EP3976838B1 - Component from a flat steel product, and method of making a method for such component - Google Patents

Description

The invention relates to a component which is produced by forming from a sheet steel blank, as well as to a method for producing such a component.

In this text, unless explicitly stated otherwise, information on alloy components is always given in mass %.

"Flat steel products" are understood here to mean rolled products whose length and width are each significantly greater than their thickness. These include in particular steel strips, steel sheets and cut pieces obtained from them, such as blanks and the like. Flat steel products of the type in question here are used for cold forming with subsequent tempering treatment to adjust the mechanical properties of the resulting component or for hot forming to form a component in order to adjust the mechanical properties of the resulting component.

"Hot forming" is also referred to as "form hardening" or "press hardening". Strictly speaking, press hardening refers to the hardening of a workpiece or component in a cooled tool, while hot forming also includes the preceding shaping in a heated state. However, the three terms mentioned are often used synonymously.

"Quenching and tempering" refers to a treatment known from the state of the art in which heating initially takes place to a temperature at which the steel of the flat steel product (component) being processed has a completely austenitic structure. This heating serves to bring the respective component to a suitable temperature. This heating is carried out as a separate work step on the component that was previously cold-formed from the flat steel product. After heating, the component is cooled at an accelerated rate so that the steel of the flat steel product from which the component is formed forms a hardened structure, with the result that the component has significantly increased strength. After quenching, the component can be subjected to tempering in order to reduce the internal stresses that can arise in the structure of the component as a result of the quenching process. In general, the highest possible temperatures are aimed for during tempering in order to shorten the cycle time and thus the costs.

The Ac3 temperatures given in this text, i.e. the temperature above which the transformation of the steel structure into the austenitic state is completed when heated, were estimated using the following formula: Ac3 [°C] = 902°C - (225 * %C - 19 * %Si + 11 * %Mn - 400 * %P - 181 * %Al + 5 * % Cr + 26 * %Cu - 13 * %Mo + 20 * %Ni - 55 * %V) * °C/mass%. where in this formula %C denotes the respective carbon content, %Si the respective silicon content, %Mn the respective manganese content, %P the respective phosphorus content, %Al the respective aluminium content, %Cr the respective chromium content, %Cu the respective copper content, %Mo the respective molybdenum content, %Ni the respective nickel content and %V the respective vanadium content of the steel composition whose Ac3 temperature is to be determined, and where the contents of the The relevant elements, if present, are inserted into the formula in mass%.

The Ar3 temperatures given in this text, i.e. the temperature at which the transformation of the previously austenitic structure of the steel begins after cooling, were estimated using the following formula: Ar3 [°C] = 910 °C - (203 * square root(%C) + 30 * %Mn - 44.7 * %Si + 11 * %Cr - 31.5 * %Mo + 15.2 * %Ni) * °C/mass% where in this formula, %C also denotes the respective carbon content, %Si the respective silicon content, %Mn the respective manganese content, %Cr the respective chromium content, %Mo the respective molybdenum content and %Ni the respective nickel content of the steel composition whose Ar3 temperature is to be determined, and where the contents of the elements in question, if present, are each inserted into the formula in mass %.

Flat steel products of the type in question here are required in particular for the manufacture of components for passenger or commercial vehicles, the mechanical properties of which have to meet the highest requirements, and of components which are exposed to high abrasive loads in practical use, such as components for machines and vehicles used in agriculture, road construction, mining or the like.

Since the beginning of the 1980s, there has been a continuously increasing demand for weight reduction from an environmental point of view, particularly in automobile bodies. Reducing the weight of the vehicle is intended to reduce the moving masses, so that less fuel is required to drive the vehicle. vehicle is required and therefore fewer climate-damaging gases are emitted.

In the steel processing industry, the trend has become established to use steel grades with ever increasing strengths to achieve a reduction in sheet thickness and thus the desired weight reduction without reducing the performance of the flat steel products in question. On the one hand, this is possible by using high-strength steels that can also be formed when cold. On the other hand, steel concepts have become established that can be hardened by heat treatment in which they undergo austenitization and subsequent controlled cooling. The strength of steels processed in this way can be further increased by martensite transformation, and internal stresses can optionally be reduced by tempering following hardening.

In addition to the state of the art explained above, it is clear from the JP 2012 180594 A a component is known which is formed from a sheet metal by hot press forming and is characterized by high strength and toughness. According to one embodiment, the sheet metal in question was produced from a steel which consisted of, in mass %, 0.31% C, 1.31% Mn, 0.009% P, 0.004% S, 0.48% Cr, 0.2% Mo, 0.022% Ti, 0.04% Nb, 0.060% Al, 0.0015% B and 0.0042% N, the remainder being iron.

The invention is based on the object of creating a weight-reduced component which, in the tempered and/or hot-formed state, has an optimal combination of strength and toughness and is, as such, suitable for uses which place the highest demands on the mechanical properties or resistance to abrasive wear.

In addition, the invention should also mention a method for producing such a component.

A component that solves this problem according to the invention has at least the features specified in claim 1.

A method that achieves the above-mentioned object according to the invention comprises at least the work steps specified in claim 8. It goes without saying that a person skilled in the art, when carrying out the method according to the invention, will add the work steps not explicitly mentioned here, which he knows from his practical experience are regularly used when carrying out such methods.

Advantageous embodiments of the invention are specified in the dependent claims and are explained in detail below.

Carbon "C" is contained in the steel of the flat steel product from which a component according to the invention is formed as a mandatory element in amounts of 0.1 - 0.6% by mass, in particular 0.10 - 0.60% by mass. The level of hardening potential is controlled by the presence of C. As the C content increases, after austenitization and accelerated cooling, both the martensite content and the hardness of the martensite obtained in the microstructure of a component according to the invention increase, with a single-phase martensite structure representing the target microstructure of the finished processed component. The increase in hardness is equivalent to an increase in strength in the tensile test. This enables a reduction in sheet thickness and thus a reduction in weight in force-transmitting component cross-sections, as is the aim in modern automotive structural engineering with regard to resource-saving lightweight body construction. In order to efficiently increase the component hardness and strength, a C content of at least 0.1% by mass, in particular at least 0.10% by mass, is required. The beneficial effects of the presence of C can be achieved particularly reliably in the flat steel product according to the invention with a C content of at least 0.12% by mass, in particular at least 0.15% by mass. With C contents of more than 0.60% by mass, the hardness or strength after accelerated quenching would be so high that the toughness of the flat steel product from which a component according to the invention is formed, which is reflected in the elongation at break or the reduction in area at break, would be significantly reduced. At the same time, the tendency to crack formation would increase and weldability would deteriorate. Negative effects of the presence of C can be particularly reliably prevented by limiting the C content to a maximum of 0.55% by mass. % by mass, in particular a maximum of 0.50 % by mass. The optimal C content is therefore 0.12 - 0.55 % by mass, in particular 0.15 - 0.50 % by mass. However, for certain applications, in particular those applications in which high abrasive loads occur and in which the potentially negative effects of higher C contents play only a minor role, it may also be sensible to provide C contents of at least 0.5 % by mass due to the associated high hardness.

Silicon "Si" can optionally be present in the steel of the flat steel product from which a component according to the invention is formed, in contents of up to 0.8 mass%. Si hinders the cementite and pearlite transformation and thereby increases the martensite hardenability of the flat steel product. Si reduces the cooling rate that is critical with regard to the desired martensite formation and thus increases the hardening of a flat steel product produced in accordance with the invention. Si also shows an inverse segregation behavior than Mn and thus improves the overall segregation resistance of the steel from which the flat steel product is made from which a component according to the invention is formed. Minimizing segregation across the cross section is of particular importance when a component according to the invention is pipes or the like. Reduced segregation sensitivity can prevent cracks in longitudinally welded pipes, particularly at higher C contents. In order to be able to use the positive effects of the presence of Si, Si contents of at least 0.1% by mass, in particular at least 0.15% by mass, can be provided. However, excessively high Si contents could impair the wetting behavior of the flat steel product from which a component according to the invention is formed, in particular if flat steel products alloyed according to the invention are to be hot-dip coated. Si tends to form external oxides during the annealing of the flat steel product carried out in this case. To prevent this, the Si content of a The Si content of the flat steel product used in the component according to the invention is at most 0.8% by mass. Negative effects of the presence of Si can be particularly reliably avoided if the Si content is limited to a maximum of 0.5% by mass. The optimal Si content is therefore 0.1 - 0.8% by mass, in particular 0.15 - 0.5% by mass.

Manganese "Mn" is present in the steel of the flat steel product from which a component according to the invention is formed, in contents of 0.1 - 2 mass%, in particular 0.10 - 2.0 mass%. Mn increases the hardenability of the steel by lowering the A3 transformation temperature (i.e. the Ac3 and/or Ar3 temperature) from ferrite to austenite. As a result, during the heat treatment of the flat steel product from which a component according to the invention is formed, the furnace temperature for complete conversion to austenite can be reduced during heating. In particular, the formation of the diffusion-controlled transformation phases ferrite, pearlite and bainite is shifted towards longer times. Therefore, in this respect, manganese is a similarly effective alloying element as carbon. Compared to carbon, manganese has the advantage of achieving a higher deformability in the hardened state, which is expressed, for example, in a higher notched impact toughness. The reduction in the critical cooling rate with increasing manganese content is also associated with an increase in the hardening capacity. Fluctuations in the cooling conditions or different contact conditions during the cooling of components made from steel material alloyed according to the invention can be better compensated and the property scatter is limited. However, excessively high Mn contents increase the C segregation behavior and can lead to inhomogeneous hardening behavior across the cross-section of the respective product and the formation of hardening cracks. Increasing Mn contents also increase the risk of external Mn oxides or Mn-based Form mixed oxides. As in the case of excessive Si contents, this would trigger the risk of a deterioration in the wetting behavior of a flat steel product made from a steel material alloyed according to the invention when hot-dip coated. In the case of bell annealing processes, excessive Mn contents on the hot or cold strip surface would also lead to undesirable discoloration or so-called "manganese haze" due to the formation of manganese oxides. In order to avoid these negative effects, the Mn content of a flat steel product intended for forming a component according to the invention is limited to a maximum of 2% by mass, in particular a maximum of 2.0% by mass, whereby unfavorable effects of the presence of Mn can be particularly reliably avoided if the Mn content is limited to a maximum of 1.5% by mass, in particular 1.50% by mass. In contrast, the positive influences of Mn on the properties of a flat steel product intended for forming a component according to the invention can be used particularly reliably if the Mn content is at least 0.4% by mass, in particular at least 0.40% by mass. The Mn content is therefore optimally 0.4 - 1.5% by mass, preferably 0.40 - 1.50% by mass, in particular 0.6 - 1.3% by mass or 0.6 - 1.2% by mass, preferably 0.60 - 1.30% by mass or 0.60 - 1.20% by mass.

Phosphorus "P" is one of the unavoidable steel components that are part of the manufacturing process. P segregates particularly at the grain boundaries and reduces the grain boundary strength. Higher P contents would therefore contribute to weakening the structure, which in turn would lead to a deterioration in the toughness of the material. The P content of a flat steel product intended for forming a component according to the invention is therefore limited to a maximum of 0.03% by mass, whereby the P content should be set as low as possible. The P content of the flat steel product is therefore preferably a maximum of 0.025% by mass, in particular a maximum of 0.02% by mass.

Sulfur "S" is also an accompanying element, the presence of which is fundamentally undesirable in the flat steel product intended for forming a component according to the invention. Due to the Mn contents provided according to the invention, non-metallic MnS precipitates would form at higher S contents, which would be present in an elongated form after rolling the flat steel product due to their low hardness and would have a negative effect on the fracture behavior. During deformation, the first microscopic material separations could form due to crack initiation and crack propagation on elongated MnS, expand and grow together until they worsen the material behavior macroscopically in the form of reduced notched impact toughness and increasing material anisotropy. In order to exclude the negative effects of the presence of S in the steel alloyed according to the invention, the S content is limited to a maximum of 0.03 mass%, with low S contents of less than 0.006 mass%, in particular less than 0.003 mass%, being particularly favorable.

Aluminum "Al" is present in the steel of the flat steel product from which a component according to the invention is formed in contents of 0.05 - 0.2 mass%, in particular 0.050 - 0.20 mass%. Al is traditionally used as a deoxidizing element, for which purpose it is typically alloyed in practice in contents of 0.02 - 0.05 mass%. According to the invention, on the other hand, increased Al contents of 0.050 - 0.20 mass%, in particular 0.050 - 0.20 mass%, in combination with optional, low Ti contents of up to 0.02 mass%, in particular up to 0.020 mass%, are provided in the steel alloyed according to the invention. In this way, the formation of AIN or NbN is promoted in competition with the nitrogen binding by TiN that is traditionally known in tempering steels and, if Ti is present in the steel of a flat steel product intended for forming a component according to the invention, the formation of comparatively coarse TiN is avoided. The aim is to prevent the formation of boron nitrides to avoid so that B, as explained below, can exert its beneficial influence on delaying the transformation in dissolved form in the crystal lattice. In addition, the presence of Al within the content limits specified in the invention results in grain refinement. At Al contents of less than 0.05% by mass, in particular less than 0.050% by mass, the precipitation pressure for the formation of AIN would be too low. In order to be able to safely utilize the positive effects of Al in the flat steel product according to the invention, the Al content can be set to at least 0.06% by mass, in particular at least 0.060% by mass, or at least 0.07% by mass, in particular at least 0.070% by mass. However, with contents of more than 0.2% by mass, in particular more than 0.20% by mass, Al there is a risk that external Al oxides will form on the surface of a product made from steel material alloyed according to the invention, which would impair the wetting behavior during hot-dip coating. In addition, with higher Al contents, the formation of non-metallic Al-based inclusions would be promoted, which essentially as alumina (Al ₂ O ₃ ) and aluminum nitride (AIN) also have a high hardness (Mohs hardness 9) and are therefore undesirable with regard to avoiding the risk of crack initiation and propagation during plastic deformation and cyclic stress. It is also disadvantageous that the oxidic Al precipitates can form conglomerates with other types of precipitates such as sulfides and silicates and thus form larger precipitates that can have a higher crack initiation and failure potential. This can prove to be particularly risky in the case of flat steel products made from steel alloyed according to the invention, which can achieve strengths of up to 2500 MPa after tempering or hot forming. In addition, high Al contents cause longitudinal cracks to form in the slabs cast from the steel alloyed according to the invention during processing. In addition, Al causes a drastic increase in the Ac3 transformation temperature, so that at higher Al contents the temperature that must be exceeded for complete austenitization is unnecessarily would be increased. According to the invention, the Al content of a flat steel product according to the invention is therefore limited to a maximum of 0.2% by mass, in particular a maximum of 0.20% by mass, whereby negative effects of the presence of Al in the flat steel product intended for forming a component according to the invention can be particularly reliably avoided by limiting the Al content to a maximum of 0.15% by mass, in particular a maximum of 0.150% by mass, or a maximum of 0.13% by mass, in particular a maximum of 0.130% by mass. The Al content of a flat steel product intended for forming a component according to the invention is therefore optimally 0.06 - 0.15% by mass, in particular 0.07 - 0.13% by mass, whereby Al contents of 0.060 - 0.150% by mass, in particular 0.070 - 0.130% by mass, have proven particularly effective.

Chromium "Cr" is present in the steel of the flat steel product from which a component according to the invention is formed as a mandatory element in contents of 0.05 - 0.8 mass% in order to increase the hardenability via the transformation-retarding effect. Chromium effectively suppresses the formation of ferrite and pearlite during accelerated cooling of the flat steel product and enables complete martensite formation even at lower cooling rates, thereby increasing the hardenability. The presence of Cr in the contents provided for in the invention thus contributes to the hardenability of the flat steel product according to the invention by means of suitable cooling and reduces the scatter of the local product properties. Cr also increases the tensile strength without significantly impairing the elongation. This is also explained by the formation of chromium carbides, which have a strength-increasing effect and can increase the tempering resistance. To achieve the desired effect on the transformation, a minimum content of 0.05 mass% Cr is required, whereby this effect is particularly likely to occur at Cr contents of at least 0.15 mass% Cr. At contents of more than 0.8 mass% Due to the overall alloy layer selected according to the invention, no increase in the positive influence of Cr can be observed, whereby the transformation-retarding effect of Cr can be used particularly effectively at contents of up to 0.55 mass%. Therefore, the Cr contents of a flat steel product intended according to the invention for forming a component according to the invention are optimally 0.15 - 0.55 mass%.

Nitrogen "N" can be present in the alloy concept according to the invention as a fundamentally undesirable accompanying element in contents of up to 0.01% by mass. At higher N contents, increased contents of nitride formers such as Ti, Nb, Al are necessary in order to be able to bind N as nitride. At the same time, the risk of the formation of coarser, toughness-impairing TiN precipitates increases in particular if Ti is optionally added to the alloy. In particular, there would then also be a risk that B would no longer be available in dissolved form. BN formation must be avoided, since otherwise the desired transformation-retarding effect of free boron could not be used. In order to keep the alloying measures required for this purpose within limits, the N content is limited to a maximum of 0.01% by mass, whereby the negative influence of N on the properties of a flat steel product intended for forming a component according to the invention can be particularly reliably avoided by limiting the N content to a maximum of 0.007% by mass, in particular a maximum of 0.005% by mass.

Niobium "Nb" is present in contents of 0.01 - 0.06 mass%, in particular 0.010 - 0.060 mass%, as a mandatory element in the steel of the flat steel product from which a component according to the invention is formed. Nb has a strong grain-refining effect because it can hinder grain growth even as a dissolved alloying element in the austenite. In addition, Nb forms fine carbide or nitride precipitations, which in the case of nitrides are significantly finer than, for example, TiN. Grain refinement and precipitation formation contribute to increasing the strength of the end product made from the steel material alloyed according to the invention and also improve the toughness. In addition, fine precipitations contribute to prevent cracks. Furthermore, fine precipitates are more beneficial than coarse precipitates in terms of preventing crack formation and crack propagation. A finer austenite grain size also reduces the martensite package size, which leads to a more homogeneous hardness and strength distribution in the product produced from a steel material alloyed according to the invention. In addition, the presence of Nb has a positive effect on the segregation behavior, since the finer formation of the grain structure already in the austenite state promotes a refinement of the segregation structure. In order to achieve these positive effects, the minimum Nb content of a flat steel product intended for forming a component according to the invention is 0.010 mass%, with Nb contents of at least 0.015 mass% or at least 0.020 mass% having proven to be particularly favorable. The upper limit of the Nb content in the flat steel product used according to the invention is 0.060 mass%, since with increasing Nb content a clogging effect can occur when casting the steels alloyed according to the invention that are melted to produce the flat steel product. In addition, especially with higher C contents at the same time, there is a risk that increased Nb contents can no longer be completely dissolved when heating slabs cast from steel material alloyed according to the invention at minimum furnace temperatures of 1100 °C. The complete dissolution of Nb-based precipitates during slab preheating is, however, advantageous in order to be able to make optimal use of grain refinement and to be able to form finely distributed, strength-relevant Nb precipitates during hot rolling or in later process phases (recrystallization annealing, hot forming furnace) of processing steel material alloyed according to the invention. Too high Nb contents can also have a negative effect on the coating behavior in the hot-dip process. The advantageous effects of the presence of Nb in the flat steel product intended for forming a component according to the invention can be particularly reliably achieved at contents of up to 0.05 mass% Nb, in particular up to 0.050 mass% Nb, or up to 0.04 mass% Nb, in particular up to 0.040 mass% Nb. The optimum Nb content of the flat steel product is therefore 0.015 - 0.05 mass%, in particular 0.015 - 0.050 mass%, with contents of 0.020 - 0.04 mass%, in particular 0.020 - 0.040 mass%, having proven particularly effective.

Titanium "Ti" can optionally be added to the steel of the flat steel product from which a component according to the invention is formed, in amounts of up to 0.02% by mass, in particular up to 0.020% by mass, in order to bind the nitrogen that is inevitably present in the steel and to ensure that B is retained in an unbound, interstitially dissolved form. At the same time, the Ti content must be limited so that the formation of coarse TiN precipitations is avoided in order to minimize as far as possible the risk of crack initiation and crack propagation in high-strength products produced from steel material alloyed according to the invention, in particular under cyclic and dynamic stress. The favorable influences of Ti on the properties of a flat steel product intended for forming a component according to the invention are certain to occur if the Ti content is at least 0.001% by mass, in particular at least 0.004% by mass or at least 0.010% by mass. For the purposes of the invention, a Ti concentration of 0.004 mass% or more, in particular of at least 0.005 mass%, is to be classified as a deliberately added element. Ti contents that are below the minimum limit of 0.004 mass%, in particular 0.005 mass%, specified in the invention for the Ti content are always regarded as unavoidable contamination that is introduced by the starting materials used in the production of the steel. At the same time, negative effects of Ti can be particularly reliably avoided by limiting the Ti content to a maximum of 0.020 mass%. The Ti content of the steel from which a component according to the invention is made is therefore ideally 0.004 - 0.016 mass%.

Optionally, the respective Ti content %Ti can be adjusted to the respective N content %N of the flat steel product from which a component according to the invention is formed, so that the ratio %Ti/%N applies: $$\%\mathrm{Ti}/\%\mathrm{N}<4.$$

According to an embodiment of the invention which is particularly advantageous in this context with regard to the desired setting of N, the Ti content %Ti, the N content %N as well as the Al content %AI and the remaining N content %Nrest which is still present after setting of N by Ti to TiN (%Nrest = %N -% Ti*14/48) are coordinated with one another in such a way that not only the condition %Ti/%N < 4 is met, but that the difference between the respective Al content %Al and the respective stoichiometrically determined size 27/14*%Nrest applies: $${\mathrm{UG}}_{\mathrm{Al\_Nrest}}\le \%\mathrm{Al}-27/14*\%\mathrm{Nrest}\le {\mathrm{OG}}_{\mathrm{Al\_Nrest}}$$

In this case, UG _{Al_Nrest} is equal to 0.070 mass%, in particular equal to 0.075 mass%, preferably equal to 0.080 mass%, in particular equal to 0.081 mass%, and OG _{Al_Nrest} is equal to 0.150 mass%, in particular equal to 0.135 mass%, preferably equal to 0.125 mass%, particularly preferably equal to 0.121 mass%. Accordingly, according to a particularly advantageous embodiment of the invention: $$\mathrm{0}\mathrm{,081}\mathrm{Masse}-\%\le \%\mathrm{Al}-27/14*\%\mathrm{Nrest}\le \mathrm{0}\mathrm{,121}\mathrm{Masse}-\%$$

Boron "B" is present in the steel of the flat steel product from which a component according to the invention is formed, in contents of 0.0005 - 0.005 mass% as a mandatory component. B is an effective hardening element that can significantly delay transformation even in very low contents and thus significantly increases hardenability. In addition, B improves grain boundary strength by primarily accumulating at grain boundaries and thus displacing harmful elements, such as P, from there. In this way, In this way, toughness and reduction in area at fracture are improved. Below 0.0005 mass% B, however, the transformation delay is too small. The beneficial effect of B in the flat steel product from which a component according to the invention is formed can therefore be used particularly reliably with B contents of at least 0.001 mass%. However, with contents of more than 0.005 mass% boron, a saturation effect occurs. The beneficial effects of the presence of B can therefore be used particularly effectively with B contents of at most 0.0035 mass%, in particular at most 0.0030 mass%. The B content of a flat steel product intended for forming a component according to the invention is therefore optimally 0.001 - 0.0035 mass%, in particular 0.001 - 0.003 mass%.

Molybdenum "Mo" can optionally be present in the steel of the steel product according to the invention in contents of up to 1.5% by mass. Like chromium, Mo suppresses the formation of ferrite and pearlite during cooling and enables increased martensite or bainite formation even at lower cooling rates, thereby increasing hardenability. The hardenability-enhancing effect of Mo is significantly higher than that of Cr. In this respect, Mo can effectively increase strength in large thicknesses and cross-sections where, due to dimensions or design, only relatively low cooling rates are possible. Mo also reduces the tempering embrittlement of tempering steels. Mo is also a strong carbide former and can therefore also contribute to increasing strength through precipitation formation. These favorable effects of Mo occur with optional Mo contents of at least 0.03% by mass, whereby the hardness-increasing contribution of Mo can be used particularly reliably with Mo contents of at least 0.1% by mass. However, if the Mo content is too high, the hot formability of the steel would be too limited. In addition, if the Mo content is too high, it could form low-melting sulphide compounds with S, which locally increase the risk of cracking during hot forming and could thus promote defects near the surface, for example. Therefore, the Mo content is limited to limited to a maximum of 1.5 mass%. Negative effects of the presence of Mo can be particularly reliably avoided by limiting the Mo content to a maximum of 0.5 mass%.

Also optionally, one or more elements from the group "Cu", "Ni", "V" and "REM" can be present in the steel of the flat steel product from which a component according to the invention is formed, in accordance with the requirements explained below:
Copper "Cu" and nickel "Ni" can optionally be provided in the steel of the flat steel product from which a component according to the invention is formed in order to increase the hardenability. Suitable Cu contents for this purpose are up to 0.5% by mass, with the effect of Cu occurring from an optional content of at least 0.1% by mass. Ni can be provided in contents of up to 1.5% by mass if not only the hardenability but also the toughness of the component made from a steel product alloyed according to the invention is to be improved. For this purpose, optional Ni contents of at least 0.15% by mass are required.

Vanadium "V" can also optionally be present in the steel of the flat steel product from which a component according to the invention is formed in order to bring about precipitation strengthening. Suitable V contents for this are up to 0.2 mass%, whereby the effect of V can be utilized through optional contents of at least 0.03 mass%.

Rare earths "REM", such as cerium and lanthanum, can cause grain refinement in the steel of the flat steel product from which a component according to the invention is formed, and thus increase toughness and strength. In order to use this effect, REM contents of at least 0.02% by mass can optionally be present. This effect can be used particularly effectively with REM contents of up to 0.05% by mass.

Calcium "Ca" is optionally present in the steel of the flat steel product from which a component according to the invention is formed in amounts of up to 0.005% by mass. Ca can be added to the steel to influence the sulfide form. It also forms sulfides, for example, in competition with manganese. Due to the higher hardness of CaS, a round precipitation form is retained in the rolling process and a smaller interface with the substrate is the result. This prevents the development of a preferred direction when cracks are initiated and propagated. In conjunction with a reduction in the sulfur content, this improves the material toughness and isotropy. In order to use this safely, the Ca content can be set to at least 0.001% by mass. If the Ca content is too high, however, the probability of other non-metallic inclusion types involving Ca forming would increase, which would impair the purity of the steel and also the toughness. For this reason, an upper limit of the Ca content of not more than 0.005 mass%, preferably not more than 0.003 mass%, should be observed.

For example, tin "Sn", arsenic "As" and cobalt "Co" and all other potential alloying elements not mentioned here are present in the steel of the flat steel product from which a component according to the invention is formed, at most as accompanying elements to be attributed to the unavoidable impurities, the contents of which must be set as low as possible, but in any case must be minimized so that they have no influence on the properties of the flat steel product and the products made from it. To this end, the following upper limits must be observed for the contents of Sn, As and Co: Sn: < 0.05 mass%, As: < 0.05 mass%, Co: < 0.05 mass%.

mit UG_%Al/%N = 8,0. Liegt das Verhältnis (%Al / %N) * 14/27 unter diesem Grenzwert, so steht nicht genügend Al zur Verfügung, um die N-Abbindung über die erstrebte Bildung von AIN gegenüber BN durch das höhere Al-Angebot ausreichend zu begünstigen. Dieses Risiko kann dadurch gemindert werden, dass die Untergrenze UG_%Al/%N für das Verhältnis (%Al / %N) * 14/27 auf UG_%Al/%N = 8,3, insbesondere UG_%Al%N = 8,6 oder, besonders bevorzugt, auf UG_%Al/%N = 9,0, angehoben wird.As already explained, B in the steel of a steel flat product intended for forming a component according to the invention makes a decisive contribution to the hardenability by delaying the microstructure transformation during cooling. At the same time, B improves the toughness and fracture reduction of the steel flat product. The B in the steel according to the invention The Al and Nb contents provided for in the flat steel product of a component according to the invention ensure that the nitrogen, which is always present in the steel in certain quantities, even if it is undesirable, is bound before boron nitrides can form. The invention provides for this purpose that in each case there is enough Al in the steel of the flat steel product that the ratio formed from the respective Al content %Al and the respective N content %N satisfies the following condition: $$\begin{array}{l}\left(\%\mathrm{Al}/\%\mathrm{N}\right)*14/27\\ \ge {\mathrm{UG}}_{\%\mathrm{Al}/\%\mathrm{N}}\end{array}$$

with UG _%Al/%N = 8.0. If the ratio (%Al / %N) * 14/27 is below this limit, there is not enough Al available to sufficiently promote N binding via the desired formation of AIN compared to BN due to the higher Al supply. This risk can be reduced by raising the lower limit UG _%Al/%N for the ratio (%Al / %N) * 14/27 to UG _%Al/%N = 8.3, in particular UG _%Al%N = 8.6 or, particularly preferably, to UG _%Al/%N = 9.0.

mit OG_%Al/%N = 15, insbesondere OG_%Al/%N = 15,0, eingehalten wird, lässt sich verhindern, dass zu viel Al im Stahl zur Verfügung steht. Dies hätte andernfalls die Gefahr zur Folge, dass gröbere, oxidische Al-Ausscheidungen des Typs Al₂O₃ sowie deren Konglomerate mit Silikaten und Sulfiden anwachsen könnten. Ist OG_%Al/%N = 13,5, insbesondere OG_%Al/%N = 12,0, so lässt sich dieses Risiko besonders sicher minimieren. Als besonders vorteilhaft erweist es sich somit, wenn für das Verhältnis %Al/%N gilt: $$\begin{array}{l}\mathrm{9,0}\le \left(\%\mathrm{Al}/\%\mathrm{N}\right)*14/27\\ \le \mathrm{12,0}\end{array}$$

By simultaneously satisfying the condition for the ratio %Al/%N $$\begin{array}{l}\left(\%\mathrm{Al}/\%\mathrm{N}\right)*14/27\\ \le {\mathrm{OG}}_{\%\mathrm{Al}/\%\mathrm{N}}\end{array}$$

with OG _%Al/%N = 15, in particular OG _%Al/%N = 15.0, it is possible to prevent too much Al from being available in the steel. This would otherwise result in the risk of coarser, oxidic Al precipitations of the type Al ₂ O ₃ and their conglomerates with silicates and sulphides growing. If OG _%Al/%N = 13.5, in particular OG _%Al/%N = 12.0, this risk can be minimised particularly safely. It is therefore particularly advantageous if the ratio %Al/%N is: $$\begin{array}{l}\mathrm{9,0}\le \left(\%\mathrm{Al}/\%\mathrm{N}\right)*14/27\\ \le \mathrm{12,0}\end{array}$$

It goes without saying that all optional elements can be present individually or in combination with one another as impurities in the steel of the flat steel product intended for forming a component according to the invention. In this case, the contents of the elements in question are so low that they are below the minimum limits above which the effect of the respective element can be used according to the above explanations. If the contents of the optionally present alloying elements are below these minimum limits, these elements have no effect on the properties of the flat steel product and can therefore be tolerated in the sense of impurities.

Endogenous or exogenous inclusions (particles, precipitates) that arise during steel production generally lead to a reduction in the degree of purity, which can lead to premature failure of components. This can be an increasing problem, especially with high-strength components. This is especially true when such components are exposed to cyclic or dynamic loads. Of interest here are the endogenous inclusions that arise during the steel production process due to the thermodynamic conditions of the chemical composition and process control. Exogenous inclusions are usually isolated cases and come from ladle slag or refractory material, for example, but play no role here and are therefore not considered here.

Based on these findings, one of the aims of the inventive adjustment of the alloy of the steel of a component according to the invention was to reduce the proportion of coarse and hard TiN, AIN and oxidic Al-based particles as well as conglomerates of these compounds for reasons of toughness and nevertheless to securely bind the nitrogen present in each case in order to achieve a complete transformation into martensite even at relatively low cooling rates of at least 30 °C/s to a maximum of 120 °C/s via the strong transformation-retarding effect of interstitially dissolved B even at relatively low cooling rates of at least 30 °C/s to a maximum of 120 °C/s. strip thicknesses and component cross-sections. As already explained above, Al and the optionally present Ti form hard precipitates which, in components formed from steel flat products alloyed according to the invention, could be the source of cracks and their propagation due to the notch effect and the stress fields surrounding the particles. The angular and cube-shaped TiN particles in particular prove to be harmful here simply because of their shape and size.

The invention has coordinated the contents of the alloying elements and the conditions during the production of flat steel products intended for forming components according to the invention in such a way that in the structure of a flat steel product according to the invention and of a component produced therefrom, a maximum of up to 150 area ppm of hard TiN particles and Al-based oxide particles as well as AIN with an average, circle-equivalent particle size of 0.2 - 10 µm are present homogeneously distributed over the strip thickness.

The definition of "hard particles" here essentially includes particles of AIN, Al ₂ O ₃ and Al ₂ O ₃ -based spinels as well as TiN particles and conglomerates formed on the basis of the above-mentioned particles. Such particles each have a high Mohs hardness of approx. 9. Due to their high hardness, they are hardly deformable in rolling or forming processes and lead to local stress fields in their environment, which can promote premature material failure. Conglomerates (mixed forms) are particularly referred to here as particle composites in which further particles form through heterogeneous nucleation on already existing particles, e.g. Al ₂ O ₃ with MnS, where the basis is one of the previously mentioned hard particle types.

The alloy concept according to the invention has also achieved that the total number of hard TiN-based precipitates and their mixed forms falling within this particle size range in a The particle size distribution in the component formed from a flat steel product alloyed according to the invention is reduced to less than 30% of the particles in the size class 0.2 - 10 µm present in the structure of a component. At the same time, the absolute number of precipitates falling into the relevant particle size range is reduced compared to conventional flat steel products, for example made from a steel with a higher Ti content, as a result of which the average distance between the 0.2 - 10 µm sized precipitates in the component formed from a flat steel product alloyed according to the invention is significantly increased. It was found that in the comparison concept the proportion of hard TiN particles and their mixed forms accounts for more than 45 - over 80% of the volume proportion of the particles present in the size class 0.2 - 10 µm. Due to this high proportion, a reduction in the Ti mass proportion makes sense, which accordingly leads to a reduction in the proportion of hard TiN particles in the inventive concept. Among other things, due to the optional addition of titanium, which is in any case limited according to the invention, coarse particles such as TiN occur much less frequently in a flat steel product intended for forming a component according to the invention than is the case with conventional concepts in which higher Ti contents are provided. By reducing the proportion of coarse precipitates, an improvement in toughness is achieved that prevents the formation and spread of cracks. Surprisingly, the moderate increase in the Al mass content does not lead to a significant increase in the proportion of similarly hard, oxidic Al-based precipitates as well as AIN and their conglomerates. As a result, the risk of premature material failure is reduced in the flat steel products intended for forming a component according to the invention. The optimization of toughness achieved by the invention is noticeable in an improvement in the fracture necking on the component according to the invention in the hot-formed, tempered state, in which it is of particular interest to the component manufacturer.

When producing a component according to the invention, a structure consisting entirely of martensite in the technical sense is produced by full austenitization with subsequent quenching and optional tempering treatment. According to the understanding of a person skilled in the art, this naturally includes the possibility that up to 5% of other components are present in the structure of a component according to the invention, which, however, are ineffective with regard to the properties of a component according to the invention determined by the martensite content.

As already mentioned, the Al and Nb contents provided in the steel according to the invention result in additional microstructure refinement. Nb and Al in dissolved and precipitated form reduce austenite grain growth during the production and heat treatment of the flat steel product made from the alloyed steel according to the invention and the component made from it, and reduce the martensite package size after the transformation. This creates other relevant precipitates in the flat steel product, such as NbN, NbC and AIN, which generally only reach a maximum size of up to approx. 100 nm as monolithic particles without nucleation on previously formed precipitates. In this way, more homogeneous precipitate fractions with narrower particle size ranges are achieved. These prove to be particularly effective with regard to controlling the austenite grain size. Thus, the steel used to produce the flat steel product from which a component according to the invention is formed has an austenite grain size during austenitization that is up to half an ASTM grain size finer than in conventional steel concepts belonging to the type of steel according to the invention. In addition, the grain sizes in a flat steel product alloyed and processed according to the invention are in a narrower range, i.e. with a reduced standard deviation. At the same time, there is a reduced variation in the former austenite grain size across the strip thickness. This leads to finer martensite packages and a high homogeneity of the martensite structure. which is advantageous for the toughness of components made from such a flat steel product in accordance with the invention in the tempered or press-hardened state. This also leads to better dimensional stability of the component, since strength fluctuations across the strip thickness can be reduced.

An important material parameter for setting the final properties is the former austenite grain size. This is the grain size of the austenite that is established after the completion of the austenitization process in the furnace as a result of recrystallization and grain growth, i.e., it is the size that prevails in the structure shortly before the start of quenching. The finer this austenite grain size on average, the finer the resulting martensite package size and the more beneficial it is for the toughness of the martensite and thus of the material or component.

It is also beneficial for the homogeneity of the local strength properties if the grain size of the austenite fluctuates only slightly and the martensite hardness therefore only shows small local fluctuations after the transformation. This also makes it possible to avoid springback effects on a press-hardened or tempered component due to locally inhomogeneous structures.

Components produced in accordance with the invention are therefore characterized after the quenching carried out during hot forming or after the quenching treatment in that the product KG = KA x Ks, referred to as "grain size quality KG", of the former austenite grain size KA, used in µm, and the simple standard deviation Ks of the former austenite grain size, also used in µm and averaged at three points over half the strip thickness, applies to the structure of the component obtained: $$\mathrm{KG}<30{\mathrm{\mu m}}^2$$

With such a qualified grain size quality KG, it is possible to achieve, compared to the state of the art, not only a small grain size that is advantageous for toughness, but also a high homogeneity of the microstructure across the component cross-section due to a small scatter of the grain size.

After a suitable heat treatment as explained below, components according to the invention achieve a tensile strength of at least 1000 MPa at C contents of 0.1% by mass, in particular 0.10% by mass, or tensile strengths of up to 2500 MPa at C contents of 0.6%, in particular 0.60% by mass.

The area of reduction in area ε(epsilon)3 was investigated here as a measure of toughness, since the investigation of the notched impact strength determined according to DIN EN ISO 148-1 according to Charpy is only limited to thicknesses of 10 mm or so-called undersize specimens (thicknesses of 2.5, 5 and 7.5 mm) and is therefore only suitable for the investigation of correspondingly thick specimens, which were not available here. The toughness properties were therefore not determined here according to DIN EN ISO 148-1 according to Charpy. The area of reduction in area was evaluated as a measure of toughness or local elongation, since contrary to the requirements of the notched impact test according to DIN EN ISO 148-1, different thicknesses and in particular thicknesses < 2.5 mm were available, so that no uniform, so-called undersize specimens could be used and the application of this standard was therefore not suitable. Although the local strain from the fracture area correlates with the hole expansion, it represents an extended description of the local deformation behavior.

Accordingly, components according to the invention are simultaneously characterized by an excellent toughness for this strength class, which also results in a percentage improvement in the tensile strength after a suitable heat treatment explained below. The reduction in area at fracture (ΔBE) is expressed by at least 5 to 45% compared to a Ti/B-based tempering concept with increasing tensile strength from 1000 to 2500 MPa. The absolute reduction in area at fracture in the thickness direction ε(epsilon)3 for components made of steel concepts according to the invention is 10 - 65%.

The components formed from flat steel products obtained in accordance with the invention in the manner explained above by processing according to the invention are in particular weight-reduced component applications in the automotive and truck sector, which include longitudinally welded tubes for stabilizers, rotor shafts, camshafts or tubular components that are used in the steering and chassis areas. In particular, a flat steel product according to the invention can be used by cold forming for a seam-welded steel tube that is suitable for use, for example, as a stabilizer for vehicle suspension, a steering shaft or a drive shaft of motor vehicles. A subsequent tempering treatment can achieve a significant increase in the strength of the formed tube.

The thickness of flat steel products intended according to the invention for forming components according to the invention is typically 1 - 16 mm, whereby sheets with a thickness of 2 - 9 mm, in particular 4 - 7 mm, can be used for automotive applications, whereby thicknesses of up to 5 mm can be of particular importance in practice. If special requirements are placed on the resistance to abrasive wear of such flat steel products, it has proven advantageous, due to the associated high hardness, if the C content of the flat steel products intended according to the invention for forming a component according to the invention is at least 0.5% by mass, in particular 0.50% by mass.

Furthermore, flat steel products obtained according to the invention can also be used as hot or cold strip for forming components according to the invention For example, structural components for automobile bodies can be hot-formed from such hot or cold strip, and their high strength can be retained by subsequent targeted cooling from the forming heat. Flat steel products intended according to the invention for forming components according to the invention, which are provided for this processing route, typically have a thickness of 0.5 - 3.5 mm, in particular 0.5 - 3 mm, 1 - 3 mm or 1.2 - 2.5 mm. Examples of components according to the invention that can be formed from flat steel products made according to the invention are supports of automobile structures subject to bending, such as the B-pillars or seat cross members of passenger vehicles, as well as longitudinal and cross members of passenger or commercial vehicle chassis, and generally body structure parts. The flat steel product according to the invention is also particularly suitable for processing into components that move during use, such as parts of shock absorbers, camshafts or parts thereof, piston rods or shafts, in particular shafts of an electric motor. It is therefore possible to obtain a high-strength component for an automobile body by hot forming or press hardening.

To protect against corrosion, flat steel products used according to the invention and components produced from them can be provided with a metallic protective layer. Metallic protective layers based on zinc or aluminum, such as AlSi coatings, which can be applied in the conventional way by hot-dip coating, are particularly suitable for this purpose. Electrolytic coatings are also conceivable.

The chemical composition of the steel of a component according to the invention, as specified in accordance with the invention, and a coordinated process control during steel production make it possible to influence the type, size and distribution of the endogenous inclusions. In addition to solidification, the influence extends in particular to the production stage of hot rolling, as explained below.

• wobei zu den Verunreinigungen Gehalte von bis zu 0,03 % P, bis zu 0,03 % S, bis zu 0,01 % N, weniger als 0,05 % Sn, weniger als 0,05 % As und weniger als 0,05 % Co zählen und
• wobei das aus dem jeweiligen Al-Gehalt %Al und dem jeweiligen N-Gehalt %N gebildete Verhältnis (%Al / %N) * 14/27 ≥ 8,0 ist,

b) die Stahlschmelze zu einem Vorprodukt vergossen wird, nämlich zu einer Bramme, einer Dünnbramme oder einem gegossenen Band,
c) das Vorprodukt, sofern erforderlich, auf eine 1100 - 1350 °C betragende Vorwärmtemperatur durcherwärmt wird,
d) das Vorprodukt zu einem Warmband mit einer Dicke von 1 - 16 mm warmgewalzt wird, wobei das Warmwalzen bei einer Warmwalzendtemperatur beendet wird, die um mindestens 50 °C und höchstens 150 °C höher ist als die Ar3-Temperatur des Stahls,
e) das erhaltene Warmband auf eine 450 - 700 °C betragende Haspeltemperatur abgekühlt wird, wobei die Abkühlung im Temperaturbereich von 800 - 650 °C mit einer Abkühlrate von 20 - 200 °C/s erfolgt,
f) das auf die Haspeltemperatur abgekühlte Warmband zu einem Coil gehaspelt wird und das Warmband im gehaspelten Zustand auf Raumtemperatur abgekühlt wird, sowie
g) optional: das im gehaspelten Zustand abgekühlte Warmband gebeizt wird und
h) ebenso optional: bei einer Kerntemperatur des Warmbands von 500 - 720 °C über eine Dauer von 5 - 50 h haubengeglüht wird.
B) Aus dem erhaltenen Warmband wird optional ein Kaltband erzeugt, indem

• i) das Warmband zu einem Kaltband mit einer Dicke von 0,5 - 3,5 mm in einem oder mehreren Kaltwalzschritten kaltgewalzt wird.
• j) Optional kann das Kaltband in einer Haubenglühe oder in einer Durchlaufglühe geglüht werden.

C) Aus dem Warmband oder dem optional daraus erzeugten Kaltband wird ein Bauteil geformt wird, indem

• k) von dem Warm- oder Kaltband eine Platine abgeteilt wird
• und gemäß Alternative 1:
  • I.1) die Platine auf eine Austenitisierungstemperatur durcherwärmt wird, die um höchstens 100 °C geringer ist als die Ac3-Temperatur des Stahls, aus dem das Warm- oder Kaltband erzeugt ist, und höchstens 950 °C beträgt,
  • l.2) innerhalb von 1 - 20 s nach dem Ende der Durcherwärmung auf die Austenitisierungstemperatur die Platine in ein gekühltes Warmumformwerkzeug eingelegt wird, in dem die Platine zu dem Bauteil warmumgeformt wird, und
  • l.3) das Bauteil durch beschleunigtes Abkühlen mit einer Abkühlrate von 30 - 120 °C/s bis Erreichen der Martensitstarttemperatur des Stahls, aus dem das jeweilige Warm- oder Kaltband besteht, pressgehärtet wird, so dass das Bauteil ein vollständig martensitisches Gefüge erhält,
• oder gemäß Alternative 2:
  • m.1) die Platine zu dem Bauteil kaltumgeformt wird,
  • m.2) das kaltgeformte Bauteil auf eine Austenitisierungstemperatur durcherwärmt wird, die um höchstens 100 °C geringer ist als die Ac3-Temperatur des Stahls, aus dem das Warm- oder Kaltband erzeugt ist, und höchstens 950 °C beträgt, und
  • m.3) das auf die Austenitisierungstemperatur durcherwärmte Bauteil mit einer Abkühlrate von 30 - 120 °C/s bis Erreichen der Martensitstarttemperatur des Stahls, aus dem das jeweilige Warm- oder Kaltband besteht, beschleunigt abgekühlt wird, so dass das Bauteil ein vollständig martensitisches Gefüge erhält.
• n) Optional kann das nach den Arbeitsschritten l.1 - l.3 oder m.1 - m.3 erhaltene Bauteil bei Temperaturen von 150 - 700 °C bei einer Glühdauer von 5 - 60 min angelassen werden.

The method according to the invention for producing a component with a structure consisting of at least 95% martensite and the remainder of other structural components therefore comprises the following work steps:
A) A hot strip is produced by
a) Steel is melted which consists of (in mass%) C: 0.1 - 0.6%, Mn: 0.1 - 2 %, Al: 0.05 - 0.2%, Nb: 0.01 - 0.06%, B: 0.0005 - 0.005%, Cr: 0.05 - 0.8%, Si: up to 0.8%, Mo: up to 1.5% Cu: up to 0.5%, Ni: up to 1.5% V: up to 0.2%, REM: up to 0.05% Ti: up to 0.02%, Ca: up to 0.005%, Rest iron and unavoidable impurities,

• where the impurities include contents of up to 0.03% P, up to 0.03% S, up to 0.01% N, less than 0.05% Sn, less than 0.05% As and less than 0.05% Co and
• where the ratio formed from the respective Al content %Al and the respective N content %N is (%Al / %N) * 14/27 ≥ 8.0,

(b) the molten steel is cast into a precursor product, namely a slab, a thin slab or a cast strip,
c) the preliminary product is, if necessary, heated to a preheating temperature of 1100 - 1350 °C,
(d) the preliminary product is hot rolled to a hot strip with a thickness of 1 - 16 mm, whereby the hot rolling is terminated at a final hot rolling temperature which is at least 50 °C and at most 150 °C higher than the Ar3 temperature of the steel,
e) the hot strip obtained is cooled to a coiling temperature of 450 - 700 °C, whereby the cooling takes place in the temperature range of 800 - 650 °C at a cooling rate of 20 - 200 °C/s,
f) the hot strip cooled to the coiling temperature is coiled into a coil and the hot strip is cooled to room temperature in the coiled state, and
g) optionally: the hot strip cooled in the coiled state is pickled and
h) also optional: the hot strip is annealed at a core temperature of 500 - 720 °C for a period of 5 - 50 hours.
B) From the hot strip obtained, a cold strip is optionally produced by

• i) the hot strip is cold rolled to a cold strip with a thickness of 0.5 - 3.5 mm in one or more cold rolling steps.
• j) Optionally, the cold strip can be annealed in a batch annealing line or in a continuous annealing line.

C) A component is formed from the hot strip or the optionally produced cold strip by

• k) a blank is cut from the hot or cold strip
• and according to Alternative 1:
  • I.1) the blank is heated to an austenitising temperature which is not more than 100 °C lower than the Ac3 temperature of the steel from which the hot or cold rolled strip is produced and does not exceed 950 °C,
  • l.2) within 1 - 20 s after the end of the heating to the austenitizing temperature, the blank is placed in a cooled hot forming tool in which the blank is hot formed into the component, and
  • l.3) the component is press hardened by accelerated cooling at a cooling rate of 30 - 120 °C/s until the martensite start temperature of the steel from which the respective hot or cold strip is made is reached, so that the component acquires a completely martensitic structure,
• or according to Alternative 2:
  • m.1) the board is cold formed into the component,
  • m.2) the cold-formed component is heated to an austenitising temperature which is not more than 100 °C lower than the Ac3 temperature of the steel from which the hot-rolled or cold-rolled strip is produced and does not exceed 950 °C, and
  • m.3) the component heated to the austenitizing temperature is cooled at an accelerated rate of 30 - 120 °C/s until the martensite start temperature of the steel from which the respective hot or cold strip is made is reached, so that the component acquires a completely martensitic structure.
• n) Optionally, the component obtained after steps l.1 - l.3 or m.1 - m.3 can be annealed at temperatures of 150 - 700 °C for an annealing time of 5 - 60 minutes.

In the method according to the invention, a melt is produced in step a) which is composed in accordance with the above explanations for alloying the steel of a flat steel product intended for forming a component according to the invention. Of course, the information given above on advantageous designs of the steel of a flat steel product intended for forming a component according to the invention applies to the alloying of this melt in the same way for the melt produced and processed in the course of the method according to the invention.

The melt produced in step a) is cast in the conventional way into slabs, thin slabs or strip (step b)). The slabs typically have thicknesses of 180 mm to 260 mm. Thin slabs are typically available in thicknesses of 40 to 60 mm, and cast strip in thicknesses of 2 to 5 mm.

In step c), the precursors are heated through for the subsequent hot rolling (step d). This heating is typically carried out in pusher or walking beam furnaces available for this purpose in the state of the art. The conditions set during the preheating of the slabs, thin slabs or strips cast from a steel melt alloyed according to the invention are of particular importance for the characteristics of a flat steel product composed and produced according to the invention. It has been shown that as a result of the inventive addition of at most a small amount of titanium to the steel processed according to the invention, relatively low furnace temperatures are also process-safe in order to produce no or little nucleation effects on coarse particles such as TiN, since these particles are much less common in the steel concepts according to the invention. occur. "Nucleation" is understood here as the formation of precipitations on previously formed precipitations based on heterogeneous nucleation, as opposed to homogeneous precipitation formation without foreign nucleation sites. The alloy of the steel specified according to the invention, from which the flat steel products processed according to the invention and the components formed from them are made, reduces on average the nucleation effects on previously formed precipitations. The nucleation of TiC, NbN, NbC, AIN on TiN would reduce the probability of these precipitations forming at lower formation temperatures, thus impairing their effectiveness with regard to the refinement of the microstructure sought by the invention. Surprisingly, it has been found that if a suitable preheating temperature is maintained, austenite grain size refinement can be achieved even at the intended, comparatively low Nb contents.

The preheating temperatures used according to the invention are 1100 - 1350 °C and preferably 1150 - 1280 °C. Below 1100 °C, coarsening and nucleation effects of the particles in the preheating must be expected. Temperatures above 1350 °C should be avoided in order to limit coarsening of the austenite grain, reduce material loss due to scaling, or from an economic point of view, reduce energy costs.

Of equal importance are the holding times during which the pre-heating of the preliminary products takes place. These are required for the complete dissolution of the precipitates present in the preliminary products to be pre-heated after casting. The total holding time for slabs provided for in the invention is 150 - 400 minutes, whereby the total holding time includes the time required for heating to the respective target pre-heating temperature and for heating the preliminary products through. With total holding times of less than 150 minutes, there is a risk that the relevant microalloy precipitate types will not dissolve completely. Holding times of However, more than 400 minutes should also be avoided in order to limit austenite grain coarsening. Thin slabs are preheated in a soaking furnace for significantly shorter times of 10 - 90 minutes.

Strips produced by strip casting are generally not preheated, but are hot rolled directly to final hot strip thicknesses of 1 - 4 mm in one or more hot strip stands.

The slabs or thin slabs heated in accordance with the invention, taking into account the above-mentioned requirements, can be hot rolled in a conventional manner in an equally conventional hot rolling mill or casting rolling mill to form a so-called "hot strip". The hot rolling can include pre-rolling, in which the slabs are rolled in a so-called "roughing stand", typically in reverse, to an intermediate thickness of approx. 35 to 60 mm. The pre-rolled slab then runs into a multi-stand finishing hot rolling mill, in which it is gradually hot rolled in a continuous run to form a hot strip.

In the case of a thin slab, pre-rolling is not necessary. It can be fed into the finishing hot rolling mill directly after preheating (if necessary).

According to the invention, hot rolling in step d) is terminated at hot rolling end temperatures that are at least 50 °C higher than the Ar3 temperature of the steel, but no more than 150 °C above this temperature. Hot rolling is thus terminated at a hot rolling end temperature at which the hot strip obtained still has a completely austenitic structure. Such a rolling strategy is referred to as "normalizing rolling". The hot rolling end temperature is selected according to the invention such that the tendency of Nb and Al to form deformation-induced precipitations is reduced and a larger proportion of Precipitation potential is available for inhibiting grain growth during austenitization in the tempering process carried out later or during hot forming. Typically, suitable rolling end temperatures for the alloy concept according to the invention are above 830 °C. The rolling end temperature is preferably at least 60 °C and at most 130 °C higher than the Ar3 temperature, with hot rolling end temperatures that are at most 110 °C above the Ar3 temperature having proven to be particularly practical for limiting austenite grain growth. Normalizing rolling is preferred here because the hot rolling forces are comparatively low and the precipitation of deformation-induced, relatively coarse precipitations is avoided. The precipitation potential for reducing the former austenite grain size can thus be maximized in later austenitization stages of tempering and hot forming. This has a positive effect on toughness.

In order to preserve the precipitation potential of Al and Nb in the hot strip obtained after hot rolling for later process steps, it is necessary in step e) to cool the hot strip after hot rolling in the temperature range from 800 °C to 650 °C with a cooling rate of more than 20 °C/s to the coiling temperature. The actual coiling temperature achieved is determined by the cooling in the cooling section. According to the invention, it is well below the A1 temperature of the steel from which the flat steel product according to the invention is produced in order to avoid relatively coarse pearlite precipitation in the hot strip. In the iron-carbon diagram, the temperature "A1" is the temperature at which austenite decomposes into pearlite from high temperatures. In the pure two-component system iron-carbon, A1 is 723 °C, whereby this conversion takes place at carbon contents > 0.02 mass%, which is the case with the steel concepts according to the invention. The A1 temperature is based on empirical formulas that describe the influence of alloying elements on A1 reproduce (see for example Hougardy, HP "Material Science Steel Volume 1: Basics", Verlag Stahleisen GmbH, Düsseldorf, 1984, p. 229 ), at 722 - 727 °C and thus in a narrow range. In the case of the invention, coiling temperatures of ≤ 720 °C are used in particular. At such low coiling temperatures, the solution state and the precipitation form of the carbon are influenced in such a way that a finely distributed C precipitation is achieved for subsequent quenching and tempering or hot forming treatments in order to accelerate the C dissolution for the hardening process. This makes flat steel products produced or obtained according to the invention particularly suitable for quenching and tempering and heating treatments in which rapid heating by means of induction heating or other rapid heating processes is to be used.

The rapid cooling of the hot strip obtained in accordance with the invention in the temperature range of 800 - 650 °C thus suppresses the precipitation of Nb and Al. This can be ensured in particular by ensuring that the cooling rate is at least 20 °C/s. It should be noted that during cooling after hot rolling, reheating of up to 30 °C can occur due to the phase transformation. In practice, the strip can be sprayed with water for the cooling controlled according to the invention after the hot rolling mill in which the hot rolling takes place. Cooling sections known in the prior art are particularly suitable for this, in which laminar and spray cooling devices are combined with one another. These should be able to achieve cooling rates of preferably more than 20 °C/s, in particular at least 50 °C/s, and a maximum of 200 °C/s, especially in the temperature range of 800 - 650 °C.

The coiling temperature to which the hot strip is cooled after hot rolling and at which the hot strip is coiled into a coil in step f) is 450 - 720 °C. The upper limit of 720 °C is advantageous, in order to be able to set a sufficiently low tensile strength for subsequent cold forming at C contents ≥ 0.4%. The coiling temperature is particularly preferably lower than 650 °C in order to further suppress the precipitation of Nb and Al and to achieve a C dissolution state that is as finely distributed as possible. An upper coiling temperature of 650 °C proves to be particularly advantageous because coarse-structured pearlite formation can then be largely avoided. At coiling temperatures of less than 450 °C, a significant increase in strength would occur in the hot strip, for which subsequent cold forming or cold rolling would represent a significant increase in the deformation forces and is therefore avoided. The hot strip is then cooled to room temperature in the coil in the conventional manner.

The flat steel product obtained according to the invention after coiling, in the form of a hot strip, typically has a tensile strength of less than 700 MPa. Only through the subsequent tempering treatment or through the processing carried out during hot forming are the largely fully martensitic structure according to the invention and, with it, the optimized mechanical properties of a component according to the invention achieved.

After coiling, in step g), which is only carried out optionally if there is a need for it, the hot strip can be subjected to pickling for further processing in order to remove any scale adhering to it. Such a processing step is advantageous if the hot strip is formed in a cold forming tool and contamination or damage to the tool can be avoided by abrasion of the scale. There are no special requirements for pickling. It can be carried out in any way known for this purpose.

The resulting hot strip has a microstructure consisting of pearlite with small amounts of ferrite (< 5%). The ferrite can be in a linear or network-like form.

The hot strip can also optionally be subjected to a bell annealing process in step h) in order to reduce the strength of the steel for subsequent cold forming. The core temperatures of the coiled steel flat product set during bell annealing are 500 - 720 °C. A core temperature of at least 500 °C is required so that a sufficient reduction in strength can occur. Annealing temperatures of more than 720 °C would, however, mean that the formation of new pearlite by exceeding the A1 temperature in all parts of the coil can be safely avoided during bell annealing. A bell annealing time at core temperature level of at least 5 h is required in order to also significantly reduce the strength level, i.e. < 700 MPa tensile strength. However, bell annealing should not last longer than 50 h, as the deformation and coagulation of the pearlite through the ongoing diffusion processes would then lead to coarse pearlite particles. Ideally, the annealing conditions for bell annealing are selected so that only partial deformation of the cementite takes place with a degree of deformation ≤ 85%. In practice, the optional bell annealing according to the invention can be carried out at core temperatures of max. 720 °C under a protective gas atmosphere. The protective gas atmosphere can consist of a pure hydrogen atmosphere (H2) or a mixture of N2 and up to 12 vol.% H2 ("HNX"). Mixtures of 95% N2 and 5% H2 are typical here. HNX annealing results in longer total annealing times of up to 50 hours, since the heat transfer is slower than with a pure H2 atmosphere. The core temperature of the bell annealing should be below 720 °C, in particular around 680 °C, but in any case below the A1 temperature of the steel from which the flat steel product is made. This restriction prevents new pearlite from being formed during the annealing process. Instead, cementite particles (carbide particles) are partially formed from the hot strip structure present at the beginning of annealing, in particular through carbon diffusion and carbon redistribution. At the same time, the structure can become coarser as a result of coagulation. In the course of the bell annealing process, which is optionally carried out according to the invention, cementite is thus formed in a partially formed, globulitic form, which is largely homogeneous and randomly distributed in a ferritic matrix, with the degree of formation being ≤ 85% according to the invention. The restriction of the bell annealing temperature and holding time serves to limit the degree of formation. A limited degree of formation reduces the time for complete C dissolution during austenitization. The structure in the hot strip bell annealed state therefore consists predominantly of partially formed cementite, pearlite in a proportion of up to 90% and a proportion of non-polygonal ferrite of up to 10%.

• i) Kaltwalzen des Warmbands zu einem Kaltband mit einer Dicke von 0,5 - 3 mm in einem oder mehreren Kaltwalzschritten;
• j) Rekristallisierendes Glühen des Kaltbandes, welches in einer Haubenglühe oder in einer Durchlaufglühe stattfinden kann. Erfolgt das Glühen in einer Haubenglühe, dann kann dies nach den oben zu Arbeitsschritt h) bereits angegebenen Bedingungen durchgeführt werden. Soll das Glühen in einer Durchlaufglüheinrichtung absolviert werden, so sind hier keine besonderen Anforderungen an die Glühparameter zu stellen. Demnach kann die Erwärmung bei Geschwindigkeiten bis 30 °C/s bis Erreichen der Glühtemperatur erfolgen, die im Bereich Ac1 bis Ac3 + 30°C liegen kann. Die Abkühlrate auf Raumtemperatur kann über Gasjet- oder Rollenkühlungen erfolgen und bei bis zu 20 °C/s liegen. In die Durchlaufglühung kann eine Schmelztauchveredelung nach dem eigentlichen Glühen integriert sein. Ergänzend kann eine Beschichtung in einer elektrolytischen Beschichtungsanlage aufgebracht werden,

If a cold-rolled flat steel product is to be produced from the hot strip produced in the manner described above, the following further work steps can be carried out following work steps a) - h), of which work steps g) and h) are only carried out as required, i.e. optionally:

• i) cold rolling of the hot strip to a cold strip with a thickness of 0.5 - 3 mm in one or more cold rolling steps;
• j) Recrystallizing annealing of the cold strip, which can take place in a bell annealer or in a continuous annealer. If the annealing takes place in a bell annealer, this can be carried out according to the conditions already specified above for step h). If the annealing is to be carried out in a continuous annealing facility, no special requirements are to be made for the annealing parameters. Accordingly, heating can take place at speeds of up to 30 °C/s until the annealing temperature is reached, which can be in the range Ac1 to Ac3 + 30°C. The cooling rate to Room temperature can be achieved via gas jet or roller cooling and can be up to 20 °C/s. A hot-dip coating can be integrated into the continuous annealing process after the actual annealing. In addition, a coating can be applied in an electrolytic coating system,

The production of the cold strip can be completed in the usual way by a skin pass with usual degrees of deformation of 0.5 - 1.5%, whereby no special requirements are imposed here either.

There are two alternative ways of forming components according to the invention from hot or cold strip produced according to the invention, which have an optimized combination of high strength and toughness. According to the first alternative, a blank cut from the respective hot or cold strip is heated and press-hardened in accordance with work steps I.1 - I.3 of the method according to the invention, whereas according to the second alternative, the blank is first cold-formed and then tempered in accordance with work steps m.1 - m.3 of the method according to the invention.

For the tempering/hot forming of components formed from flat steel products according to the invention by cold forming or for hot forming into components, the respective component (tempering) or flat steel product (hot forming or press hardening) is first heated to a suitably high austenitizing temperature ("austenitizing"). In practice, this can be done, for example, in a manner known per se, first in a furnace in which the respective flat steel product (work step 1.1) of the first alternative) or the component (work step m.2) of the second alternative) is heated to the respective austenitizing temperature over a sufficient total time, in particular (including the heating time).

In the first alternative, the blank is heated in step I.1 and in the second alternative, the component is heated in step m.2 to an austenitizing temperature that is no more than 100 °C below the Ac3 temperature of the steel from which the hot or cold strip is produced and from which the blank or component is made (austenitizing temperature ≥ (Ac3 - 100 °C)). Austenitizing temperatures that are no more than 75 °C lower than the Ac3 temperature (austenitizing temperature ≥ (Ac3 - 75 °C)), in particular no more than 50 °C lower than the Ac3 temperature of the steel of the hot- or cold-rolled sheet from which the blank or component is made (austenitizing temperature ≥ (Ac3 - 50 °C)), lead in practice to the desired result in a particularly reliable manner. Austenitizing temperatures that are at least equal to the Ac3 temperature of the steel from which the respective blank or component is made are particularly suitable. The austenitizing temperature is limited to a maximum of 950 °C. The austenitizing temperature maintained in work steps I.1 and m.2 is accordingly in a range from (Ac3 - 100 °C) to 950 °C, in particular (Ac3 - 75 °C) to 950 °C or, particularly advantageously, from (Ac3 - 100 °C) to 950 °C, with austenitizing temperatures of Ac3 - 950 °C being particularly practical.

A total time of typically 1 second to 20 minutes is required for heating the board or component, although in practice total times of at least 10 seconds, in particular at least 1 minute, are suitable to achieve reliable heating. The total heating time includes the time required to heat up to the austenitizing temperature.

For heating circuit boards in step I.1, total times (including heating time) of 1 - 20 minutes are particularly suitable.

In the case of tempering a component (work steps m.2 and m.3 of the second alternative), the component is heated piece by piece to the Austenitizing temperature typically provides total times of 1 - 20 min or 2 - 10 min, in particular 5 - 10 min.

Inductive continuous flow heating devices available on the market can be used to heat the component more quickly. The component to be heated passes through these devices in a continuous flow, so that the section of the component that is in the effective range of an electromagnetic field induced by the heating device is heated within a short time. In this way, the component is gradually heated to the austenitizing temperature along its length. Continuous flow heating devices of this type are particularly suitable for the continuous heating of components such as pipes or profiles that require a high degree of dimensional accuracy.

In the case of hot forming according to the first alternative, the respective flat steel product is placed after austenitizing within a transfer time of 1 - 20 seconds into a hot forming device known for this purpose from the state of the art, in which it is then press-hardened into a component in an equally known manner, the average cooling rate to room temperature being 30 - 120 °C/s.

In the case of tempering according to the second alternative, the component heated to the austenitizing temperature is also quenched to room temperature after austenitizing at an average cooling rate of 30 - 120 °C/s. For this purpose, the component can be immersed in a suitable quenching medium in a manner known per se or can be exposed to the quenching medium using equally known devices such as nozzles or jet devices. If a continuous heating device of the type described above, particularly an inductive one, is used to heat the component, the section of the component heated to the austenitizing temperature can be The blank must also be cooled in the process by means of a suitable quenching device as it exits the heating device in question.

Quenching takes place within 1 - 20 seconds after removal from the device used to heat to the austenitizing temperature (quenching and tempering) or through contact with the tool at the end of the press hardening process (hot forming). In practice, an oil bath can be used for quenching during tempering, in which the respective component is quenched to room temperature within 1 - 30 seconds while moving. Typical transfer times between the furnace in which heating to the austenitizing temperature takes place and the oil bath are 1 - 20 seconds.

Due to their special property profile, flat steel products processed in accordance with the invention are particularly suitable for the production of highly stressed components for vehicle bodies, in particular for supports, structural parts, frames, bumpers, battery boxes and the like. In particular, the components according to the invention are tubular components, in the production of which blanks of hot or cold strip produced according to the invention are formed into a tubular body and then welded longitudinally.

The structure of the flat steel products in question here and the components according to the invention made from them was investigated as follows:
The proportions of hard oxide and nitride particles in the microstructure of a flat steel product are given in area ppm, unless otherwise stated. The exact procedure for determining them is described below. According to ASTM E2142 from 2008, the area proportion of inclusions can be equated to the volume proportion. Likewise, the phase proportions of the structure given in this text refer to the evaluated ground surface and are therefore given in area %.

The inclusions were examined on longitudinal sections across the strip thickness using a scanning electron microscope (SEM) from Zeiss (model GeminiSEM 500), equipped with the EDX system "Oxford Xmax" from the manufacturer "Oxford instr." for energy-dispersive element analysis. The data was evaluated using the software "Aztec 3.3 SPI, Feature Analysis" from "Oxford instr." Inclusions with a size of approx. 0.2 µm or larger were recorded. The element content in the precipitates was determined using calibration samples. The inclusions were classified based on the stoichiometry of the known precipitates, with a division into oxides, sulfides and TiN. The measured elements were quantified and standardized without Fe, C and Ag. The recorded elements were converted into oxides (without S, P, Cl, F) and normalized to 100%. In addition, the subsystem <Al ₂ O ₃ -SiO ₂ -CaO> was calculated and normalized to 100%. Computer-aided classification tables of the analyzed inclusions were then created from the raw data obtained in this way. Inclusions that could not be clearly classified were listed separately. These inclusions were checked individually. The particle size was idealized as a circle-equivalent diameter, regardless of the particle shape.

The homogeneity of the microstructure of the former austenite and the distribution of the components contained in it was determined using electron backscatter diffraction (EBSD) in the fully martensitic state after tempering or press hardening on longitudinal sections across the strip thickness. The samples were polished using the polishing agent "OP-S Suspension" from the manufacturer "Struers". For this purpose, a measuring field with the dimensions 140 µm x 140 µm was positioned in different layers across the strip thickness and scanned with a step size of 0.15 µm. In addition, several layers across the strip thickness were examined (1/6, 1/3, 1/2) in order to obtain a statement about the homogeneity of the The data obtained from the EBSD investigations on the martensitic microstructure were then used to reconstruct the austenitic initial microstructure using the software "ARPGE 2.0, Reconstruction of Parent Grains from EBSD Data" (described in C. Cayron, Acta Cryst. A62 (2006) 21-40 ; C. Cayron, J. Appl. Cryst. 40 (2007), p. 1183-1188 ) using the Nishiyama-Wassermann orientation relationship for the transformation of body-centered cubic crystals into face-centered cubic crystals ( Z. Nishiyama, Sci. Rep. Res. Inst. Tohuku Univ. Vol. 23 (1934-1935), p. 647 ).

The quantitative estimation of the microstructure components ferrite, pearlite, cementite, bainite and martensite, which goes beyond the inclusion structure, was carried out using light microscopy on the basis of longitudinal sections in the 1/3 zone in the strip thickness direction at 500 - 1000 times magnification.

The mechanical characteristics of flat steel products or components made from them mentioned in this text are the tensile test values (tensile strength, yield strength, modulus of elasticity, uniform elongation and elongation at break), which were determined according to DIN EN ISO 6892-1.

As a measure of toughness, the area of reduction at fracture, also called "absolute strain in the thickness direction" ε(epsilon)3 = (t0 - tf)/t0, expressed in % was used ( Larour, P., Freudenthaler, J., Weissböck, T.: Reduction of cross section area at fracture in tensile test: measurement and applications for flat sheet steels, IDDRG 2017, 36th International Deep Drawing Research Group Conference, Materials Modeling and Testing for Sheet Metal Forming, Munich, DE, Jul 2-6, 2017, Volume 896 (2017), p. 012073/1-8 ). Here, to is the initial thickness of the sample, tf is the thickness of the thinnest points in the area of the necking strain of the fracture cross-section determined from four measurements across the sample width. The "absolute strain in the thickness direction" or "necking at fracture" was measured on tensile specimens according to the Tempering treatment was measured using an optical system (microscope). To do this, the thickness tf was determined in the fracture cross-section at four points across the width (1 mm to the right of the left edge, middle, minimum, 1 mm to the left of the right edge). Three parallel tensile specimens were tested in each case in order to obtain a representative statement for the condition examined. In total, six fracture cross-sections were measured. The average value for a specimen was calculated from the six measured values. The tensile specimens were oriented lengthways to the rolling direction.

The invention is explained in more detail below using exemplary embodiments:
To demonstrate the effectiveness of the invention, three steels 1 - 3a according to the invention were melted, the compositions of which are given in Table 1. For comparison, three further steels 4 - 6 were melted which were not alloyed according to the invention and the compositions of which are also given in Table 1.

The production of flat steel products from steels 1 - 6 was carried out in a conventional manner in an integrated steelworks in which the process chain "pig iron and crude steel production", "steel production" and the various stages of semi-finished product manufacture, such as "preheating" and "hot rolling" and optionally "pickling" and "batch annealing" for the hot strip stage, as well as "pickling", "cold rolling", "continuous annealing" and optionally "AISi coating" and "skin passing" for the cold strip stage according to the invention are represented. The inventive specifications and measures apply to the production of uncoated hot or cold strips for tempering or hot forming processes, as well as to the production of an AlSi-coated, alloyed and cold-rolled sheet for hot forming.

Steels 1 - 6 were each melted and cast into slabs. The slabs were then heated to a preheating temperature and then hot rolled into a hot strip. The hot strips obtained during hot rolling were cooled to a coiling temperature at which they were coiled into a coil. The coil was then cooled to room temperature.

The flat steel product thus produced as unpickled hot strip from steel 2 was made available for further processing without any further treatment.

After cooling in the coil, the hot strips produced from steels 1 and 3 - 6 were subjected to a pickling treatment in order to remove any scale adhering to them.

The hot strips produced from steels 1 and 4 were then cold rolled into cold strips without intermediate annealing. The cold strips obtained in this way were each subjected to continuous annealing, given an AISi coating by hot dip coating and finally skin pass rolled. The flat steel products available as cold strips with an AlSi coating were made available for further processing into components.

The hot strips produced from steels 3, 3a and 6 were subjected to batch annealing and were made available in this condition as flat steel products for further processing.

The flat steel product produced from steel 5 as hot-rolled strip was made available for further processing after pickling.

Table 1 lists the chemical compositions of steels 1 - 6. The contents of the metals present during production but not The elements P, S and N which are attributable to impurities are indicated here because they are of particular importance for the quality of the steels produced according to the invention and, in particular in the case of steels 1 - 3a according to the invention, it must be ensured that the contents of these elements correspond to the requirements of the invention.

The final thickness D of the flat steel products produced from steels 1 - 6 is given in Table 2. This means that for the steel strips produced from steels 1 and 4, the thickness D is given in the finished cold-rolled state with the AISi coating, and for the hot-rolled steel strips produced from steels 2, 3, 3a, 5 and 6, the thickness is given after coiling (hot strip produced from steel 2) or after descaling (hot strip produced from steels 3, 3a, 5 and 6).

Likewise, Table 2 shows for steels 1 - 6 the result of the equation %Ti-48/14*%N, the ratio %Ti/%N, the content %Nrest of the nitrogen not bound by Ti, the result of the equation %Al-27/14*%Nrest, the ratio %Al/%N and the result of the equation (%Al / %N) * 14/27, where %Ti is the Ti content, %N is the N content and %Al is the Al content of the respective steel.

Each of the steels was alloyed with B, with the B contents being at least 0.001 mass%.

The steels 1 - 3a according to the invention each have Ti contents that are not sufficient, or at best only just sufficient, to bind the N content present in the respective steel. The stoichiometric ratio %Ti/%N that must be maintained for theoretically complete binding of the nitrogen present by Ti is 48/14 = 3.43. In the case of the steel 1 according to the invention, the ratio %Ti/%N is significantly below this value. Likewise, in the case of the steels 2, 3, 3a according to the invention, the ratio %Ti/%N is still below the stoichiometric ratio. of 3.43. In any case, the %Ti/%N ratio of the steels according to the invention was less than 4. In contrast, all comparison steels 4 - 6 had a %Ti/%N ratio > 5.

To compensate for the low Ti contents, the Al content of the steels 1 - 3a according to the invention has been increased in order to achieve AlN precipitation through the higher Al contents, i.e. through a higher precipitation pressure, and to avoid BN formation. The %Al/%N ratios of the steels 1 - 3 according to the invention are significantly higher than those of the comparison steels 4 - 6 and are each more than 15. They are therefore also significantly above the stoichiometric %Al/%N ratio, which is 27/14 = 1.93. In the comparison variants 4 - 6, the %Al/%N ratio reaches a maximum of 12.3.

The melts composed according to the invention were cast in a conventional continuous casting plant to form slabs which, after being heated to a preheating temperature "VWT" over a holding time "LIZ", were first pre-rolled in a conventional hot strip mill to form a pre-strip with a thickness "VBD". The pre-strips leaving the roughing stand at a temperature "VBT" were then hot-rolled in a continuous, conventional hot-rolling process to form hot strips with a hot strip thickness "WBD". The hot-rolled hot strips leaving the hot rolling mill were cooled to a coiling temperature HT of less than 650 °C, with a cooling rate "ABK" of at least 50 °C/s being set in the temperature range of 800 - 650 °C.

The exemplary process parameters set for the production and further processing of slabs made of steels 1 - 6, "preheating temperature VWT", "holding time LIZ", "pre-strip temperature VBT", "pre-strip thickness VBD", "hot rolling end temperature WET", "cooling rate ABK", "coiling temperature HT" and "hot strip thickness WBD" are shown in Table 3. plus the temperature "Ar3" calculated using the formula given above.

The hot strips produced from the steel 1 according to the invention and the comparative steel 4 were rolled to their final thickness "D" in cold rolling mills. The degree of cold rolling achieved by cold rolling is not a decisive factor here. It is determined solely by the given hot strip thickness and the respective required cold strip thickness, so that cold rolling can be carried out according to the usual procedure in the state of the art. Cold rolling causes the strip to undergo plastic deformation, which results in strong hardening in terms of the material and results in a reduction in further deformability. Therefore, after cold rolling, recrystallizing annealing is also carried out in a conventional manner, which softens the respective strip and makes it suitable for forming into a component again. In the event that hot-dip coating is to be carried out, as in the example of the cold strip produced from steel 1, the annealing can be integrated in a similarly known manner into the hot-dip coating process that is usually carried out in a continuous process. Alternatively, bell annealing can also be used. Likewise, electrolytic coating can be used instead of hot-dip coating.

In order to determine the behavior of the flat steel products manufactured in the manner explained above from the steels 1 - 3a according to the invention and from the comparative steels 4 - 6 during tempering or during a hot forming process, samples of the flat steel products in question were subjected to a simulation of a conventional tempering or hot forming process. The samples were each heated to an austenitizing temperature "T_aust" which was approximately 60 °C higher than the Ac3 temperature of the respective steel 1 - 6. The austenitizing time required for heating and heating through at the austenitizing temperature T_aust was 7-10 min. including heating time in a salt bath furnace. Following austenitization, the samples were quenched in oil at an average cooling rate of 70 - 120 °C/s to room temperature. These process parameters correspond to the usual conditions that prevail in practice when hardening components that have been cold-formed from flat steel products of the type made from steels 1 - 6, or when press-hardening such flat steel products to form components. The austenitization parameters are listed in Table 4.

After quenching, the samples were tempered at 170 - 200 °C for a period of 20 minutes. This tempering corresponds to both a heat treatment typically carried out at the end of tempering and the conditions that prevail during cathodic dip painting in the typical automotive painting process. Temperatures of 150 - 700 °C in combination with holding times of 5 - 60 minutes are also common in industrial practice for tempering a component.

From the individual values of the reduction in area at fracture ("absolute elongation in the thickness direction") ε(epsilon)3 determined in the manner explained at the beginning, a linear regression calculation for ε(epsilon)3 as a function of the tensile strength was created for the inventive concept and the comparison concept. With a coefficient of determination of > 90%, statistically significant influences were thus obtained. The value ΔBE = (s(epsilon)3_Erf. - ε(epsilon)3_Vergl.)/ s(epsilon)3_Erf. was defined as a function of the tensile strength of both regression calculations as the improvement in toughness. Erf. here means inventive, vergl. means comparative. This value, like the value ε(epsilon)3 of the reduction in area at fracture determined in each case, is listed as the value ΔBE in Table 4.

After these final heat treatment processes, the samples produced from steels 1 - 6 in the manner described above were subjected to The mechanical tensile test parameters "E-modulus", "yield strength Rp0.2", "tensile strength Rm", "uniform elongation Ag" and "elongation at break A" were determined in accordance with DIN EN ISO 6892-1. For cold strip thicknesses ≤ 3 mm, the elongation at break A refers to sample form 2 with cross sections 20 mm wide and an initial measuring length of 80 mm. For thicknesses > 3.0 mm, an initial measuring length of 50 - 65 mm (proportional tensile specimens) was used. The determination was carried out at three points on the samples examined in the longitudinal direction at room temperature. The results of these tests, averaged from the three measurements in each case, are also summarized in Table 4.

The inclusion state in the structure of the samples produced from steels 1 - 6 in the manner described above was also examined, firstly in the condition as delivered, i.e. before tempering, and secondly after tempering. No significant differences in the non-metallic precipitates were found, so an average was calculated from both measurements. The test methods explained above were used for this purpose. These tests confirmed that the samples made from the steels 1 - 3a alloyed according to the invention had a significantly reduced proportion of less than 150 area ppm of hard TiN, AIN and Al-based, oxidic precipitates in their structures, the average particle size of which was 0.2 - 10 µm. The precipitates were also distributed homogeneously across the strip thickness. The measuring area covered by the respective investigation, the corresponding microstructure parameters "number of TiN per cm ² ", average diameter of the TiN precipitates "TiN-Ø", % share of TiN precipitates (including TiN particles as a conglomerate with softer particles) in the total sum of hard particles "TiN share+conglomerates" as well as the total sum of hard particles of the TiN, AIN and Al ₂ O ₃ precipitates and their conglomerates with softer particles "TiN, Al ₂ O ₃ , AIN+conglomerates" are listed in Table 5.

It was also shown that the samples produced and processed in accordance with the invention from the steels 1 - 3a according to the invention have a reduced former austenite grain size compared to the non-inventive variants produced from the steels 4 - 6, in conjunction with a likewise reduced scatter of the austenite grain size across the strip thickness, also averaged at three points 1/6, 1/3 and 1/2 across the strip thickness. The parameters determined in these tests, mean diameter of the former austenite grains "KA=Ø former A-KG", standard deviation (related to the sample size) of the diameter of the former austenite grains "Ks=σ former A-KG" and the derived size of the grain size quality "KG quality" are summarized in Table 6. The KG quality is obtained by multiplying the mean former austenite grain size by the standard deviation of the diameter of the former austenite grain size. The smaller the KG grade, the more favorable the effects on toughness and local elongation are. It is known that toughness improves with decreasing grain size. In addition, a smaller dispersion of grain size ensures increased homogeneity of the deformation behavior and thus a delayed onset of instability due to fracture reduction, since there are fewer local differences. Table 1 Steel C Si Mn Al Cr Nb Ti B Ca P*) S*) N*) V Cu Mo Ni According to the invention? 1 0.13 0.20 1.17 0.130 0.21 0.021 0.004 0.0021 0.0012 0.011 0.005 0.0059 0.001 0.04 0.026 0.09 YES 2 0.15 0.21 1.06 0.092 0.21 0.024 0.015 0.0023 0.0008 0.014 0.002 0.0046 0.006 0.04 0.009 0.05 YES 3 0.45 0.19 0.67 0.090 0.30 0.026 0.016 0.0027 0.0005 0.010 0.001 0.0048 0.002 0.01 0.005 0.02 YES 3a 0.47 0.27 1.27 0.081 0.32 0.025 0.013 0.0023 0.0005 0.008 0.001 0.0038 0.001 0.02 0.005 0.03 YES 4 0.21 0.27 1.17 0.048 0.21 0.001 0.034 0.0025 0.0004 0.015 0.001 0.0039 0.003 0.08 0.027 0.05 NO 5 0.41 0.22 1.16 0.036 0.41 0.001 0.029 0.0023 0.0005 0.017 0.003 0.0051 0.003 0.01 0.002 0.02 NO 6 0.38 0.19 1.25 0.036 0.31 0.001 0.029 0.0012 0.0011 0.020 0.002 0.0057 0.002 0.01 0.003 0.02 NO *) contamination
Information in mass%, remainder iron and other unavoidable impurities Steel D %Ti-48/14*%N %Ti/%N %Nrest %Al-27/14*%Nrest %Al/%N %Al/%N*14/27 [mm] [%] [%] [%] 1 1.45 -0.0162 0.678 0.0047 0.1209 22,034 11,425 2 3.03 -0.0008 3,261 0.0002 0.0916 20,000 10,370 3 4.04 -0.0005 3,333 0.0001 0.0897 18,750 9,722 3a 4.02 -0.0004 3,333 0.0001 0.0808 21,316 11,053 4 2.20 0.0206 8,718 0.0000 0.0480 12,308 6,382 5 4.05 0.0115 5,686 0.0000 0.0360 7,059 3,660 6 3.10 0.0095 5,088 0.0000 0.0360 6,316 3,275 %Ti = Ti content, %N = N content, %Nrest = N content not bound by Ti, %Al = Al content of the respective steel Steel VWT LIZ VBT VBD WET Ar3 ABK* HT WBD [°C] [min.] [°C] [mm] [°C] [°C] [°C/s] [°C] [mm] 1 1287 341.6 1084 37.4 869 808 61 579 3.30 2 1265 222.2 1093 39.6 892 807 76 592 3.03 3 1205 216.3 1072 40.3 861 758 53 612 4.04 3a 1261 224.0 1036 38.1 859 741 35 612 4.02 4 1273 236.9 1090 40.5 840 791 41 586 3.11 5 1260 247.6 1080 40.2 849 751 35 634 4.05 6 1259 359.4 1088 39.7 910 752 47 717 3.10 *: average cooling rate in the temperature range between 800 °C and 650 °C Steel Ac3 T_aust E-modulus Rp0.2 Rm Ag A ε(Epsilon)3 ΔBE [°C] [°C] [MPa] [MPa] [MPa] [%] [%] [%] [%] 1 888 940 193721 930 1183 3.6 5.3 59.9 7.5 2 881 940 195068 937 1213 3.5 6.1 58.5 7.8 3 815 875 173333 1515 2243 4.9 9.1 16.7 53.0 3a 803 870 192402 1555 2159 4.5 8.1 13.5 42.2 4 857 915 196419 1105 1455 3.8 6.8 43.5 -11.5 5 812 875 203333 1462 2188 3.9 5.0 8.0 -45.5 6 818 880 192333 1312 1862 4.0 8.7 27.5 -22.6 Steel measuring surface number of TiN per cm ² TiN-Ø TiN content + conglomerates TiN, Al ₂ O ₃ , AIN + conglomerates - [mm ² ] [N/cm ² ] [µm] [% share of area ppm] [area ppm] 1 22.8 62 0.4 0 115 2 39.0 5221 1.0 24 121 3 39.1 1408 1.1 11 116 3a 39.2 3235 1.1 16 135 4 35.2 13450 1.2 79 326 5 35.9 13298 1.0 48 357 6 32.7 12502 1.2 63 373 Steel Position KA= Ø former A-KG Ks= σ former A-KG KG quality [µm] [µm] [µm ² ] 1 1/2 5.6 2.4 13.2 1 1/6 6.8 3.3 22.4 1 1/3 7.2 3.6 25.7 2 1/2 8.3 3.4 28.6 3 1/6 6.4 3.1 19.6 3 1/3 6.7 4.0 26.7 3 1/2 7.5 3.8 28.4 3a 1/6 7.1 3.3 23.4 3a 1/3 7.3 3.2 23.4 3a 1/2 6.9 3.5 24.2 4 1/2 9.1 6.0 54.5 5 1/6 7.7 4.3 32.7 5 1/3 7.5 5.7 42.5 5 1/2 8.5 4.5 38.1 6 1/2 8.2 4.7 38.6

Claims (10)

1. Component produced by being formed from a sheet steel blank,

   - consisting of a steel which, in wt.%, consists of C: 0.1 - 0.6%, Mn: 0.1 - 2%, Al: 0.05 - 0.2%, Nb: 0.01 - 0.06%, B: 0.0005 - 0.005%, Cr: 0.05 - 0.8%, Si: up to 0.8%, Mo: up to 1.5%, Cu: up to 0.5%, Ni: up to 1.5%, V: up to 0.2%, REM: up to 0.05% Ti: up to 0.02%, Ca: up to 0.005%, the remainder being iron and unavoidable impurities,

   - wherein the impurities include contents of up to 0.03% P, up to 0.03% S, up to 0.01% N, less than 0.05% Sn, less than 0.05% As and less than 0.05% Co,

   - wherein the ratio formed from the corresponding Al content %Al and the corresponding N content %N is (%Al / %N) * 14/27 ≥ 8.0, and

   - wherein the component has a structure which consists of at least 95% by area of martensite and other structural components as the remainder and in which a maximum of 150 ppm by area of particles are present in a homogeneous distribution over the strip thickness, of which the average circular equivalent particle size is 0.2 - 10 µm and which consist of oxide-based Al compounds, of AIN, TiN or of conglomerates formed on the basis of these particles, measured as described in the description.
2. Component according to claim 1, characterized in that its Al content is 0.07 - 0.13 wt.%.
3. Component according to any of the preceding claims, characterized in that its B content is 0.001 - 0.0035 wt.%.
4. Component according to any of the preceding claims, characterized in that for its structure, KA x Ks of the former austenite grain size KA, used in µm, and the simple standard deviation Ks of the former austenite grain size, also used in µm and averaged over the austenite grain sizes determined at one sixth, one third and half of the thickness of the relevant wall portion of the component under consideration, applies for the product, measured as described in the description: $$\mathrm{KA}\times \mathrm{Ks}<30{\mathrm{\mu m}}^2.$$

   
5. Component according to any of the preceding claims, characterized in that in the tempered or press-hardened state, it has a tensile strength Rm of 1000 - 2500 MPa.
6. Component according to any of the preceding claims, characterized in that in the tempered or press-hardened state, it has a fracture reduction in the thickness direction ε3 of 10 - 65%.
7. Component according to any of the preceding claims, characterized in that it is a body structural part, a stabilizer for a vehicle suspension, a steering shaft or a drive shaft of motor vehicles.
8. Method for producing a component according to claim 1 having a structure consisting of at least 95% by area of martensite and other structural components as the remainder, wherein

   A) a hot strip is produced, by

   a) melting steel, which, in wt.%, consists of C: 0.1 - 0.6%, Mn: 0.1 - 2%, Al: 0.05 - 0.2%, Nb: 0.01 - 0.06%, B: 0.0005 - 0.005%, Cr: 0.05 - 0.8%, Si: up to 0.8%, Mo: up to 1.5%, Cu: up to 0.5%, Ni: up to 1.5%, V: up to 0.2%, REM: up to 0.05% Ti: up to 0.02%, Ca: up to 0.005%, the remainder being iron and unavoidable impurities,

   - wherein the impurities include contents of up to 0.03% P, up to 0.03% S, up to 0.01% N, less than 0.05% Sn, less than 0.05% As and less than 0.05% Co and

   - wherein the ratio formed from the corresponding Al content %Al and the corresponding N content %N is (%Al / %N)*14/27 ≥ 8.0,

   b) casting the molten steel into a preliminary product, such as a slab, a thin slab or a cast strip,

   c) heating the preliminary product, if necessary, through to a preheating temperature of 1100 - 1350°C,

   d) hot-rolling the preliminary product to form a hot strip with a thickness of 1 - 16 mm, wherein the hot-rolling is terminated at a final hot-rolling temperature which is at least 50°C and at most 150°C higher than the Ar3 temperature of the steel,

   e) cooling the resulting hot strip to a coiling temperature of 450 - 700°C, wherein the cooling takes place in the temperature range of 800 - 650°C at a cooling rate of 20 - 200°C/s,

   f) coiling the hot strip cooled to the coiling temperature to form a coil and cooling to room temperature in the coiled state, and

   g) optionally: pickling the hot strip cooled in the coiled state and

   h) optionally: carrying out batch annealing at a core temperature of the hot strip of 500 - 720°C for a period of 5 - 50 hours;

   B) wherein a cold strip is optionally produced from the hot strip obtained by

   i) cold-rolling the hot strip into a cold strip with a thickness of 0.5 - 3.5 mm in one or more cold-rolling steps;

   j) optionally: annealing the cold strip in a batch annealing furnace or in a continuous annealing furnace;

   C) a component is formed from the hot strip or the optionally produced cold strip by

   k) cutting a blank from the hot or cold strip

   and according to alternative 1:

   I.1) the blank is heated through to an austenitizing temperature which is at most 100°C lower than the Ac3 temperature of the steel from which the hot-rolled or cold-rolled strip is produced and at most 950°C,

   I.2) within 1 - 20 s after the end of the through-heating to the austenitizing temperature, the blank is placed in a cooled hot-forming tool in which the blank is hot-formed to form the component, and

   I.3) the component is press-hardened by accelerated cooling at a cooling rate of 30 - 120°C/s until the martensite start temperature of the steel from which the respective hot or cold strip is made is reached, so that the component acquires a completely martensitic structure,

   or according to alternative 2:

   m.1) the blank is cold-formed to form the component,

   m.2) the cold-formed component is heated to an austenitizing temperature which is at most 100°C lower than the Ac3 temperature of the steel from which the hot-rolled or cold-rolled strip is produced and is at most 950°C, and

   m.3) the component heated through to the austenitizing temperature is cooled at an accelerated rate of 30 - 120°C/s until the martensite start temperature of the steel from which the respective hot or cold strip is made is reached, so that the component acquires a completely martensitic structure;

   and

   n) optionally: the component obtained after steps I.1 - I.3 or m.1 - m.3 is annealed at temperatures of 150 - 700°C for an annealing time of 5 - 60 min.
9. Method according to claim 8, characterized in that the coiling temperature is 550 - 630°C.
10. Method according to either claim 8 or claim 9, characterized in that the total time provided for through-heating the blank in step I.1 or for through-heating the component in step m.2 is between 1 second and 20 minutes, wherein the total time includes the heating time required for heating to the relevant austenitizing temperature.