EP3807429A1 - Flat steel product and method for the production thereof - Google Patents

Abstract

The present invention relates to a flat steel product which has good deep-drawing ability, low edge-crack sensitivity and good bending behaviour. To this end, the flat steel product contains a steel which consists of (in wt%) 0.1-0.5% C, 1.0-3.0% Mn, 0.9-1.5% Si, up to 1.5% AI, up to 0.008% N, up to 0.020% P, up to 0.005% S, 0.01-1% Cr and optionally one or more of the following elements: up to 0.2% Mo, up to 0.01% B, up to 0.5% Cu, up to 0.5% Ni and optionally a total of 0.005-0.2% microalloying elements, the remainder being iron and unavoidable impurities, wherein 75 < (Mn2 + 55*Cr)/Cr < 3000 where Mn is the Mn content of the steel in wt% and Cr is the Cr content of the steel in wt%. The steel has a structure which consists of at least 80 area % martensite, of which at least 75 area % is tempered martensite and at most 25 area % is non-tempered martensite, at least 5 volume % residual austenite, 0.5 to 10 area % ferrite and at most 5 area % bainite, wherein in the region of the phase boundary between tempered martensite and residual austenite there is a low-Mn ferrite seam which has a width of at least 4 nm and at most 12 nm and the Mn content of which is at most 50% of the average Mn content of the flat steel product. The flat steel product contains carbides with a length of less than or equal to 250 nm.The invention also relates to a method for producing a flat steel product according to the invention, in which method the structural characteristics of the flat steel product according to the invention are set by suitable heat treatment.

Description

 Flat steel product and process for its manufacture

The present application relates to a cold-rolled flat steel product, in particular a cold-rolled flat steel product for automobile construction, which has good deep-drawing properties, low sensitivity to edge cracking and good bending behavior, and a method for producing such a flat steel product.

High-strength and ultra-high-strength steels are preferably used to reduce the vehicle's weight in automotive engineering, which should also have good formability in addition to high strength. On sheets that have been subjected to shear cutting processes, the shape-changing capacity in the edge area is greatly reduced, so that the risk of edge tears occurring during further processing is increased. One method for characterizing the edge crack sensitivity is the hole expansion test according to ISO 16630. In contrast, the bending test determines the bending strength and the maximum deflection until a first crack. The angle obtained after the springback of the bent sample is referred to as the bending angle and is a measure of the suitability of the tested material for forming. Particularly for complex component geometries, high demands are placed on the deep-drawing ability of the steels A method for assessing the deep-drawing ability is offered by the cup drawing test according to DIN 8584-3, which provides information on the deep-drawing ability of the material by determining the maximum deep-drawing ratio (limit drawing ratio ß _ma x) the elongation at break and the maximum deep-drawing ratio decrease with increasing strength.

When we are talking about flat steel products, this means steel strips, steel sheets or blanks produced from them, such as blanks.

WO 2012/156428 A1 discloses a method for producing flat steel products, in which the flat steel products are subjected to a heat treatment, in which the flat steel products are cooled to a cooling stop temperature after austenitizing, held and then reheated in one stage to a temperature TP at a heating rate Theta_Pl become. The flat steel products have a yield strength of 600 to 1400 MPa, a tensile strength of at least 1200 MPa, an elongation A50 of 10 to 30%, a hole expansion of 50 to 120% and a bending angle of 100 to 180 °. The flat steel products consist of 0.10-0.50% by weight of C, 0.1-2.5% by weight of Si, 1.0-3.5% by weight of Mn, up to 2.5% by weight. % AI, up to 0.020% by weight P, up to 0.003% by weight S, up to 0.02% by weight N, and optionally 0.1-0.5% by weight Cr, 0, 1 - 0.3% by weight Mo, 0.0005 - 0.005% by weight % B, up to 0.01 wt% Ca, 0.01-0.1 wt% V, 0.001-0.15 wt% Ti, 0.02-0.05 wt% Nb , the sum of the V, Ti and Nb contents being less than or equal to 0.2% by weight. The structure of the flat steel products has less than 5% ferrite, less than 10% bainite, 5-70% non-tempered martensite, 5-30% residual austenite and 25-80% tempered martensite. From WO 2012/156428 A1, however, it is not known how high strength and good deep-drawing ability can be achieved at the same time.

If information on alloy contents and compositions is given here, these relate to the weight or the mass, unless otherwise expressly stated. Unless otherwise stated, the information on the structural components for the structural components martensite, ferrite and bainite in the present case is based on area%. and based on vol.% for residual austenite.

Against the background of the prior art, the object of the invention was to provide a high-strength flat steel product with optimized mechanical properties, in particular with very good forming properties, in particular a good deep-drawing capability with high strength at the same time.

Another object of the invention was to provide a method for producing such a flat steel product. This process should be particularly suitable for being able to be integrated into a process for hot-dip coating.

With regard to the flat steel product, the object was achieved by a product which has at least the features specified in claim 1. With regard to the method, the object was achieved in that at least the method steps specified in claim 9 are carried out in the production of a flat steel product according to the invention.

A flat steel product according to the invention contains a steel which consists of (in% by weight)

0.1 - 0.5% C,

 1.0 - 3.0% Mn,

 0.9 - 1.5% Si,

 up to 1.5% AI,

 up to 0.008% N,

up to 0.020% P, up to 0.005% S

 0.01 - 1% Cr, and optionally from one or more of the following elements up to 0.2% Mo,

 up to 0.01% B,

 up to 0.5% Cu,

 up to 0.5% Ni and optionally consisting of a total of 0.005 - 0.2% microalloying elements and the rest of iron and unavoidable impurities, whereby:

75 S (Mn ² + 55 * Cr) / Cr ä 3000

 with Mn: Mn content of the steel in% by weight, Cr: Cr content of the steel in% by weight.

A flat steel product according to the invention has a structure that

 at least 80 area% martensite, of which at least 75 area% is tempered martensite and at most 25 area% is not tempered martensite,

 at least 5% by volume of residual austenite,

 0.5 to 10 area% ferrite and

 at most 5 area% bainite

consists.

It is essential for the good mechanical properties that in the area of the phase boundaries between tempered martensite and residual austenite, a low-Mn ferrite seam is preferred. In this ferrite seam, the Mn content is at most 50% of the average total Mn content of the flat steel product. The width of the Mn-poor ferrite seam is at least 4 nm, preferably more than 8 nm, and at most 12 nm, preferably less than 10 nm. In addition, a steel flat product according to the invention contains carbides whose length is equal to or less than 250 nm, preferably less than Is 175 nm.

A flat steel product according to the invention is characterized by a tensile strength Rm of 900 to 1500 MPa, a proof stress Rp02 which is equal to or more than 700 MPa and less than the tensile strength of the flat steel product is, an elongation A80 7-25%, a bending angle which is greater than 80 °, a hole expansion which is greater than 25% and a maximum drawing ratio ß _ma x, for which the following applies: ß _max> -1 9 IO ⁶ x (ß _m ) ² + 3.5 IO ^"3 x R _m + 0.5 with im: tensile strength of the flat steel product in MPa, whereby the tensile strength, the yield strength and the elongation in the tensile test according to DIN EN ISO 6892 -1 (sample form 2) from 02/2017, the bending angle according to VDA238- 100 from 12/1010, the hole expansion according to ISO 16630 from 10/2017 and the maximum deep-drawing ratio, ßmax according to DIN 8584-3 from 09/2003.

The carbon content of the steel of a flat steel product according to the invention is 0.1-0.5% by weight. In the steel of a flat steel product according to the invention, the carbon contributes to the formation and stabilization of the austenite. C contents of at least 0.1% by weight, preferably at least 0.12% by weight, contribute to the stabilization of the austenitic phase, especially during the first cooling after austenizing and during the subsequent partitioning annealing, thereby it is possible to ensure a residual austenite content of at least 5% by volume in the flat steel product according to the invention. In addition, the C content has a strong influence on the strength of the martensite. This applies both to the strength of the martensite that arises during the first quenching and to the strength of the martensite that is formed during the second quenching that begins after the partitioning annealing. In order to use the influence of carbon on the strength of the martensite, the C content is at least 0.1% by weight. With increasing C content, the martensite start temperature Ms is shifted to lower temperatures. A C content above 0.5% by weight could therefore result in insufficient martensite being formed during quenching. In addition, a high C content can lead to the formation of large brittle carbides. The processability, in particular the weldability, is also impaired at higher C contents, which is why the C content should be at most 0.5% by weight, preferably at most 0.4% by weight.

Manganese (Mn) is important as an alloying element for the hardenability of the steel as well as for avoiding the formation of the structural component perlite during cooling. The Mn content of the steel of a steel flat product according to the invention is at least 1.0% by weight, preferably at least 1.9% by weight, so that after the first quenching a pearlite-free structure consisting of martensite and residual austenite is available for the further process steps put. A too low Mn content would also lead to the fact that no Mn-poor ferrite seam can form. The positive influences of Mn can be used particularly safely with contents of preferably at least 1.9% by weight. With increasing Mn content, however, the weldability of a deteriorates Flat steel product according to the invention and the risk of the occurrence of strong segregations increases. Segregations are chemical formed during the solidification process

Inhomogeneities of the composition in the form of macroscopic or microscopic

Segregations. In order to reduce segregation and ensure good weldability, the Mn content of the steel of a flat steel product according to the invention is limited to at most 3.0% by weight, preferably to at most 2.7% by weight.

Silicon (Si) as an alloying element supports the suppression of cementite formation. Cementite is an iron carbide. Due to the formation of cementite, carbon is bound in the form of iron carbide and is no longer available as interstitially dissolved carbon to stabilize the remaining austenite. As a result, the elongation of the flat steel product deteriorates, since residual austenite helps to improve the elongation. A similar effect with regard to the stabilization of the residual austenite can also be achieved by alloying aluminum. In order to use the positive effect of Si, at least 0.9% by weight of Si should be present in the steel of the flat steel product according to the invention. However, since a high Si content can have a negative effect on the surface quality of the flat steel product, the steel should not contain more than 1.5% by weight, preferably less than 1.5% by weight, of Si.

Aluminum (Al) can be added to the steel of a steel product according to the invention for deoxidation and for setting nitrogen, insofar as nitrogen is present in the steel, in a content of up to 1.5% by weight. Aluminum can also be added to suppress the formation of cementite. However, AI increases the austenitizing temperature of the steel. If higher annealing temperatures are to be set for austenitizing, Al up to 1.5% by weight can be alloyed. Since aluminum increases the annealing temperature required for complete austenitization and complete austenitization is difficult with Al contents above 1.5% by weight, the Al content of the steel of a flat steel product according to the invention is at most 1.5% by weight. -%, preferably at most 1.0% by weight, limited. If a low austenitizing temperature is to be set, Al contents of at least 0.01% by weight, in particular from 0.01 to 0.1% by weight, have proven to be expedient.

Phosphorus (P), sulfur (S) and nitrogen (N) have a negative effect on the mechanical-technological properties of flat steel products according to the invention. P has an unfavorable effect on weldability, which is why the P content should be at most 0.02% by weight, preferably less than 0.02% by weight. S leads to the formation of MnS or formation at higher concentrations of (Mn, Fe) S, which have a negative effect on the elongation. The S content is therefore limited to values of at most 0.005% by weight, preferably less than 0.005% by weight. Nitrogen bound to nitrides can have a negative effect on the formability, which is why the N content should be limited to at most 0.008% by weight, preferably to less than 0.008% by weight.

Chromium (Cr) is present in steel from 0.01 to 1.0% by weight. Chromium is an effective inhibitor of pearlite and contributes to strength. Therefore, at least 0.01% by weight of Cr, preferably at least 0.1% by weight of Cr, should be contained in the steel according to the invention. If the Cr content exceeds 1.0% by weight, the weldability of a steel flat product according to the invention deteriorates and the risk of a pronounced grain boundary oxidation occurring, which leads to a deterioration in the surface quality, is increased. The Cr content is therefore limited to at most 1.0% by weight, preferably at most 0.50% by weight, particularly preferably to less than 0.2% by weight.

Furthermore, the invention is based on the knowledge that compliance with a specific ratio of Mn and Cr has a favorable effect on the formation of a ferrite fringe low in Mn along the phase boundary from residual austenite to tempered martensite. A Mn-poor ferrite fringe can be set along the phase boundary from residual austenite to tempered martensite if the following condition is met:

75 <(Mn ² + 55 * Cr) / Cr <3000 with Mn: Mn content of the steel in% by weight and Cr: Cr content of the steel in% by weight. If the chromium content is too high compared to the Mn content, the grain boundaries can be covered with chromium carbides. This is undesirable since the formation of the Mn-poor ferrite seam would then be prevented by a reduced mobility of the phase boundary. However, if the Mn content is chosen too high compared to the chromium content, the austenite saturates early on Mn and the diffusion of the manganese ceases. A low Mn ferrite seam cannot be formed due to the still high local Mn concentration. The absence of the ferrite seam would degrade the forming properties and in particular the maximum deep-drawing ratio ßmax.

Optionally, one or more elements from the group of molybdenum (Mo), boron (B) and copper (Cu) can be present in the steel of a flat steel product according to the invention in order to improve the mechanical-technological properties. Molybdenum (Mo) can also optionally be contained in the steel of a steel flat product according to the invention in contents of up to 0.2% by weight, preferably less than 0.2% by weight, in order to prevent the formation of Perl it.

Boron (B) can be present as an optional alloying element in contents of up to 0.01% by weight in the steel of a flat steel product according to the invention. Boron segregates on the phase boundaries and thus blocks their movement. This supports the formation of a fine-grained structure, which improves the mechanical properties of the flat steel product. When adding boron, enough Ti should be available to set N, which prevents the formation of harmful boron nitrides, namely Ti> 3.42 * N. From a technical point of view, the lower limit for boron is 0.0003%.

Copper (Cu) can be contained as an optional alloying element in contents of up to 0.5% by weight in the flat steel product according to the invention. Cu can increase the yield strength and strength. In order to effectively use the strength-increasing effect of Cu, Cu can preferably be added in contents of at least 0.03% by weight. At these levels, the resistance to atmospheric corrosion is also increased. At the same time, however, there is a significant decrease in elongation at break with increasing Cu contents. In addition, the weldability at Cu contents of greater than 0.5% by weight is significantly reduced and the tendency towards red brittleness increases, which is why the Cu content is up to 0.5% by weight, preferably 0.2% by weight. % is.

Nickel (Ni) can be contained as an optional alloying element in contents of up to 0.5% by weight in the steel of a flat steel product according to the invention. Like chromium, it is an inhibitor of pearlite and is effective even in small amounts. With optional alloying with nickel of preferably at least 0.02% by weight, in particular at least 0.05% by weight, this supporting effect can be achieved. In view of the desired setting of the mechanical properties, it is also expedient to limit the Ni content to 0.5% by weight, with Ni contents of at most 0.2% by weight, in particular 0.1% by weight. -% have shown to be particularly practical.

Steels of steel products according to the invention can optionally contain one or more microalloying elements. In the present case, microalloying elements are understood to mean the elements titanium (Ti), niobium (Nb) and vanadium (V). Titanium and / or niobium are preferably used. The microalloying elements can form carbides with carbon, which in the form of very finely divided precipitates contribute to greater strength. With a total content of microalloying elements of at least 0.005% by weight, excretions can occur arise, which lead to the freezing of grain and phase boundaries during austenitizing. At the same time, however, carbon, which is favorable in atomic form for the stabilization of the residual austenite, is bound as carbide. In order to ensure adequate stabilization of the residual austenite, the total concentration of the microalloying elements should not be more than 0.2% by weight. In a preferred embodiment, the sum of Ti and / or Nb is 0.005-0.2% by weight.

In a preferred embodiment, the flat steel product according to the invention is a cold-rolled flat steel product.

In a further preferred embodiment, the flat steel products can optionally be provided with a metallic coating for the purpose of corrosion protection. Zn-based coatings are particularly suitable for this. The coating can be applied in particular by hot dip coating.

The method according to the invention for producing a high-strength flat steel product comprises at least the following steps: a) Providing a slab which consists of a steel which, in addition to iron and unavoidable impurities, comprises (in% by weight)

 0.1-0.5% C, preferably 0.12-0.4% C, 1.0-3.0% Mn, preferably 1.9-2.7% Mn, 0.9-1.5% Si, up to 1.5% AI, up to 0.008% N, up to 0.020% P, up to 0.005% S, 0.01 to 1% Cr and optionally from one or more of the following elements: up to 0.2 % Mo, up to 0.01% B, up to 0.5% Cu, up to 0.5% Ni and optionally from a total of 0.005 - 0.2%

Microalloying elements, preferably consisting of a total of 0.005 - 0.2% Ti and / or Nb, where: 75 S (Mn ² + 55 * Cr) / Cr <3000, with Mn: Mn content of the steel in wt. %, Cr: Cr content of the steel in% by weight; b) heating the slab to temperatures of 1000-1300 ° C and hot rolling the slab to a hot strip, the final rolling temperature T_ET being greater than 850 ° C; c) cooling the hot strip within a maximum of 25 s to a reel temperature T_HT, which is 400 to 620 ° C and winding the hot strip into a coil; d) pickling the hot rolled flat steel product; e) cold rolling the hot rolled flat steel product; f) heating the cold-rolled flat steel product to a holding zone temperature T_HZ which is at least 15 ° C above the A3 temperature of the steel and is at most 950 ° C, the heating either fl) in one stage with an average heating rate of 2 - 10 K / s or f2) in two stages with a first heating rate Theta.H 1 of 5 - 50 K / s up to a turning temperature T_W of 200 - 400 ° C and above the turning temperature T_W with a second heating rate Theta_H2 of 2 - 10 K / s; g) holding the flat steel product at the holding zone temperature T_HZ for a period t_HZ of 5-15 s; h) cooling the flat steel product from the holding zone temperature T_HZ to a cooling stop temperature T_Q, which lies between the martensite start temperature T_MS and a temperature which is 175 ° C. lower than T_MS, with either hl) a cooling rate Theta_Q 1, which is at least 30 K / s ; or h2) a first cooling rate Theta_LK of less than 30 K / s for a first cooling to an intermediate temperature TJLK, which is not lower than 650 ° C., and a second cooling rate Theta_Q2 for a second cooling from T_LK to T_Q, where Theta_Q2 is at least 30 K / s; i) holding the steel flat product at the cooling stop temperature T_Q for 1-60 seconds; j) heating the flat steel product with a first heating rate Theta_Bl, which is between 5 and 100 K / s, to a first treatment temperature T_B1, which is at least T_Q + 10 ° C and at most 450 ° C, keeping the flat steel product at the first treatment temperature TJBl for a duration t_Bl from 8.5 s to 245 s, heating the flat steel product with a second heating rate Theta_B2, which is between 2 and 50 K / s, to a second treatment temperature T_B2, which is at least T_B1 + 10 ° C and at most 500 ° C optional holding of the flat steel product at the treatment temperature T_B2 for a duration t_B2 of up to 34 s, the total treatment time t_BT for the heating and the isothermal holding is between 10 and 250 s in total; k) optional coating of the flat steel product in a Zn-based coating bath;

L) cooling the flat steel product to room temperature with a cooling rate Theta_B3 of at least 5 K / s.

In step a), a slab produced in a conventional manner is provided, which consists of a steel of the composition mentioned in step a).

In step b) the slab is heated to temperatures of 1000-1300 ° C and rolled out to a hot strip. Hot rolling is carried out with a final rolling temperature T_ET greater than 850 ° C in an otherwise customary manner. The final rolling temperature TJET should be higher than 850 ° C in order to avoid the formation of coarse, polygonal ferrite grains during the rolling process.

In step c), the hot strip is cooled after hot rolling and before coiling and then wound into a coil at the coiling temperature T_HT. In order to reduce or preferably completely suppress the formation of polygonal ferrites, the cooling takes place within a time period t_RG equal to or less than 25 s, that is to say within a maximum of 25 s. Here, t_RG is the time period which begins after the rolling process, that is to say after the last rolling pass, and ends after the cooling process has ended, that is to say when the reel temperature T_HT has been reached. The formation of polygonal ferrites can be minimized particularly effectively if t_RG is at most 18 s, preferably at most 15 s. Typically, for process reasons, t_RG is at least 2 s, usually at least 5 s.

In order to prevent the formation of the undesirable structural component perlite, the coiling takes place at a coiling temperature T_HT of at most 620 ° C. In a preferred embodiment, the reel temperature T_HT is set to at most 600 ° C., which additionally has a positive effect on the avoidance of polygonal ferrites. Coil temperatures of at most 580 ° C. are particularly preferred in order to increase the proportion of bainite in the structure of the hot strip. If the reel temperature is selected so that it is between 620 ° C to 580 ° C, the proportion of bainite and bainitic ferrite increases with decreasing reel temperature. This means that a uniform structure can be achieved without large differences in hardness, which enables the adherence to narrow thickness and width tolerances in the subsequent cold rolling step. Another positive effect of the low Reel temperatures represent the reduced sensitivity to grain boundary oxidation. In general, the higher the reel temperature, the more likely oxygen-affine elements, such as Si, Cr or Mn, will diffuse to the grain boundary and form stable oxides there, which reduce the surface quality and make an optional later coating difficult. However, the reel temperature T_HT should also not be chosen lower than 400 ° C, since at low reel temperatures the cold rollability is impaired due to extensive formation of martensite. Martensite is a particularly hard and brittle phase that has a negative impact on cold rollability. In addition, at lower reel temperatures there is no longer enough thermal energy available to redistribute the Mn.

If the cooling time t_RG according to the invention and the reel temperature T_HT are observed, a largely bainitic structure is generated in the first minutes of reeling. This consists mainly of very finely divided bainitic ferrite and very finely divided austenite, the grain sizes of the ferrite and austenite each being in the nanometer range. The shortest distance between two phases is typically less than or equal to 20 pm. Mn is a strong austenite former, which is why there is a driving force for a shift of Mn atoms from the ferritic structure components into the austenite grains. During the cooling in the coil, which is very slow, Mn diffuses from the ferrite into the austenite. As a result, the ferritic structural components of Mn become poor in an area which lies directly behind the phase interface from ferrite to austenite. This area depleted in Mn is a few nanometers wide. At the same time, Mn accumulates in the austenite grains directly behind the phase boundary. The diffusion process is locally limited to a few nanometers wide around the phase boundary between austenite and ferrite, since the volume diffusion of Mn takes place very slowly in a temperature range between 620 ° C and 400 ° C. As cooling to temperatures below 400 ° C progresses, the austenite partly breaks down into iron carbides. However, this has no influence on the redistribution of Mn, since the diffusion rate of Mn is too low below 400 ° C and there is no thermodynamic driving force available for homogenization.

The diffusion process of the Mn is supported by very low cooling rates and correspondingly long holding times. In a preferred embodiment, low cooling speeds can be set by cooling the hot strip in the coil in air, in particular in standing air. In a further preferred embodiment, the coil weight can be used to influence the cooling in the coil. The heavier a coil is, the slower it cools down because the ratio of the coil mass to the coil surface increases. A slow cooling and thus a redistribution of Mn in the hot strip can be supported if the coil mass m_CG is at least 101, particularly preferably at least 15 1, very particularly preferably at least 20 1

After cooling in the coil, the hot-rolled flat steel product is pickled in a conventional manner (step d)) and then subjected to cold rolling in a conventional manner (step e)).

The cold rolled flat steel product is heated in step f) to an annealing temperature T_HZ, which can also be referred to as the holding zone temperature. The heating takes place either in one stage with an average heating rate of 2-10 K / s, preferably 5-10 K / s. Alternatively, the heating can also be done in two stages. The flat steel product is first heated up to a turning temperature T_W, which is 200 - 400 ° C, with a heating rate Theta_H 1 of 5 - 50 K / s. Above the turning temperature T_W the heating takes place until the holding zone temperature T_HZ is reached with a heating speed Theta_H2 from 2 - 10 K / s. In two-stage heating, the first heating rate Theta_Hl is not the same as the second heating rate Theta_H2. Preferably, Theta_H2 is less than ThetaJHl.

In a preferred embodiment, the flat steel product is heated in a continuous furnace. In a particularly preferred embodiment, the flat steel product is heated in an oven which is equipped with ceramic radiant tubes, which is particularly advantageous for reaching strip temperatures above 900 ° C.

The holding zone temperature TJfZ is at least 15 ° C., preferably more than 15 ° C., above the A3 temperature of the steel in order to enable a complete structural transformation into the austenite. The A3 temperature is analysis dependent and can be estimated using the following empirical equation:

A3 [° C] = 910-15.2% Ni + 44.7% Si + 31.5% Mo-21, l% Mn-203 * V% C with% C = C content of the steel in weight %,% Ni = Ni content of the steel in% by weight,% Si = Si content of the steel in% by weight,% Mo = Mo content of the steel in% by weight,% Mn = Mn content of the steel in% by weight. The holding zone temperature T_HZ is limited to a maximum of 950 ° C., since at higher temperatures and longer holding times the Mn enrichment already produced in the hot strip could be homogenized again in the austenite and the Mn depletion in the ferrite. Operating costs can also be saved by annealing temperatures limited to 950 ° C.

The flat steel product is held in step g) for a holding time t_HZ of 5-15 s at the holding zone temperature T_HZ. The holding time t_HZ should not exceed 15 seconds in order to avoid the formation of a coarse austenite grain as well as an irregular austenite grain growth and thus negative effects on the formability of the flat steel product. The holding time should last at least 5 s in order to achieve a complete transformation into austenite and a homogeneous C distribution in austenite. The formation of the Mn-poor zone is also negatively influenced by a long t_HZ and the associated Mn homogenization. A too long holding time t_HZ leads to an even distribution of the manganese and thus not to the formation of the Mn-poor ferrite seam.

In step h), the flat steel product is cooled from the holding zone temperature T_HZ to a cooling stop temperature T_Q. The cooling in step h) produces martensite, which is also referred to as primary martensite. The cooling can take place either in one stage or in two stages. In both cases, rapid cooling with a cooling rate Theta_Q of at least 30 K / s takes place over at least part of the temperature range between T_HZ and T_Q. The rapid cooling rate Theta_Q is referred to as Theta_Ql in the case of a one-step cooling and in the case of a two-step cooling as Theta_Q2 for easier distinction between one-step and two-step cooling. In the case of one-step cooling, the flat steel product is called the Theta_Ql with a cooling rate of at least 30 K / s T_HZ cooled to T_Q. The maximum value for Theta_Q1 is 1000 K / s, preferably at most 500 K / s, particularly preferably at most 200 K / s, in order to ensure a uniform temperature distribution. The cooling takes place with at least 30 K / s in order to avoid the conversion into bainite and ferrite contents of more than 10%.

In a two-stage cooling, the flat steel product is first cooled to an intermediate temperature T_LK at a first cooling rate Theta _LK, which is less than 30 K / s. In a preferred embodiment, Theta_LK is greater than 0.1 K / s in order to avoid the formation of ferrite fractions of more than 10% as far as possible. TJLK is smaller than T_HZ and not lower than 650 ° C in order to avoid the formation of ferrite components of more than 10%. After the intermediate temperature T_IK has been reached, the cooling to the cooling stop temperature T_Q takes place without interruption second cooling rate Theta_Q2, which is at least 30 K / s. The maximum value for Theta_Q2 is 1000 K / s _» preferably a maximum of 500 K / s _» particularly preferably a maximum of 200 K / s _»in order to ensure a uniform temperature distribution. The two-stage cooling is also carried out in the temperature range below 650 ° C. at at least 30 K / s in order to avoid both the formation of ferrite fractions of more than 10% and a bainitic conversion. The ferritic and bainitic transformations can be restricted with particular certainty if the time tJLK for the cooling from T_HZ to T_LK is not more than 30 seconds.

To control the formation of martensite, the cooling stop temperature TJJ is selected such that TJJ lies between the martensite start temperature, T_MS, and a temperature which is up to 175 ° C lower than T_MS. The following applies:

 (T_MS-175 ° C) <T_Q <T_MS.

In a preferred embodiment, T_Q can be selected such that T_Q lies between a temperature that is 75 ° C lower than T_MS and a temperature that is 150 ° C lower than T_MS: (T_MS- 150 ° C) <T_Q <( T_MS-75 ° C).

The martensite start temperature TJVIS is the temperature at which the transformation from austenite to martensite begins. The martensite start temperature can be estimated using the following equation:

T_MS [° C] = 539 ° C + (- 423% C-30.4% Mn-7.5% Si + 30% AI) ° C / wt .-% with% C = C content of the steel in wt. %,% Mn = Mn content of the steel in% by weight,% Si = Si content of the steel in% by weight,%,% AI = Al content of the steel in% by weight.

Manganese lowers the martensite start temperature because Mn, as an austenite former, inhibits the thermodynamic driving force for martensite formation. The formation of martensite is thus promoted by reduced Mn contents. For this reason, the first martensite lancets form preferentially in areas that are depleted in Mn, whereas areas with increased Mn content remain primarily austenitic. The phase boundaries from austenite to martensite are therefore preferably located at local Mn enrichment and local Mn depletion points. These points of local Mn enrichment and local Mn depletion were already created during the hot strip manufacturing process and are finely distributed in the material. Typically the sites are local Mn enrichments and more local Mn depletions at intervals of less than 5 pm, preferably less than 1 pm, distributed in the material.

The steel flat product cooled to T_Q is held in work step i) for a duration t_Q, which is 1-60 seconds, at the cooling stop temperature T_Q in order to achieve a homogenization of the temperature distribution in the steel flat product both over the thickness and over the width. A homogeneous distribution of the temperature over the thickness and width of the flat steel product favors the formation of a particularly fine structure. The average grain size is typically less than 20 pm. In some cases, structures with average grain sizes of less than 15 pm or even less than 10 pm can also occur. Typically, there is a uniform structure of primary martensite and residual austenite across the thickness and width of the flat steel product, which has a favorable effect on the formability of the cold-rolled and annealed end product, here the coil or the sheet metal. The temperature distribution can be achieved particularly reliably if the flat steel product is kept at T_Q for at least 5 s, particularly preferably at least 10 s.

After holding at T_Q, the flat steel product is reheated in step j). During the heating, the flat steel product is first heated to a first treatment temperature TJBi, which is at least 10 ° C. above the cooling stop temperature T_Q, at a first heating rate Theta_Bl, which is between 5 and 100 K / s. The treatment temperature T_B1 is at least T_Q + 10 ° C, preferably T_Q + 15 ° C, particularly preferably T_Q + 20 ° C, and at most 450 ° C. The flat steel product is then heated at a second heating rate Theta_B2, which is between 2 and 50 K / s, to a second treatment temperature T_B2, which is at least 10 ° C. above the first treatment temperature T_B1. The second treatment temperature T_B2 is at least T_B1 + 10 ° C, preferably at least T_B1 + 15 ° C, particularly preferably at least TJ31 + 20 ° C. The second treatment temperature TJ32 is at most 500 ° C. In a subsequent optional treatment step, the flat steel product can be kept isothermal at the second treatment temperature T_B2 for a period t_B2 of up to 34 s. The total treatment time t_BT, which includes heating to TJBI, isothermal holding to T_B1, heating to TJB2 and optional holding to T_B2, is between 10 and 250 seconds.

During the heating to the first treatment temperature TJBI, the remaining austenite is enriched with carbon from the supersaturated primary martensite. In a preferred embodiment the ratio of primary martensite to residual austenite is greater than 2: 1, since such a ratio has proven to be particularly favorable for achieving good forming behavior. If the ratio of primary martensite to residual austenite is greater than 2: 1, the effect of an increased thermodynamic driving force can be used to support the carbon shift into the remaining austenite. Because of the comparatively low atomic mass and the high diffusibility of the carbon, especially in the cubic, body-centered lattice of martensite, the diffusion process begins at the Köhl stop temperature T_Q and thus with the start of the martensitic transformation. Since the diffusibility of carbon in the face-centered cubic lattice of austenite is much lower than in martensite, carbon atoms accumulate at the phase boundary between the primary martensite and the austenite. This enrichment leads to a local increase in the C concentration at this point, which can be several percent by weight. In order to ensure sufficient accumulation of carbon atoms at the phase boundary between the primary martensite and the austenite, the first treatment temperature T_B1 should be at least 10 ° C, preferably at least 15 ° C, particularly preferably at least 20 ° C above the cooling stop temperature T_Q. In order to prevent an excessive local increase in the C concentration at this point, TJBi should not exceed 450 ° C, preferably not above 430 ° C, and the duration of the isothermal hold on TJBl should not exceed 245 s, preferably at most 200 s particularly preferably be at most 150 s.

By heating to the second treatment temperature T_B2, the thermodynamic stability of the residual austenite is increased to such an extent that the austenite phase is locally expanded. The pent-up carbon atoms are initially taken up by the residual austenite. In the course of the heating, the diffusion of carbon in the residual austenite increases with a further increase in temperature. This reduces the concentration gradient of the C content at the phase boundary from primary martensite to austenite, so that the carbon in the remaining austenite is approximately evenly and homogeneously distributed. In order to ensure adequate homogenization, the second treatment temperature T_B2 is at least 10 ° C., preferably at least 15 ° C., particularly preferably at least 20 ° C. above the first treatment temperature TJBI and is at most 500 ° C. With the homogenization of the carbon, the grain boundaries of the residual austenite recede, so that the proportion of residual austenite formed during the isothermal holding at the treatment temperature T_B1 decreases. Due to the moving phase boundary, the carbon is transported in the residual austenite which has receded during the heating to the second treatment temperature T_B2. At the same time, the heating increases the diffusibility of the manganese in the area of the phase boundary, which leads to an accumulation of manganese in the receded residual austenite. Has also been found to be advantageous for carbon and manganese diffusion optional hold at treatment temperature TJB2 has been demonstrated for up to 34 s. Along the receding austenite phase boundary, a seam is formed from low manganese ferrite, which has a width of a few nanometers, in particular equal to or less than 12 nm. The low-Mn ferrite seam is formed especially in the low-Mn areas already formed during the production of the hot binders in steps b) and c), since ferrite formation is particularly favored in these areas. The Mn-poor ferrite seam is much more ductile than the other structural components. In the end product, this ductile ferrite serves as a compensation zone between differently plasticizing structural components, such as tempered and non-tempered martensite. The low-Mn ferrite seam works together with the residual austenite to prevent microcracks from spreading, which in particular improves hole expansion.

The duration of the heating up to TJBi is referred to here as t_BRi t_BRi can be determined from the quotient of the difference between the treatment temperature T_B1 and the cooling stop temperature T_Q by the heating rate Theta_Bi: tJBRi = (T_B 1 - T_Q) / Theta_B 1 with tJBRi = heating up time in s; T_B1 = treatment temperature in ° C; T_Q = cooling stop temperature in ° C; Theta_Bl = heating rate in K / s.

When heating up faster with heating rates Theta_Bl greater than 100 K / s, the uniform setting of the treatment temperature T_B1 over the range is difficult to achieve in terms of process and control technology. With a very slow heating up with heating rates Theta_Bl less than 5 K / s, the process runs very slowly and carbides are formed. However, the carbides bind carbon, which is then no longer available to stabilize the remaining austenite. In addition, these carbides are brittle, which prevents flow in the material, which in turn deteriorates the later macroscopic properties, e.g. of the deep-drawing ratio, the elongation at break and the hole expansion.

A complete avoidance of carbide formation is usually not possible in terms of process technology. However, the length of the carbide, which influences the mechanical-technological properties of the flat steel product, can be influenced via the heating rate. The heating rate Theta_Bl is between 5 and 100 K / s, the length of the carbides to at most 250 nm, preferably at most 175 nm. In the present case, the length of the carbides is understood to mean the longest axis of the carbides.

The average heating rate ThetaJB2 _» with which the flat steel product is brought from the first treatment temperature T_B1 to the second treatment temperature T_B2 during two-stage heating is 2 to 50 K / s. The duration in which the flat steel product is brought from T_B1 to T_B2 is referred to here as t_BR2. t_BR2 is 0 to 35 s. The average heat treatment rate Theta_B2 can be determined as

Theta_B2 = (G_B2 - T_B1) / t_BR2 with Theta_B2 = heat treatment rate in K / s; t_BR2 = duration in which the flat steel product is brought from T_B1 to T_B2, in s; T_Bl or T_B2 = treatment temperature in ° C.

Basically, the heating can be carried out by means of conventional heating devices. However, the use of jet pipes or a booster has proven to be particularly effective.

In step j), the flat steel product is kept isothermally at the treatment temperature TJBi and optionally at the treatment temperature T_B2. An isothermal hold on T_B1 and optionally on T_B2 can be used to support the redistribution of carbon. The flat steel product is kept at the treatment temperature T_B1 for a duration t_Bl between 8.5 to 245 s and optionally at a treatment temperature T_B2 for a duration t_B2 of up to 34 s. In a preferred embodiment, the duration of heating to T_B2 and the holding time at temperature T_B2 total at most 35 s, that is to say (t_B2 + t_BR2) <35 s, preferably less than 25 s and particularly preferably less than 20 s.

The total treatment time tJBT, during which the flat steel product is heated to T_B1, held to TJül, heated to T_B2 and optionally held to T_B2, should be between 10 and 250 s. Treatment times shorter than 10 s adversely affect the redistribution of carbon. Treatment times longer than 250 s promote undesired carbide formation.

During the holding or directly during the heating in work step j), the steel flat product can in an optional work step k) a hot-dip coating in a Zn-based Be subjected to coating bath. The duration with which the flat steel product is passed through the coating bath is included in the holding time t_B2 or in the heating-up time t_BR2.

In order to avoid loss of strength, it has proven advantageous to keep the duration t_BR2 for heating to the second treatment temperature T_B2 and the holding time t_B2 short. In particular, it has proven to be advantageous if the holding time t_B2 is zero seconds, so that the flat steel product passes directly into the coating bath from the second heating phase t_BR2. High strength values can thus be achieved particularly reliably if the duration t_BR2 for heating to T_B2 and the optional holding time t_B2 together are at most 35 s, preferably less than 25 s and particularly preferably less than 20 s.

Coating baths suitable for hot dip coating have the following composition:

> 96 wt% Zn, 0.5-2 wt% AI, 0-2 wt% Mg.

The coating baths typically have temperatures of 450 - 500 ° C.

After the optional coating in step k) or - if step k) is omitted - after heating and optional maintenance at treatment temperature T_B2 in step j), the flat steel product is in a further step I) with a cooling rate Theta_B3, which is more than 5 K / s is cooled. The cooling rate should be more than 5 K / s to allow the formation of secondary martensite. Secondary martensite is understood to mean the martensite that is formed during the cooling in step I). Since the secondary martensite does not undergo any heat treatment, it is also referred to here as non-tempered martensite.

The flat steel product produced according to the invention has a particularly fine-grained structure with an average grain size of less than 20 pm, which has a total martensite content of at least 80 area%, of which at least 75 area% is tempered martensite and at most 25 area% is not tempered martensite, at least 5 vol .-% residual austenite, 0.5 to 10 area% ferrite and at most 5 area% bainite.

The structure contains carbides with a length of equal to or less than 250 nm, in particular less than 250 nm, and preferably less than 175 nm. The residual austenite is surrounded by a low-Mn ferrite seam. This seam forms in an area of the phase boundary between tempered martensite and Austenite is a low-Mn zone whose Mn content is at most 50%, in particular less than 50% of the mean total Mn content of the flat steel product, preferably at most 30%, in particular less than 30% of the mean total Mn content of the flat steel product , The width of the Mn-poor ferrite seam is at least 4 nm, in particular more than 4 nm, and preferably at least 8 nm, in particular more than 8 nm. The width of the Mn-poor ferrite seam is at most 12 nm, in particular less than 12 nm, and preferred at most 10 nm, in particular less than 10 nm.

In the present case, the mean total Mn content of the flat steel product is equated with the average Mn content of the molten steel from which the flat steel product was produced

Martensite: The total martensite content in the structure of a flat steel product according to the invention is at least 80% by area. The martensite present in the structure of a steel flat product according to the invention is formed in step h) during the first cooling and in step I) during the second cooling. The martensite formed during the first cooling is also called primary martensite, the martensite formed during the second cooling is also referred to as secondary martensite. The primary martensite is heated in step j). The heated primary martensite is also called tempered martensite or primary tempered martensite. The sum of the martensite components of the tempered and the secondary martensite is also referred to as the total martensite component. As a hard structural component, martensite contributes significantly to the strength of the flat steel product. The total martensite content is at least 80% by area in order to obtain a flat steel product with a tensile strength Rm of at least 900 MPa.

Annealed martensite: The primary martensite, which is formed before the heating, which is carried out in step j), is the source of the carbon which diffuses into the residual austenite during the heating treatment and stabilizes it. After the heat treatment, this martensite is called tempered martensite. Its share should amount to at least 75% of the total area to ensure a bending angle that is greater than 80 ° and a hole widening that is greater than 25%.

Secondary martensite: The secondary martensite arises from the residual austenite insufficiently stabilized in treatment step j) and contributes to the strength. In proportions greater than 25 area% of the total martensite content, the secondary martensite leads to premature crack formation during the forming process and must therefore be kept below 25 area%. Residual austenite In the structure of a steel flat product according to the invention, residual austenite is present at room temperature. Residual austenite contributes to the improvement of the elongation properties. To ensure sufficient elongation, the proportion of residual austenite should be at least 5% by volume.

Ferrite: Ferrite has a lower strength than martensite, but can support formability in small quantities. The proportion of ferrite in the structure of a steel flat product according to the invention is therefore limited to 0.5 to 10 area%. Due to the Mn-poor ferrite seam formed during reheating, step j), there is a minimum ferrite content of 0.5 area% in the structure.

Bainite: In principle, bainite can also form during the phase transformation of austenite. When austenite is converted into bainite, part of the dissolved carbon is incorporated into the bainite and is therefore no longer available for enrichment of the carbon in the austenite. In order to provide as much carbon as possible for enriching the austenite, the bainite content should be limited to a maximum of 5 area%. The lower the bainite content, the more reliably the mechanical properties of the flat steel product can be achieved. The mechanical properties can be achieved particularly reliably if the formation of the bainite can be completely suppressed and the bainite content is reduced to up to 0% by area.

Mn-poor ferrite seam: In the flat steel product according to the invention, the residual austenite grains are surrounded by a narrow Mn-poor ferrite seam. During the heating to treatment temperature T_B1 or T_B2 and while holding on T_B1 or T_B2, a low-Mn zone is formed around the remaining austenite grains, which zone consists of a low-Mn ferrite seam. The Mn-poor ferrite hem is much more ductile than the surrounding structural components. It represents a compensation zone between differently plasticizing structural components and thus counteracts the spread of microcracks. This leads to an improvement in the forming behavior, in particular the hole expansion and the maximum deep-drawing ratio, of the end product. The Mn content in the Mn-poor zone is at most 50%, in particular less than 50% of the mean total Mn content of the flat steel product in order to achieve a hole expansion of more than 25% and a bending angle of more than 80 °. This effect can be achieved particularly reliably if the Mn content in the Mn poor zone is at most 30%, in particular less than 30%, of the average Mn content of the flat steel product. The width of the Mn-poor ferrite seam is at least 4 nm, in particular more than 4 nm, since ductile compensation can only arise from a width of 4 nm. Would be the Mn-poor ferrite hem narrower, the zone would no longer effectively contribute to ductility compensation, but the forming would already be influenced by congruence effects. The ductility compensation can be achieved particularly safely if the Mn-poor ferrite seam is preferred to be at least 8 nm, in particular more than 8 nm wide. The width of the Mn-poor ferrite seam wobbles with increasing treatment time during treatment step j). Since the positive contribution of the seam has been increased from 12 nm and the risk of carbide formation increases with the duration of treatment during step j), the width of the seam should be at most 12 nm, in particular less than 12 nm. This effect can be achieved particularly reliably if the Mn-poor ferrite seam is preferably at most 10 nm, in particular less than 10 nm wide.

Carbides: Carbides bind carbon. The carbon bound in carbide form is not available for redistribution into the austenite. In addition, carbides show brittle fracture behavior. The brittle behavior of the carbides prevents plastic flow in the material, which leads to a deterioration in the macroscopic properties, such as the maximum deep-drawing ratio and / or the hole expansion. The maximum length of the carbides should be equal to or less than 250 nm in order to avoid deterioration of the elongation at break and / or the hole expansion. The mechanical-technological properties can be achieved particularly safely if the length of the carbides is preferably less than 175 nm. The length of a carbide is understood to mean its longest axis. In the present case, the term “carbides” generally means carbon deposits. These are precipitations in which carbon forms compounds, such as iron carbides, chromium carbides, titanium carbides, niobium carbides or vanadium carbides, together with elements present in the flat steel product.

The method according to the invention enables the production of a flat steel product with a tensile strength Rm of 900 to 1500 MPa, an elastic limit Rp02 which is equal to or more than 700 MPa and less than the tensile strength of the flat steel product, an elongation A80 of 7 to 25%, a bending angle, which is greater than 80 °, a hole expansion which is greater than 25%, and a maximum deep-drawing ratio ßmax, for which the following relationship applies: ßmax> -1.9 IO ⁶ x (Ä _m ) ² + 3.5 10 ^{~ 3} x R _m + 0.5 with Rm: tensile strength of the flat steel product in MPa, In a preferred embodiment, the flat steel product has a balanced ratio of high tensile strength and good deep-drawing behavior. In this case, the maximum deep-drawing ratio ßmax is at least 1.475. A flat steel product according to the invention thus has both good strength and forming properties.

The method according to the invention enables in particular the production of a flat steel product according to one of claims 1 to 8.

The invention is explained in more detail below on the basis of exemplary embodiments and figures.

Figure 1 shows schematically a possible variant of the method according to the invention. The cold-rolled and uncoated flat steel product is heated to a holding temperature T_HZ and held before it is cooled in one stage to a cooling stop temperature T_Q with a cooling rate Theta_Ql. After an isothermal hold at T_Q, the flat steel product is heated in a first heating step to the treatment temperature T_B1, at which it is kept isothermal. It is then heated to a second treatment temperature TJB2, at which it is held again before it is cooled to room temperature.

Figure 2 shows schematically a further variant of the method according to the invention. The cold-rolled and uncoated flat steel product is also heated and held at a holding temperature TJHZ, before it is first cooled to an intermediate temperature T_LK with a first, slower cooling rate Theta_LK and then to the cooling stop temperature T_Q with a second, faster cooling rate. The flat steel product is then, as already explained for FIG. 1, heated in two stages and then cooled to room temperature.

Each of the variants described can also be combined with a hot-dip coating treatment. Hot-dip coating is included in the isothermal hold at treatment temperature T_B2 or in the period t_BR2 during heating to treatment temperature T_B2 before the flat steel product is cooled to room temperature.

The invention was tested using several exemplary embodiments. 14 tests were carried out for this. Samples of 14 cold-rolled and coated steel strips, which were produced from the steels A - G shown in Table 1, were examined. For that were initially, slabs were produced in a conventional manner from melts of the compositions given in Table 1. Before the hot rolling, the slabs were each heated to a temperature of 1000-1300 ° C. and, with the conditions given in Table 2, rolled in a conventional manner to hot binders and wound into hot strip coils. The hot binders were subjected to pickling in a conventional manner and then also in a conventional manner Way cold rolled. Table 3 shows the conditions under which the samples were each heat treated. The cold-rolled flat steel products were each heated in one stage to the holding zone temperature T_HZ using the heating rate ThetaJHi given in Table 3 and held at the temperature TJHZ for 5 to 15 s. The flat steel products were then cooled in two stages, first with an initial cooling rate Theta_LK, which was more than 0.1 K / s and equal to or less than 30 K / s, to the intermediate temperature T_LK and then with a second cooling rate Theta_Q2 to the cooling stop temperature T_Q. The flat steel products were held at> 1 second and <60 seconds at T_Q and then heated to a first treatment temperature T_B1 at a first heating rate Theta_Bl for a period t_BRi. After the heating, the flat steel products were kept at T_Bt for a period t_Bl and then heated at a second heating rate Theta_B2 over a period tJBR2 to the second treatment temperature T_B2, with which they were introduced directly into a Zn-based coating bath. The flat steel products were continuously passed through a coating bath which had a composition of> 96% Zn, 0.5-2% Al, 0-2% Mg. The time t_B2, which also includes the passage of the flat steel products through the coating bath, and the total treatment time are also given in Table 3. After coating, the flat steel products were cooled at a cooling rate Theta_B3 of more than 5 K / s.

After cooling, samples were taken for structural analysis and for determining the mechanical properties. The structure was examined on three cross sections, which were taken equidistantly across the width of the flat steel product. The structural examination was carried out in each case over the thickness of the flat steel product at at least three equidistantly spaced locations. Due to the very fine-grained structure, it was not possible to assess the structure using conventional light-optical examination methods. Therefore, the proportions of the primary, tempered martensite (M (PRI) M_l), the secondary martensite (M (SEK) M_2), the ferrite (F) and the bainite (B) were determined using a scanning electron microscope (SEM) at a minimum of 5000- magnification examined. The quantitative determination of the residual austenite content was carried out by means of X-ray diffraction (XRD) according to ASTM E975. The description of the Mn-poor ferrite fringe and the measurement of the Mn-

The content of the Mn-poor ferrite seam was determined using a tomographic atom probe (Atom-Probe-Tomography, API). In this way, the width of the Mn-poor ferrite seam, which is referred to in Table 4 as the Mn edge, was also determined The Mn content of the Mn-poor ferrite was determined in a defined volume element, for example a cylinder or a cuboid, to determine the number of atoms, in order to determine the width of the Mn-poor ferrite fringe a width measurement of the The individual values were arithmetically averaged and represent the size referred to as the width of the Mn-poor ferrite seam. The Mn content of the Mn-poor ferrite is referred to in Table 4 as the Mn content edge. The length of the carbides was determined by TEM. The results of the structural studies are shown in Table 4.

The results of the testing of the mechanical properties are shown in Table 5. The mechanical properties were investigated in each case on samples which were taken from three points distributed equidistantly over the length of the flat steel product, each in the middle of the width of the flat steel product. The yield strength Rp02, the tensile strength Rm and the elongation ASO were determined in the tensile test according to DIN EN ISO 6892-1 (sample form 2) from 02/2017. The bending angle (bending) was determined according to VDA238-100 of 12/1010, the hole expansion (HER) was measured according to ISO 16630 of 10/2017 and the maximum drawing ratio ß _ma x was determined from 09/2003 in accordance with DIN 8584-3.

The results show that tests with the method carried out according to the invention lead to high strengths and at the same time good forming behavior. Samples B2, B3, D7, Di, F12, F13 and G14 show bending angles greater than 80 ° and hole expansion values greater than 25%. Test A1 shows that the structure of the invention could not be set if the silicon content was not according to the invention. The high proportion of secondary martensite and the high proportion of ferrite led to a comparatively low yield strength and tensile strength. Furthermore, there was only a very small Mn-poor ferrite seam, so that only a small bending angle and a low hole expansion were achieved.

Test B4 shows that despite the steel composition according to the invention, the formability is impaired if the end roll temperature TJET and the cooling stop temperature T_Q are not according to the invention and the Mn-poor ferrite seam is too narrow. The yield strength and tensile strength are sufficiently high, but the bending angle and the hole expansion are too small due to the insufficient Mn depletion in the Mn-poor ferrite fringe or the insufficient Mn enrichment in the zone adjacent to the Mn-poor ferrite fringe.

Experiments C5 and C6 show that if the carbon and silicon content is too low, the proportion of bainite (experiment C5) or of secondary martensite and ferrite (experiment C6) is too high and the width of the Mn-poor ferrite seam is too small, in order to be able to achieve a sufficiently high hole expansion (test C5) or a sufficient yield strength, bending angle and hole expansion (test C6).

 Test D8 shows that despite the steel composition according to the invention, the formability is impaired by carbides that are too long if the reel temperature T_HT is too high, the heating rate Theta_Bl is too low and the overall heat treatment time tJBT is too long. If the t_BT is selected too long, the maximum carbide length is exceeded, which has a negative effect on the hole expansion.

Test E10 shows that if the silicon content is too low and the time for cooling down after hot rolling to coiling temperature, t_RG, the proportion of secondary martensite and the proportion of ferrite increases, which leads to an inhomogeneous structure and thus to an insufficient bending angle and an insufficient one Hole expansion leads.

Test Eli shows that if the silicon content is too low and the reel temperature is not according to the invention, the proportion of secondary martensite increases and the carbides become too long, which worsens the elongation A80 and the hole widening. Experiment Eli also shows that both a too low reel temperature and an excess of the Treatment time on T_B2, i.e. tJBR2 + t_B2> 35 seconds, has a negative effect on the properties of the flat steel product If it is not possible to suppress carbide formation sufficiently, carbides that are too long form, leading to premature crack formation and correspondingly poor values for hole expansion.

Figures in% by weight, remainder iron and unavoidable impurities. Underlined values lie outside the specifications according to the invention.

Table 1

 Underlined values lie outside the specifications according to the invention.

Table 2 O

n H

Underlined values lie outside the specifications according to the invention.

 bo o o

Table 3

t u

Underlined values lie outside the specifications according to the invention.

Table 4

 Underlined values lie outside the specifications according to the invention.

Table 5

Claims

claims 1. Flat steel product, characterized in that it contains a steel which consists of (in% by weight) 04 - 0.5% C,  1.0 - 3.0% Mn,  0.9 - 1.5% Si,  up to 1.5% AI,  up to 0.008% N,  up to 0.020% P,  up to 0.005% S  0.01-1% Cr,  and optionally one or more of the following elements  up to 0.2% Mo,  up to 0.01% B,  up to 0.5% Cu,  up to 0.5% Ni  and optionally from a total of 0.005 - 0.2% microalloying elements  and the remainder consists of iron and unavoidable impurities, whereby: 75 <(Mn2 + 55 * Cr) / Cr <3000 with Mn: Mn content of the steel in% by weight, Cr: Cr content of the steel in% by weight  and has a structure that consists of  at least 80 area% martensite, of which  is at least 75% by area tempered martensite and at most 25% by area non-tempered martensite,  at least 5% by volume of residual austenite,  0.5 to 10 area% ferrite and  at most 5 area% bainite  exists in the area of the phase boundary between tempered martensite and Residual austenite is a Mn-poor ferrite seam which has a width of at least 4 nm and at most 12 nm and whose Mn content is at most 50% of the mean total Mn content of the flat steel product and the flat steel product has carbides whose length is equal to or less than than 250 nm. 2. Flat steel product according to claim 1, characterized in that it has a tensile strength Rm of 900 to 1500 MPa, an elastic limit Rp02 of more than 700 MPa, an elongation A80 of 7 to 25%, a bending angle greater than 80 °, a hole expansion greater than 25% and a maximum Thermoforming _ratio ß _m® for which the following applies: 3.5 IO ^{"" 3} XR _m + 0.5 with Rm: tensile strength of the flat steel product in MPa. 3. Flat steel product according to claim 1 or 2, characterized in that the width of the Mn-poor ferrite hem is at least 8 nm. 4. Flat steel product according to one of the preceding claims, characterized in that the width of the Mn-poor ferrite hem is at most 10 nm. 5. Flat steel product according to one of the preceding claims, characterized in that the Mn content of the Mn-poor ferrite seam is at most 30% of the average total Mn content in the flat steel product. 6. Flat steel product according to one of the preceding claims, characterized in that the length of the carbides is less than 175 nm. 7. Flat steel product according to one of the preceding claims, characterized in that the flat steel product has a maximum deep-drawing ratio βmax of at least 1.475. 8. Flat steel product according to one of the preceding claims, characterized in that the flat steel product is provided with a metallic coating. 9. A process for producing a high-strength flat steel product comprising at least the following steps: a) Providing a slab which consists of a steel which, in addition to iron and unavoidable impurities, consists of (in% by weight) 0.1 - 0.5% C, 1.0 - 3.0% Mn, 0.9 - 1.5% SI, up to 1.5% AI, up to 0.008% N, up to 0.020% P, up to 0.005% S, 0.01 to 1% Cr, and optionally from one or more of the following elements: up to 0.2% Mo, up to 0.01% B, up to 0.5% Cu, up to 0.5% Ni and optionally consists of a total of 0.005 - 0.2% microalloying elements, where: 75 <(Mn ² + 55 * Cr) / Cr <3000 with Mn: Mn content of the steel in% by weight, Cr Cr content of the steel in% by weight; b) heating the slab to temperatures of 1000-1300 ° C and hot rolling the slab to a hot strip, the final rolling temperature T_ET being greater than 850 ° C; c) cooling the hot strip within a maximum of 25 s to a reel temperature T_HT which is 400 to 620 ° C. and winding the hot strip into a coli; d) pickling the hot rolled flat steel product; e) cold rolling the hot rolled flat steel product; f) heating the cold-rolled flat steel product to a holding zone temperature T_HZ which is at least 15 ° C above the A3 temperature of the steel and is at most 950 ° C, the heating either fl) in one stage with an average heating rate of 2 - 10 K / s or f2) in two stages with a first heating rate Theta_H 1 of 5 - 50 K / s up to one 200 - 400 ° C turning temperature TJV and above the turning temperature T_W with a second heating rate Theta_H2 of 2 - 10 K / s; g) holding the steel product for a duration t_HZ of 5 - 15 s at the holding zone temperature T_HZ; h) cooling the steel product from the holding zone temperature T_HZ to a cooling stop temperature T_Q, which lies between the martensite start temperature T_MS and a temperature which is 175 ° C lower than T_MS, with either hl) a cooling rate Theta_Ql which is at least 30 K / s; or h2) a first cooling rate ThetaJLK of less than 30 K / s for a first cooling to an intermediate temperature T_LK, which is not lower than 650 ° C, and a second cooling rate Theta_Q2 for a second cooling from T_LK to T_Q, where Theta_Q2 is at least 30 K. / s is; i) holding the flat steel product on the coulter istop ptem peratu r T_Q for 1-60 seconds; j) heating the flat steel product with a first heating rate Theta_Bl, which is between 5 and 100 K / s, to a first treatment temperature T_B1, which is at least T_Q + 10 ° C and at most 450 ° C, keeping the flat steel product at the first treatment temperature T_B1 for a duration t_Bl from 8.5 s to 245 s, heating the flat steel product with a second heating rate Theta_B2, which is between 2 and 50 K / s, to a second treatment temperature T_B2, which is at least T_B1 + 10 ° C and at most 500 ° C optionally holding the steel flat product at the treatment temperature T_B2 for a duration t_B2 of up to 34 s, the total treatment time t_BT for the heating and the isothermal holding being between 10 and 250 s in total; k) optional coating of the flat steel product in a Zn-based coating bath; L) cooling the flat steel product to room temperature with a cooling rate Theta_B3 of at least 5 K / s; 10. The method according to claim 9, characterized in that the weight of the hot strip wound into a coil is at least 101. 11. The method according to claim 9 or 10, characterized in that the cooling of the hot strip in work step c) takes place within a maximum of 18 s. 12. The method according to any one of claims 9 to 11, characterized in that the reel temperature T_HT is at most 600 ° C. 13. The method according to any one of claims 9 to 12, characterized in that the holding time t_Q is at least 5 s. 14. The method according to any one of claims 9 to 13, characterized in that the cooling stop temperature T_Q between a temperature which is 75 ° C lower than that Martensite start temperature T_MS actual and a temperature that is 150 ° C lower than T_MS. 15. The method according to any one of claims 9 to 14, characterized in that the duration t_BR2 for heating up to T_B2 and the duration for the optional holding on T_B2 together is at most 35 s.