WO2025078050A1 - High-tensile steel exhibiting improved resistance to hydrogen embrittlement - Google Patents

Abstract

The invention relates to a flat steel product to be subjected to hot forming, to a shaped sheet metal blank, to a process for manufacturing the flat steel product, and to a process for manufacturing the shaped sheet metal blank, wherein the flat steel product and the sheet metal blank, in particular in combination with an aluminum-based corrosion-resistant coating, exhibit improved resistance to hydrogen embrittlement.

Description

High-strength steel with improved resistance to hydrogen embrittlement

The invention relates to a flat steel product for hot forming and a method for producing such a flat steel product. Furthermore, the invention relates to a sheet metal part with improved properties and a method for producing such a sheet metal part from a flat steel product.

When reference is made below to a "flat steel product" or a "sheet metal product," this refers to rolled products, such as steel strips or sheets, from which "sheet metal blanks" (also called blanks) are cut for the production of, for example, body components. "Formed sheet metal parts" or "sheet metal components" of the type according to the invention are made from such sheet metal blanks, whereby the terms "formed sheet metal part" and "sheet metal component" are used synonymously here.

All information regarding the contents of the steel compositions specified in this application is based on weight, unless expressly stated otherwise. All unspecified "%" data relating to a steel alloy are therefore to be understood as data in "wt. %." With the exception of the data relating to the residual austenite content of the microstructure of a sheet metal part according to the invention, which is based on volume (specified in "vol. %"), data on the contents of the various microstructure constituents refer to the area of a microsection of a sample of the respective product (specified in area percent, "area %"), unless expressly stated otherwise. Information provided in this text regarding the contents of the constituents of an atmosphere refers to the volume (specified in "vol. %").

Mechanical properties reported here, such as tensile strength, yield strength, and elongation, were determined in tensile tests according to DIN EN ISO 6982-1, specimen form 2 (Appendix B Table B1) (as of 2020-06), unless explicitly stated otherwise. The bending angle is determined according to VDA standard 238-100.

The microstructure was determined on longitudinal sections etched with 3% Nital (alcoholic nitric acid). The content of retained austenite was determined by X-ray diffraction. WO 2019/223854 A1 discloses a sheet metal part and a method for producing such a sheet metal part, which has a tensile strength of at least 1000 MPa. The sheet metal part consists of a steel which, in addition to iron and unavoidable impurities, is composed of (in wt. %) 0.10-0.30% C, 0.5-2.0% Si, 0.5-2.4% Mn, 0.01-0.2% Al, 0.005-1.5% Cr, 0.01-0.1% P, and optionally further optional elements, in particular 0.005-0.1% Nb. Furthermore, the sheet metal component comprises a corrosion protection coating containing aluminum.

EP 2 553 133 B1 also discloses a sheet metal part and a method for producing such a sheet metal part.

As delivered, all manganese-boron steel grades are low in hydrogen. Their diffusible hydrogen contents are below the current detection limit of 0.1 ppm. As a result, MnB steels generally exhibit only a low tendency toward hydrogen-induced delayed cracking. However, practice has shown that during hot forming of manganese-boron steels in humid furnace atmospheres, hydrogen enrichment in the steel substrate occurs. A metal-water vapor reaction has been identified as the cause of this.

This reaction occurs when the steel flat product is heated to elevated temperatures in a heating furnace under a steam-containing atmosphere for hot forming. The steam in the furnace atmosphere reacts on the material surface to form hydrogen and a metal oxide. The resulting hydrogen diffuses into the steel material and can then lead to delayed failure by preferentially concentrating in areas of high tensile residual stress. If a very high local hydrogen concentration is reached, this weakens the bond at the grain boundaries of the steel substrate structure to such an extent that, during use, the resulting stress causes a crack to develop along the grain boundary.

To prevent hydrogen ingress due to surface reactions in the furnace, dew point control devices are often used. The goal here is to limit the water vapor supply in the furnace atmosphere.

From EP 2 993 248 B1 it is known to use an admixture of alkali or alkaline earth metals in flat steel products with an aluminum-based corrosion protection coating, which are very quickly form a protective oxide layer on the surface, thus preventing surface reaction with water vapor. This results in less free hydrogen being produced, which can diffuse into the steel substrate.

Furthermore, corrosion occurs during typical automotive use of sheet metal parts. This is particularly the case with uncoated substrates. This corrosion also produces hydrogen, which diffuses into the steel material and can lead to delayed failure by preferentially concentrating in areas of high tensile residual stress. If a very high local hydrogen concentration is reached, this weakens the bonding at the grain boundaries of the steel substrate structure to such an extent that, during use, the resulting stress causes a crack along the grain boundary.

The object of the present invention is to further reduce the content of free hydrogen in sheet metal parts, in particular in uncoated substrates.

Another object of the present invention is to reduce the susceptibility to hydrogen embrittlement under corrosion.

This object is achieved by a steel flat product for hot forming, comprising a steel substrate made of steel which, in addition to iron and unavoidable impurities (in wt.%), consists of:

C: 0.12-0.30%,

Si: 0.02-1.2%,

AI: 0.01-1.0%,

B: 0.0005-0.01%,

P: < 0.05%,

S: < 0.02%,

N: < 0.02%,

Sn: < 0.03%,

As: < 0.010%,

Sb: < 0.02%, at least one of the elements from the group comprising Ti, Nb and V, with the proviso that the contents are as follows:

Ti: 0.008-0.10%, Nb: 0.01-0.08%,

V: 0.01-0.4%, and optionally one or more of the elements “Cr, Mn, Cu, Mo, Ni, Ca, W” in the following contents:

Cr: 0.01-1.0%,

Mn: 0.2-3.3%,

Cu: up to 0.2%,

Mo: 0.002-0.5%,

Ni: 0.01-2.0%,

Ca: 0.0005-0.01%,

W: 0.001-1.0%, where the following applies to a degree of dispersion Di of precipitates in the near-surface third of the steel substrate:

1 1

25 ■ 10“ ⁶ — < D < 5 ■ 10“ ³ — nm nm

von Ausscheidungen definiert als das Verhältnis aus Flächenanteil der Ausscheidungen im metallographischen Schliff zum Durchmesser der Ausscheidungen, wobei der Flächenanteil der Ausscheidungen das Verhältnis der Gesamtfläche der Ausscheidungen in einem Messfeld zur Größe des Messfeldes ist. Bei dem Durchmesser handelt es sich um den mittleren Durchmesser über das Messfeld. The degree of dispersion

Precipitation area is defined as the ratio of the area of precipitates in the metallographic section to the diameter of the precipitates, where the area of precipitates is the ratio of the total area of precipitates in a measurement field to the size of the measurement field. The diameter is the average diameter across the measurement field.

According to the invention, the precipitates act as traps for free hydrogen. Free hydrogen penetrating from the outside is thus localized at the precipitates. Consequently, high concentrations in regions with high tensile residual stress are avoided. It has been shown that this effect is practically nonexistent when the dispersion D is too low. This means that the precipitates are too large and/or the proportion of precipitates is too small. There is then insufficient specific interface between the precipitates and the substrate to effectively localize hydrogen. If, on the other hand, the dispersion D is too large, the precipitates themselves lead to a reduction in mechanical strength. This is caused by a repulsive force of the precipitates on the movement of dislocations in the crystal, the so-called Zener drag. This Zener drag is directly proportional to the dispersion Di. The dislocation movement, in turn, causes plastic deformation and influences This directly reduces the achievable bending angle. For example, the bending angle is significantly reduced. In particular, the dispersion D is at least 35 ^10-6 ", particularly preferably at least ^{15 10-5} nm . Furthermore, the dispersion D is preferably at most 2 ^10-3 _nm , particularly at most 9.0 ^10-4 nm.

The distribution of the precipitates and thus the dispersion D _± is measured using electron-optical images in combination with X-ray microanalysis (TEM and EDX) based on carbon extraction replicas (known in the specialist literature as "carbon extraction replicas"). The carbon extraction replicas are created from longitudinal sections. The magnification of the measurement is between 10,000x and 200,000x. Based on these images, the mean diameter and the area fraction of the precipitates in the measurement field can be calculated using computer-assisted image analysis. Five measurement fields are measured for each of these. The results of the five measurement fields are then averaged. The size of the measurement fields depends on the selected magnification and ranges from 18.5 pm x 14.5 pm at 10,000x magnification to 0.925 pm x 0.725 pm at 200,000x magnification.

For example, dispersion can be determined using the following steps:

- Making carbon impressions of a longitudinal section of the flat steel product or sheet metal part;

- Determination of the mean diameter and the area fraction of the precipitates on 5 different measuring fields of size l.85pm x l.45pm using TEM and computer-assisted image analysis at a magnification of 100,000x;

- Calculate the dispersion in each of the 5 measuring fields from the mean diameter and the area fraction;

- Determination of the dispersion of the flat steel product or sheet metal part as the mean value of the dispersion over the 5 measuring fields.

For the dispersion D _± of flat steel products or sheet metal parts, the 5 measuring fields are placed in the near-surface third of the steel substrate.

For dispersion D ₂ in the alloy layer of the sheet metal part, which will be explained later, the corresponding measuring procedure applies, except that the 5 measuring fields are placed in the alloy layer. The precipitates are primarily carbides and/or carbonitrides of one or more of the elements from the group vanadium, titanium, niobium, chromium, and molybdenum. Vanadium, titanium, and niobium are so-called microalloying elements, which exert an effect even in the smallest quantities.

The near-surface third of the steel substrate is the area of the steel substrate that is at a distance from any surface of the steel substrate that corresponds to a maximum of one-third of the thickness of the steel substrate. This means that the steel substrate is conceptually divided into three slices of equal thickness, parallel to the surface. The middle slice then contains the central region of the steel substrate, in which the midplane lies. The other two slices are each bounded by one of the two surfaces of the steel substrate. These two slices form the two near-surface thirds of the steel substrate.

Carbon ("C") is present in the steel substrate of the flat steel product in concentrations of 0.12–0.30 wt.%. Such adjusted C contents contribute to the hardenability of the steel by delaying ferrite and bainite formation and stabilizing the residual austenite in the microstructure. A carbon content of at least 0.06 wt.% is required to achieve sufficient hardenability and the associated high strength.

However, high carbon contents can negatively affect weldability. To improve weldability, the carbon content can be adjusted to a maximum of 0.28 wt.%, preferably to a maximum of 0.25 wt.%.

To ensure the most reliable use of the positive effects of the presence of C, C contents of at least 0.15 wt.%, preferably at least 0.20 wt.%, can be provided. At these contents, tensile strengths of the sheet metal part of at least 1100 MPa, in particular at least 1250 MPa, in particular at least 1400 MPa, can be reliably achieved after hot press forming, subject to the further provisions of the invention.

Silicon (“Si”) is used to further increase the hardenability of the steel flat product as well as the strength of the press-hardened product through solid solution strengthening. Silicon also enables the use of ferro-silicon-manganese as an alloying agent, which has a beneficial effect on production costs. Silicon is present in the steel substrate of the steel flat product in amounts of 0.02-1.2 wt.%. A hardening effect is already present at an Si content of 0.05 wt.%. A significant increase in strength occurs at a Si content of at least 0.15 wt.%, in particular at least 0.20 wt.%. Si contents above 0.65 wt.% have a detrimental effect on the coating behavior, especially in the case of Al-based coatings. Si contents of at most 0.55 wt.%, in particular at most 0.35 wt.%, are preferred in order to improve the surface quality of the coated flat steel product.

Aluminum ("Al") is known to be added as a deoxidizer during steel production. The content of the steel substrate is at least 0.01 wt.%, in particular at least 0.02 wt.%, preferably at least 0.11 wt.%, in particular at least 0.15 wt.%. The maximum aluminum content is 1.0 wt.%, preferably a maximum of 0.7 wt.%, in particular a maximum of 0.25 wt.%, in particular a maximum of 0.20 wt.%. In certain embodiments, the maximum aluminum content is 0.10 wt.%, preferably a maximum of 0.05 wt. To reliably bind the oxygen contained in the molten steel, at least 0.01 wt.% Al is required. In addition, Al can be used to bind undesirable, but unavoidable, levels of N during production. Comparatively high aluminum contents have been avoided to date, since the Ac3 temperature also shifts upwards with the aluminum content. This has a negative impact on austenitization, which is important for hot forming. However, it has been shown that increased aluminum contents surprisingly lead to positive effects when combined with an aluminum-based corrosion protection coating.

Surprisingly, it has been shown that by increasing the aluminum content ("Al") in the steel substrate to the described lower limits and beyond, a significant reduction in pore formation can be achieved during coating with an aluminum-based corrosion protection coating and subsequent hot forming. Particularly in the transition region between the steel substrate and the corrosion protection coating, the locally higher aluminum consumption during the formation of denser iron-aluminide compounds can be at least partially compensated by the aluminum content of the steel substrate, thus suppressing the formation of pores, especially a band of pores. If the Al content is too high, especially at contents of more than 1.0 wt% Al, there is a risk of Al oxides forming on the surface of a product made from steel alloyed according to the invention, which would impair the wetting behavior during hot-dip coating. Furthermore, higher Al contents promote the formation of non-metallic Al-based inclusions, which, as coarse inclusions, negatively impact crash behavior. Therefore, the Al content is preferably selected below the upper limits already mentioned.

Furthermore, the steel comprises at least one of the elements from the group comprising Ti, Nb, and V. This means that the steel contains at least one of the elements Ti, Nb, or V. The steel can also preferably contain two elements from the group (Ti, Nb, or Ti, V, or Nb, V) or, in particular, all three elements from the group (Ti, Nb, and V). The limits listed below, with their preferred ranges, apply in all cases to the contents of the three elements.

The microalloying elements vanadium, niobium, and titanium all form precipitates that contribute to grain refinement and act as traps for free hydrogen. However, due to their different solubilities, the precipitates occur at different temperatures. Ti precipitates have the lowest solubility in austenite and therefore precipitate at very high temperatures, thereby reducing austenite grain growth. Nb precipitates at medium temperatures, while vanadium only precipitates below approximately 900°C. This means that vanadium leads to particularly fine precipitates. Otherwise, the same mechanism occurs for all three elements. The following example explains the mechanism using niobium. The same applies to vanadium and titanium.

The niobium content, particularly in the process described below for producing a flat steel product for hot forming with a corrosion protection coating, leads to a distribution of niobium carbides and niobium carbonitrides, which results in a particularly fine hardened microstructure during subsequent hot forming. During cooling after hot-dip coating, the coated flat steel product is kept for a certain time in a temperature range of 400 °C to 300 °C. In this temperature range, a certain diffusion rate of carbon still exists in the steel substrate, while the thermodynamic solubility is very low. Thus, carbon diffuses to lattice defects and accumulates there. Lattice defects are caused in particular by dissolved niobium atoms, which, due to their significantly higher atomic volume, expand the atomic lattice and thus enlarge the tetrahedral and octahedral gaps in the atomic lattice, thus increasing the local solubility of carbon. Consequently, clusters are formed. of C and Nb in the steel substrate, which then transform into very fine precipitates in the form of niobium carbides and niobium carbonitrides in the subsequent austenitizing step of hot forming and act as additional austenite nuclei. This results in a refined austenite microstructure with smaller austenite grains and thus also a refined hardening microstructure. Furthermore, these precipitates form traps for free hydrogen and thus promote the inventive resistance to hydrogen embrittlement.

This particularly applies to the ferritic interdiffusion layer that forms during hot forming. The refined ferritic microstructure in the interdiffusion layer helps reduce crack initiation under bending loads, and the precipitates in the ferritic interdiffusion layer trap free hydrogen before it can concentrate inside the substrate.

For the effect described above, the Nb content is at least 0.01 wt.%, preferably at least 0.010 wt.%, in particular at least 0.02 wt.%. In certain embodiments, the niobium content is at least 0.03 wt.%. The maximum niobium content is 0.08 wt.%, in particular a maximum of 0.07 wt.%, in particular a maximum of 0.05 wt.%, preferably 0.03 wt.%, in particular a maximum of 0.02 wt.%.

The titanium content for the effect described above is at least 0.008 wt.% Ti, in particular at least 0.010 wt.%, preferably at least 0.015 wt.% Ti, which should be added to ensure sufficient availability. From 0.10 wt.% Ti, cold rollability and recrystallizability deteriorate significantly, which is why higher Ti contents should be avoided. To improve cold rollability, the Ti content can preferably be limited to 0.08 wt.%, in particular to a maximum of 0.050 wt.%, particularly preferably to a maximum of 0.040 wt.%, in particular a maximum of 0.030 wt.%. Titanium also has the effect of binding nitrogen, thus enabling boron to exert its strong ferrite-inhibiting effect. Therefore, in a preferred development, the titanium content is more than 3.42 times the nitrogen content in order to achieve sufficient binding of nitrogen.

For the effect described above, the V content is at least 0.01 wt.%, in particular at least 0.02 wt.%, preferably at least 0.04 wt.%, in particular at least 0.10 wt.%. For cost reasons, a maximum of 0.4 wt.%, preferably a maximum of 0.3 wt.%, in particular a maximum of 0.25 wt.%, particularly preferably a maximum of 0.15 wt.%, preferably a maximum of 0.10 wt.%, particularly preferably a maximum of 0.05 wt.%, particularly preferably a maximum of 0.03 wt.%, in particular a maximum of 0.02 wt.%.

Excessive proportions of the microalloying elements Nb, V, and Ti can lead to excessive dispersion, which in turn negatively affects the mechanical properties (such as the bending angle). Therefore, in a preferred embodiment, the sum of the Nb, V, and Ti contents is a maximum of 0.13 wt.%, in particular a maximum of 0.12 wt.%, preferably a maximum of 0.10 wt.%, in particular a maximum of 0.08 wt.%, preferably a maximum of 0.06 wt.%. The following therefore preferably applies to the element contents (in wt.%):

V + Nb + Ti < 0.13

For the same reasons, the individual contents of Nb, V and Ti are also preferably limited, as described in the previous paragraphs.

The elements Nb, V and Ti can also occur as impurities in steel below the minimum contents mentioned above.

Boron ("B") is added to improve the hardenability of the flat steel product by reducing the grain boundary energy through boron atoms or boron precipitates deposited on the austenite grain boundaries, thereby suppressing the nucleation of ferrite during press hardening. A significant effect on hardenability occurs at contents of at least 0.0005 wt.%, preferably at least 0.0007 wt.%, in particular at least 0.0010 wt.%, in particular at least 0.0020 wt.%. At contents above 0.01 wt.%, however, increased formation of boron carbides, boron nitrides, or boron nitrocarbides occurs, which in turn represent preferred nucleation sites for the nucleation of ferrite and reduce the hardening effect. For this reason, the boron content is limited to at most 0.01 wt.%, in particular at most 0.010 wt.%, preferably at most 0.0100 wt.%, preferably at most 0.0050 wt.%, in particular at most 0.0035 wt.%, in particular at most 0.0030 wt.%, preferably at most 0.0025 wt.%.

Phosphorus ("P") and sulfur ("S") are elements that are introduced into steel as impurities by iron ore and cannot be completely eliminated in the industrial steelmaking process. The P and S content should be kept as low as possible, since the mechanical properties, such as the notched-bar impact energy, deteriorate with increasing P content or S content. Furthermore, from P contents of 0.03 wt.%, embrittlement of the martensite begins to occur, which is why the P content of a flat steel product according to the invention is preferably limited to a maximum of 0.05 wt.%, in particular to 0.03 wt.%, preferably to a maximum of 0.02 wt.%, in particular to a maximum of 0.015 wt.%. The S content of a flat steel product according to the invention is limited to a maximum of 0.02 wt.%, preferably to a maximum of 0.0020 wt.%, in particular to a maximum of 0.0010 wt.%.

Nitrogen ("N") is also present in small amounts as an impurity in steel due to the steelmaking process. The N content should be kept as low as possible and should not exceed 0.02 wt.%. Nitrogen is particularly harmful to alloys containing boron, as it inhibits the transformation-retarding effect of boron by forming boron nitrides. Therefore, the nitrogen content in this case should preferably not exceed 0.010 wt.%, and in particular not exceed 0.007 wt.%.

Other typical impurities are tin (“Sn”), arsenic (“As”), cobalt (“Co”), and antimony (“Sb”). The Sn content is a maximum of 0.03 wt.%, preferably a maximum of 0.02 wt.%. The As content is a maximum of 0.010 wt.%, in particular a maximum of 0.005 wt.%. The Co content is a maximum of 0.01 wt.%, in particular a maximum of 0.005 wt.%. The Sb content is a maximum of 0.02 wt.%, in particular a maximum of 0.01 wt.%, in particular a maximum of 0.005 wt.%.

In addition to the impurities P, S, N, Sn, As, and Sb explained above, other elements may also be present as impurities in the steel. These additional elements are summarized under the term "unavoidable impurities." The total content of these "unavoidable impurities" is preferably a maximum of 0.2 wt.%, preferably a maximum of 0.1 wt.%. The optional alloying elements Cr, Mn, Cu, Mo, Ni, V, Ti, Ca, and W described above and below, for which a lower limit is specified, may also be present as unavoidable impurities in the steel substrate in contents below the respective lower limit. In this case, they are also counted as "unavoidable impurities," whose total content is limited to a maximum of 0.2 wt.%, preferably a maximum of 0.1 wt.%.

Chromium, manganese, cobalt, copper, molybdenum, nickel, calcium and tungsten can optionally be alloyed to the steel of a flat steel product according to the invention, either individually or in combination with one another. Chromium (“Cr”) suppresses the formation of ferrite and pearlite during accelerated cooling of a flat steel product according to the invention and enables complete martensite formation even at lower cooling rates, thereby increasing hardenability.

These effects occur from a content of 0.01 wt.% upwards, whereby a content of at least 0.08 wt.%, in particular of at least 0.10 wt.%, preferably at least 0.20 wt.%, has proven to be effective in practice for reliable process control. However, excessive Cr contents impair the coatability of the steel. Therefore, the Cr content of the steel or the steel substrate is limited to a maximum of 1.0 wt.%, in particular a maximum of 0.75 wt.%, preferably a maximum of 0.50 wt.%.

Molybdenum (Mo) can optionally be added to improve process stability, as it significantly slows ferrite formation. Starting at a content of 0.002 wt.%, dynamic molybdenum-carbon clusters form along the grain boundaries, extending to ultrafine molybdenum carbides. These clusters significantly slow grain boundary mobility and thus diffusive phase transformations. Furthermore, molybdenum reduces grain boundary energy, which slows the nucleation rate of ferrite. The Mo content is preferably at least 0.004 wt.%, in particular at least 0.01 wt.%. Due to the high costs associated with a molybdenum alloy, the content should be at most 0.5 wt.%, in particular at most 0.30 wt.%, preferably at most 0.10 wt.%, and in particular at most 0.05 wt.%.

Chromium and molybdenum also form precipitates in a similar manner to the microalloying elements Nb, V, and Ti, but with a lesser effect. Therefore, the addition of one or both of these elements in the above-mentioned amounts supports the formation of precipitates with the dispersion according to the invention.

Manganese (“Mn”) acts as a hardening element by significantly delaying the formation of ferrite and bainite. At manganese contents of less than 0.2 wt.%, significant amounts of ferrite and bainite are formed during press hardening, even at very rapid cooling rates, which should be avoided. Mn contents of at least 0.5 wt.%, preferably at least 0.8 wt.%, in particular of at least 1.0 wt.%, particularly preferably of at least 1.10 wt.%, are advantageous if a martensitic microstructure is to be ensured, particularly in areas of greater deformation. Manganese contents of more than 3.3 wt.% have a detrimental effect on the Processing properties are affected, which is why the Mn content of flat steel products according to the invention is limited to a maximum of 3.3 wt.%, preferably a maximum of 2.5 wt.%. Weldability, in particular, is severely restricted, which is why the Mn content is preferably limited to a maximum of 1.6 wt.%, and in particular to a maximum of 1.40 wt.%, preferably to a maximum of 1.3 wt.%. Manganese contents of less than or equal to 1.6 wt.% are also preferred for economic reasons.

Copper (Cu) can optionally be added to the alloy to increase hardenability with additions of at least 0.01 wt.%, preferably at least 0.010 wt.%, in particular at least 0.015 wt.%. Furthermore, copper improves the resistance to atmospheric corrosion of uncoated sheets or cut edges. If the Cu content is too high, hot-rollability deteriorates significantly due to low-melting Cu phases on the surface, which is why the Cu content is limited to a maximum of 0.2 wt.%, preferably a maximum of 0.1 wt.%, in particular a maximum of 0.10 wt.%.

Nickel (Ni) stabilizes the austenitic phase and can optionally be added to the alloy to reduce the Ac3 temperature and suppress the formation of ferrite and bainite. Nickel also has a positive influence on hot rollability, particularly when the steel contains copper. Copper impairs hot rollability. Furthermore, it is known from EP 3 175 006 A1 that significant nickel contents can lead to a near-surface layer with an increased Ni content, which also hinders the penetration of free hydrogen. The measures described there can therefore preferably be combined with the improvements explained here to further reduce the free hydrogen content. Therefore, 0.01 wt.% nickel can be added to the steel; the Ni content is preferably at least 0.010 wt.%, in particular at least 0.020 wt.%. For economic reasons, the nickel content should be limited to a maximum of 2.0 wt.%, preferably a maximum of 1.0 wt.%, in particular a maximum of 0.60 wt.%. Furthermore, the Ni content is preferably a maximum of 0.5 wt.%, in particular a maximum of 0.50 wt.%.

Calcium (Ca) is used in steels to form non-metallic inclusions, particularly manganese sulfides. The rounded shape significantly reduces the negative effect of the inclusions on hot formability, fatigue strength, and toughness. To utilize this effect in a flat steel product according to the invention, a flat steel product according to the invention can optionally contain at least 0.0005 wt.% Ca, in particular at least 0.0010 wt.% % by weight, preferably at least 0.0020 wt. %. The maximum Ca content is 0.01 wt. %, in particular a maximum of 0.007 wt. %, preferably a maximum of 0.005 wt. %. Excessively high Ca contents increase the likelihood of non-metallic inclusions involving Ca forming, which impair the purity of the steel and also its toughness. For this reason, an upper limit of the Ca content of a maximum of 0.005 wt. %, preferably a maximum of 0.003 wt. %, should be maintained.

Tungsten (W) can optionally be added to the alloy in amounts of 0.001–1.0 wt.% to slow ferrite formation. A positive effect on hardenability is already achieved at W contents of at least 0.001 wt.%. For cost reasons, a maximum of 1.0 wt.%, and in particular a maximum of 0.30 wt.%, of tungsten is added.

In a preferred variant of the flat steel product, the elements Mo, Cr, Ti, V and Nb are present in the steel and the following applies to the element contents:

0,7 0.23

0.7

The elements Mo, Cr, Ti, V and Nb are the main contributors to the precipitates. The geometric mean of the contents (in wt. %) of these elements given here is particularly suitable for describing the superposition of properties which change according to a linear relationship depending on the concentration. Since the carbon content is also important for the formation of carbides, (Mo ■ Cr - Ti - V ■ Nb)s is particularly suitable for describing carbide distributions. It has been shown that particularly suitable carbide distributions are achieved when the above relationship is satisfied. The value c Mo ■ Cr - Ti - V ■ Nb)s is preferably greater than 0.30, in particular greater than 0.35. It is also preferred that it is less than 0.6, in particular less than 0.55.

In another preferred variant of the flat steel product, the element contents are as follows:

Ni + Cr < 1.1 Due to health concerns and, in particular, the European Union's REACH regulation, the Ni and Cr contents must be kept low. Therefore, the above ratio is advantageously met. In particular, the sum of the contents is less than 0.7 wt.%, and especially less than 0.5 wt.%.

The flat steel product preferably comprises a corrosion protection coating on at least one side to protect the steel substrate from oxidation and corrosion during hot forming and during use of the produced sheet metal part.

In a specific embodiment, the flat steel product preferably comprises an aluminum-based anti-corrosive coating. The anti-corrosive coating can be applied to one or both sides of the flat steel product. The two large, opposing surfaces of the flat steel product are referred to as the two sides of the flat steel product. The narrow surfaces are referred to as the edges.

Such a corrosion protection coating is preferably produced by hot-dip coating the flat steel product. The flat steel product is passed through a liquid melt consisting of 0.1-15 wt.% Si, preferably more than 1.0 wt.% Si, optionally 2-4 wt.% Fe, optionally up to 5 wt.% alkali or alkaline earth metals, preferably up to 1.0 wt.% alkali or alkaline earth metals, and optionally up to 15 wt.% Zn, preferably up to 10 wt.% Zn and optionally further components, the total contents of which are limited to a maximum of 2.0 wt.%, with aluminum as the remainder. The optional content of alkali or alkaline earth metals is preferably at least 0.1 wt.%.

In a preferred variant, the Si content of the melt is 0.5-3.5 wt% or 7-12 wt%, in particular 8-10 wt%.

In a preferred variant, the optional content of alkali or alkaline earth metals in the melt comprises 0.1-1.0 wt.% Mg, in particular 0.1-0.7 wt.% Mg, preferably 0.1-0.5 wt.% Mg. Furthermore, the optional content of alkali or alkaline earth metals in the melt can comprise in particular at least 0.0015 wt.% Ca, in particular at least 0.01 wt.% Ca. Further preferably, the optional content of alkali or alkaline earth metals in the melt consists of 0.1-1.0 wt.% Mg, in particular 0.1-0.7 wt.% Mg, preferably 0.1-0.5 wt.% Mg and optionally at least 0.0015 wt.% Ca, in particular at least 0.01 wt.% Ca. During hot-dip coating, iron diffuses from the steel substrate into the liquid coating, so that the corrosion protection coating of the flat steel product has, in particular, an alloy layer and an Al base layer upon solidification.

The alloy layer lies on the steel substrate and is directly adjacent to it. The alloy layer is essentially made of aluminum and iron. The alloy layer preferably consists of 25-50 wt.% Fe, 5-20 wt.% Si, optional further components whose total content is limited to a maximum of 5.0 wt.%, preferably 2.0 wt.%, and the remainder aluminum. The optional further components include, in particular, the remaining components of the melt (i.e., optionally alkali or alkaline earth metals, in particular Mg or Ca) and the remaining components of the steel substrate in addition to iron. In a further variant (variant with an Si content in the melt of 0.5-3.5 wt.%), the alloy layer consists of 25-50 wt.% Fe, 0.5-5.0 wt.% Si, optional further components whose total content is limited to a maximum of 5.0 wt.%, preferably 2.0 wt.%, and the remainder aluminum. The optional additional components also include in particular the remaining components of the melt (i.e. alkali or alkaline earth metals, in particular Mg or Ca) and the remaining components of the steel substrate in addition to iron.

In a preferred variant of the flat steel product, the degree of dispersion D ₂ of precipitates in the alloy layer is:

1 1

25 ■ 10 ^-6 — < D ₂ < 5 ■ IO ^-3 — nm nm

The precipitates are in particular carbides and/or carbonitrides of one or more of the elements from the group vanadium, titanium, niobium, chromium and molybdenum.

The adjustment of the precipitates in the steel substrate of the flat steel product results in a corresponding degree of dispersion D ₂ in the preferably ferritic alloy layer. Here, too, the precipitates act as traps for free hydrogen. Free hydrogen penetrating from the outside is thus localized at the precipitates. This causes the free hydrogen to collect in the alloy layer. This is preferably ferritic and thus softer than the steel substrate. Hydrogen embrittlement therefore does not occur as quickly, since the free hydrogen collects in the softer alloy layer and does not enter the more brittle steel substrate.

The Al base layer lies on top of the alloy layer and directly adjoins it. The composition of the Al base layer preferably corresponds to the composition of the melt of the molten bath. This means that it consists of 1.0-15 wt.% Si, optionally 2-4 wt.% Fe, optionally 5 wt.% alkali or alkaline earth metals, preferably up to 1.0 wt.% alkali or alkaline earth metals, optionally up to 15 wt.% Zn and optionally further components, the total contents of which are limited to a maximum of 2.0 wt.%, and the remainder being aluminum. Preferred compositions of the Al base layer correspond to the preferred melt compositions.

In a preferred variant of the Al base layer, the optional content of alkali or alkaline earth metals comprises 0.1-1.0 wt.% Mg, in particular 0.1-0.7 wt.% Mg, preferably 0.1-0.5 wt.% Mg. Furthermore, the optional content of alkali or alkaline earth metals in the Al base layer can in particular comprise at least 0.0015 wt.% Ca, in particular at least 0.1 wt.% Ca. Further preferably, the optional content of alkali or alkaline earth metals consists of 0.1-1.0 wt.% Mg, in particular 0.1-0.7 wt.% Mg, preferably 0.1-0.5 wt.% Mg and optionally at least 0.0015 wt.% Ca, in particular at least 0.1 wt.% Ca.

In a further preferred variant of the corrosion protection coating, the Si content in the alloy layer is lower than the Si content in the Al base layer.

trägt das Auflagengewicht des Korrosionsschutzüberzuges 100-200^ bei beidseitigen Überzügen bzw. 50- für einseitige Überzüge. Besonders bevorzugt beträgt das Auflagengewicht des Korrosionsschutzüberzuges 120-180^ bei beidseitigen Überzügen bzw. 60-90^ für einseitige Überzüge. Die Dicke der Legierungsschicht ist bevorzugt kleiner als 20 pm, besonders bevorzugt kleiner 16 pm, besonders bevorzugt kleiner 12 pm, insbesondere kleiner 10 pm. Die Dicke der Al-Basisschicht ergibt sich aus der Differenz der Dicken von Korrosionsschutzüberzug und Legierungsschicht. Bevorzugt beträgt die Dicke der Al-Basisschicht auch bei dünnen Korrosionsschutzüberzügen mindestens 1 pm. The corrosion protection coating preferably has a thickness of 5-60 μm, in particular 10-40 μm. The coating weight of the corrosion protection coating is in particular 30-360 μm for double-sided corrosion protection coatings or 15-180 μm for the single-sided variant.

The coating weight of the corrosion protection coating is 100-200^ for double-sided coatings or 50^ for single-sided coatings. The coating weight of the corrosion protection coating is particularly preferably 120-180^ for double-sided coatings or 60-90^ for single-sided coatings. The thickness of the alloy layer is preferably less than 20 μm, particularly preferably less than 16 μm, particularly preferably less than 12 μm, and especially less than 10 μm. The thickness of the Al base layer results from the difference between the thicknesses of the anti-corrosive coating and the alloy layer. The thickness of the Al base layer is preferably at least 1 μm, even with thin anti-corrosive coatings.

In a preferred variant, the flat steel product comprises an oxide layer arranged on the corrosion protection coating. The oxide layer is located in particular on the aluminum base layer and preferably forms the outer edge of the corrosion protection coating.

The oxide layer consists in particular of more than 80 wt.% oxides, with the majority of the oxides (i.e., more than 50 wt.% of the oxides) being aluminum oxide. Optionally, in addition to aluminum oxide, hydroxides and/or magnesium oxide are present in the oxide layer, alone or as a mixture. Preferably, the remainder of the oxide layer not occupied by the oxides and optionally present hydroxides consists of silicon, aluminum, iron, and/or magnesium in metallic form. For the optional embodiment with zinc as a component of the aluminum base layer, zinc oxide components are also present in the oxide layer.

Preferably, the oxide layer of the flat steel product has a thickness greater than 50 nm. In particular, the thickness of the oxide layer is a maximum of 500 nm.

In an alternative design, the flat steel product includes a zinc-based corrosion protection coating. The corrosion protection coating can be applied to one or both sides of the flat steel product. The two large, opposing surfaces of the flat steel product are referred to as the two sides. The narrow surfaces are referred to as the edges.

Such a zinc-based corrosion protection coating preferably comprises 0.2-6.0 wt.% Al, 0.1-10.0 wt.% Mg, optionally 0.1-40 wt.% manganese or copper, optionally 0.1-10.0 wt.% cerium, optionally at most 0.2 wt.% other elements, unavoidable impurities, and the remainder zinc. In particular, the Al content is a maximum of 2.0 wt.%, preferably a maximum of 1.5 wt.%. The Mg content is in particular a maximum of 3.0 wt.%, preferably a maximum of 1.0 wt.%. The corrosion protection coating can be applied by hot-dip coating or by physical vapor deposition or by electrolytic processes. The above explanations regarding element contents and their preferred limits apply accordingly to the process described below for producing a flat steel product, for the sheet metal part and for the process for producing a sheet metal part.

The method according to the invention for producing a flat steel product for hot forming (optionally with a corrosion protection coating) comprises the following working steps: a) Providing a slab or a thin slab made of steel which, in addition to iron and unavoidable impurities (in wt. %), consists of

C: 0.12-0.30%,

Si: 0.02-1.2%,

AI: 0.01-1.0%,

B: 0.0005-0.01%,

P: < 0.05%,

S: < 0.02%,

N: < 0.02%,

Sn: < 0.03%,

As: < 0.010%,

Sb: < 0.02%, at least one of the elements from the group comprising Ti, Nb and V, with the proviso that the contents are as follows:

Ti: 0.008-0.10%,

Nb: 0.01-0.08%,

V: 0.01-0.4%, and optionally one or more of the elements “Cr, Mn, Cu, Mo, Ni, Ca, W” in the following contents:

Cr: 0.01-1.0%,

Mn: 0.2-3.3%,

Cu: up to 0.2%,

Mo: 0.002-0.5%,

Ni: 0.01-2.0%,

Ca: 0.0005-0.01%, W: 0.001-1.0%; b) through-heating the slab or thin slab at a temperature (TI) of 1100-1400 °C; c) rough-rolling the through-heated slab or thin slab to a transfer strip with a transfer strip temperature (T2) of 800-1200 °C; d) hot-rolling the transfer strip to a hot-rolled flat steel product using a rolling mill, wherein before and during the hot-rolling of the transfer strip the local temperature of the transfer strip is adjusted so that the temperatures of the individual sections of the transfer strip do not vary by more than 60 K during hot rolling and wherein the final rolling temperature (T3) is 750-910 °C; e) cooling the hot-rolled flat steel product to a coiling temperature (T4) of 500-670 °C; f) coiling the hot-rolled flat steel product; g) descaling the hot-rolled flat steel product; h) optionally cold rolling the flat steel product, wherein the cold rolling degree is at least 30%; i) annealing the flat steel product at an annealing temperature (T5) of 650-900 °C; j) cooling the flat steel product to an intermediate temperature (T6) which is 650-800 °C, preferably 670-800 °C; k) optionally coating the flat steel product cooled to the intermediate temperature with a corrosion protection coating by hot-dip coating in a molten bath with a melt temperature (T7) of 660-800 °C, preferably 680-740 °C; l) cooling the flat steel product to room temperature, wherein the first cooling time T in the temperature range between 600 °C and 450 °C is more than 10 s, in particular more than 14 s, and the second cooling time t _n in the temperature range between 400 °C and 300 °C is more than 8 s, in particular more than 12 s; m) optionally skin-passing the flat steel product. In step a), a slab or thin slab composed according to the alloy specified for the flat steel product according to the invention is provided. This is done by means of conventional continuous slab casting or thin slab casting.

In step b), the slab or thin slab is thoroughly heated to a temperature (TI) of 1100-1320 °C. If the slab or thin slab has cooled after casting, it is first reheated to 1100-1320 °C for thorough heating. The thorough heating temperature should be at least 1100 °C to ensure good formability for the subsequent rolling process. The thorough heating temperature (TI) is preferably at least 1200 °C, more preferably at least 1250 °C. The thorough heating temperature should not exceed 1320 °C, preferably not exceeding 1300 °C. This avoids the presence of molten phases in the slab or thin slab. On the other hand, it has been shown that the dispersion D _± of the precipitates is too high when thorough heating is carried out at higher temperatures. Furthermore, higher temperatures lead to even greater reactions with the environment (e.g., decarburization in the near-surface area). Grain growth also occurs in the slab, resulting in poorer and more uneven product properties. Furthermore, the resulting increased scaling would also reduce yield.

In step c), the slab or thin slab is pre-rolled into a roughing strip. In this case, the temperature of the roughing strip (T2) at the end of roughing should be at least 800 °C so that the intermediate product contains sufficient heat for the subsequent finish-rolling step. However, high rolling temperatures can also promote grain growth during the rolling process, which has a detrimental effect on the mechanical properties of the flat steel product. To keep grain growth during the rolling process low, the temperature of the intermediate product at the end of roughing should not exceed 1200 °C. Preferably, the roughing strip temperature (T2) is at least 850 °C. Further preferably, the roughing strip temperature is a maximum of 1150 °C, in particular a maximum of 1100 °C, preferably a maximum of 1050 °C. For the purposes of this application, the roughing strip temperature (T2) refers to the temperature at the beginning of the roughing strip upon entry into the (downstream) rolling mill.

In step d), the pre-rolled strip is rolled into a hot-rolled flat steel product using a rolling mill. Finish rolling preferably begins no later than 90 seconds after the end of pre-rolling. Before and during hot rolling of the pre-rolled strip, the local temperature of the pre-rolled strip is measured. set so that the temperatures of the individual sections of the transfer strip do not vary by more than 60 K during hot rolling. To date, only one temperature for hot rolling has been specified in the literature. This is either the temperature at which the beginning of the transfer strip reaches during hot rolling or an average temperature of the transfer strip during hot rolling. However, since the strips are of significant length, large temperature differences can occur along the transfer strip. Even if a transfer strip that is uniformly tempered to a certain temperature enters the rolling mill, only the beginning of the transfer strip is rolled at this temperature. While the beginning of the transfer strip is being rolled, significant cooling already occurs in the upstream sections of the transfer strip. This effect is generally known and leads to an increase in the required rolling force during finish rolling due to the falling temperature and thus to an increase in strength. However, such variations can be implemented without problems in modern rolling stands, so that it has not seemed necessary to compensate for this until now. However, the invention recognized that this effect can also impact the final material properties. For example, in the well-known finish rolling of a homogeneously tempered preliminary strip, a temperature gradient inevitably arises along the steel flat product (also referred to as the finished strip) after finish rolling. This is because the preliminary strip is significantly thicker before rolling and therefore cools more slowly than the finished strip after rolling. While the beginning of the strip is rolled first and then immediately cools quickly as the finished strip, the end of the strip is rolled last and therefore initially cools slowly until it is rolled. The beginning and end of the strip therefore experience different temperature control. Surprisingly, this affects the final properties of the entire strip. The reason is probably that the temperature distribution along the strip during subsequent coiling leads to a temperature gradient along the coil radius, which in turn influences the cooling behavior of the coil. As a result, the precipitates are distributed with a dispersion D _± outside the desired range, which is beneficial for the reduction of free hydrogen.

The above-mentioned problems are solved according to the invention by adjusting the local temperature of the roughing strip before and during hot rolling of the roughing strip so that the temperatures of the individual sections of the roughing strip do not vary by more than 60 K during hot rolling. This refers to the temperature that each section of the roughing strip has at the time at which this section is rolled. This tempering ensures that each section of the pre-strip has essentially the same temperature when rolled. This means that from this point onward, every section of the pre-strip also experiences the same temperature control. This results in the desired distribution of the precipitates.

Local temperature adjustment can be achieved through various temperature control measures. Possible temperature control measures include:

- Thermal insulation of the transfer strip before the rolling mill, preferably by guiding the transfer strip through an enclosed area;

- Sectional reheating of the pre-strip.

The thermal insulation of the transfer strip upstream of the rolling mill is preferably achieved by guiding the transfer strip through an enclosed area. A section of the strip path upstream of the rolling mill is thermally insulated by an enclosure to reduce cooling of the transfer strip in this area. In the case of slight cooling, the high thermal conductivity of the strip results in a relatively homogeneous temperature, so that the variation is less than 60 K. The enclosed area can be linear or consist of several linear segments. Alternatively, this can also be achieved by first coiling the transfer strip into a spiral, and then holding this spiral within thermal insulation to compensate for temperature differences.

The section-by-section reheating of the pre-strip can be realized in particular as follows:

- Local, particularly inductive, reheating of the pre-strip before the rolling mill;

- Varying the forming speed during rough rolling or hot rolling so that the energy introduced by the forming varies.

It is known that rolling not only reduces the thickness of a strip, but also introduces local heat energy. The heat energy introduced depends on the forming speed during rolling. By varying the forming speed along the strip, more heat energy can be introduced in specific sections of the strip. In this case, for example, the forming speed during pre-rolling can be varied to achieve a targeted temperature distribution that compensates for the subsequent cooling behavior and thus ensures that temperatures do not fluctuate excessively during hot rolling. Likewise, the forming speed can be varied directly during hot rolling to compensate for the previous cooling behavior.

Thermal insulation of the transfer strip before the rolling mill, local reheating of the transfer strip before the rolling mill, and variation of the forming speed during pre-rolling ensure that the transfer strip already has a relatively homogeneous temperature upon entering the rolling mill. This means that the temperatures of the individual sections of the transfer strip upon entering the rolling mill do not vary by more than 60 K. This refers to the local temperature along the passing transfer strip at a fixed point, namely the entry to the rolling mill.

In contrast, by varying the forming speed during hot rolling, the inhomogeneous temperature distribution is at least partially compensated directly during hot rolling.

The final rolling temperature (T3), i.e., the temperature of the finished hot-rolled flat steel product at the end of the hot rolling process, is 750-910 °C. The final rolling temperature (T3) is preferably at least 800 °C, in particular at least 825 °C. The final rolling temperature is limited to values of no more than 910 °C to prevent coarsening of the austenite grains. Furthermore, final rolling temperatures of no more than 910 °C are relevant from a process engineering perspective for setting coiler temperatures (T4) below 670 °C.

In a preferred variant, the ratio T2/T3 of the pre-strip temperature T2 to the final rolling temperature T3 is at least 1.0. Furthermore, the ratio T2/T3 is preferably at most 1.35, preferably at most 1.25, in particular at most 1.15.

In the following step e), the hot-rolled flat steel product is cooled to a coiling temperature (T4) of 500-670 °C.

Step f) involves coiling the hot-rolled flat steel product. For this purpose, the flat steel product is cooled after hot rolling to a coiling temperature (T4), particularly within less than 50 seconds. The cooling medium used for this purpose can be water, air, or a combination of both. The coiling temperature (T4) should not exceed 670 °C. preferably a maximum of 650 °C. In principle, there is no lower limit on the coiling temperature. However, coiling temperatures of at least 550 °C have proven favorable for cold rolling. The coiled flat steel product is then cooled to room temperature in air using the conventional method.

In step g), the hot-rolled flat steel product is descaled in a conventional manner by pickling or other suitable treatment.

The scale-cleaned hot-rolled flat steel product can optionally be subjected to cold rolling prior to annealing in step h), for example, to meet higher thickness tolerance requirements for the flat steel product. The cold rolling degree (KWG) should be at least 30% to inject sufficient deformation energy into the flat steel product for rapid recrystallization. The cold rolling degree KWG is defined as the quotient of the thickness reduction during cold rolling (AdKW) divided by the hot strip thickness d:

KWG = AdKW/d where AdKW = thickness reduction during cold rolling in mm and d = hot strip thickness in mm, whereby the thickness reduction AdKW results from the difference between the thickness of the flat steel product before cold rolling and the thickness of the flat steel product after cold rolling. The flat steel product before cold rolling is usually a hot strip with a hot strip thickness of d. The flat steel product after cold rolling is usually also referred to as cold strip. The cold rolling degree can, in principle, assume very high values of over 90%. However, cold rolling degrees of no more than 80% have proven to be advantageous for preventing strip breakage.

The further process is described below in a first variant in which the flat steel product is coated:

In step i), the flat steel product is subjected to an annealing treatment at annealing temperatures (T5) of 650-900 °C. For this purpose, the flat steel product is first heated to the annealing temperature within 10 to 120 seconds and then held at the annealing temperature for 30 to 600 seconds. The annealing temperature is at least 650 °C, preferably at least 720 °C. Annealing temperatures above 900 °C are undesirable for economic reasons. In step j), the flat steel product is cooled to an intermediate temperature (T6) after annealing to prepare it for subsequent coating treatment. The intermediate temperature T6 can also be referred to as the immersion temperature T6. The intermediate temperature is lower than the annealing temperature and is adjusted to the temperature of the molten bath (T7). The intermediate temperature is 600-800 °C, preferably at least 660 °C, more preferably at least 670 °C, more preferably at most 740 °C, in particular at most 700 °C. For particularly homogeneous boundary layer formation, it is important that sufficient thermal energy is present in the boundary layer between the steel substrate and the aluminum melt. This is not the case at temperatures lower than 600 °C, so that undesirable compounds can form, the subsequent reconversion of which can lead to pores. From the preferred intermediate temperatures, the diffusion rate of iron into aluminum increases significantly again, allowing increased iron diffusion into the still-liquid boundary layer right at the beginning of the coating process. The cooling time of the annealed flat steel product from the annealing temperature T5 to the intermediate temperature T6 is preferably 10-180 s. In particular, the intermediate temperature T6 deviates from the temperature of the molten bath T7 by no more than 30 K, in particular no more than 20 K, and preferably no more than 10 K.

The flat steel product is subjected to a coating treatment in step k). The coating treatment is preferably carried out by continuous hot-dip coating. The coating can be applied to just one side, both sides, or all sides of the flat steel product. The coating treatment is preferably carried out as a hot-dip coating process, in particular as a continuous process. The flat steel product usually comes into contact with the molten bath on all sides, so that it is coated on all sides. The molten bath, which contains the alloy to be applied to the flat steel product in liquid form, typically has a temperature (T7) of 660-800 °C, preferably 680-740 °C. Aluminum-based alloys have proven particularly suitable for coating ageing-resistant flat steel products with a corrosion-protective coating. In such a case, the molten bath contains 0.1-15 wt.% Si, preferably more than 1.0%, optionally 2-4 wt.% Fe, optionally up to 5 wt.% alkali or alkaline earth metals, preferably up to 1.0% wt.% alkali or alkaline earth metals, and optionally up to 15% Zn and optional further components, the total contents of which are limited to a maximum of 2.0 wt.%, and the remainder aluminum. In a preferred variant, the Si content of the melt is 1.0-3.5 wt.% or 7-12 wt.%, in particular 8-10 wt.%. In a preferred variant, the optional content of alkali or alkaline earth metals in the melt comprises 0.1-1.0 wt.% Mg, in particular 0.1-0.7 wt.% Mg, preferably 0.1-0.5 wt.% Mg. Furthermore, the optional content of alkali or alkaline earth metals in the melt can comprise, in particular, at least 0.0015 wt.% Ca, in particular at least 0.01 wt.% Ca. Further variants of the melt were explained above in connection with the flat steel product.

After the coating treatment, the coated flat steel product is cooled to room temperature in step I). A first cooling time t _m y in the temperature range between 600 °C and 450 °C (medium temperature range mT) is more than 10 s, in particular more than 14 s, and a second cooling time t _n in the temperature range between 400 °C and 300 °C (low temperature range nT) is more than 8 s, in particular more than 12 s.

The first cooling time t _mi can be achieved in the temperature range between 600 °C and 450 °C (medium temperature range mT) by slow, continuous cooling or by holding the temperature at a certain time within this temperature range. Intermediate heating is even possible. The only important thing is that the flat steel product remains in the temperature range between 600 °C and 450 °C for at least a cooling time t _m y. In this temperature range, on the one hand, there is a significant diffusion rate of iron into aluminum and, on the other hand, the diffusion of aluminum into steel is inhibited because the temperature is below half the melting temperature of steel. This allows diffusion of iron into the corrosion protection coating without strong diffusion of aluminum into the steel substrate.

The diffusion of iron into the corrosion protection coating has several advantages:

Firstly, the melting of the corrosion protection coating is delayed during austenitization prior to press hardening. Secondly, the thermal expansion coefficients of the corrosion protection coating and the substrate are homogenized. This means that the transition area between the substrate's thermal expansion coefficient and the surface becomes wider, which reduces thermal stresses during reheating.

At the same time, the diffusion of aluminum into the steel substrate would have significant disadvantages: Due to the very high affinity of aluminum to nitrogen, a high aluminum content can lead to nitrogen dissolving from fine precipitates, such as niobium carbonitrides or titanium carbonitrides, and instead coarse precipitates, such as aluminum nitrides, preferentially depositing on the Grain boundaries would form. These would impair crash performance and reduce the bending angle. Furthermore, this would destabilize the fine precipitates (e.g., niobium-containing precipitates) in the uppermost substrate region, which, according to the invention, are important for reducing free hydrogen. Furthermore, the inhomogeneous diffusion rate of aluminum in the steel substrate in ferrite compared to pearlite/bainite/martensite would lead to an uneven distribution of Al in the surface layer of the steel substrate. This should also be prevented to improve crash and bending performance. These disadvantages of aluminum diffusion into the steel substrate are therefore reduced or avoided by inhibition.

Due to the preferred initial cooling time t _my (more than 14 s), the iron concentration in the transition boundary layer increases to such an extent that the activity of aluminum in the coating directly at the substrate boundary is further reduced. This then leads to an even further reduced aluminum uptake into the substrate during austenitization prior to press hardening, with the associated advantages described above.

The second cooling time t _n in the temperature range between 400 °C and 300 °C (low temperature range nT) can also be achieved by slow, continuous cooling or by holding the product at a temperature within this temperature range for a certain period of time. Intermediate heating is even possible. The only important thing is that the flat steel product remains in the temperature range between 400 °C and 300 °C for at least the cooling time t _n .

With the preferred second cooling time t _n y of more than 12s, very fine iron carbides (so-called transition carbides) are also formed, which in turn dissolve very quickly during austenitization and lead to additional austenite nuclei and thus an even finer austenite structure and thus also a hardening structure.

The coated flat steel product can optionally be subjected to skin passing with a skin passing degree of up to 2% in the subsequent step m) in order to improve the surface roughness of the flat steel product.

As an alternative to the described process variant, the flat steel product can also be manufactured in an uncoated version. In this case, process steps i)-m) are designed as follows: In step i), the flat steel product is subjected to an annealing treatment at annealing temperatures (T5) of 650-900 °C. For this purpose, the flat steel product is first heated to the annealing temperature within 10 to 120 seconds and then held at the annealing temperature for 30 to 600 seconds. The annealing temperature is at least 650 °C, preferably at least 720 °C. Annealing temperatures above 900 °C are undesirable for economic reasons.

In step j), the flat steel product is cooled to an intermediate temperature (T6) after annealing. The cooling time of the annealed flat steel product from the annealing temperature T5 to the intermediate temperature T6 is preferably 10-180 seconds.

Work step k) is not implemented in this uncoated variant.

Subsequently, the flat steel product is cooled to room temperature in step I). A first cooling time t _m y in the temperature range between 600 °C and 450 °C (medium temperature range mT) is more than 10 s, in particular more than 14 s, and a second cooling time tn? in the temperature range between 400 °C and 300 °C (low temperature range nT) is more than 8 s, in particular more than 12 s.

The initial cooling time tmT can be achieved in the temperature range between 600 °C and 450 °C (medium temperature range mT) through slow, continuous cooling or by holding at a temperature within this temperature range for a certain period of time. Intermediate heating is even possible. The only important thing is that the flat steel product remains in the temperature range between 600 °C and 450 °C for at least the cooling time tmT.

The second cooling time t _n T in the temperature range between 400 °C and 300 °C (low temperature range n T) can also be achieved by slow, continuous cooling or by holding at a temperature within this temperature range for a certain period of time. Intermediate heating is even possible. The only important thing is that the flat steel product remains in the temperature range between 400 °C and 300 °C for at least the cooling time t _n T.

With the preferred second cooling time t _n T of more than 12s, very fine iron carbides (so-called transition carbides) are formed, which in turn degrade very quickly during austenitization. dissolve and lead to additional austenite nuclei and thus an even finer austenite structure and thus also a hardening structure.

The uncoated flat steel product thus obtained can optionally be subjected to skin passing with a skin passing degree of up to 2% in the subsequent step m) in order to improve the surface roughness of the flat steel product.

The invention further relates to a sheet metal part formed from a flat steel product comprising a steel substrate as described above and optionally a corrosion protection coating. The corrosion protection coating has the advantage of preventing scale formation during austenitization during hot forming. Furthermore, such a corrosion protection coating protects the formed sheet metal part against corrosion.

For the sheet metal part, as well as for the flat steel product, the following applies for a degree of dispersion Di of precipitates in the near-surface third of the steel substrate:

25 ■ IO“ ⁶ — < Di < 5 ■ IO“ ³ — . nm

In the forming processes explained below, the temperatures and holding times are not high enough to cause a significant change in the precipitates. Consequently, this property of the flat steel product, which was established during hot rolling, is passed on to the formed sheet metal part. The precipitates can therefore assume the desired function in the formed sheet metal part and reduce free hydrogen.

In a specific embodiment, the sheet metal part preferably comprises an aluminum-based corrosion protection coating. The corrosion protection coating of the sheet metal part preferably comprises an alloy layer and an aluminum base layer. In the sheet metal part, the alloy layer is also often referred to as an interdiffusion layer.

The thickness of the corrosion protection coating is preferably at least 10 pm, particularly preferably at least 20 pm, in particular at least 30 pm. The thickness of the alloy layer is preferably less than 30 pm, particularly preferably less than 20 pm, in particular less than 16 pm, particularly preferably less than 12 pm. The thickness of the Al base layer results from the difference between the thicknesses of the corrosion protection coating and the alloy layer.

The alloy layer lies on the steel substrate and is directly adjacent to it. The alloy layer of the sheet metal part preferably consists of 35-90 wt.% Fe, 0.1-12 wt.% Si and optional further components, the total contents of which are limited to a maximum of 3.5 wt.%, preferably 2.0 wt.%, with aluminum as the remainder. The optional further components are preferably the elements present in the steel of the steel substrate alongside iron and the remaining elements from the melt, such as Zn and alkali or alkaline earth metals. These elements from the melt only accumulate in the alloy layer to a very small extent.

The alloy layer preferably has a ferritic structure in the area close to the substrate.

The Al base layer of the sheet metal part lies on top of the alloy layer and directly borders it. The Al base layer of the sheet metal part preferably consists of 35-55 wt.% Fe, 0.4-10 wt.% Si, optionally up to 3 wt.% alkali or alkaline earth metals, preferably up to 1.0 wt.% alkali or alkaline earth metals, optionally up to 10% Zn, and optional further components, the total contents of which are limited to a maximum of 2.0 wt.%, with aluminum as the remainder. The optional content of alkali or alkaline earth metals is preferably at least 0.1 wt.%.

In a preferred variant of the Al base layer, the optional content of alkali or alkaline earth metals comprises 0.1-1.0 wt.% Mg, in particular 0.1-0.7 wt.% Mg, preferably 0.1-0.5 wt.% Mg. Furthermore, the optional content of alkali or alkaline earth metals in the Al base layer can in particular comprise at least 0.0015 wt.% Ca, in particular at least 0.1 wt.% Ca. Further preferably, the optional content of alkali or alkaline earth metals consists of 0.1-1.0 wt.% Mg, in particular 0.1-0.7 wt.% Mg, preferably 0.1-0.5 wt.% Mg and optionally at least 0.0015 wt.% Ca, in particular at least 0.1 wt.% Ca.

The Al base layer can have a homogeneous element distribution, with local element contents varying by no more than 10%. Preferred variants of the Al base layer, however, have silicon-poor phases and silicon-rich phases. Silicon-poor phases are areas whose average Si content is at least 20% less than the average Si content of the Al Base layer. Silicon-rich phases are areas whose average Si content is at least 20% higher than the average Si content of the Al base layer.

In a preferred variant, the silicon-rich phases are arranged within the silicon-poor phase. In particular, the silicon-rich phases form at least a 40% continuous layer bordered by silicon-poor regions. In an alternative embodiment, the silicon-rich phases are arranged in island-like patterns within the silicon-poor phase.

For the purposes of this application, “island-shaped” means an arrangement in which discrete, unconnected areas are enclosed by another material - i.e., “islands” of a particular material are located within another material.

In a preferred variant, the sheet metal part comprises an oxide layer arranged on the corrosion protection coating. The oxide layer is located in particular on the aluminum base layer and preferably forms the outer edge of the corrosion protection coating.

The oxide layer of the sheet metal part consists, in particular, of more than 80 wt.% oxides, with the majority of the oxides (i.e., more than 50 wt.% of the oxides) being aluminum oxide. Optionally, in addition to aluminum oxide, hydroxides and/or magnesium oxide are present in the oxide layer, alone or as a mixture. Preferably, the remainder of the oxide layer not occupied by the oxides and optionally present hydroxides consists of silicon, aluminum, iron, and/or magnesium in metallic form.

The oxide layer preferably has a thickness of at least 50 nm, in particular of at least 100 nm. Furthermore, the thickness is a maximum of 4 pm, in particular a maximum of 2 pm.

In a special design, the sheet metal part includes a zinc-based anti-corrosion coating.

Such a zinc-based corrosion protection coating preferably comprises up to 80 wt.% Fe, 0.2-6.0 wt.% Al, 0.1-10.0 wt.% Mg, optionally 0.1-40 wt.% manganese or copper, optionally 0.1-10.0 wt.% cerium, optionally at most 0.2 wt.% other elements, unavoidable impurities, and the remainder zinc. In particular, the Al content is a maximum of 2.0 wt.%, preferably a maximum of 1.5 wt.%. The Fe content, which results from diffusion, is preferably more than 20 wt.%, in particular more than 30 wt.%. In addition, the Fe content is in particular a maximum of 70 wt.%, in particular a maximum of 60 wt.%. The Mg content is in particular a maximum of 3.0 wt.%, preferably a maximum of 1.0 wt.%. The corrosion protection coating can be applied by hot-dip coating, by physical vapor deposition, or by electrolytic processes.

In a specific development, the steel substrate of the sheet metal part has a structure with at least partially more than 80% martensite, preferably at least partially more than 90% martensite, in particular at least partially more than 95%, particularly preferably at least partially more than 98%. In this context, "partially having" means that there are regions of the sheet metal part that have the mentioned structure. In addition, there may also be regions of the sheet metal part that have a different structure. The sheet metal part therefore has the mentioned structure in sections or in regions.

The high martensite content allows very high tensile strengths and yield points to be achieved.

In an alternative development, the steel substrate of the sheet metal part has a microstructure with a ferrite content of more than 5%, preferably more than 10%, in particular more than 20%. Furthermore, the ferrite content is preferably less than 85%, in particular less than 70%. The martensite content is less than 80%, in particular less than 50%. In addition, the microstructure can optionally contain bainite and/or pearlite. The exact ratio of the microstructure components depends on the level of the C content and the Mn content as well as on the cooling conditions during forming. The microstructure designed in this way has greater ductility and therefore leads to improved forming behavior. A corresponding sheet metal part therefore preferably has an elongation at break A30 in a range of 8% to 25%, preferably between 10% and 22%, in particular between 12% and 20%.

In a further developed variant, the sheet metal part has at least partially a tensile strength of at least 1100 MPa, in particular at least 1250 MPa, preferably at least 1400 MPa.

In particular, the sheet metal part has at least partially an elongation at break A80 of at least 4%, preferably at least 5%, particularly preferably at least 6%. In addition, in a preferred variant, the sheet metal part can at least partially have a bending angle of at least 50°, in particular at least 55°, preferably at least 60°. The bending angle here is understood to be the bending angle corrected for the sheet thickness. The corrected bending angle results from the determined bending angle at the maximum force (measured according to VDA standard 238-100) (also referred to as the maximum bending angle) from the formula

_Corrected bending angle = _determined bending angle /sheet thickness, where the sheet thickness in mm must be entered into the formula. This applies to sheet thicknesses greater than 1.0 mm. For sheet thicknesses less than 1.0 mm, the corrected bending angle corresponds to the determined bending angle.

In this context, "partially exhibit" means that there are areas of the sheet metal part that exhibit the stated mechanical property. In addition, there may also be areas of the sheet metal part whose mechanical properties are below the limit value. The sheet metal part therefore exhibits the stated mechanical property in sections or regions. This is because different areas of the sheet metal part can undergo different heat treatments. For example, individual areas can be cooled more quickly than others, resulting in more martensite, for example, forming in the faster-cooled areas. This also results in different mechanical properties in the different areas.

The mechanical characteristics mentioned have proven to be particularly advantageous in ensuring use in a vehicle with good crash performance.

In a preferred embodiment, the sheet metal part has a free hydrogen content (Hdiff) of not more than 0.30 ppm, preferably not more than 0.25 ppm, in particular not more than 0.20 ppm.

In a particularly preferred variant, the sheet metal part has an aluminum-based corrosion protection coating comprising an alloy layer and an Al base layer. The following applies to a degree of dispersion D ₂ of precipitates in the alloy layer: 1 1

25 ■ IO ^-6 - < D ₂ < 5 ■ 10“ ³ - nm nm

By adjusting the precipitates in the steel substrate of the flat steel product, a corresponding degree of dispersion D ₂ is also present in the preferably ferritic alloy layer. This is particularly advantageous because it allows incoming hydrogen to be trapped in the soft alloy layer before it penetrates the steel substrate. This further reduces the risk of hydrogen embrittlement, as the hydrogen cannot penetrate the hard substrate, which is therefore susceptible to embrittlement.

The sheet metal part according to the invention is preferably a component for a land vehicle, marine vehicle, or aircraft. It is particularly preferably an automotive part, in particular a body part. The component is preferably a B-pillar, longitudinal member, A-pillar, sill, or cross member.

In the method according to the invention for producing a shaped sheet metal part according to the invention, which is designed in the manner explained above, at least the following work steps are carried out: a) providing a sheet metal blank from a previously described flat steel product; b) heating the sheet metal blank in such a way that the AC3 temperature of the blank is at least partially exceeded and the temperature T _{E inig} of the blank when inserted into a forming tool intended for hot press forming (work step c)) at least partially has a temperature above Ms+100°C, where Ms denotes the martensite start temperature; c) inserting the heated sheet metal blank into a forming tool, wherein the transfer time t _T rans required for removing the blank from the heating device and inserting it is at most 20s, preferably at most lös; d) hot-press forming the sheet metal blank into the formed sheet metal part, wherein the blank is cooled during the hot-press forming over a period twz of more than 1 s at a cooling rate r _W z which is at least partially more than 30 K/s to the target temperature T _Z iei and is optionally held there; e) removing the formed sheet metal part cooled to the target temperature Tziei from the tool. In the method according to the invention, a blank is provided which consists of a steel suitably composed in accordance with the above explanations (working step a)), which is then heated in a manner known per se such that the AC3 temperature of the blank is at least partially exceeded and the temperature T _{E inig} of the blank when placed in a forming tool intended for hot press forming (working step c)) is at least partially above Ms+100 °C. Partially exceeding a temperature (here AC3 or Ms+100 °C) is understood in the context of this application to mean that at least 30%, in particular at least 60%, of the volume of the blank exceeds a corresponding temperature. When placed in the forming tool, therefore, at least 30% of the blank has an austenitic structure, i.e. the transformation from the ferritic to the austenitic structure does not have to be complete when placed in the forming tool. Rather, up to 70% of the volume of the blank when placed in the forming tool can consist of other microstructure components, such as tempered bainite, tempered martensite and/or non- or partially recrystallized ferrite. For this purpose, certain areas of the blank can be deliberately kept at a lower temperature level than others during heating. To do this, the heat supply can be specifically directed only to certain sections of the blank, or the parts that are to be heated less can be shielded from the heat supply. In the part of the blank material whose temperature remains lower, no or only significantly less martensite is formed during forming in the tool, so that the microstructure there is significantly softer than in the other parts that have a martensitic microstructure. In this way, a softer area can be deliberately set in the respective formed sheet metal part, for example by ensuring optimal toughness for the respective intended use, while the other areas of the sheet metal part have maximized strength.

Maximum strength properties of the obtained sheet metal part can be achieved by ensuring that the temperature reached at least partially in the sheet metal blank is between Ac3 and 1000 °C, preferably between 850 °C and 950 °C.

The minimum temperature Ac3 to be exceeded is determined according to the formula given by HOUGARDY, HP. in Werkstoffkunde Stahl, Volume 1: Grundlagen, Verlag Stahleisen GmbH, Düsseldorf, 1984, p. 229:

Ac3 = (902 - 225*%C + 19*%Si - ll*%Mn - 5*%Cr + 13*%Mo - 20*%Ni + 55*%V) °C with %C = respective C content, %Si = respective Si content, %Mn = respective Mn content, %Cr = respective Cr content, %Mo = respective Mo content, %Ni = respective Ni content and %V = respective V content of the steel from which the blank is made.

An optimally uniform distribution of properties can be achieved by heating the blank completely in step b).

In a preferred embodiment, the average heating rate _r of the sheet metal blank during heating in step b) is at least 3 K/s, preferably at least 5 K/s, in particular at least 10 K/s, preferably at least 15 K/s. The average heating rate r is to be understood as the average heating rate from 30 °C to 700 °C.

In a preferred embodiment, the heating takes place in a furnace with a furnace temperature of at least 850 °C, preferably at least 880 °C, particularly preferably at least 900 °C, in particular at least 920 °C, and at most 1000 °C, preferably at most 950 °C, particularly preferably at most 930 °C.

The dew point of the furnace atmosphere in the furnace is preferably at least -20 °C, preferably at least -15 °C, in particular at least -5 °C, particularly preferably at least 0 °C and at most +25 °C, preferably at most +20 °C, in particular at most +15 °C.

In a special embodiment, the heating in step b) takes place stepwise in areas with different temperatures. In particular, the heating takes place in a roller hearth furnace with different heating zones. Here, the heating takes place in a first heating zone at a temperature (so-called furnace inlet temperature) of at least 650 °C, preferably at least 680 °C, in particular at least 720 °C. The maximum temperature in the first heating zone is preferably 900 °C, in particular a maximum of 850 °C. Furthermore, the maximum temperature of all heating zones in the furnace is preferably a maximum of 1200 °C, in particular a maximum of 1000 °C, preferably a maximum of 950 °C, particularly preferably a maximum of 930 °C.

The total time in the oven, which consists of a heating time and a holding time, is preferably at least at least 2 minutes, in particular at least 3 minutes, preferably at least 4 minutes. Furthermore, the total furnace time for both variants is preferably a maximum of 20 minutes, in particular a maximum of 15 minutes, preferably a maximum of 12 minutes, in particular a maximum of 8 minutes. Longer total furnace times have the advantage of ensuring uniform austenitization of the sheet metal blank. On the other hand, holding the furnace for too long above Ac3 leads to grain coarsening, which has a negative impact on the mechanical properties.

The blank heated in this way is removed from the respective heating device, which can be, for example, a conventional heating furnace, an equally known induction heating device or a conventional device for keeping sheet metal parts hot, and transported into the forming tool so quickly that its temperature upon arrival in the tool is at least partially above Ms+100 °C, preferably above 600 °C, in particular above 650 °C, particularly preferably above 700 °C. Here, Ms denotes the martensite start temperature. In a particularly preferred variant, the temperature is at least partially above the ACl temperature. In all of these variants, the temperature is in particular a maximum of 900 °C. These temperature ranges ensure good formability of the material overall.

In step c), the transfer of the austenitized blank from the heating device used to the forming tool is completed within preferably a maximum of 20 seconds, in particular within a maximum of 15 seconds. Such rapid transport is necessary to avoid excessive cooling prior to forming.

When the blank is inserted, the tool typically has a temperature between room temperature (RT) and 200 °C, preferably between 20 °C and 180 °C, in particular between 50 °C and 150 °C. Optionally, in a special embodiment, the tool can be tempered at least partially to a temperature Twz of at least 200 °C, in particular at least 300 °C, in order to only partially harden the component. Furthermore, the tool temperature Twz is preferably a maximum of 600 °C, in particular a maximum of 550 °C. It only needs to be ensured that the tool temperature Twz is below the desired target temperature Tziei. The residence time in the tool twz is preferably at least 2s, in particular at least 3s, particularly preferably at least 5s. The maximum residence time in the tool is preferably 25s, in particular a maximum of 20s. The target temperature Tziei of the sheet metal part is at least partially below 400°C, preferably below 300°C, in particular below 250°C, preferably below 200°C, particularly preferably below 180°C, in particular below 150°C. Alternatively, the target temperature Tael of the sheet metal part is particularly preferably below Ms-50°C, where Ms denotes the martensite start temperature. Furthermore, the target temperature of the sheet metal part is preferably at least 20°C, particularly preferably at least 50°C.

The martensite start temperature of a steel within the scope of the invention is according to the formula:

Ms [°C] = (490.85 - 302.6 %C - 30.6 %Mn - 16.6 %Ni - 8.9 %Cr + 2.4 %Mo - 11.3 %Cu + 8.58 %Co + 7.4 %W - 14.5 %Si) [°C/wt.%], where C% is the C content, %Mn is the Mn content, %Mo is the Mo content, %Cr is the Cr content, %Ni is the Ni content, %Cu is the Cu content, %Co is the Co content, %W is the W content and %Si is the Si content of the respective steel in wt.%.

The ACl temperature and the AC3 temperature of a steel within the scope of the invention specifications are according to the formulas:

AC1[°C] = (739 - 22*%C - 7*%Mn + 2*%Si + 14*%Cr + 13*%Mo - 13*%Ni + 20*%V) [°C/wt%] and

AC3[°C] = (902 - 225*%C + 19*%Si - 11*%Mn - 5*%Cr + 13*%Mo - 20*%Ni +55*%V) [°C/wt.%], where %C denotes the C content, %Si the Si content, %Mn the Mn content, %Cr the Cr content, %Mo the Mo content, %Ni the Ni content and +%V the vanadium content of the respective steel (Brandis H 1975 TEW-Techn. Ber. 1, pp. 8-10).

In the tool, the blank is not only formed into the sheet metal part, but is also quenched to the target temperature at the same time. The cooling rate in the tool r _W z to the target temperature is in particular at least 20 K/s, preferably at least 30 K/s, in particular at least 50 K/s, in a special embodiment at least 100 K/s.

After removal of the sheet metal part in step e), the sheet metal part is cooled to a cooling temperature TAB of less than 100 °C within a cooling time _tAB of 0.5 to 600 s. This is usually done by air cooling.

In the following, the invention is explained in more detail using exemplary embodiments.

To demonstrate the effectiveness of the invention, several tests were conducted. Slabs with the compositions listed in Table 1, thicknesses of 200-280 mm and widths of 1000-1200 mm, were produced, heated in a pusher-type furnace to a respective temperature TI, and held at TI for between 30 and 450 minutes until the temperature TI in the core of the slabs was reached and the slabs were thus thoroughly heated. The production parameters are listed in Table 2. The slabs were discharged from the pusher-type furnace at their respective through-heating temperatures TI and subjected to hot rolling. The tests were carried out as continuous hot strip rolling. For this purpose, the slabs were first pre-rolled to an intermediate product with a thickness of 40 mm. The intermediate products, which in hot strip rolling can also be referred to as pre-strips, were fed to the finish rolling immediately after pre-rolling. The pre-strip temperature T2 corresponds to the temperature of the strip start during hot rolling. In all cases except test 10, the local temperature of the transfer strip was adjusted so that the temperatures of the individual sections of the transfer strip did not vary by more than 60 K during hot rolling. Table 2 shows AT2 as the difference between the maximum and minimum temperature of the individual sections of the transfer strip as the corresponding sections entered the rolling mill. In test 10, no such tempering measure was used, whereupon the temperature along the strip dropped with a temperature variation of 95 K. In all other cases, suitable tempering measures were used so that the temperature variation did not exceed 44 K. Specifically, the transfer strips were passed through an enclosed area that reduced cooling. The transfer strips were rolled into hot strips with a final thickness of 3-7 mm and the respective final rolling temperatures T3 specified in Table 2, cooled to the coiling temperature T4 and wound into coils at the respective coiling temperatures T4 and then cooled in still air. The hot strips were descaled in a conventional manner by pickling before They were subjected to cold rolling with the cold rolling grades specified in Table 2. The thickness of the steel strips produced was between 1.0 mm and 2.0 mm in all tests.

The cold-rolled flat steel products 1-12 were heated in a continuous annealing furnace to a respective annealing temperature T5 (see Table 3) and held at annealing temperature Tö for 100 s each before being cooled to their respective intermediate temperature T6 at a cooling rate of 1 K/s. The cold-rolled strips were passed through a molten coating bath at temperature T7 at their respective intermediate temperature T6 (or immersion temperature T6). The composition of the coating bath is given in Table 3. After coating, the coated strips were blown off in a conventional manner, creating deposits with varying layer thicknesses (see Table 3). The strips were first cooled to 600 °C at an average cooling rate of 10-15 K/s. During the further cooling process between 600 °C and 450 °C and between 400 °C and 300 °C, the strips were cooled for the cooling times T _m y and T _n specified in Table 2. Between 450 °C and 400 °C and below 220 °C, the strips were cooled at a cooling rate of 5-15 K/s.

For comparison, steel flat products 1, 6, and 8 were subjected to a corresponding annealing treatment in a continuous annealing furnace without coating. These are tests 13, 14, and 15. Therefore, the manufacturing conditions for strips 1 and 13, as well as 6 and 14, and 8 and 15 in Table 2 are identical. Cold-rolled steel flat products 13-15 were heated in a continuous annealing furnace to their respective annealing temperatures Tö (see Table 3) and held at annealing temperature T5 for 100 s each before being cooled to their respective intermediate temperatures T6 at a cooling rate of 1 K/s. The strips were then cooled to 600 °C at an average cooling rate of 10-15 K/s. In the further cooling process between 600 °C and 450 °C and between 400 °C and 300 °C, the strips were cooled for the cooling times T _m y and T„T specified in Table 2. Between 450 °C and 400 °C and below 220 °C, the strips were cooled at a cooling rate of 5-15 K/s each.

After cooling to room temperature, longitudinal sections were taken of the resulting flat steel products. In longitudinal sections, the plane of the section is perpendicular to the strip surface and parallel to the main deformation direction (rolling direction). These longitudinal sections depicted the entire thickness of the flat steel product (substrate and corrosion protection coating). Carbon pull-out impressions were taken from these longitudinal sections. To determine D _± , These carbon pull-out impressions were made in the near-surface third. The carbon pull-out impressions were examined in a transmission electron microscope (TEM) and evaluated using computer-assisted image analysis.

For variants 1-9 and 13-15 according to the invention, the results were all between 25 ■ IO ^-6 — and 5 ■ IO ^-3 — . nm nm

In comparison variant 10, the precipitation process varies locally due to the large temperature fluctuation at the pre-strip AT2. This results in a heterogeneous dispersion. Furthermore, large amounts of precipitation-forming elements also lead to a strong driving force for precipitation formation. Furthermore, the temperature variation AT2 leads to a different cooling behavior of the coiled coil. As a result, a high dispersion occurs throughout the strip. Both this high dispersion and its non-uniformity jointly lead to a deterioration in the mechanical properties, particularly the low bending angle, of the final sheet metal part.

In comparison variant 11, the temperature TI is too low, so that an insufficient amount of the mostly coarse precipitates already present in the slab could be dissolved. Therefore, the driving force for precipitation formation and thus the amount of precipitates is low. Furthermore, coarse, undissolved precipitates remain, increasing the mean diameter of the precipitates. This leads to insufficient dispersion and insufficient hydrogen binding in the sheet metal part.

In comparison variant 12, the temperature T2 is so high that almost all of the precipitates present in the slab are dissolved. Due to the large amount of precipitate-forming elements (e.g., demonstrated by the value for (Mo • Cr - Ti - V • c

Nb)s), the driving force for precipitation formation is high, resulting in the formation of a large amount of predominantly fine precipitates and thus a high degree of dispersion. While this has a positive effect on the very low diffusible hydrogen content, it also results in a deterioration in the mechanical properties of the sheet metal part, which can be seen in the bending angle. For the comparison variants 10-12, the non-inventive entries are underlined in the tables.

From the 15 steel strips produced in this way, blanks were cut and used for further tests. The data for this is given in Table 4. In these tests, sheet metal part samples 1-15 in the form of 200 x 300 ^mm2 plates were hot-pressed from the respective blanks. For this purpose, the blanks were heated in a heating device, for example in a conventional heating furnace, from room temperature at an average heating rate r _{of furnace} (between 30 °C and 700 °C) in a furnace with a furnace temperature Toten. The total time in the furnace, which includes heating and holding, is designated toten. The dew point of the furnace atmosphere is given in Table 4. The blanks were then removed from the heating device and placed in a forming tool which has the temperature Twz. At the time of removal from the furnace, the blanks had reached the furnace temperature. The transfer time t _Tr ans, comprising the time required for removal from the heating device, transport to the tool and insertion into the tool, was 8 s. The temperature T _{E inig} of the blanks when inserted into the forming tool was above the respective martensite start temperature of +100 °C in all cases. The blanks were formed into the respective sheet metal parts in the forming tool, with the sheet metal parts being cooled in the tool at a cooling rate r _W z. The residence time in the tool is designated twz. Finally, the samples were cooled in air to room temperature. Table 4 shows the parameters mentioned for various variants, where “RT” stands for room temperature.

Table 5 summarizes the overall results for the obtained sheet metal parts. The first columns indicate the sample number and steel grade according to Table 1. The remaining columns indicate the free hydrogen content, the bending angle, and the tensile strength. Methods for determining the free hydrogen content are known to those skilled in the art, for example, thermal desorption mass spectrometry (TDMS) with heated samples.

In the uncoated samples, the amount of free hydrogen was below the detection limit. The degree of dispersion D _± was comparable in the uncoated samples 13-15 to the respective coated samples 1, 6, and 8 due to the comparable annealing treatment. Because the same inventive distribution of the precipitates was observed in the uncoated The resulting hydrogen trapping effect is also achieved by the samples. This results in improved resistance to hydrogen embrittlement under corrosion conditions when using the sheet metal part.

wobei die Blechdicke in mm in die Formel einzusetzen ist. Die Zugfestigkeit wurde gemäß DIN EN ISO 6892-1 Probenform 2 (Anhang B Tab. Bl) an Proben quer zur Walzrichtung ermittelt. The maximum bending angle was determined according to VDA standard 238-100 with a bending axis perpendicular to the rolling direction. The maximum bending angle is calculated from the punch travel according to the formula specified in the standard (the maximum bending angle is the bending angle at which the force reaches its maximum in the bending test). To eliminate the influence of sheet thickness on the bending angle, the corrected bending angle was calculated from the maximum bending angle according to the formula:

where the sheet thickness in mm is to be inserted into the formula. The tensile strength was determined according to DIN EN ISO 6892-1, specimen shape 2 (Appendix B, Table B1) on specimens transverse to the rolling direction.

After cooling to room temperature, longitudinal sections were taken of the sheet metal parts produced in this way. In longitudinal sections, the section plane is perpendicular to the strip surface and parallel to the main deformation direction (rolling direction). These longitudinal sections showed the entire thickness of the sheet metal part (substrate and corrosion protection coating). Carbon pull-out impressions were taken from these longitudinal sections. To determine D _±, these carbon pull-out impressions were taken in the third closest to the surface. To determine D _2, the carbon pull-out impressions were taken in the alloy layer. The carbon pull-out impressions were examined under a transmission electron microscope (TEM) and evaluated using computer-assisted image analysis.

For the samples 1-9 according to the invention, the degree of dispersion D ₂ was between 25 ■ 1°~ ⁶ “ and

5 ■ IO ^-3 — (see Table 5). nm

The uncoated samples naturally do not have an alloy layer, so the degree of dispersion D ₂ is not defined. The degree of dispersion D was found to be consistent, within the measurement uncertainty, with the degree of dispersion D ₁ of the flat steel product before forming. Therefore, the described forming process does not result in any significant change in the degree of dispersion D ₁ . This is presumably due to the fact that the heating is not high enough or does not last long enough to lead to the dissolution of the precipitates. The degree of dispersion D _± is therefore not shown again in Table 5, but corresponds to the values from Table 3.

Furthermore, the microstructure of the sheet metal part was determined in each case. In all cases, a martensitic microstructure was found, with a martensite content of more than 95% by area.

The remainder is iron and unavoidable impurities. All values are in wt.%.

* non-inventive reference examples Table 1 (steel grades)

Table 2 (Manufacturing conditions for flat steel products)

Table 3 (Coating manufacturing conditions)

Table 4 (Hot forming parameters)

Table 5 (Sheet metal parts)

Claims

Patent claims 1. A steel flat product for hot forming, comprising a steel substrate made of steel which, in addition to iron and unavoidable impurities (in % by weight), consists of: C: 0.12-0.30%, Si: 0.02-1.2%, AI: 0.01-1.0%, B: 0.0005-0.01%, P: <0.05%, S: <0.02%, N: <0.02%, Sn: <0.03%, As: <0.010%, Sb: <0.02%, at least one of the elements from the group comprising Ti, Nb and V with the proviso that the following applies to the contents: Ti: 0.008-0.10%, Nb: 0.01-0.08%, V: 0.01-0.4%, and optionally one or more of the elements “Cr, Mn, Cu, Mo, Ni, Ca, W” in the following contents: Cr: 0.01-1.0%, Mn: 0.2-3.3%, Cu: up to 0.2%, Mo: 0.002-0.5%, Ni: 0.01-2.0%, Ca: 0.0005-0.01%, W: 0.001-1.0%, where the dispersion degree Di of precipitates in the near-surface third of the steel substrate is: 1 1 25 ■ 10“ ⁶ — < D < 5 ■ 10“ ³ — nm nm 2. Steel flat product according to claim 1, wherein the elements Mo, Cr, Ti, V and Nb are present in the steel and wherein the element contents are: 0.23

0.7 3. Flat steel product according to one of claims 1 to 2, wherein the element contents are: Ni + Cr < 1.1 4. Flat steel product according to one of claims 1 to 3, wherein the element contents are: V + Nb + Ti < 0.13 5. Flat steel product according to one of the preceding claims, characterized in that it has a corrosion protection coating on at least one side. 6. Steel flat product according to claim 5, characterized in that the corrosion protection coating is an aluminum-based corrosion protection coating and has an alloy layer and an Al base layer. 7. Flat steel product according to claim 6, characterized in that the alloy layer consists of 25-50 wt.% Fe, 5-20 wt.% Si, optional further components whose total contents are limited to a maximum of 5.0 wt.%, and the remainder aluminum and/or the Al base layer consists of 1.0-15 wt.% Si, optionally 2-4 wt.% Fe, optionally up to 5.0 wt.% alkali or alkaline earth metals, optionally up to 10% Zn and optional further components whose total contents are limited to a maximum of 2.0 wt.%, and the remainder aluminum. 8. Steel flat product according to one of claims 6 to 7, characterized in that for a degree of dispersion D ₂ of precipitates in the alloy layer, the following applies: 1 1 25 ■ 10 ^-6 — < D ₂ < 5 ■ IO ^-3 — nm nm 9. A method for producing a flat steel product for hot forming, comprising the following steps: a) Provision of a slab or a thin slab made of steel which, in addition to iron and unavoidable impurities (in % by weight), consists of: C: 0.12-0.30%, Si: 0.02-1.2%, AI: 0.01-1.0%, B: 0.0005-0.010%, P: < 0.05%, S: < 0.02%, N: < 0.02%, Sn: < 0.03%, As: < 0.010%, Sb: < 0.02%, at least one of the elements from the group comprising Ti, Nb and V, with the proviso that the contents are as follows: Ti: 0.008-0.10%, Nb: 0.01-0.08%, V: 0.01-0.4%, and optionally one or more of the elements “Cr, Mn, Cu, Mo, Ni, Ca, W” in the following contents: Cr: 0.01-1.0%, Mn: 0.2-3.3%, Cu: up to 0.2%, Mo: 0.002-0.5%, Ni: 0.01-2.0%, Ca: 0.0005-0.01%, W: 0.001-1.0%; b) heating the slab or thin slab at a temperature (TI) of 1100-1320 °C; c) rough rolling the heated slab or thin slab into a pre-strip with a pre-strip temperature (T2) of 800-1200 °C; d) hot rolling of the preliminary strip into a hot-rolled flat steel product by means of a rolling mill, wherein before and during the hot rolling of the preliminary strip the local temperature of the preliminary strip is adjusted so that the temperatures of the individual sections of the preliminary strip do not vary by more than 60 K during hot rolling and wherein the final rolling temperature (T3) is 750-910 °C; e) cooling the hot-rolled flat steel product to a coiling temperature (T4) of 500-670 °C; f) coiling the hot-rolled flat steel product; g) descaling the hot-rolled flat steel product; h) optionally cold rolling the flat steel product, wherein the cold rolling degree is at least 30%; i) annealing the flat steel product at an annealing temperature (T5) of 650-900 °C; j) Cooling the flat steel product to an intermediate temperature (T6) which is 650-800 °C, preferably 670-800 °C; k) Optionally coating the flat steel product cooled to the intermediate temperature with a corrosion protection coating by hot-dip coating in a molten bath with a melt temperature (T7) 660-800 °C, preferably 680-740 °C; l) Cooling the flat steel product to room temperature, wherein the first cooling time tmT in the temperature range between 600 °C and 450 °C is more than 10 s, in particular more than 14 s, and the second cooling time t _n y in the temperature range between 400 °C and 300 °C is more than 8 s, in particular more than 12 s; m) Optionally skin passing the flat steel product. 10. A process according to claim 9, wherein the elements Mo, Cr, Ti, V and Nb are present in the steel and wherein the element contents are: 0.03

0.4 11. A process according to any one of claims 9 to 10, wherein the element contents are: Ni + Cr < 1.1 12. A process according to any one of claims 9 to 11, wherein the element contents are: V + Nb + Ti < 0.13 13. A method according to any one of claims 9 to 12, characterized in that during hot-dip coating, a molten bath is used which contains the corrosion protection to be applied to the flat steel product in liquid form, which consists of 0.1-15 wt.% Si, optionally 2-4 wt.% Fe, optionally up to 5 wt.% alkali or alkaline earth metals and optionally up to 15 wt.% Zn, and optionally further components, the total contents of which are limited to a maximum of 2.0 wt.%, and the remainder being aluminum. 14. Sheet metal part formed from a flat steel product, comprising a steel substrate made of steel which, in addition to iron and unavoidable impurities (in % by weight), consists of: C: 0.12-0.30%, Si: 0.02-1.2%, AI: 0.01-1.0%, B: 0.0005-0.01%, P: < 0.05%, S: < 0.02%, N: < 0.02%, Sn: < 0.03%, As: < 0.010%, Sb: < 0.02%, at least one of the elements from the group comprising Ti, Nb and V, with the proviso that the contents are as follows: Ti: 0.008-0.10%, Nb:0.01-0.08%, V: 0.01-0.4%, and optionally one or more of the elements “Cr, Mn, Cu, Mo, Ni, Ca, W” in the following contents: Cr: 0.01-1.0%, Mn: 0.2-3.3%, Cu: up to 0.2%, Mo: 0.002-0.5%, Ni: 0.01-2.0%, Ca: 0.0005-0.01%, W: 0.001-1.0%, where the following applies to a degree of dispersion Di of precipitates in the near-surface third of the steel substrate: 25 ■ IO ^-6 — < D, < 5 ■ IO ^-3 —, nm nm and optionally a corrosion protection coating. 15. Sheet metal part according to claim 14, wherein the elements Mo, Cr, Ti, V and Nb are present in the steel and wherein the elements Mo, Cr, Ti, V and Nb are present in the steel and wherein the element contents are: 0.03

0.4 16. Sheet metal part according to one of claims 14 to 15, wherein the element contents are: Ni + Cr < 1.1 17. Sheet metal part according to one of claims 14 to 16, wherein the element contents are: V + Nb + Ti < 0.13 18. Sheet metal part according to one of claims 14 to 17, the sheet metal part at least partially has a tensile strength of at least 1100 MPa, preferably at least 1450 MPa. 19. Sheet metal part according to one of claims 14 to 18, the corrosion protection coating is an aluminum-based corrosion protection coating and comprises an alloy layer and an Al base layer. 20. Sheet metal part according to claim 19, characterized in that the alloy layer consists of 35-90 wt.% Fe, 0.1-12 wt.% Si and optional further components, the total contents of which amount to at most 3.5 wt.%, and the remainder being aluminum and/or the Al- Base layer consisting of 35-55 wt% Fe, 0.4-10 wt% Si, optionally 3 wt% alkali or alkaline earth metals, optionally up to 10% Zn and optional further components, the total contents of which are limited to a maximum of 2.0 wt%, and the remainder being aluminum. 21. Sheet metal part according to one of claims 19 to 20, characterized in that for a degree of dispersion D ₂ of precipitates in the alloy layer, the following applies: 1 1 25 ■ 10“ ⁶ — < D ₂ < 5 ■ 10“ ³ — nm nm 22. A method for producing a shaped sheet metal part, comprising the following work steps: a) providing a sheet metal blank from a flat steel product according to one of claims 1 to 8; b) heating the sheet metal blank such that the AC3 temperature of the blank is at least partially exceeded and the temperature T _{E inig} of the blank when inserted into a forming tool provided for hot press forming (work step c)) at least partially has a temperature above Ms+100 °C, where Ms denotes the martensite start temperature; c) inserting the heated sheet metal blank into a forming tool, wherein the transfer time t _T rans required for removing the blank from the heating device and inserting it is at most 20 s, preferably at most lös; d) hot-press forming the sheet metal blank to form the sheet metal part, wherein the blank is cooled during the hot-press forming process over a period twz of more than 1s at a cooling rate r _W z of at least partially more than 30 K/s to the target temperature Tael and is optionally held there; e) removing the sheet metal part cooled to the target temperature T _Z iei from the tool.