US20240141450A1

US20240141450A1 - High-strength hot-rolled flat steel product having high local cold formability and a method of producing such a flat steel product

Info

Publication number: US20240141450A1
Application number: US18/547,444
Authority: US
Inventors: Ingwer Denks; Djordje Mirkovic; Marco WITTE
Original assignee: Salzgitter Flachstahl GmbH
Current assignee: Salzgitter Flachstahl GmbH
Priority date: 2021-02-25
Filing date: 2022-02-24
Publication date: 2024-05-02
Also published as: WO2022180146A1; CN116888283A; KR20230148167A; EP4298255A1; MX2023007872A; DE102021104584A1

Abstract

The present disclosure provides a high-strength hot-rolled flat steel product and a method of producing such a product. This is achieved by a high-strength hot-rolled flat steel product with high local cold formability, having a tensile strength Rm of at least 760 MPa, a yield point ratio of at least 0.8 and a hole expansion ratio of at least 30%, an elongation at break of at least 10, a measure of cold formability of at least 0.12, and a ratio of local and global cold formability of at least 5 and at most 13, and a microstructure consisting of more than 50% by volume of bainite and up to 10% by volume, of carbon-rich microstructure constituents

Description

CROSS REFERENCE TO RELATED APPLICATION

The present application is a national stage application of and claims the priority benefits of International Patent Application No. PCT/EP2022/054616, filed Feb. 24, 2022, and claims benefit of German patent application no. 10 2021 104 584.3, filed on Feb. 25, 2021.

BACKGROUND AND FIELD OF THE INVENTION

The invention relates to a high-strength, hot-rolled flat steel product having high local cold-formability. Furthermore, the invention relates to a method for producing such a flat steel product.
In the present case, cold-formability is understood to mean the formability at a temperature in the range from 10° C. to 700° C., preferably between 10° C. and 200° C., particularly preferably between 10 and 80° C. and especially preferably at room temperature between 15 and 40° C.
In particular, the invention relates to a high-strength, microalloyed, predominantly bainitic hot strip having an optimised alloy composition and microstructure which is used e.g. as a chassis component in the automotive industry.
The invention also relates to high-strength hot strips having tensile strengths of at least 760 MPa with simultaneously high cold-formability.
The description of cold-formability is complex and can only be sufficiently quantified by an AND-linked combination of characteristic values.
Therefore, the following characteristic values are used to describe the cold-formability of the flat steel product in accordance with the invention:

- 1. Elongation at fracture (A)
- 2. Hole expansion ratio (LA)
- 3. Measure of cold-formability (FL)
- 4. Ratio of local to global cold-formability (LFR) Characteristic values, such as elongation at fracture, uniform elongation and hole expansion ratio are established characteristic values for describing cold-formability.

In order to quantify high global cold-formability and high local cold-formability, it is necessary to use a local cold-formability characteristic value and a global cold-formability characteristic value. Within the scope of the invention, the characteristic value of hole expansion ratio is selected as being representative of the local formability and the characteristic value of uniform elongation is selected as being representative of the global cold-formability. The true variables are indicated and not the technical (percentage) variables, the determination of which is indicated in the description of the exemplified embodiments.
The invention includes, in particular, flat steel products consisting of steels which have a multiphase microstructure which contains substantially, i.e. a proportion of more than 50 vol. %, bainite and which have a yield strength ratio of at least 0.8. In addition to a high tensile strength of at least 760 MPa and an elongation at fracture A of at least 10%, the flat steel product also has a high hole expansion capability with a hole expansion ratio LA of at least 30%, a measure of cold-formability FL of at least 0.12 and a ratio of local to global cold-formability LFR in the range of at least 5 and at most 13.
It is known that bainitic steels are steels which are characterised by a comparatively high yield strength and tensile strength with sufficiently high elongation for cold-forming processes. An effective welding capability is provided by reason of the chemical composition. The microstructure typically consists of bainite as the predominant constituent with proportions of ferrite. The microstructure can contain in isolation small proportions of other phases, such as e.g. martensite and residual austenite.
Such a steel is disclosed e.g. in the laid-open document DE 10 2012 002 079 A1. However, a disadvantage with this is a still insufficiently high hole expansion capability.
The fiercely competitive car market means that producers are constantly forced to find solutions for reducing fleet fuel consumption and CO2 exhaust emissions whilst maintaining the highest possible level of comfort and passenger protection. On the one hand, the weight reduction of all of the vehicle components plays a decisive role as does, on the other hand, the most favourable possible behaviour of the individual components under the high static and dynamic stresses both during use of an automobile and also in the event of a crash.
Through the provision of high-strength to ultra high-strength steels with strengths of up to 1050 MPa or more and through the reduction in the sheet thickness which can be achieved with these steels, it is possible to reduce the weight of vehicles whilst having, at the same time, improved deformation behaviour of the steels used during manufacture and in operation.
Therefore, high-strength steels must meet comparatively high demands with respect to their strength, ductility and energy absorption, without disadvantages occurring during processing thereof, such as e.g. during stamping, hot and cold forming, during thermal tempering (e.g air hardening, press hardening), welding and/or a surface treatment, e.g. a metallic finishing, organic coating or lacquering, in comparison with conventional steels.
Thus, in addition to the required weight reduction by reduced sheet thicknesses, newly developed steels must meet the increasing material requirements for elasticity limit, tensile strength, hardening behaviour and elongation at fracture while having good processing properties such as formability and weldability.
Therefore, in order to guarantee the required reduction in sheet thickness, a high-strength steel with a single-phase or multi-phase microstructure must be used in order to ensure sufficient strength for the motor vehicle components and in order to meet the high deformation and component demands with respect to toughness, lack of sensitivity to edge cracking, improved bending angle and bending radius, energy absorption.
Improved joining suitability, in the form of better general welding capability, expressed by a larger usable welding area for resistance spot welding and improved failure behaviour of the weld seam (fracture pattern) under mechanical stress, and sufficient resistance to delayed crack formation owing to hydrogen embrittlement are also required to an increasing extent.
The hole expansion capability is a material property which describes the resistance of the material to crack initiation and crack propagation in deformation operations in regions close to the edge and previously shear-cut regions, such as e.g. during plunging.
The hole expansion test is regulated e.g. in the ISO 16630 standard. Accordingly, holes which are stamped into a metal sheet are expanded by means of a mandrel.
The measurement variable is the change in hole diameter, related to the initial diameter, with respect to the diameter, at which the first crack through the metal sheet occurs at the edge of the hole.
Improved edge crack insensitivity signifies an increased deformation capability of the sheet edges and is described by an increased hole expansion capability. This situation is known under the synonyms “Low Edge Crack” (LEC) or “High Hole Expansion” (HHE) and Xpand®.
Patent specification EP 3 516 085 B1 discloses a method for producing a high-strength, hot-rolled steel strip having a tensile strength of at least 570 MPa, preferably at least 780 MPa, with which good cold-formability of the steel strip is to be achieved.
The method comprises the following steps:

- casting a slab, followed by the step of reheating the solidified slab to a temperature between 1050 and 1260° C.;
- hot rolling the steel slab at an entry temperature in the last final roll stand between 980 and 1100° C.;
- finish-rolling at a finish-rolling temperature between 950 and 1080° C.;
- cooling the hot-rolled steel strip at a primary cooling rate between 50 and 150° C./s to an intermediate temperature on the ROT (outlet roller conveyor) between 600 and 720° C. and followed by:
- mildly heating the steel between 0 and +10° C./s due to latent heat resulting from the phase conversion from austenite to ferrite; or
- isothermally holding the steel; or
- mildly cooling the steel, resulting overall in a rate of temperature change into the secondary stage of ROT of −20 to 0° C./s; —in order to reach the winding temperature between 580 and 660° C.

The steel disclosed therein consists of (in wt. %): between 0.015 and 0.15 C; at most 0.5 Si; between 1.0 and 2.0 Mn; at most 0.06 P; at most 0.008 S; at most 0.1 Al sol; at most 0.02 N; between 0.02 and 0.45 V; optionally one or more of: at least 0.05 and/or at most 0.7 Mo; at least 0.15 and/or at most 1.2 Cr; at least 0.01 and/or at most 0.1 Nb; optionally Ca in an amount consistent with calcium treatment for inclusion control; the remainder being Fe and unavoidable impurities; and wherein the steel has a substantially single-phase ferritic microstructure which contains a mixture of polygonal ferrite (PF) and needle-shaped/bainitic ferrite (AF/BF) and wherein the total volume proportion of the sum of the ferrite constituents is at least 95% and wherein the ferrite constituents are precipitation-hardened with fine composite carbides and/or carbonitrides of V and optionally of Mo and/or Nb.
However, it has been found that the cold forming capability, in particular the local cold-formability of this flat steel product, is not yet sufficiently high. Moreover, the alloy concept is comparatively cost-intensive.
Furthermore, patent specification EP 3 492 611 B1 discloses a method for producing hot-rolled steel having a tensile strength of at least 950 MPa and a microstructure which comprises bainite having an area ratio of 70% or more, wherein the difference is one or both of the following: martensite having an area ratio of 30% or less, and optionally ferrite having an area ratio of 20% or less, wherein the method comprises the steps of:

- heating steel with the chemical composition to a temperature of at least 1250° C.,
- hot rolling the steel at a final rolling temperature of 850-930° C.,
- quenching the steel to a reeling temperature of 450-575° C.,
- reeling the steel at the reeling temperature,
- cooling the steel, and
- skin pass rolling.

The steel has the following alloy composition in mass %:

- C 0.07-0.10, Si 0.01-0.25, Mn 1.5-2.0, Cr 0.5-1.0, Ni 0.1-0.5, Cu 0.1-0.3, Mo 0.01-0.2,
- Al 0.01-0.05, Nb 0.015-0.04, V 0-0.1, Ti 0-0.1, the difference being Fe and unavoidable impurities. The steel used is comparatively cost-intensive by reason of the addition by alloying of chromium, copper and nickel and likewise still does not have sufficiently high cold-formability. Local cold-formability will not be discussed.

SUMMARY OF THE INVENTION

The present invention provides a high-strength, hot-rolled flat steel product, as well as a method for producing such a flat steel product, and thus, in relation to the steel, achieving a combination of high strength with simultaneously high local cold-formability and a high level of economic viability.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 is a graph of a range of cold-formability; and

FIG. 2 is a schematic drawing of the positioning of the hardness impressions.

In accordance with an embodiment, a high-strength, hot-rolled flat steel product with high local cold-formability, having a tensile strength Rm of at least 760 MPa, a yield strength ratio of at least 0.8 and a hole expansion ratio of at least 30%, advantageously at least 40%, particularly advantageously at least 50%, an elongation at fracture of at least 10%, preferably at least 16%, a measure of cold-formability of at least 0.12, advantageously at least 0.17, and a ratio of local to global cold-formability of at least 5 and at most 13, and a microstructure consisting of more than 50 vol. % bainite, up to 10 vol. %, advantageously up to 5 vol. %, carbon-rich microstructure constituents, such as e.g. martensite, residual austenite, pearlite, residual austenite, with the remainder being precipitation-hardened ferrite, with the following chemical composition of the steel (in wt. %):

- C: 0.04 to 0.08
- Si: 0.1 to 0.6
- Mn: 1.0 to 2.0
- P: max. 0.06
- S: max. 0.01
- N: max. 0.012
- Al: up to 0.06
- Ti: up to 0.18 and/or Nb: up to 0.08
- Mo: up to 0.35

with Ti+Nb more than 0.06 wt. %, wherein a hyperstoichiometric proportion of carbon and nitrogen is present according to the following formula 1:
1.0<(C/12+N/14)/(Ti/48+Nb/93+Mo/96),

- with the remainder being iron including unavoidable, steel-associated elements, with optional addition by alloying of one or more elements of Cr, Ni, V, B or Ca,
- wherein the microstructure distribution over the thickness of the flat steel product in the three regions near the surface, ¼ thickness and ½ thickness of the flat steel product, is characterised by
  - an absolute deviation of a maximum of 12, advantageously a maximum of 7 vol. % of the proportion of ferrite in the regions near the surface or ¼ thickness of the flat steel product, in relation to the region ½ thickness of the flat steel product, and/or
  - the deviation in aspect ratio in the rolling direction in the three regions of the flat steel product with respect to the mean value is less than 0.3 in each of the three regions, and/or
  - in the three regions the difference in hardness HV0.1 is in each case a maximum of 20 HV 0.1, advantageously a maximum of 15 HV 0.1, even more advantageously a maximum of 10 HV 0.1 in comparison with the mean value over the entire thickness of the flat steel product, an excellent combination of strength, elongation and forming properties.

For the measurements of the microstructure distribution and the hardness over the thickness of the flat steel product, it is irrelevant from which surface side these measurements are conducted.
In the case of the optional addition by alloying of one or more of the elements Cr, Ni, V, B and Ca, provision is made in particular that Cr is added by alloying up to 0.6 wt. %, Ni up to 0.6 wt. %, V up to 0.2 wt. %, B: up to 0.01 wt. % and Ca up to 0.01 wt. %.
The microstructure consists preferably of more than 50 vol. % bainite and the remainder is precipitation-hardened ferrite.
In particular, the flat steel product is characterised by a combination of high strength and excellent cold-formability at the same time. Furthermore, the production of this flat steel product in accordance with the invention based on the alloy elements C, Si, Mn, Nb and/or Ti is comparatively inexpensive.
The flat steel product in accordance with the invention is specifically characterised by a high elongation at fracture A of at least 10%, a high hole expansion ratio (LA) of at least 30%, advantageously at least 40%, particularly advantageously at least 50%, a measure of cold-formability (FL) of at least 0.12, advantageously at least 0.17, and a ratio of local to global cold-formability (LFR) of at least 5 and at most 13, whilst at the same time having a high tensile strength of at least 760 MPa.
In one advantageous development of the invention, in order to set optimised combinations of properties, the steel alloy optionally also comprises one or more elements of Cr, Ni, V or B with the following contents in wt. %: Cr: more than 0.1 up to 0.6, Ni: more than 0.1 up to 0.6, V: more than 0.01 up to 0.2 and B: more than 0.0005 up to 0.01, wherein a hyperstoichiometric proportion of carbon and nitrogen is present according to formula 2 below:
1.0<(C/12+N/14)/(Ti/48+Nb/93+Mo/96+V/51)
In another advantageous development of the invention, Ca is added by alloying to the steel for inclusion control. As a result, the inclusions of MnS and Al₂O₃which are unfavourable with regard to the final properties are replaced by Ca-containing inclusions, which are less detrimental, in particular with regard to the morphology. The addition by alloying to the steel is a maximum of 0.01 wt. %.
In a further advantageous development of the invention, in order to achieve particularly favourable combinations of properties, the flat steel product contains the following alloy composition in wt. %:
Ti: at least 0.02, advantageously at least 0.04, even more advantageously at least 0.06, Nb: at least 0.01, Mo: at least 0.05 and Ti+Nb: up to 0.2.
The microstructure consists predominantly of bainite and, in smaller proportions, ferrite. The bainite is a mixture of constituents which is characterised by a main constituent of at least 50 vol. % and secondary constituents, wherein the main constituent consists of bainitic ferrite hardened by precipitations of (Ti, Nb, Mo)(C,N) or V(C,N) and the secondary constituents consist of carbon-richer constituents such as martensite, residual austenite, lower bainite, and pearlite. In an advantageous manner, the microstructure consists of more than 75 vol. % bainite.
In addition, the microstructure may contain carbon-rich microstructure constituents.
Particularly favourable properties are achieved only if the microstructure contains a maximum of 10%, advantageously a maximum of 5%, of carbon-rich microstructure constituents (such as e.g. martensite, residual austenite, pearlite).
It has likewise been proven to be advantageous if the grain extension of all microstructure constituents in the rolling direction, measured at a position at ½ thickness of the flat steel product below the surface of the flat steel product, characterised by the area-averaged aspect ratio of all microstructure constituents in the rolling direction of at most 2.0, and/or the mean value over the three regions near the surface, ¼ thickness and ½ thickness of the flat steel product is at most 2.0, advantageously 1.6.
It has likewise been proven to be advantageous for high cold-formability if half of the precipitations of (Ti, Nb, Mo)(C,N) or V(C,N) which harden the ferrite and the main constituent of bainitic ferrite have, on average, a diameter of less than 10 nm and/or the precipitations have an average spacing of less than 750 nm.
It is also advantageous if the ratio of shear texture components to rolling texture components increases towards the surface and has the following values:

- near the surface: at least 0.9
- ½ thickness of the flat steel product: a maximum of 0.1

The hot-rolled flat steel product in accordance with the invention can be provided with a metallic or non-metallic coat. The metallic coat can be applied to the flat steel product electrolytically or by means of hot-dipping and is advantageously zinc-based.
Such a hot-rolled flat steel product is advantageously used in the automotive industry for the production of components, in particular chassis components. Hot-rolled flat steel products in accordance with the invention have thicknesses of 1.6 to 6.0 mm. However, the invention includes thicknesses less than 1.6 mm or thicknesses greater than 6.0 mm.
In an advantageous manner, the flat steel product in accordance with the invention has, along the rolling direction, a tensile strength Rm of at least 760 MPa, a yield strength ratio of at least 0.8, an elongation at fracture A of at least 10%, preferably at least 16%, a hole expansion ratio of at least 30%, advantageously at least 40% or even at least 50%. The degree of cold-formability is at least 0.12, advantageously at least 0.17 with a ratio of local and global cold-formability of at least 5 and at most 13.
Alloy elements are generally added to the steel in order to influence specific properties in a targeted manner. An alloy element can thereby influence different properties in different steels. The effect and interaction generally depend greatly upon the quantity, presence of further alloy elements and the solution state in the material. The correlations are varied and complex. The effect of the alloy elements in the alloy in accordance with the invention will be discussed in greater detail hereinafter.
In the case of the numerical values indicated hereinafter and provided in the claims for alloy element contents and all other numerical values, the numbers are also to be included as limit values. The use of the term “to” in the content ranges, e.g. 0.01 to 1 wt. %, means that the limit values, in this case 0.01 and 1, are also included.
Carbon C: is required to form carbides, in particular in conjunction with the so-called microalloy elements Nb, V and Ti, requires the formation of martensite and bainite, stabilises the austenite and generally increases strength. Higher contents of C impair the welding properties and result in the impairment of the elongation and toughness properties, for which reason a maximum content of at most 0.08 wt. % is set. In order to achieve a sufficient strength for the material, a minimum addition of 0.04 wt. % is required.
Manganese Mn: stabilises the austenite, increases the strength and toughness. Higher contents of >2.0 wt. % Mn increase the risk of middle segregations which significantly reduce the ductility and therefore the product quality. Lower contents <1.0 wt. % do not allow the required strength and toughness to be achieved at the desired moderate analysis costs. Therefore, the content of Mn is fixed to 1.0 to 2.0 wt. %.
Aluminium Al: is used for deoxidation in the steel works process. The quantity of Al used is process-dependent. Thus no minimum Al content is given. An Al content of greater than 0.06 wt. % considerably impairs the casting behaviour in the continuous casting process. This gives rise to increased effort when casting. Therefore, the content of Al is fixed to a maximum of 0.06 wt. %.
Silicon Si: belongs to the elements which permit an increase in the strength of steel by mixed crystal hardening in an inexpensive manner. However, Si reduces the quality of the surface of the hot strip by the conveying of firmly adhering scale on the reheated slabs, which, in the case of high Si contents, can be removed only with a considerable effort or may only be removed to an insufficient degree. This is disadvantageous particularly in the case of subsequent galvanising. Therefore, the Si content is limited to a maximum of 0.6 wt. %. For the efficacy of Si, a lower limit of 0.1 wt. % can be considered reasonable.
Calcium Ca: is added by alloying to the steel for inclusion control in order to prevent the formation of unfavourable inclusions of MnS and Al2O3 and to form with these elements less detrimental Ca-containing inclusions in terms of morphology. The addition by alloying to the steel is a maximum of 0.01 wt. %.
Microalloy elements are generally added only in very small amounts (<0.2 wt. % per element). In contrast to the alloy elements, they mainly act by precipitate formation but can also influence the properties in the dissolved state. Despite the small amounts added, microalloy elements greatly influence the target-orientated production conditions and the processing properties and final properties of the product.
Typical microalloy elements are e.g. niobium and titanium. These elements can be dissolved in the iron lattice and form carbides, nitrides and carbonitrides with carbon and nitrogen. Since the microalloy elements are comparatively cost-intensive, the alloyed proportion is kept as low as possible. On the other hand, the carbon, which is hyperstoichiometric and therefore not bound in precipitations of the microalloy elements, in carbon-rich microstructure components contributes to the cost-effective and necessary increase in strength. Therefore, the hyperstoichiometric proportion of carbon and nitrogen is calculated according to formula 1: (C/12+N/14)/(Ti/48+Nb/93+Mo/96) set to >1.
The effect of Nb and Ti depends in particular on how the processing is carried out during hot-rolling and subsequent cooling. The addition of microalloy elements seeks to achieve grain refinement during the process and to produce precipitations in the nanometre size range. Therefore, a Nb+Ti content of more than 0.06 wt. % is a prerequisite for achieving the desired strength and good elongation properties. In contrast, a total value of more than 0.2 wt. % no longer has any effect on improving the properties of the steel, since contents above the indicated total value can no longer be dissolved during the reheating of the slabs when the stated analysis is carried out and conventional furnaces are used, and thus do not demonstrate any positive effect.
Niobium Nb: the addition by alloying of niobium acts in a grain-refining manner in particular by forming carbides in the rolling process, whereby at the same time the strength, toughness and elongation properties are improved. In addition, very fine Nb-containing precipitations can be formed after phase conversion which contribute significantly to the strength of the product. In the case of contents of over 0.08 wt. %, a saturation behaviour sets in, for which reason a maximum content of less than or equal to 0.08 wt. % is provided. A minimum content of 0.01 wt. % is provided for sufficient efficacy.
Titanium Ti: acts in a grain-refining manner as a carbide-forming agent, whereby at the same time the strength, toughness and elongation properties are improved. Contents of Ti of more than 0.18 wt. % impair the ductility and hole expansion capability by the formation of very coarse primary TiN precipitations, for which reason a maximum content of 0.18 wt. % is set. For sufficient efficacy, a minimum content of 0.02, advantageously 0.04, even more advantageously 0.06 wt. %, is provided.
Molybdenum Mo: increases the hardenability or decreases the critical cooling rate and thus promotes the formation of fine bainite microstructures. Furthermore, the use of small quantities of Mo already delays the coarsening of fine precipitations which should be as fine as possible in order to increase the strength of microalloyed microstructures. A minimum content of 0.05 wt. % is provided for sufficient efficacy and is limited to a maximum of 0.35 wt. % for cost reasons.
Phosphorus P: is a trace element from iron ore and is dissolved in the iron lattice as a substitution atom. Phosphorus increases hardness by means of mixed crystal hardening and improves hardenability. However, attempts are generally made to lower the phosphorus content as much as possible because inter alia it exhibits a strong tendency towards segregation and greatly reduces the level of toughness. The attachment of phosphorus to the grain boundaries can cause cracks along the grain boundaries during hot rolling. Moreover, phosphorus increases the transition temperature from tough to brittle behaviour by up to 300° C. However, by targeted measures, which are precisely controlled in terms of processing, the use of low quantities of P also makes possible an inexpensive increase in strength. For the aforementioned reasons, the phosphorus content is limited to at most 0.06 wt. %.
Sulphur S: like phosphorous, is bound as a trace element in the iron ore. It is generally not desired in steel since it leads to undesired inclusions of MnS, whereby the elongation and toughness properties are adversely affected. It is thus attempted to achieve the lowest possible quantities of sulphur in the melt and possibly to transform the elongated inclusions by so-called Ca treatment into a more favourable geometric form. For the aforementioned reasons, the sulphur content is limited to at most 0.01 wt. %.
Nitrogen N: is likewise an associated element from steel production. Steels with free nitrogen tend to have a strong ageing effect. The nitrogen diffuses even at low temperatures to dislocations and blocks same. It thus produces an increase in strength associated with a rapid loss of toughness. Binding of the nitrogen in the form of nitrides is possible e.g. by addition by alloying of aluminium, niobium or titanium. However, the stated alloy elements are subsequently no longer available later in the process for new formation of small precipitations which are very efficient with respect to strength. For the aforementioned reasons, the nitrogen content is limited to at most 0.012 wt. %.
Chromium Cr: as an element optionally added by alloying, Cr improves the strength and reduces the corrosion rate and delays the formation of ferrite and pearlite. The maximum content is set to at most 0.6 wt. %, since higher contents result in an impairment in ductility. A content of more than 0.1 wt. % is provided for sufficient efficacy.
Nickel Ni: the optional use of already low quantities of Ni promotes ductility while leaving the strength unchanged. Owing to the comparatively high costs, the content of Ni is limited to at most 0.6 wt. %. A content of more than 0.1 wt. % is provided for sufficient efficacy.
Vanadium V: in the case of the present alloy concept, an addition of vanadium is not absolutely necessary. The content of vanadium is limited to at most 0.2 wt. % for cost reasons. However, if an addition of V is provided, the hyperstoichiometric proportion of carbon and nitrogen is calculated according to formula 2: (C/12+N/14)/(Ti/48+Nb/93+Mo/96+V/51) set to >1. Then, a content of V of more than 0.01 wt. % is provided for sufficient efficacy.
Boron B: boron is an effective element for increasing hardenability which becomes effective even in very small quantities. The martensite starting temperature remains unaffected thereby. In order to become effective, boron must be present in a solid solution. Since it has a high affinity to nitrogen, the nitrogen must initially be removed, preferably by the stoichiometrically required quantity of titanium. By reason of its low solubility in iron, the dissolved boron preferably becomes attached to the austenite grain boundaries. At this location, it partially forms Fe—B carbides which are coherent and decrease the grain boundary energy. Both effects act in such a way as to delay the formation of ferrite and perlite and thus increase the hardenability of the steel.
However, excessively high contents of boron are hazardous because iron boride can form which has a negative effect upon hardenability, deformability and toughness of the material.
For the aforementioned reasons, the boron content for the alloy concept in accordance with the invention is limited to values of a maximum of 0.01 wt. %. A content of more than 0.0005 wt. % is provided for sufficient efficacy.
A method in accordance with the invention for producing a hot-rolled flat steel product with high local cold-formability, having a tensile strength Rm of at least 760 MPa, a yield strength ratio of at least 0.8 and a hole expansion ratio of above 30%, advantageously at least 40%, particularly advantageously at least 50%, a measure of cold-formability of at least 0.12, advantageously at least 0.17, and a ratio of local and global cold-formability of at least 5 and at most 13, comprising the steps of:

- melting a steel melt containing (in wt. %):

C: 0.04 to 0.08

Si: 0.1 to 0.6

Mn: 1.0 to 2.0

P: max. 0.06

S: max. 0.01

N: max. 0.012

Al: up to 0.06

Ti: up to 0.18 and/or

Nb: up to 0.08

Mo: up to 0.35

with Ti+Nb more than 0.06, and wherein a hyperstoichiometric ratio of carbon and nitrogen is present according to the following formula 1:1.0<(C/12+N/14)/(Ti/48+Nb/93+Mo/96) is set, with optional addition by alloying of one or more elements of Cr, Ni, V, B or Ca, with the remainder being iron including unavoidable steel-associated elements

- casting the steel melt to form a slab or thin slab by means of a horizontal or vertical slab or thin slab casting process,
- reheating the slab or thin slab to 1100° C. to 1270° C. and then hot-rolling the slab or thin slab with the following directly consecutive steps of:
- rolling in the last rolling pass to a hot strip to the required final thickness at a final rolling temperature EWT, wherein the following applies:

EWT≥EWTmin=682° C.+464 C+6445 Nb−644×Nb0.5+732 V−230 V0.5+890 Ti+363 Al−36 Si (formula 3)

- cooling at an average cooling rate of 30 K/s to 150 K/s
- reeling the hot strip into a coil at a reeling temperature HT which is low enough to set the advantageous microstructure constituents with

HT≥HTmax=761° C.−217×C−77×Mn+97×Si−47×Mo−53×Cr−34×Ni−21×V (formula 4)
and on the other hand is suitable to provide sufficient precipitation-hardening in the subsequent time-dependent cooling process T(t), defined by (formula 5) 17000≤HP≤18800 with HP(T, t)=T(t)×(ln(t)+20), wherein the temperature T is indicated in K and the duration t is indicated in h.

- cooling in a cooling process T(t) at an average cooling rate of 5 K/h to 50 K/h between reeling temperature and 100° C., with subsequent cooling in still air to room temperature.

In the case of the optional addition by alloying of one or more of the elements Cr, Ni, V, B and Ca, provision is made in particular that Cr is added by alloying up to 0.6 wt. %, Ni up to 0.6 wt. %, V up to 0.2 wt. %, B: up to 0.01 wt. % and Ca up to 0.01 wt. %.
The idea forming the basis of the invention will be explained hereinafter and described in greater detail using examples.
Nowadays, thermomechanical rolling is typically used in order to produce high-strength, microalloyed hot strip. In this process, finish-rolling is effected in a low temperature range of less than EWTmin, in which the austenite no longer recrystallises and, as a result, the accumulated dislocations lead to a high nucleation density with the beginning of the phase conversion and thus produce a fine hot strip microstructure. An aim of the thermomechanical rolling is to increase the strength and ductility through a small grain size of the hot strip microstructure.
In the case of alloys having a hyperstoichiometric ratio of carbon and nitrogen to microalloy elements, the carbon, in contrast to the nitrogen, is not completely brought to precipitation in the form of strength-increasing microalloy precipitations. The carbon not precipitated in microalloy precipitations leads to the formation of carbon-rich microstructure constituents and to different carbon-rich constituents of the bainite. For the cold-formability, it is crucial that the carbon-rich microstructure components and the carbon-rich components of the bainite are advantageously present in terms of size and distribution. “Advantageously” means that there is a small size and as uniform a distribution as possible.
In order to achieve a balanced ratio of local and global cold-formability, different processes are also used in addition to different alloy compositions.
It is possible to distinguish between three process routes, in which the type and distribution of the carbon-rich microstructure constituents and the precipitation state of the microalloy elements are set. The type and distribution of the carbon-rich microstructure constituents influences the cold-formability and the precipitation state of the microalloy elements influences the strength.
The process routes are:
1. Low reeling temperatures of e.g. 450<HT<550° C. lead to the setting of a low-temperature bainite having constituents of carbon-rich constituents which are very finely distributed, e.g. lower bainite. The resulting product exhibits high cold-formability with a pronounced proportion of local cold-formability (“high local cold-formability”). However, the strength is comparatively low, since at the low reeling temperatures only a small proportion of microalloy elements are brought to precipitation and the contribution to precipitation-hardening is correspondingly low.
2. High reeling temperature of e.g. HT>650° C. for setting a ferritic microstructure. The carbon is present in the form of hard microstructure constituents, such as carbides, pearlite or martensite. The resulting product exhibits high cold-formability with a lower proportion of local cold-formability. The strength is greater because a higher proportion of microalloy elements are brought to precipitation.
3. An average reeling temperature of e.g. 550<HT<650° C. to produce a mixed microstructure consisting of high-temperature bainite (e.g. upper bainite and granular bainite) and ferrite with both high local cold-formability and high strength by reason of a high precipitation proportion has hitherto not been expedient. Either only high-cold formability or only high strength was achieved by reason of the high precipitation proportion.
It was found in tests that the predominantly bainitic, microalloyed hot strip has both high strength and high local cold-formability when, in combination with the alloy composition and a hyperstoichiometric ratio of 1.0<(C/12+N/14)/(Ti/48+Nb/93+Mo/96), the flat steel product is finish-rolled at a final rolling temperature of at least EWTmin according to formula 3 and subsequently reeled and cooled in a temperature time window which is characterised by a maximum reeling temperature HTmax according to formula 4 and by 17000≤HP≤18800, wherein HP is calculated according to formula 5.
It has been found in tests that the hot-rolled flat steel product has a strength contribution to the tensile strength by precipitation-formation of at least 80 MPa or more when reeled and cooled in a temperature time window which is characterised by 17000≤HP≤18800 compared to a temperature time window which is characterised by HP≤15990. The strength contribution is necessary to achieve a high tensile strength and a high yield strength ratio in a cost-effective manner. At the same time, if HTmax is maintained within the temperature time window, the microstructure which is favourable for the local formability and strength is formed.
The flat steel product in accordance with the invention is characterised, when said temperature time window is maintained, in that half of the precipitations of (Ti, Nb, Mo)(C,N) and/or V(C,N) which harden the ferrite and the main constituent of bainitic ferrite have a diameter of less than 10 nm and/or the precipitations have an average spacing of less than 750 nm.
Surprisingly, it has been found that, when said temperature time window is maintained in combination with a final rolling temperature of at least EWTmin, the local cold-formability is high, whereas, when said temperature time window is maintained in combination with a final rolling temperature of less than EWTmin, the local cold-formability is comparatively low.
In addition to a cost-effective alloy concept, the flat steel product produced in accordance with the invention has high strength with high local cold-formability at the same time. Moreover, the production method in accordance with the invention is characterised by high process stability.
In contrast to ferrite, the bainite consists typically of different constituents. The different constituents of the bainite are formed from the austenitic phase during the production of the hot strip after final rolling at a decreasing temperature. In comparison with ferrite, the bainite is formed at lower temperatures and the bainite has, on average, a higher dislocation density.
Only with a predominantly bainitic microstructure are both the high strength and the high local cold-formability achieved. The reason is that the bainitic microstructure has a high dislocation density and a small grain size.
The high local cold-formability is not achieved with a predominantly ferritic microstructure. The reason is that the grain size of the ferrite is comparatively large and the hyperstoichiometric carbon precipitates in the form of very hard and comparatively coarse carbides at phase boundaries. These carbides lead to premature material failure during cold forming by reason of local stress concentrations.
With previously known solutions, it is possible to achieve hot strips with either the property combination of comparatively low strengths with comparatively high local cold-formability or the property combination of comparatively high strength with comparatively low local cold-formability.
In contrast, the invention allows the combination of properties of high strength and high local cold-formability to be achieved.
The reason is that the accumulated local damage during the forming procedure, in particular in the case of large elongation differences, can only be appropriately taken into account when considering the true sizes.

- Definition of true uniform elongation: ln(1+Ag/100), where Ag is the technical uniform elongation.
- Definition of true hole expansion ratio: ln(1+LA/100), where LA is the technical hole expansion ratio.

The measure of cold-formability is described by the geometric mean value of local and global formability:

- Definition of measure of cold-formability (Formability Level “FL”): (true uniform elongation×true hole expansion ratio)^0.5

The ratio of local and global cold-formability is defined as: Ratio of local and global cold-formability (Local Formability Ratio “LFR”)=(true hole expansion ratio/true uniform elongation)
The following criteria are required for high local cold-formability, in particular for the field of application of the flat steel product in accordance with the invention:

- A≥10%
- LA≥30%
- Measure of cold-formability≥0.12
- 5≤ratio of local and global cold-formability ≤13

The test results illustrated in the appendix cover exemplified embodiments having a tensile strength of at least 760 MPa. For the exemplified embodiments, the following criteria are required for high local cold-formability, in particular for the field of application:

- A≥16%
- LA≥50%
- Measure of cold-formability≥0.17
- 5≥ratio of local to global cold-formability≥13

Within the scope of the tests, the mechanical-technological properties as well as the microstructures of the hot-rolled flat steel products produced were tested. In addition to tensile tests according to ISO 6892-1 to determine tensile strength Rm, yield strength Rp0.2 and elongation at fracture A and uniform elongation Ag, hole expansion tests were carried out according to ISO 16630.
The following abbreviations are used in the tables listed in the appendix and in the following description:

- EWT=Final Rolling Temperature
- HT=Reeling Temperature
- MW=Mean Value
- Leg.=Alloy
- GOS=Grain Orientation Spread
- KAM=Kernel Average Misorientation
- IQ=Image Quality
- AR=Aspect Ratio
- SGV=Yield Strength Ratio
- S_P=Strength contribution by precipitation-formation
- S_M=Strength in a predominantly bainitic microstructure which has no precipitations by reason of the low value of the parameter HP.

For hot strip thicknesses >3 mm, the proportional sample shape with the elongation at fracture designation A was used. In contrast, for hot strip thicknesses of ≥3 mm the non-proportional sample shape with an initial measurement length of 80 mm was used. For better comparability, when using the non-proportional sample shape, the values for the elongation at fracture were converted from the uniform elongation according to A=AG×b, where b was determined beforehand to be 2.254 using reference samples. In the hole expansion test, the mean value of at least 3 individual tests is always indicated.
For the metallographic evaluation of the microstructure of the hot strip, the following regions were defined in different thickness ranges of the sample:

- Near the surface: measuring field with 100 μm×100 μm with a spacing from the sample surface of 0.1 mm
- ¼ Thickness: measuring field with 100 μm×100 μm centrally between surface and sample centre
- ½ Thickness: measuring field with 100 μm×100 μm with a spacing from the sample centre of 0.1 mm

The positioning of the measuring fields can be seen from the sketch in FIG. 2 .
The metallographic tests were conducted on samples longitudinally with respect to the rolling direction.
Electron backscatter diffraction (EBSD) images were taken in the measuring fields defined above in order to characterise the microstructure. For this purpose, longitudinal sections were produced which were mechanically ground and polished to 1 μm. Subsequently, the samples were polished with OP-S for ca. 10 min in order to produce a surface which was prepared to be as free of deformation as possible. An EDAX DigiView 5 EBSD camera with a binning of 10×10 and an acquisition rate of 140 Hz was used for the measurements, the acceleration voltage was 15 kV. The step size between the individual measuring points was 0.1 μm in each case. The parameters which are relevant to the invention were determined as follows: GOS (Grain Orientation Spread): average misorientation of all measuring points within a grain with respect to the average orientation of the grain. A segmentation angle of 15° is used to determine the grains.
KAM (Kernel Average Misorientation) and GKAM: in order to calculate the KAM values, the average misorientation of an EBSD measuring point is determined with respect to its next but one neighbouring measuring points. The maximum permissible misorientation is 4°. For the GKAM, the KAM values of all measuring points of a grain are averaged, wherein a segmentation angle of 15° is used in order to determine the grains.
The occurring microstructure types are defined in terms of metallography as follows:

- Ferrite consists of polygonal and quasi-polygonal ferrite and the grains are delimited by grain boundaries with misorientation angles >15°. No small angle grain boundaries <15° occur in the grain interior of the ferrite, the Grain Orientation Spread (GOS) values are <2° and the Grain Kernel Average Misorientation (GKAM) values are typically <0.4°. TEM images show a high density of (Ti, Nb, Mo)(C,N) precipitationss in the grain interior. In particular, (Fe,Mn) carbides can be present in the region of the grain triple junctions.
- The grains of granular bainite are delimited by grain boundaries of >15°. By reason of the displacive phase conversion of the austenite into the bainite, small angle grain boundaries occur in the grain interior, the GOS values are 2° and the GKAM values are typically 0.4°. The EBSD IPF (Inverse Pole Figure) map typically shows lancets of different orientation in the grain interior. Lancets which do not show a second phase in the EBSD Image Quality Map are designated hereinafter as “bainitic ferrite”. A carbon-rich second phase in the form of martensite, MA phase, lower bainite or pearlite is incorporated between the grains of the bainitic ferrite. In particular, (Fe,Mn) carbides can be present in the region of the grain triple junctions. The surface proportion of the second phase decreases as the reeling temperature increases and can be 0-10%.

Aspect ratio: The samples were oriented for the EBSD measurement in the electron microscope such that the rolling direction corresponds with the Y-direction of the measuring field. With the aid of the Matlab Toolbox MTEX, ellipses were adapted to the shape of the individual grains (segmentation angle 15°) and parameterised via their long and short semi-axis, as well as the orientation of the long axis. The ellipses were calculated from these parameters for each grain and then the intersection of these ellipses with the X- and Y-axes of the coordinate system was determined. The ratio of the intersection points of the grain ellipses with the X-axis to the intersection points of the grain ellipses with the Y-axis corresponds to the aspect ratio of the grains in the rolling direction to the sheet normal. This calculation method ensures that the extension of grains, of which the long axis does not point exactly in the rolling direction is determined only in the rolling or sheet normal direction.
The hardness test HV0.1 was carried out on the polished sample in points with different spacings with respect to the surfaces. No measurement is taken at a spacing from the surfaces and the centre of 0.1 mm. The following also applies:
The hardness values are indicated as the mean value from 6 individual measurements.
In each case, 3 hardness impressions for the near-surface position are positioned between 0% and 10% as well as a 90% and 100% spacing from the surface in relation to the thickness of the sheet.
In each case, 3 hardness impressions for the ¼ position are positioned between 20% and 30% as well as a 70% and 80% spacing from the surface in relation to the thickness of the sheet.
In each case, 3 hardness impressions for the ½ position are positioned between 40% and 50% as well as a 50% and 60% spacing from the surface in relation to the thickness of the sheet.
The positioning of the hardness impressions can be seen from the sketch in FIG. 2 .
The alloy compositions of two exemplified embodiments are summarised in Table 1. Alloys A and B are single casts and so all examples A1-A14 and B1-B20 have the same compositions. Likewise, Table 1 shows the calculated values for the hyperstoichiometric ratio of carbon and nitrogen to microalloy elements (formula 2), i.e.
1.0<(C/12+N/14)/(Ti/48+Nb/93+Mo/96+V/51)
Table 1 illustrates the alloy compositions of two exemplified embodiments.
Tables 2 and 3 illustrate the results on different exemplified embodiments. Also illustrated is an evaluation of the results with regard to the achievement of the required characteristic values with Y (achieved) and N (not achieved). If the specifications in accordance with the invention are not met, according to the 2nd line in the tables, this is indicated by underlining. The values listed are rounded in accordance with standard commercial practice.
For alloys A and B, Table 2 lists the results for the mechanical characteristic values with different process conditions are listed. The values which are underlined are outside the required mechanical properties or outside the expedient process conditions.
With regard to the final rolling temperature (EWT), it must be ensured that complete recrystallisation is achieved over strip thickness in every thickness range. This is achieved when EWT-EWTmin≥0, where EWTmin=682° C.+464 C+6445 Nb−644×Nb^0.5+732 V−230 V^0.5+890 Ti+363 Al−36 Si (formula 3). All element specifications are given in wt. %.
During subsequent reeling of the strip, it is necessary to ensure that a microstructure consisting of more than 50 vol. % bainite is produced. This is achieved when HT −HTmax <0, where HTmax=761° C.−217×C−77×Mn+97×Si−47×Mo−53×Cr−34×Ni−21×V (formula 4). All element specifications are given in wt. %. If the conditions for EWT and HT are met, the characteristic values for formability indicated in Table 2 can be achieved. It is shown that a cost-effective increase in strength can be achieved only if a sufficient contribution to the tensile strength Rm is provided by precipitation-hardening S_Pafter reeling during the subsequent cooling process. To this end, it is necessary that a suitable temperature T prevails for a suitable duration t during the cooling process T(t). This is achieved (formula 5) when 17000≤HP≤18800 with HP(T, t)=T(t)×(ln(t)+20), where T is always indicated in K and t is always indicated in h when calculating HP.
In order to calculate the parameter HP, the following procedure is performed:
1. The cooling process T(t) is divided into n equal time periods ti with the corresponding temperatures Ti, wherein n is to be selected to be sufficiently large so that the result remains almost the same when divided into significantly more time periods.
2. Calculation of the individual parameters HP_i=HP_i(t_i, T_i)=T_i×(ln (t_i)+20).
3. Calculation of the time periods t_i*=exp (HP_i/T*−20), where T* represents any temperature, e.g. the reeling temperature HT.
4. Calculation of the parameter HP with HP=T*×(ln(t₁*+t₂*+ . . .+t_n*)+20)
The strength contribution S_Pdue to precipitation-formation is determined in the following steps:
1. Determining the strength S_Min a predominantly bainitic microstructure which has no precipitations by reason of the low value of the parameter HP. With the aid of TEM tests, it was established that the state is present irrespective of the alloy at HP=15990. As a first step, the data Rm is plotted over HP in the range 16080<HP<18000 for all application examples and then a linear regression is performed. In the second step, the strength at HP=15990 is determined with the aid of the regression line. In the present case, this is the strength contribution S_M,A(15990)=804 MPa for alloy A and the strength contribution S_M,B(15990)=762 MPa for alloy B.
2. Calculation of the theoretical strength of a precipitation-free microstructure in dependence upon HP by S_M(HP)=S_M(15990)−0.0495 (HP-15990)
3. Calculation of the strength contribution S_P(HP) by S_P(HP)=Rm (HP)−S_M(HP)
In the case of the alloy compositions A and B, high strengths Rm are achieved in a particularly cost-effective manner, since the cooling process at the indicated ranges of HP allows a strength contribution S_Pby precipitation-formation ≥80 MPa.
Furthermore, high local-cold-formability with high strength can be observed only at 17000≥HP≥18800, but not at HP>18800 or HP<17000 (Table 2).
The causes in terms of materials science for the different local cold-formability at high strength were analysed on the basis of microstructure constituents and features in longitudinal sections.
The bainitic microstructure of the hot strip produced in accordance with the invention consists of a main constituent of 50% and secondary constituents, wherein the main constituent is formed from bainitic ferrite hardened by precipitations (Ti, Nb, Mo)(C,N).
Transmission-electronic tests of the precipitations on individual representative samples have shown that half of the precipitations of (Ti, Nb, Mo)(C,N) which harden the main constituent of bainitic ferrite have a diameter of <10 nm and/or the precipitations have an average spacing of less than 750 nm. The secondary constituents consist of carbon-richer constituents such as e.g. martensite, MA phase, lower bainite, and pearlite. Since the main constituent has a higher formability than the secondary constituents, a minimum proportion of the main constituent of 50% is advantageous.
Table 3 illustrates the results of the microstructure tests for alloy A for different final rolling temperatures according to formula 3, different reeling temperatures according to formula 4 and HP values according to formula 5.
The following applies to samples not processed in accordance with the invention: the microstructure of the hot strip samples are inhomogeneous and anisotropic over strip thickness. The inhomogeneity and anisotropy can be described for the two samples A2 and A6 with HP values 17232 and 18380 as follows:
a. The samples consist of a ferritic-bainitic microstructure. The ferrite proportion is 48% to 66%.
b. The deviation in the proportion of the ferrite in the near-surface position and the % thickness position in relation to the ½ thickness position is a maximum of 59% and 17%.
c. The microstructure extension is comparatively strong. This applies in particular to the ½ thickness position; in this case, the aspect ratio is 2.9 and 2.5.
d. The hardness over thickness varies comparatively greatly. Particularly at the surface, the hardness is lower and deviates by −24 HV0.1 and −26 HV0.1 from the mean value over the sample thickness.
e. The shear texture components vary comparatively greatly and are 0.92 and 0.96 in the near-surface position and are 0.01 and 0.01 in the ½ thickness position.
It is known that the local material behaviour during cold forming with its high demands on local cold forming capability is negatively influenced by microstructure inhomogeneities, as the damage is localised early and leads to material failure. In the present case, the features listed in a.-d. lead directly and indirectly to elongated regions of reduced formability, e.g. regions with an increased proportion of carbon-rich second phase of the granular bainite. However, the influence of the feature listed in e. on the local cold-formability is not known.
The cause for the microstructure inhomogeneity over hot strip thickness was identified as the different completed recrystallisation of the austenite over strip thickness immediately after the last rolling step and before cooling. With an increased temperature range of the finish-rolling to more than EWTmin but at a constant average reeling temperature, it has been possible within the scope of the invention to achieve complete recrystallisation over strip thickness in every thickness range and thus to achieve a homogeneous ferritic-bainitic microstructure over strip thickness.
The result is surprising, since the absence of recrystallisation in the last passes of the rolling process leads to a coarser microstructure in hyperstoichiometric bainitic hot strip grades, as expected, but the local cold-formability is positively influenced, against expectation.
For complete recrystallisation over strip thickness in every thickness range, an EWT of at least EWTmin according to formula 3 is necessary.
The causes in terms of materials science for the high local cold-formability were analysed on the basis of microstructure constituents and features in longitudinal sections. The result of the microstructure analysis of samples A7 and A9 according to Table 3, which were processed according to the invention, is:

- The microstructure of the hot strip samples are comparatively homogeneous and isotropic over strip thickness. The homogeneity and isotropy can be described for the two samples with HP values 18380 (sample A7) and 17232 (sample A9) as follows:

a. The samples consist of a predominantly bainitic microstructure. The ferrite proportion is 22% and 49%.
b. The deviation in the proportion of the ferrite in the near-surface position and the % thickness position in relation to the ½ thickness position is a maximum of 7% and −2%.
c. The microstructure extension is comparatively weak. This applies in particular to the ½ thickness position; in this case, the aspect ratio is 1.5 and 1.5.
d. The hardness over thickness varies comparatively little. This applies in particular to the near-surface position. The hardness on the surface deviates by −10 HV0.1 and 0 HV0.1 from the mean value.
e. The shear texture components are 0.98 and 0.98 in the near-surface position and 0.01 and 0.03 in the ½ thickness position.
Through further tests, the target microstructure was further characterised in the following point:
At least half of the precipitations of (Ti, Nb, Mo)(C,N) which harden the main constituent of bainitic ferrite have a diameter of <10 nm and/or the precipitations have an average spacing of less than 750 nm.
FIG. 1 summarises once again the range of cold-formability which is claimed in accordance with the invention and is limited by the specifications for FL and LFR.
FIG. 2 shows the positioning of the hardness impressions with a spacing from the surfaces (0% and 100%): 0.1 mm and spacing from centre (50%): 0.1 mm and the EBSD measuring fields with a spacing from the surface (0%): 0.1 mm Spacing from centre (50%): 0.1 mm.

TABLE 1

Alloy	C	Si	Mn	N	V	Ti	Nb	Mo	Cr	Ni	Al	B	Formula 1

A	0.064	0.564	1.890	0.006	0.006	0.094	0.041	0.188	0.036	0.036	0.066	0.0001	1.29
B	0.056	0.314	1.634	0.006	0.004	0.106	0.038	0.182	0.046	0.042	0.059	0.0001	1.10

TABLE 2

									Uni-
									form	Elonga-				In accor-
	Thick-	EWT −	HT −						elonga-	tion at				dance with
	ness	EWTmin	HTmax		Rm	S_P	Rp0.2	LA	tion	fracture A				the
Alloy	[mm]	[K]	[K]	HP	[MPa]	[MPa]	[MPa]	[%]	[%]	[%]	FL	LFR	SGV	invention

			≤0	≤18800	≥760	≥80		≥30		≥10	≥0.12	≥5	≥0.8
				≥17000
A1	4.0	−40	−54	17805	897	183	790	39	10.2	21.9	0.18	3.4	0.88	N
A2	4.0	−40	−84	17232	899	156	712	27	9.2	20.2	0.15	2.8	0.79	N
A3	4.0	−40	−124	16468	836	55	676	55	9.0	19.8	0.19	5.1	0.81	N
A4	4.0	−40	−164	15705	833	15	703	69	8.4	19.6	0.21	6.5	0.84	N
A5	4.0	−40	−194	15134	818	−29	686	81	8.7	19.7	0.22	7.1	0.84	N
A6	4.0	−40	−24	18380	889	203	792	30	10.3	21.7	0.16	2.6	0.89	N
A7	4.0	60	−24	18380	906	220	793	59	9.5	21.0	0.21	5.1	0.88	Y
A8	4.0	60	−54	17805	896	182	793	76	9.0	21.1	0.22	6.5	0.88	Y
A9	4.0	60	−84	17232	914	172	740	66	7.5	18.4	0.19	7.1	0.81	Y
A10	4.0	60	−124	16468	823	43	682	72	7.2	17.8	0.19	7.8	0.83	N
A11	4.0	60	−144	16086	833	33	654	56	7.0	17.6	0.17	6.6	0.79	N
A12	4.0	50	−114	16659	803	32	640	61	8.2	18.3	0.19	6.1	0.80	N
A13	4.0	30	−124	16468	840	60	630	64	7.9	18.8	0.19	6.5	0.75	N
A14	4.0	20	−124	16468	808	28	643	76	8.0	18.7	0.21	7.3	0.80	N
B1	4.0	64	−58	17652	869	189	753	73	7.4	18.9	0.20	7.7	0.87	Y
B2	4.0	51	−46	17882	866	198	794	83	7.9	19.5	0.21	7.9	0.92	Y
B3	4.0	−21	−69	17442	828	138	694	50	9.5	21.3	0.19	4.4	0.84	N
B4	4.0	−33	−38	18035	836	175	768	51	9.0	20.5	0.19	4.8	0.92	N
B5	4.0	56	−43	17939	862	196	786	95	7.7	19.7	0.22	9	0.91	Y
B6	4.0	−49	18	19107	818	210	741	46	11.5	23.2	0.20	3.5	0.91	N
B7	2.1	13	−61	17595	858	175	750	101	8.7	19.6	0.24	8.4	0.87	Y
B8	2.1	33	−81	17213	850	149	717	79	7.8	17.6	0.21	7.7	0.84	Y
B9	2.1	28	−21	18360	852	207	793	78	9.8	22.2	0.23	6.1	0.93	Y
B10	2.1	44	−70	17423	871	180	789	103	8.1	18.3	0.23	9.1	0.91	Y
B11	4.0	57	−70	17423	852	161	727	69	7.1	19.1	0.19	7.6	0.85	Y
B12	4.0	57	−63	17557	846	162	731	75	7.5	19.2	0.20	7.8	0.86	Y
B13	4.0	47	−1	18743	836	210	756	71	7.6	19.1	0.20	7.3	0.90	Y
B14	4.0	69	−80	17232	828	127	686	73	6.5	18.2	0.19	8.7	0.83	Y
B15	2.1	40	−52	17767	866	192	821	79	7.8	17.6	0.21	7.8	0.95	Y
B16	2.1	30	−48	17844	879	209	806	84	7.0	15.7	0.20	9.1	0.92	Y
B17	2.1	58	−65	17519	867	181	787	79	8.7	19.5	0.22	7.0	0.91	Y
B18	4.0	60	−98	16888	795	77	628	68	7.5	19.0	0.19	7.2	0.79	N
B19	4.0	55	−61	17595	861	178	733	74	7.2	18.5	0.20	8	0.85	Y
B20	3.0	−59	−17	18437	898	257	860	57	10.6	20.3	0.21	4.5	0.96	N

TABLE 3

	In accordance					MW		MW
	with the	EWT − EWTmin	HT − HTmax			Bainite	Bainite	Ferrite
Alloy	invention	[K]	[K]	HP	Position	[%]	[%]	[%]

A2	N	−40	−84	17232	Surf.	52	18	48
					¼		59
					½		77
A6	N	−40	−24	18380	Surf.	34	23	66
					¼		40
					½		40
A7	Y	60	−24	18380	Surf.	51	51	49
					¼		51
					½		50
A9	Y	60	−84	17232	Surf.	78	75	22
					¼		79
					½		81

		Δ Ferrite in					Δ Hardness
		relation to	MW		MW		in relation	Shear
	Ferrite	½ thickness	AR		Hardness	Hardness	to MW	texture
Alloy	[%]	[%]	[ ]	AR	[HV 0.1]	[HV 0.1]	[HV 0.1]	components

A2	82	59	2.3	1.6	294	271	−24	0.92
	41	18		2.3		305	11	0.33
	23	0		2.9		284	−10	0.01
A6	77	17	2.2	2.0	295	269	−26	0.96
	60	−1		2.1		304	9	0.08
	61	0		2.5		296	1	0.01
A7	49	−1	1.3	1.2	306	296	−10	0.98
	49	−2		1.2		310	4	0.21
	50	0		1.5		305	−1	0.03
A9	26	7	1.5	1.7	307	307	0	0.98
	21	3		1.4		311	4	0.28
	19	0		1.5		303	−3	0.01

Claims

1. A high-strength, hot-rolled flat steel product with high local cold-formability, having a tensile strength Rm of at least 760 MPa, a yield strength ratio of at least 0.8 and a hole expansion ratio of at least 30%, an elongation at fracture of at least 10%, a measure of cold-formability of at least 0.12, and a ratio of local and global cold-formability of at least 5, and a microstructure consisting of more than 50 vol. % bainite, carbon-rich microstructure constituents, with the remainder being ferrite, said flat steel product having a chemical composition comprising (in wt. %):

C: in a range of 0.04 to 0.081

Si: in a range of −0.1 to 0.6;

Mn: in a range of 1.0 to 2.0;

P: max. 0.06;

S: max. 0.01;

N: max. 0.012;

Al: up to 0.06;

T_i: up to 0.18 and/or Nb: up to 0.08;

Mo: up to 0.35;

with T_i+Nb more than 0.06, wherein a hyperstoichiometric proportion of carbon and nitrogen is present according to the following formula:

1.0<(C/12+N/14)/(T_i/48+Nb/93+Mo/96),

with the remainder being iron including unavoidable, steel-associated elements, with optional addition by alloying of one or more elements of Cr, Ni, V, B or Ca,

wherein the microstructure distribution over the thickness of the flat steel product in the three regions near the surface, ¼ thickness and ½ thickness of the flat steel product comprises:

an absolute deviation of a maximum of 12 vol. %, advantageously a maximum of 7 vol. % of the proportion of ferrite in the region near the surface or in the region of ¼ thickness of the flat steel product, or in to the region of ½ thickness of the flat steel product and/or

the deviation in aspect ratio of the grains in the rolling direction in the three regions of the flat steel product with respect to the mean value is less than 0.3 in each of the three regions,

and/or

in the three regions the difference in hardness HV0.1 is in each case a maximum of 20 HV 0.1, advantageously a maximum of 15 HV 0.1, even more advantageously a maximum of 10 HV 0.1 in comparison with the mean value over the entire thickness of the flat steel product.

2. The flat steel product as claimed in claim 1, wherein an addition by alloying of Ca of a maximum of 0.01 wt. %.

3. The flat steel product as claimed in claim 1 or 2, wherein the steel contains (in wt. %):

T_i+Nb: max. 0.2.

4. The flat steel product as claimed in at least one of claims 1 to 3, wherein the steel contains (in wt. %):

T_imin.: 0.02, optionally 0.04, even more optionally 0.06

Nb min.: 0.01

Mo min.: 0.05.

5. The flat steel product as claimed in at least one of claims 1 to 4, wherein in that the steel contains (in wt. %):

Cr: up to 0.6

Ni: up to 0.6

V: up to 0.2

B: up to 0.01

wherein a hyperstoichiometric proportion of carbon and nitrogen is present according to the following formula:

1.0<(C/12+N/14)/(T_i/48+Nb/93+Mo/96+V/51).

6. The flat steel product as claimed in claim 5, wherein the steel contains (in wt. %):

Cr: more than 0.1

Ni: more than 0.1

V: more than 0.01

B: more than 0.0005

7. The flat steel product as claimed in claim 1, wherein the bainite is a mixture of constituents which is includes by a main constituent of at least 50 vol. % and secondary constituents, wherein the main constituent consists of bainitic ferrite hardened by precipitations of (Ti, Nb, Mo)(C,N) and/or V(C;N) and the secondary constituents consist of carbon-richer constituents, such as e.g. martensite, residual austenite, lower bainite, and pearlite.

8. The flat steel product as claimed in claim 1, wherein the microstructure consists of more than 50 vol. % bainite and the remainder consists of ferrite.

9. The flat steel product as claimed in claim 8, wherein the microstructure consists of more than 75 vol. % bainite and the remainder consists of ferrite.

10. The flat steel product as claimed in claim 1, wherein the grain extension of all microstructure constituents in a position ½ thickness of the flat steel product and the aspect ratio of all microstructure constituents in the rolling direction of at most 2.0 and/or the mean value over the three regions near the surface, ¼ thickness and ½ thickness of the flat steel product is at most 2.0, and optionally 1.6.

11. The flat steel product as claimed in claim 1, wherein half of the precipitations of (Ti, Nb, Mo)(C,N) and/or V(C,N) which harden the ferrite and the main constituent of bainitic ferrite have a diameter of less than 10 nm and/or the precipitations have an average spacing of less than 750 nm.

12. The flat steel product as claimed in claim 1, wherein the ratio of shear texture components to rolling texture components increases towards the surface and has the following values:

near the surface: min. 0.9; and

½ thickness: max. 0.1.

13. A method for producing a hot-rolled flat steel product with high local cold-formability, having a tensile strength Rm of at least 760 MPa, a yield strength ratio of at least 0.8 and a hole expansion ratio of above 30%, a measure of cold-formability of at least 0.12, and a ratio of local and global cold-formability of at least 5 and at most 13, said method comprising the steps of:

melting a steel melt containing (in wt. %):

C: in a range of 0.04 to 0.08:

Si: in a range of 0.1 to 0.6;

Mn: in a range of 1.0 to 2.0;

P: max. 0.06;

S: max. 0.01;

N: max. 0.012;

Al: up to 0.06;

T_i: up to 0.18 and/or Nb: up to 0.08;

Mo: up to 0.35;

with T_i+Nb more than 0.06, and wherein a hyperstoichiometric ratio of carbon and nitrogen is present according to the following formula: 1.0<(C/12+N/14)/(T_i/48+Nb/93+Mo/96) is set, with optional addition by alloying of one or more elements of Cr, Ni, V, B or Ca, with the remainder being iron including unavoidable steel-associated elements;

casting the steel melt to form a slab or thin slab by means of a horizontal or vertical slab or thin slab casting process;

reheating the slab or thin slab to 1100° C. to 1270° C. and then hot-rolling the slab or thin slab with the following directly consecutive steps of:

rolling in the last rolling pass to a hot strip to the required final thickness at a final rolling temperature EWT, wherein the following applies:

EWT≥EWTmin=682° C.+464 C+6445 Nb−644×Nb^0.5+732 V−230 V^0.5+890 T_i+363 Al

36 Si

cooling at an average cooling rate of 30 K/s to 150 K/s;

reeling the hot strip into a coil at a reeling temperature HT which is low enough to set the advantageous microstructure constituents with

HT≤HTmax=761° C.−217×C−77×Mn+97×Si−47×Mo−53×Cr−34×Ni−21×V

and on the other hand is suitable to provide sufficient precipitation-hardening in the subsequent cooling process T(t), defined by

17000≤HP<18800 with HP(T, t)=T(t)×(ln(t)+20), where T is indicated in K and t is indicated in h; and

cooling in a cooling process T(t) at an average cooling rate of 5 K/h to 50 K/h between reeling temperature and 100° C., with subsequent cooling in still air to room temperature.

14. The method as claimed in claim 13, wherein said hot-rolling the slab includes hot-rolling a flat steel product with a thickness of 1.6 mm to 6.0 mm.

15. The method as claimed in claim 13, further comprising providing the flat steel product with a metallic coat electrolytically or by means of hot-dipping.

16. The method as claimed in claim 15, wherein said providing the flat steel product comprises providing the flat steel product with a zinc-based metallic coat.

17. The method as claimed in claim 13 for producing a flat steel product as claimed in claim 1.

18. A method of using a flat steel product as claimed in claim 1 for producing a component in the automotive industry.

19. The method as claimed in claim 18 for producing a chassis component.

20. The flat steel product according to claim 1, wherein the microstructure distribution over the thickness of the flat steel product in the three regions comprises:

an absolute deviation of a maximum of 7 vol. % of the proportion of ferrite in the three regions; and/or

the deviation in aspect ratio of the grains in the rolling direction in the three regions of the flat steel product with respect to the mean value is less than 0.3 in each of the three regions; and/or

in the three regions the difference in hardness HV0.1 is in each case a maximum of 15 HV 0.1, and optionally a maximum of 10 HV 0.1 in comparison with the mean value over the entire thickness of the flat steel product.