High-strength hot dip-coated steel strip with plasticity brought about by microstructural transformation and method for production thereof

Abstract

The invention relates to a method of producing a hot dip-coated high-strength steel strip with plasticity brought about by microstructural transformation, comprising the following steps: (i) producing a hot-rolled steel strip consisting of the following elements in % by weight: C: from 0.15 to 0.205, Mn: from 1.9 to 2.6, Al: from 0.2 to 0.7, Si: from 0.5 to 0.9, Cr: from 0.2 to 0.5, Nb: from 0.01 to 0.06, Mo: < 0.15, B: ≤ 0.001, P: ≤ 0.02, S: ≤ 0.005, balance: iron, including customary steel-accompanying elements, where, for a value μ = 4.5 x ([Si] + 0.9 x [Al] + [Cr]) + 200 x [Nb], where [Si], [Al], [Cr] and [Nb] are the proportions of the corresponding elements in % by weight, 8 ≤ µ ≤ 16, (ii) etching and optionally cold rolling the hot-rolled steel strip to give a cold-rolled steel strip, (iii) subsequently continuously annealing in the course of a continuous process of hot dip coating the cold- or hot-rolled steel strip at a maximum temperature between 750°C and 950°C for a total duration of 10 s to 1200 s, (iv) subsequently cooling the cold- or hot-rolled steel strip to an intermediate temperature within a temperature range from 620 to 760°C at an average cooling rate CR1 of up to 10 K / s, (v) subsequently further cooling the cold- or hot-rolled steel strip from the intermediate temperature to a cooling stop temperature within a temperature range between 200°C and 450°C, especially between 280°C and 450°C, at an average cooling rate CR2 > CR1 and not more than 150 K / s, and then keeping the temperature within a temperature range between 200°C and 450°C, especially between 280°C and 450°C, for 25 to 500 s, (vi) then hot dip coating the cold- or hot-rolled steel strip at a temperature between 380 and 500°C and (vii) then cooling the hot dip coated cold- or hot-rolled steel strip at an average cooling rate of 1 K / s to 50 K / s to ambient temperature. The invention further relates to a corresponding hot dip-coated high-strength steel strip with plasticity brought about by microstructural transformation.

Description

[0001] High-strength hot-dip coated steel strip with plasticity induced by microstructural transformation and process for its production

[0002] The present invention is based on a process for producing a hot-dip coated high-strength steel strip with plasticity caused by microstructural transformation, comprising the following steps:

[0003] (i) producing a hot-rolled steel strip,

[0004] (ii) pickling and optional cold rolling of the hot-rolled steel strip into a cold-rolled steel strip,

[0005] (iii) Subsequent continuous annealing as part of a continuous hot-dip coating process of the cold-rolled or hot-rolled steel strip,

[0006] (iv) subsequently cooling the cold-rolled or hot-rolled steel strip to an intermediate temperature,

[0007] (v) subsequently further cooling the cold-rolled or hot-rolled steel strip from the intermediate temperature to a cooling stop temperature and then maintaining the temperature,

[0008] (vi) subsequent hot-dip coating of the cold-rolled or hot-rolled steel strip and

[0009] (vii) subsequent final cooling of the hot-dip coated cold-rolled or hot-rolled steel strip to ambient temperature.

[0010] In addition to this method, the invention also relates to a corresponding hot-dip coated high-strength steel strip with plasticity caused by microstructural transformation.

[0011] In high-strength dual-phase and multi-phase steels, the achievable elongation or ductility, and thus the formability, decreases with increasing strength class. However, the automotive industry demands steel grades that still offer high formability at high strengths in order to meet the requirement for high elongation in complex forming operations. This high formability can be achieved with plasticity induced by microstructural transformation. This requires a sufficient content of retained austenite in the microstructure.

[0012] A higher Si content can stabilize the proportion of residual austenite in the microstructure, resulting in higher elongations due to the well-known TRIP effect (TRIP: Transformation Induced Plasticity). EP2439291 B1 describes a multi-phase steel with 1.4 - 2.0 wt.% Si, whereby the high Si content suppresses carbide formation and achieves sufficient stability of the residual austenite. Similarly, EP3024951 B1 describes a steel sheet with a tensile strength > 900 MPa with 1 < Si < 3 wt.%. EP3128023B1 describes a high-strength steel strip with a high yield strength ratio and Si contents between 1.2 and 2.2 wt.%. The same range of Si between 1.2 and 2.2 wt.%. is also used in EP2707514B1 to produce a steel strip with a tensile strength > 1000 MPa and a uniform elongation > 12%.

[0013] However, during continuous galvanizing, Si significantly impairs the galvanizability by impairing the galvanizing reaction when the steel strip is immersed in the molten zinc. High Si contents can also lead to the formation of highly adhesive scale in the hot-rolled strip, complicating surface quality and further processing. Si can sometimes be substituted by increased Al contents to stabilize retained austenite. For example, EP3394300B1 and EP3394297B1 describe the use of 1 < Si+Al < 2 wt.% and 1 < Si+Al < 2.2 wt.%, respectively, to achieve sufficient amounts of retained austenite. Similarly, EP2439290B1 discloses a multi-phase steel with a tensile strength of at least 950 MPa and a residual austenite content of at least 6% with Si between 0.2 and 0.7 wt% and Al between 0.5 and 1.5 wt%, with the Cr content being < 0.1 wt%.

[0014] However, excessively high Al contents can be detrimental to hot ductility or castability during continuous casting. Furthermore, Al increases the Ac3 transformation temperature, necessitating high annealing temperatures during continuous galvanizing, which should be avoided for process and cost reasons. The sum of Si+Al should therefore be as low as possible when producing hot-dip galvanized steel strip. Low Si+Al contents are also advantageous for avoiding liquid metal embrittlement. In the context of the invention, the terms "hot-dip galvanized" and "hot-dip coated" are used synonymously, unless it is explicitly stated that the purpose of hot-dip galvanizing is to produce a coating of zinc or a zinc alloy.

[0015] In addition to the effect of Si and Al, EP2831299B2 describes the advantageous use of Cr in combination with Si and Al to produce higher contents of residual austenite with a minimum content of Si+0.8Al+Cr of 1.4 wt.% for 5-20 vol.% residual austenite. In addition to high costs and restrictions in weldability, a high Cr content has the disadvantage that, under conventional process control, the hot strip can become too strong due to increased hardenability to subsequently be cold rolled into a cold strip. By increasing hardenability, Cr also has a major impact on the microstructure development during continuous hot-dip galvanizing. Therefore, Cr cannot be added to any desired amount if a specific target microstructure is to be achieved. A predominantly bainitic microstructure requires a longer overaging zone in continuous hot-dip galvanizing with increasing Cr content, which is not always feasible on an industrial scale.

[0016] Document EP3730635 A1 describes a process for producing a hot-dip coated high-strength steel sheet, as well as a corresponding hot-dip coated high-strength steel sheet, in which the steel sheet itself has a carbon (C) content of 0.04 to a maximum of 0.15 wt.% and an antimony (Sb) content of 0.05 wt.% or less – excluding 0 wt.%. Antimony (Sb) is added to the alloy so that, distributed in grain boundaries, it retards the diffusion of oxidizing elements such as Mn, Si, Al, and the like through the grain boundaries. Furthermore, the steel sheet itself is characterized by a titanium (Ti) and / or niobium (Nb) content – ​​each of 0.003 to a maximum of 0.06 wt.% – which serves to increase strength and refine grains.

[0017] The invention is based on the object of providing a method for producing a hot-dip coated high-strength steel strip and a corresponding hot-dip coated high-strength steel strip, in which the hot-dip coated steel strip has overall high strength in combination with high uniform elongation with good adhesion of the coating to the actual steel strip.

[0018] The problem is solved by the subject matter of the independent patent claims.

[0019] Advantageous further developments of the invention result from the features of the subclaims.

[0020] The invention thus relates to a method for producing a hot-dip coated high-strength steel strip with plasticity caused by microstructural transformation, comprising the following steps:

[0021] (i) producing a hot-rolled steel strip, in particular from a slab heated to over 1200 °C, consisting of the following elements in % by weight: C: from 0.15 to 0.205; Mn: from 1.9 to 2.6; Al: from 0.2 to 0.7; Si: from 0.5 to 0.9; Cr: from 0.2 to 0.5; Nb: from 0.01 to 0.06; Mo: < 0.15; B: < 0.001; P: < 0.02; S: < 0.005, and optionally one or more of the following elements in % by weight:

[0022] Ti: 0.005 to 0.060; V: 0.001 to 0.060; N: 0.0001 to 0.016; Ni: 0.01 to 0.5 and Cu: 0.01 to 0.3, the remainder being iron, including common steel-accompanying elements, where 8 < p < 16 applies for a value p = 4.5 x ([Si] + 0.9 x [Al] + [Cr]) + 200 x [Nb], where [Si], [Al], [Cr] and [Nb] are the proportions of the corresponding elements in weight %,

[0023] (ii) pickling and optional cold rolling of the hot-rolled steel strip into a cold-rolled steel strip,

[0024] (iii) Subsequently, continuous annealing as part of a continuous hot-dip coating process of the cold-rolled or hot-rolled steel strip at a maximum temperature between 750 °C and 950 °C inclusive, in particular between 800 and 870 °C, for a total duration of 10 s to 1200 s, in particular from 50 s to 650 s,

[0025] (iv) subsequent cooling of the cold-rolled or hot-rolled steel strip to an intermediate temperature in a temperature range of 620 to 760 °C with an average cooling rate CRi of up to 10 K / s,

[0026] (v) subsequently further cooling the cold-rolled or hot-rolled steel strip from the intermediate temperature to a cooling stop temperature in a temperature range between 200 °C and 450 °C inclusive, in particular between 280 °C and 450 °C inclusive, with an average cooling rate CR2 > CRi and a maximum of 150 K / s and then maintaining the temperature in the temperature range between 200 °C and 450 °C inclusive, in particular between 280 °C and 450 °C inclusive, for 25 to 500 s,

[0027] (vi) subsequent hot-dip coating of the cold-rolled or hot-rolled steel strip at a temperature between 380 and 500 °C and

[0028] (vii) subsequent final cooling of the hot-dip coated cold- or hot-rolled steel strip to ambient temperature at an average cooling rate of 1 K / s to 50 K / s. A key aspect of the invention arises from the proportions of the elements Si; Al; Cr; and Nb, as well as the resulting value p in combination with the described continuous hot-dip coating process. The proportion of the element Nb is particularly noteworthy.

[0029] It is known that Nb, as a microalloying element, increases strength through strong grain refinement and the formation of fine Nb(C,N) precipitates. This effect is particularly advantageous in steels with low C content, since the solubility limit of Nb shifts to higher concentrations with decreasing C content, thus requiring a larger amount of Nb. The beneficial effect of Nb in an Al-free dual-phase steel with retained austenite and an Nb content of approximately 300 ppm is described in the scientific article "Mohrbacher, J.-R. Yang, Y.-W. Chen, J. Rehrl; Metals 10 (2020) 504." Advantages such as increased strength and toughness with an increase in R p o,2-yield strength and tensile strength (R m) through grain refinement and nano-sized Nb precipitates, more homogeneous strain distribution in the microstructure, and finer MA grain sizes to improve crack resistance. Furthermore, a significant delay in the recrystallization of the cold-rolled microstructure was observed. Regarding the Nb content of 300 ppm, the authors point out that this represents the maximum possible soluble Nb content for the dual-phase steel investigated in the industrial-scale process, since the maximum holding temperature during heating of the slabs prior to hot rolling is technologically limited. Publication EP 3394300 B1 also describes the optional use of Nb in steels with a TRIP effect for grain refinement of the austenite during hot rolling and strength enhancement through precipitation hardening during the final annealing treatment.

[0030] Contrary to this common use of Nb to increase strength, an unexpected effect was observed in the steel strip produced according to the invention. In combination with Si, Al, and iron, an increase in the Nb content in the steel according to the invention leads to lower strengths, but at the same time to an increase in the content of retained austenite with suitable processing (Fig. 1). Despite a decrease in tensile strength with increasing Nb content, the product of tensile strength and uniform elongation R increased. m x A gdue to the higher proportion of residual austenite. In the steel according to the invention, the contents of Si and Al in particular can be kept low compared to the achieved contents of residual austenite. The steel can therefore be produced using the continuous hot-dip galvanizing process and is characterized by very good coating adhesion, in particular zinc adhesion. It was determined that the value p = 4.5 x ([Si] + 0.9 [Al] + [Cr]) + 200 x [Nb] in wt.% should be between 8 and 16, in particular between 8.0 and 16.0, in order to stabilize residual austenite, whereby for the reasons described the sum [Si] + 0.9 x [Al] is limited to < 1.2 wt.%, preferably < 1.0 wt.% in order to guarantee good galvanizability. In addition, it was also found that Nb contents above 300 ppm are metallurgically effective and increase the residual austenite content in the final microstructure, allowing the highest elongations to be achieved.By avoiding excessively high Cr contents (< 0.5 wt.%, more precisely < 0.50 wt.%), different microstructures can be flexibly adjusted, in which, for example, more bainite is present than tempered martensite or vice versa.

[0031] In the step of producing a hot-rolled steel strip, in particular from a slab heated to over 1200 °C, it is preferably provided that the hot-rolled steel strip consists of the following elements in weight %: C: from 0.15 to 0.205; Mn: from 1.90 to 2.60; Al: from 0.20 to 0.70; Si: from 0.50 to 0.90; Cr: from 0.20 to 0.50; Nb: from 0.0100 to 0.0600; Mo: < 0.15; B: < 0.0010; P: < 0.02; S: < 0.005, and optionally one or more of the following elements in weight %:

[0032] Ti: 0.005 to 0.060; V: 0.001 to 0.060; N: 0.0001 to 0.016; Ni: 0.01 to 0.5; and Cu: 0.01 to 0.3, the remainder being iron, including common steel-accompanying elements, where 8.0 < p < 16.0 applies for a value p = 4.5 x ([Si] + 0.9 x [Al] + [Cr]) + 200 x [Nb], where [Si], [Al], [Cr], and [Nb] are the proportions of the corresponding elements in weight percent. In this specification of the composition of the steel strip, selected limits are specified more precisely with regard to their accuracy.

[0033] It is preferably provided that 10.0 < p < 16.0 applies to the value p and the expression [Si] + 0.9 x [AI] < 1.2, in particular [Si] + 0.9 x [AI] < 1.0, where [Si] and [AI] are the proportions of the corresponding elements in % by weight of the hot-rolled steel strip.

[0034] According to a preferred embodiment of the method according to the invention, the proportion of C in the hot-rolled steel strip is at least 0.16% by weight, i.e. the C content in weight% is in the range from 0.16 to 0.205. It is preferably provided that the proportion of Mn in the hot-rolled steel strip is between 1.95 and 2.4% by weight, in particular between 1.95 and 2.40% by weight, the proportion of C in the hot-rolled steel strip is at least 0.16% by weight and preferably the expression (100[C] + 10[Mn]) / (4.5 x ([Si] + 0.9[Al] + [Cr]) + 200x[Nb]) < 4.5 applies. Where [Si], [Al], [Cr], [C] and [Mn] are the proportions of the corresponding elements in the hot-rolled steel strip in weight %.

[0035] Furthermore, it is advantageous that the sum of the proportions of the elements Cr and Mo in the hot-rolled steel strip in weight % is less than 0.5, i.e. [Cr] + [Mo] < 0.5.

[0036] According to a further preferred embodiment of the method according to the invention, the intermediate temperature is in a temperature range of 650 to 730 °C and, when this temperature is reached, the steel strip has a structure with at least 10 volume % ferrite.

[0037] According to yet another preferred embodiment of the method according to the invention, the upper limit of the cooling stop temperature is 400 °C, preferably 350 °C, i.e. cooling stop temperature < 400 °C, preferably < 350 °C, wherein after the final cooling to ambient temperature more than 8 volume % austenite is present in the microstructure and wherein the temperature at which the cold-rolled or hot-rolled steel strip is kept before hot-dip coating is < 400 °C, preferably < 350 °C.

[0038] Further advantages of Nb for large-scale production in the continuous hot-dip galvanizing process include higher process speeds and wider cooling-stop temperature processing windows. The Nb content of the steel according to the invention also allows for the targeted adjustment of the ferrite content during the continuous hot-dip galvanizing process, allowing the control of technological properties such as tensile strength, yield strength, yield ratio, and uniform elongation without changing the annealing cycle. The wider cooling-stop temperature processing window allows for greater temperature fluctuations to be tolerated in large-scale production.Thus, in a preferred embodiment with p > 10, a fluctuation of the cooling stop temperature around AT in the range of 315 to 400 °C over the length of the steel strip leads only to a maximum fluctuation of the tensile strength in the rolling direction of the hot-dip coated steel strip of 0.47 x A in MPa (where the numerical value A corresponds to the numerical value in Kelvin of AT), which results in a high process stability of the process.

[0039] The invention relates to a high-strength hot-dip coated steel strip with a dual or multi-phase structure with retained austenite for plasticity caused by microstructural transformation, as well as excellent formability with A g> 8% and a low content of Si and Al, in combination with the appropriate process for producing the steel strip with high process stability. Furthermore, the steel strip should preferably exhibit good weldability and a low tendency to liquid metal and hydrogen embrittlement. In particular, the resulting hot-dip coated high-strength steel strip has a tensile strength R m of at least 900 MPa and uniform elongation A g of at least 8%.

[0040] According to a further preferred embodiment of the method according to the invention, the hot-dip coated steel strip is skin-rolled during or after the last process step with a rolling degree of maximum 2% in such a way that the R p o,2-yield strength increases by at least 20 MPa.

[0041] Finally, in the method according to the invention, it is advantageously provided that in the hot-rolled steel strip, the Ti content is at least 0.005% by weight, the N content is at most 0.008% by weight, the Al content is at most 0.50% by weight and TiN and TiAIN particles with a diameter of > 0.96 pm in total in a surface area of ​​at least 1 pm 2 / mm 2 on a measuring area of ​​at least 100 mm 2 in the slab before reheating and on a measuring surface of at least 20 mm 2 in the hot-dip coated high-strength steel strip.

[0042] Higher aluminum contents, such as those used in the steel according to the invention, can lead to the formation of harmful plate-like AlN at the primary grain boundaries during or immediately after continuous casting, making the slabs susceptible to cracking (so-called AlN embrittlement). It has been found that the number of these harmful AlN precipitates in the steel according to the invention can be reduced by TiN and TiAlN precipitates, which are not critical for crack susceptibility. TiN partially forms in the melt and thus binds the nitrogen before it can react with Al to form AlN. Thus, to reduce the susceptibility to cracking of the slabs used, the steel according to the invention advantageously contains TiN and TiAlN precipitates with a diameter of > 0.96 pm in total, covering an area of ​​at least 1 pm. 2 / mm 2 on a measuring area of ​​at least 100 mm 2in the slab before reheating. Since TiN and TiAlN do not dissolve during subsequent annealing processes, they can also be detected in the hot-dip coated high-strength steel strip. The TiN and TiAlN content can be quantitatively determined using energy-dispersive X-ray spectroscopy (EDX). This involves measuring the element contents in at.% in particles and the area A. pof the particles is measured, from which the diameter is given as (4Ap / TT). The concentration of elements in the particle (elements N, O, F, Na, Mg, Al, Si, S, K, Ca, Ti, V, Cr, Mn, Zn, Zr, Nb and Pb) minus the contents of the elements Fe and C results in a total of 100 at.%. TiN is characterized by the fact that the elements with the highest and second highest concentration in the particle in at.% are from the group [Ti, N], additionally the Ti and N contents are each > 5 at.% and the Al content is < 5 at.%. For TiAIN it applies that the elements with the highest and second highest concentration in the particle in at.% are from the group [N, Al, Ti] and the Ti, Al and N contents are each > 5 at.%.

[0043] The production of the high-strength hot-dip coated steel strip with plasticity caused by microstructural transformation and a microstructure of ferrite, martensite, bainite and retained austenite usually takes place as follows:

[0044] (i) Production of hot-rolled strip of varying thicknesses, typically between 1.7 and 6 mm, from slabs with the steel chemistry composition according to the invention, which were necessarily reheated to a temperature > 1200 °C prior to hot rolling to dissolve Nb. The final rolling temperature during hot rolling is between 860 and 960 °C and is typically > 900 °C to avoid thermomechanical rolling and microstructural inhomogeneities. The coiling temperature can be between 420 and 750 °C and is typically > 600 °C to achieve a predominantly ferritic-pearlitic hot strip structure.

[0045] (ii) Pickling and optional batch annealing of the hot strip at a maximum holding temperature between 400 and 700°C for an annealing time of 12 hours to 6 days, wherein the annealing time includes the time for heating and cooling to room temperature. Batch annealing can primarily be carried out if the hot strip was coiled at temperatures < 600°C in order to reduce resistance during cold rolling. In the context of the present invention, room temperature or ambient temperature is understood to mean a temperature between 10 and 40°C, preferably 15 to 25°C. Optional cold rolling of the hot strip into a cold strip or sheet to a thickness of typically 0.5 mm to 2.5 mm. (iii) Continuous annealing during the continuous hot-dip galvanising of the cold-rolled or hot-rolled strip at a temperature between 750 and 950 °C, preferably > 800 °C, in order to achieve a high degree of austenitisation.

[0046] (iv) Initial slow cooling in the slow cooling section of the plant (so-called slow cooling section) to an intermediate temperature of 620 to 760 °C, preferably 650 to 730 °C, to form ferrite. In addition to the residual austenite, ferrite, as a soft phase, contributes to increased ductility of the steel strip. The intermediate temperature must be selected so that the final microstructure contains at least 10% ferrite. Due to the high C and Mn content of the steel according to the invention, the hardenability is increased to such an extent that, in the region of the slow cooling section, with prior almost complete austenitization and higher process speeds or higher cooling rates in the slow cooling section, the formation of ferrite only occurs to an insufficient extent if < 100 ppm Nb is alloyed. Alloying the steel according to the invention with Nb enables the formation of ferrite in the region of the slow cooling section, even at higher process speeds or the resulting high cooling rates in the slow cooling section, such as 4 K / s.For this reason, the Nb content of at least 100 ppm Nb is mandatory and allows plant configurations with a short segment of the slow cooling section or high process speeds to increase throughput.

[0047] (v) Second rapid cooling in the rapid cooling segment of the plant from the intermediate temperature to a cooling stop temperature between 200°C and 450°C, in particular between 280°C and 450°C. Holding at an overaging temperature between 200°C and 450°C, in particular between 280°C and 450°C, preferably approximately equal to the cooling stop temperature, thereby increasing the carbon concentration in the retained austenite and ensuring that the retained austenite remains stable in the final microstructure even at room temperature. An advantage of approximately equal cooling stop and overaging temperatures is the use of hot-dip galvanizing plants that do not have a reheating unit immediately after reaching the cooling stop temperature. This allows flexible production of the steel according to the invention with various plant configurations. Lower cooling stop and overaging temperatures typically lead to higher strengths, but at the expense of the retained austenite content.In this context, it was found that with the steel according to the invention, with a higher Nb content > 100 ppm, in particular 200-600 ppm, at a low cooling stop temperature and the same overaging temperature of < 400 °C, in particular < 350 °C, a high retained austenite content can still be achieved, which is not possible without Nb. Furthermore, the achievable retained austenite contents for different cooling stop temperatures equalize at higher Nb contents, enabling larger process windows and consistent technological properties in terms of plasticity induced by microstructural transformation. Low cooling stop temperatures of < 400 °C, in particular < 350 °C, are also advantageous for forming tempered martensite and establishing an associated higher yield strength, as well as for allowing harmful hydrogen from the production process to diffuse out of the steel when held at this temperature.At excessively high cooling stop temperatures of > 450 °C, the martensite transformation only occurs after the hot-dip coating, and the hydrogen is trapped by the hot-dip coating. Excessively low cooling stop and overaging temperatures of < 200 °C also lead to hydrogen entrapment, as hydrogen diffusion is inhibited at low temperatures. Likewise, excessively low cooling stop temperatures of < 200 °C lead to an excessively high proportion of tempered martensite and thus to an excessively high yield strength ratio of the steel strip, requiring very high cooling rates. Therefore, the minimum cooling stop and overaging temperature is 200 °C, preferably 280 °C, in order to achieve optimal proportions of bainite and tempered martensite and to enable hydrogen diffusion from the steel strip.

[0048] (vi) optionally reheating the steel strip and performing a hot-dip coating at a temperature between 380 and 500 °C. The hot-dip coating step is an integral part of the manufacturing process of the steel according to the invention, since it not only applies a hot-dip coating to the steel surface, but the temperature treatment during this step also influences the final technological characteristics of the hot-dip coated steel strip. If martensite is already formed in the microstructure of the steel strip before reheating, this martensite is tempered during the hot-dip coating and is therefore referred to as tempered martensite.

[0049] (vii) The coated steel strip is then cooled to ambient temperature. During the final cooling to ambient temperature, the insufficiently stabilized austenite transforms into martensite (fresh martensite). The remaining austenite is referred to as retained austenite. Optional skin-passing and / or stretch-straightening of the coated steel strip to achieve the R P o,2-yield strength to be finally set.

[0050] The alloying concept and processing of the steel strip according to the invention aims to achieve high tensile strengths of > 900 MPa and a residual austenite content of > 8 vol.% to guarantee excellent formability. The microstructure of the steel strip according to the invention is composed of 8-16 vol.% residual austenite, > 10 and < 40 vol.% ferrite, and a total of at least 50 vol.% bainite, tempered martensite, and fresh martensite. In a preferred embodiment, the total amount of bainite and fresh martensite is more than tempered martensite. The percentages of ferrite, bainite, and martensite given for the microstructure components were determined in a longitudinal section perpendicular to the rolling surface and refer to area proportions (area spanned by the sheet normal and rolling direction), which are usually also adopted as volume proportions. Furthermore, the microstructure proportions refer to % position above thickness.

[0051] The retained austenite content can be measured using a magnetic induction method using a magnetizing yoke. Alternatively, the retained austenite content can also be determined using X-ray diffraction or electron backscatter diffraction (EBSD) on electropolished samples.

[0052] Due to its formation mechanism, fresh martensite possesses a high dislocation density and high hardness. In electron backscatter diffraction, such regions appear darker in the Kikuchi band contrast than other microstructure components, as the diffraction condition is violated by a disrupted crystal lattice. This allows the proportion of fresh martensite to be quantitatively determined. Alternatively, the formation of fresh martensite can be determined using dilatometry based on the volume change upon cooling of a sample.

[0053] The invention further relates to a hot-dip coated high-strength steel strip with plasticity brought about by microstructural transformation, in particular produced according to the above-mentioned process, consisting of the following elements in % by weight: C: from 0.15 to 0.205; Mn: from 1.9 to 2.6; Al: from 0.2 to 0.7; Si: from 0.5 to 0.9; Cr: from 0.2 to 0.5; Nb: from 0.01 to 0.06; Mo: < 0.15; B: < 0.001; P: < 0.02; S: < 0.005, and optionally one or more of the following elements in % by weight: Ti: 0.005 to 0.060; V: 0.001 to 0.060; N: 0.0001 to 0.016; Ni: 0.01 to 0.5 and Cu: 0.01 to 0.3, the remainder being iron, including common steel-associated elements. It is intended that for a value p = 4.5 x ([Si] + 0.9 x [Al] + [Cr]) + 200 x [Nb], where [Si], [Al], [Cr] and [Nb] are the proportions of the corresponding elements in weight %, 8 < p < 16 applies, whereby the steel strip is a product of R m -Tensile strength and uniform elongation A gof greater than 8000 MPa %, in particular greater than 9000 MPa %, and particularly advantageously between 9900 and 13000 MPa %,.

[0054] The steel of the actual steel strip is preferably made up of the following elements in weight %: C: from 0.15 to 0.205; Mn: from 1.90 to 2.60; Al: from 0.20 to 0.70; Si: from 0.50 to 0.90; Cr: from 0.20 to 0.50; Nb: from 0.0100 to 0.0600; Mo: < 0.15; B: < 0.0010; P: < 0.02; S: < 0.005, and optionally one or more of the following elements in weight %: Ti: 0.005 to 0.060; V: 0.001 to 0.060; N: 0.0001 to 0.016; Ni: 0.01 to 0.5 and Cu: 0.01 to 0.3, the remainder iron, including common steel elements. It is intended that for a value p = 4.5 x ([Si] + 0.9 x [Al] + [Cr]) + 200 x [Nb], where [Si], [Al], [Cr], and [Nb] are the proportions of the corresponding elements in weight %, 8.0 < p < 16.0 applies. In this specification of the composition of the steel strip, selected limits are specified more precisely with regard to their accuracy.

[0055] The preferred embodiments mentioned in connection with the process for producing a hot-dip coated high-strength steel strip shall also apply analogously to the hot-dip coated high-strength steel strip. For example, for the steel of the actual steel strip, it is specifically intended that the carbon content of this steel strip be at least 0.16% by weight, i.e., the carbon content in weight % is in the range of 0.16 to 0.205.

[0056] The coating created by hot-dip coating / hot-dip galvanizing is, for example, a coating made of zinc or a zinc alloy such as zinc-aluminum or zinc-aluminum-magnesium. Such coatings are well known and will not be discussed further here. The information mentioned in connection with the invention regarding the composition and structure of the hot-dip coated high-strength steel strip refers only to the actual steel strip serving as the substrate for the coating.

[0057] According to a preferred embodiment of the hot-dip coated high-strength steel strip according to the invention, the area proportion of special Z3 grain boundaries with a maximum deviation of 10° to the Z3 orientation relationship of 60° <111> , based on the total grain boundary area for large angle grain boundaries with a disorientation angle > 15°, less than 30% of the total grain boundaries.

[0058] Furthermore, it is advantageously provided that the hot-dip coated high-strength steel strip has a yield strength ratio R p o.2 / R m of < 0.87 and has a bake hardening value BH2 of > 25 MPa.

[0059] According to a further preferred embodiment of the hot-dip coated high-strength steel strip according to the invention, the microstructure of the actual steel strip comprises at least the following components: 8-16% by volume of residual austenite, >10 and <40% by volume of ferrite, and at least 50% by volume of bainite, tempered martensite, and fresh martensite. The "actual steel strip" refers to the steel strip serving as the substrate for the coating.

[0060] According to yet another preferred embodiment of the hot-dip coated high-strength steel strip according to the invention, the structure of the actual steel strip has at least two of the following properties:

[0061] - the proportion of bainite and fresh martensite in volume% is in total greater than the proportion of tempered martensite in volume%,

[0062] - in which the proportion of granular bainite in volume % is higher than the proportion of lower bainite in volume % and

[0063] - there is a proportion of at least 2% by volume of fresh martensite in the total microstructure.

[0064] The effect of the elements in the inventive high-strength steel strip with plasticity induced by microstructural transformation is described in more detail below. Accompanying elements are unavoidable and are taken into account in the analysis concept with regard to their effect, where necessary.

[0065] Trait elements are elements that are already present in iron ore or that enter the steel during production. Due to their predominantly negative effects, they are generally undesirable. Attempts are made to remove them down to a tolerable level or to convert them into less harmful forms. Nitrogen (N) is a tramp element from steel production. Steels with free nitrogen are prone to severe aging. The nitrogen diffuses onto dislocations even at low temperatures and blocks them. It thus causes an increase in strength combined with a rapid loss of toughness. Binding the nitrogen in the form of nitrides is possible by alloying with aluminum or titanium. For the reasons stated above, the optional nitrogen content is limited to < 0.016 wt.% or to amounts unavoidable during steel production.In particular, a nitrogen content of maximum 0.008 wt% is advantageous to avoid harmful AIN excretions.

[0066] Like phosphorus, sulfur (S) is bound as a trace element in iron ore. It is undesirable in steel (with the exception of free-cutting steels) because it is prone to strong segregation and has a strong embrittling effect. Attempts are therefore made to keep the sulfur content in the melt as low as possible (e.g., through deep vacuum treatment). Furthermore, the existing sulfur is converted into the relatively harmless compound manganese sulfide (MnS) by adding manganese. The manganese sulfides are often rolled out in rows during the rolling process and act as nuclei for the transformation. This leads to a pronounced, row-like microstructure, particularly in the case of diffusion-controlled transformation, and can, in the case of pronounced rowing, lead to impaired mechanical properties (e.g., pronounced martensite rows instead of distributed martensite islands, anisotropic material behavior, reduced elongation at break). For the reasons stated above, the sulfur content is limited to < 0.005 wt.% orlimited to quantities unavoidable in steel production.

[0067] Phosphorus (P) is a trace element from iron ore and is dissolved in the iron lattice as a substitution atom. Phosphorus increases hardness through solid solution strengthening and improves hardenability. However, attempts are generally made to reduce the phosphorus content as much as possible because, among other things, its low diffusion rate leads to a strong tendency to segregation and significantly reduces toughness. The accumulation of phosphorus at grain boundaries causes grain boundary fractures. In addition, phosphorus raises the transition temperature from tough to brittle behavior to as much as 300 °C. During hot rolling, near-surface phosphorus oxides at the grain boundaries can lead to fracture cracking. The negative effects of phosphorus can be partially compensated by adding small amounts of boron to the alloy. It is assumed that boron increases grain boundary cohesion and reduces phosphorus segregation at the grain boundaries.However, in some steels, it is used in small amounts (<0.1%) as a microalloying element due to its low cost and significant strength enhancement, for example, in high-strength IF (interstitial-free) steels. For the reasons stated above, the phosphorus content is limited to <0.02% or to amounts unavoidable during steel production.

[0068] Alloying elements are usually added to steel to specifically influence certain properties. A single alloying element can influence different properties in different steels. The relationships are diverse and complex. The effect of alloying elements will be discussed in more detail below.

[0069] Carbon (C) is considered the most important alloying element in steel. Its targeted addition of up to 2.06% iron is what makes it steel. The carbon content is often drastically reduced during steelmaking. Due to its comparatively small atomic radius, carbon is dissolved interstitially in the iron lattice. Its solubility is a maximum of 0.02% in alpha iron and a maximum of 2.06% in y-iron. In dissolved form, carbon significantly increases the hardenability of steel. Due to the varying solubility, pronounced diffusion processes are necessary during phase transformation, which can lead to very different kinetic conditions. Furthermore, carbon increases the thermodynamic stability of austenite, which is reflected in the phase diagram by an expansion of the austenite region at lower temperatures and enables the stabilization of higher contents of retained austenite in the microstructure at room temperature.As the forced carbon content in martensite increases, so does the lattice distortion and, with it, the strength of the diffusion-free phase. To ensure sufficient strength and retained austenite content, the minimum C content is set at 0.15 wt.%. A minimum C content of 0.16 wt.% is particularly advantageous for elongation, increasing the C content in the retained austenite and providing greater stability. Since the solubility and thus the effectiveness of Nb decreases with increasing C content, the maximum C content in the steel according to the invention is limited to 0.205 wt.%. Excessively high C contents also typically prove detrimental to weldability and liquid metal embrittlement. Aluminum (Al) is typically added to steel to bind the oxygen and nitrogen dissolved in the iron. The oxygen and nitrogen are thus converted into aluminum oxides and aluminum nitrides.These precipitates can cause grain refinement by increasing the nucleation sites, thus improving toughness and strength properties. In the dissolved state, aluminum, like silicon, shifts ferrite formation to shorter times, thus enabling the formation of sufficient amounts of ferrite. It also suppresses carbide formation, thus leading to a delayed transformation of austenite. For this reason, Al is also used as an alloying element in retained austenitic steels to replace some of the silicon with aluminum. The reason for this approach is that Al is less critical for the galvanizing reaction than Si. However, Al can be detrimental to hot ductility or castability during continuous casting. Al also causes an undesirable increase in the Ac3 transformation temperature. The Al content is therefore limited to 0.2 wt% to a maximum of 0.7 wt%, in particular 0.20 wt% to a maximum of 0.70 wt%.In particular, the AI ​​content may be limited to a maximum of 0.5 wt%, especially 0.50 wt%, in order to avoid harmful AIN precipitations that can lead to AIN embrittlement.

[0070] Silicon (Si) increases the strength and yield strength ratio of ferrite through solid solution strengthening with only a slight decrease in elongation at break. Another important effect is that silicon shifts the formation of ferrite to shorter times, thus enabling ferrite to form before quenching. Ferrite formation enriches and stabilizes the austenite with carbon. At higher contents, silicon noticeably stabilizes the austenite in the lower temperature range, particularly in the area of ​​bainite formation, by preventing carbide formation. During hot rolling, highly adhesive scale can form at high silicon contents, which can impair further processing. During continuous galvanizing, silicon can diffuse to the surface during annealing and form film-like oxides, either alone or together with manganese.These oxides impair galvanizability by impairing the galvanizing reaction (iron dissolution and inhibition layer formation) when the steel strip is immersed in the zinc melt. This manifests itself in poor zinc adhesion and ungalvanized areas. However, good galvanizability of the steel strip and good zinc adhesion can be ensured by appropriate furnace operation with an adjusted moisture content in the annealing gas and / or a low Si / Mn ratio and / or the use of moderate amounts of silicon. For the reasons stated above, the minimum Si content is set at 0.50 wt.% and the maximum Si content at 0.90 wt.%.

[0071] Manganese (Mn) is added to almost all steels for desulfurization to convert harmful sulfur into manganese sulfides. Furthermore, manganese increases the strength of the ferrite through solid solution strengthening and shifts the transformation to lower temperatures. A primary reason for alloying with manganese is the significant improvement in hardenability. Due to the diffusion hindrance, the pearlite and bainite transformation is shifted to longer times and the martensite initiation temperature is lowered. Like silicon, manganese tends to form oxides on the steel surface during annealing. Depending on the annealing parameters and the contents of other alloying elements (particularly Si and Al), manganese oxides (e.g., MnO) and / or mixed Mn oxides (e.g., Mn2SiO4) may occur. However, manganese is considered less critical at a low Si / Mn or Al / Mn ratio, since globular oxides rather than oxide films are more likely to form.However, high manganese contents can negatively affect the appearance of the zinc layer and zinc adhesion. The Mn content is therefore set at 1.9 wt% to 2.6 wt%, particularly 1.90 wt% to 2.60 wt%, and preferably only up to 2.40 wt%, to avoid Mn lineation.

[0072] Chromium (Cr): The addition of chromium primarily improves hardenability. In the dissolved state, chromium postpones the pearlite and bainite transformation to longer times and simultaneously lowers the martensite initiation temperature. Another important effect is that chromium significantly increases tempering resistance, so that there is almost no loss of strength in the zinc bath. Chromium is also a carbide former. If chromium is present in carbide form, the austenitizing temperature before hardening must be set high enough to dissolve the chromium carbides. Otherwise, the increased nuclei count can impair hardenability. Chromium also tends to form oxides on the steel surface during annealing, which can deteriorate the galvanizing quality. The Cr content is therefore set at values ​​of 0.2 to 0.5 wt.%, in particular 0.20 to 0.50 wt.%.

[0073] Molybdenum (Mo): Molybdenum is added in a similar way to chromium to improve hardenability. The pearlite and bainite transformations are accelerated, and the martensite initiation temperature is lowered. Molybdenum also significantly increases tempering resistance, so that no loss of strength is to be expected in the zinc bath. It also increases the strength of the ferrite through solid solution strengthening. The Mo content is alloyed depending on the dimensions, the plant configuration, and the microstructure. By slowing carbon diffusion, Mo can also counteract the enrichment of carbon in the residual austenite. High Mo contents also lead to high strength in the hot-rolled strip, which negatively affects cold-rollability. For these reasons, the Mo content is set at up to 0.15 wt.%.

[0074] Copper (Cu): The addition of copper can increase tensile strength and hardenability. In combination with nickel, chromium, and phosphorus, copper can form a protective oxide layer on the surface, which can significantly reduce the corrosion rate. In combination with oxygen, copper can form harmful oxides at the grain boundaries, which can have negative effects, especially in hot forming processes. The optional copper content is therefore limited to 0.01 to 0.3 wt.%.

[0075] Nickel (Ni): Nickel can increase tensile strength and hardenability. However, in combination with oxygen, nickel can form harmful oxides at the grain boundaries, which can have negative effects, especially on hot forming processes. The optional nickel content is therefore limited to 0.01 to 0.5 wt.%.

[0076] Microalloying elements are generally added in very small quantities (<0.1%). Unlike alloying elements, they act primarily through precipitation, but can also influence properties in dissolved form. Despite the small amounts added, microalloying elements strongly influence the manufacturing conditions as well as the processing and final properties. Carbide and nitride formers that are soluble in the iron lattice are generally used as microalloying elements. The formation of carbonitrides is also possible due to the complete solubility of nitrides and carbides in each other. The tendency to form oxides and sulfides is generally most pronounced with microalloying elements, but is usually deliberately prevented by the addition of other alloying elements. This property can be used positively by binding the generally harmful elements sulfur and oxygen.However, setting can also have negative effects if there are no longer enough microalloying elements available for carbides to form. Typical microalloying elements are vanadium, titanium, niobium, and boron. These elements can dissolve in the iron lattice and form carbides or nitrides with carbon and nitrogen.

[0077] Titanium (Ti) forms very stable nitrides (TiN) and sulfides (TiS2) even at high temperatures. Depending on the nitrogen content, some of these only dissolve in the melt. If the resulting precipitates are not removed with the slag, the high temperature at which they form causes them to form coarse particles in the material, which are generally not beneficial for the mechanical properties. A positive effect on toughness is achieved through the binding of free nitrogen and oxygen. Titanium protects other dissolved microalloying elements such as niobium from binding by nitrogen, allowing them to optimally develop their effect. Titanium also helps prevent harmful AlN precipitates, which can lead to AlN embrittlement in the steel according to the invention due to the comparatively high Al content.Unbound titanium forms titanium carbides at temperatures above 1150 °C and can thus cause grain refinement (inhibition of austenite grain growth, grain refinement through delayed recrystallization, and / or increased nuclei during a- / y-transformation) as well as precipitation hardening. The optional Ti content therefore ranges from 0.005 to 0.060 wt.%.

[0078] Niobium (Nb) typically results in strong grain refinement, as it is the most effective of all microalloying elements in delaying recrystallization and also inhibiting austenite grain growth. A further effect of niobium is the delay of the a- / y-transformation and the lowering of the martensite initiation temperature in the dissolved state. In principle, the addition of niobium is limited until its solubility limit is reached. While this limits the amount of precipitates, exceeding it primarily results in early precipitation formation with relatively coarse particles. Precipitation hardening can therefore be particularly effective in steels with low C content (higher supersaturation possible) and in hot forming processes (deformation-induced precipitation). As previously described, it was found that with the hot-dip coated steel strip according to the invention, Nb can stabilize higher contents of residual austenite.The special effect of Nb in the steel according to the invention will be explained in more detail below. The Nb content is therefore limited to values ​​of 0.01 to 0.06 wt.%, in particular 0.0100 to 0.0600 wt.%, and is particularly effective in the present invention at contents of 0.0200 wt.% and above, and even more advantageously at 0.0300 wt.%. Nb is therefore not optional in the present invention and must be added as an alloy.

[0079] Vanadium (V): Carbide and nitride formation in vanadium only begins at temperatures around 1000 °C or even after the a / y transformation, i.e., much later than in titanium and niobium. Due to the small number of precipitates present in the austenite, vanadium has little grain-refining effect. Austenite grain growth is also not inhibited by the late precipitation of vanadium carbides. Thus, the strength-enhancing effect is based almost exclusively on precipitation hardening. However, in dissolved form, vanadium also delays transformation. One advantage of vanadium is its high solubility in austenite and the large volume fraction of fine precipitates caused by the low precipitation temperature. The optional V content is therefore limited to values ​​of 0.001 to 0.060 wt.%.

[0080] Boron (B) forms nitrides or carbides with nitrogen as well as with carbon; however, this is generally not the desired result. Due to their low solubility, only a small amount of precipitates forms, and these are mostly precipitated at the grain boundaries. An increase in surface hardness is not achieved (except for boronizing, which forms FeB and Fe2B in the surface zone of a workpiece). To prevent nitride formation, attempts are generally made to bind the nitrogen with elements with higher affinity. Titanium, in particular, can ensure the binding of all the nitrogen. In very small amounts in the dissolved state, boron leads to a significant improvement in hardenability. The mechanism of action of boron can be described as follows: under suitable temperature control, boron atoms accumulate at the grain boundaries and, by reducing the grain boundary energy, significantly impede the formation of growth-capable ferrite nuclei.When controlling the temperature, care must be taken to ensure that boron is predominantly distributed atomically within the grain boundary and not present in the form of precipitates due to excessively high temperatures. The effectiveness of boron decreases with increasing grain size and rising carbon content (> 0.8%). A quantity above 60 ppm also causes a decrease in hardenability, as boron carbides act as nuclei at the grain boundaries. Due to its small atomic diameter, boron diffuses extremely well and has a very high affinity for oxygen, which can lead to a reduction in the boron content in areas close to the surface (up to 0.5 mm). In this context, annealing above 1000 °C is not recommended. This is also recommended, as boron can lead to severe coarse grain formation at annealing temperatures above 1000 °C.Boron is an extremely critical element in the continuous hot-dip zinc plating process, as even the smallest amounts, either alone or in combination with manganese, can form film-like oxides on the steel surface during annealing. These oxides passivate the strip surface and prevent the galvanizing reaction (iron dissolution and inhibition layer formation). Whether film-like oxides form depends on both the amount of free boron and manganese and the annealing parameters used (e.g., moisture content in the annealing gas, annealing temperature, annealing time). Higher manganese contents and longer annealing times tend to lead to globular and less critical oxides. By increasing the moisture content in the annealing gas, it is also possible to reduce the amount of boron-containing oxides on the steel surface. For the reasons stated above, the boron content is kept as low as possible and limited to values ​​of up to 0.0010 wt.%.

[0081] A comparison of the alloy composition of the reference steels and the example steels according to the invention is shown in Table 1. The respective composition refers to the actual steel strip in the final product, i.e., the substrate for the coating.

[0082]

[0083]

[0084] Table 1 : Chemical composition of the steels tested-A - U

[0085] Table 2 contains the process parameters for different annealing cycles of the continuous annealing with hot-dip coating process step, in which the technological parameters (in the longitudinal direction / rolling direction) are set.

[0086] Table 2: Process parameters for continuous annealing

[0087] The following parameters are noted in Table 2:

[0088] T An : Annealing temperature tAn : Annealing time

[0089] T m : Intermediate temperature

[0090] CR average cooling rate when cooling from T An on T m

[0091] T q : Cooling stop temperature

[0092] TOA : Overaging temperature

[0093] CR2: average cooling rate when cooling from T m on T q toA : Hold time on TOA

[0094] THD : Temperature hot dip coating

[0095] CR3: average cooling rate after hot-dip coating

[0096] Corresponding technological parameters and microstructure composition for the respective steels with different annealing cycles are summarized in Table 3.

[0097] Table 3: Comparison of parameters of the respective hot-dip coated steel strip for the individual steels and the annealing process parameters used

[0098] The following parameters are noted in Table 3:

[0099] R m : Tensile strength

[0100] R P O,2° : Yield strength before tempering

[0101] Rpo,2° / Rm: Yield strength ratio before skin passing

[0102] A g : Uniform stretching

[0103] RA : proportion of retained austenite in the structure

[0104] *A g was determined too low (no standard fracture)

[0105] Fresh and tempered martensite are abbreviated as “fresh M” and “tempered M”.

[0106] To determine the technological parameters, tensile tests were carried out in accordance with DIN EN ISO 6892-1:2020-06 on specimens with a gauge length of 80 mm, which were taken longitudinally to the rolling direction.

[0107] Reference steel A has the required contents of Si, Al and Cr, but is not alloyed with Nb and does not meet the requirement of p = 4.5 x (Si + 0.9Al + Cr) + 200 x Nb > 8. With this steel, particularly at low cooling stop temperatures T qsufficient contents of retained austenite (> 8 vol.%) and ferrite (> 10 vol.%) cannot be achieved.

[0108] The steels B to G, I, and J according to the invention have Si, Al, and Cr contents comparable to those of A and are additionally alloyed with Nb, with Nb contents of 130 to 410 ppm. In these steels, p > 8. These steels are also characterized by a Si+0.9 x Al content of < 1 wt.%. High contents of retained austenite > 8 vol.% and ferrite > 10 vol.% were achieved in these steels. R m x A g was sufficiently high for all annealing cycles with > 8000 MPa %. In particular, steels C to F with p > 10 showed a very high product of tensile strength and uniform elongation of R in all cycles. m x A g > 9000 MPa %.

[0109] Steel E, with the highest Nb content of 410 ppm and p = 13.8, stabilized the highest levels of retained austenite. Consequently, this steel also achieved the highest values ​​for Rm x A g . Furthermore, the achieved tensile strength and the content of retained austenite were almost the same for all annealing cycles with correspondingly different cooling stop temperatures, which proves to be advantageous in large-scale production with regard to process stability and process fluctuations. In particular, the retained austenite content, at 13.7%, 13.8%, and 13.3%, is unexpectedly stable in the cooling stop temperature range from 315 °C to 450 °C (annealing cycles Ia, II, and III), enabling consistent plasticity induced by microstructural transformation. Even at even lower cooling stop temperatures of 290 °C and 260 °C, very high values ​​for R were achieved with a sufficiently high Nb content and p = 12.8. m x A g possible (steel U with annealing cycle Id and le).

[0110] Reference steel H is also alloyed with Nb, but has a C content of > 0.205 wt.%, which is considered detrimental to weldability and LME (liquid metal embrittlement). Furthermore, due to the excessive C content, hardenability was increased to such an extent that a sufficient amount of ferrite (> 10 vol.%) could not be formed during the slow cooling stage.

[0111] Reference steel K has an excessively high Al content of > 0.7. For this reason, sufficient austenitization could not be ensured for this steel during continuous hot-dip galvanizing.

[0112] Reference steel L does not meet the requirement of p = 4.5 x (Si + 0.9Al + Cr) + 200 x Nb > 8. A sufficient amount of retained austenite could not be generated in the microstructure of this steel. The tensile strength with annealing cycle 1a was also too low (< 900 MPa).

[0113] In contrast to the steels according to the invention, reference steels M and N were alloyed with > 0.15 wt.% Mo. A sufficiently high content of residual austenite could not be stabilized in these steels, as Mo significantly slows the diffusion of carbon and counteracts the enrichment of the residual austenite with C. Furthermore, the steels were unsuitable for cold rolling due to high rolling forces and high susceptibility to edge cracking. The Mo content of the steels according to the invention is therefore limited to < 0.15 wt.%.

[0114] The inventive steels O and P have a higher Si content than the inventive steels B to F, G, I, and J. For these steels, Si + 0.9 x Al < 1.2 wt.% applies. Sufficient residual austenite contents and the required technological parameters were achieved for steels O and P.

[0115] Reference steels Q to S are classic complex phase steels with usual low

[0116] Contains (< 5 vol%) of retained austenite in the structure before skin-passing. For these steels, < 6 and R m x A g is well below 8000 MPa%.

[0117] Reference steel T falls below the Al content of > 0.2 weight % required for the inventive design and does not reach the required proportion of residual austenite.

[0118] These and other features and advantages of the present invention will become apparent from the following examples with reference to the accompanying drawings, in which:

[0119] Figure 1 shows a graphical representation of the dependence of the technological parameters (Rm-tensile strength, R p 0.2° proof strength, A g -uniform elongation) and the volume content of retained austenite (RA) from the Nb content (steels A to E) after the process step of continuous annealing and hot-dip coating with temperature cycle la,

[0120] Figure 2 shows a graphical representation of the dependence of the volume content of a) retained austenite (RA) and b) the A g -Uniform elongation of 4.5 x (Si + 0.9 x AI + Cr) + 200 x Nb (steels A to E) after the process step of continuous annealing and hot-dip coating with temperature cycles Ia and II,

[0121] Figure 3 is a graphical representation of a relative change in length dL / LO (elongation) measured with a dilatometer as a result of the phase transformation of the austenite for the steel F according to the invention and reference steel H during continuous annealing in the context of continuous hot-dip galvanizing for the annealing cycle 1a (cooling in the slow cooling section X1 / rapid cooling section X2) as a function of the temperature T and

[0122] Figure 4 Section of a band contrast map (Kikuchi band contrast) measured by electron backscatter diffraction for an area with a) lower bainite from steel A with annealing cycle II and b) granular bainite from steel D with annealing cycle II.

[0123] Figure 1 shows the dependence of the technological parameters R on the example of steels A to E with annealing cycle la m , R P 0.2°, A g and the content of retained austenite from the Nb content or 4.5 x (Si + 0.9 x AI + Cr) + 200 x Nb. With increasing Nb content in combination with Si, AI and Cr there is a decrease of R m , Rpo,2° and an increase in RA and A g The plasticity caused by structural transformation (increase in RA and A g) increases approximately linearly with increasing content of p = 4.5 x (Si + 0.9 x AI + Cr) + 200 x Nb. From a content of Nb > 100 ppm and p > 8, it is ensured that the steel has sufficient plasticity caused by microstructural transformation.

[0124] Figure 2 shows the inventive effect of the content of p = 4.5 x (Si + 0.9 x AI + Cr) + 200 x Nb on increasing the proportion of residual austenite (Fig. 2a) and the uniform elongation (Fig. 2b) for annealing cycles 1a and II of continuous hot-dip galvanizing. As shown in Table 2, the annealing cycles differ in the cooling stop and overaging temperatures of 315 °C (cycle 1a) and 400 °C (cycle II), respectively. A linearly increasing proportion of residual austenite and the uniform elongation was observed with increasing content of p = 4.5 x (Si + 0.9 x AI + Cr) + 200 x Nb. For p > 8, the uniform elongation is still > 8% even for annealing cycle 1a with the lower cooling stop temperature of 315 °C. With increasing content of p = 4.5 x (Si + 0.9 x AI + Cr) + 200 x Nb, the proportions of retained austenite RA for different cooling stop temperatures approach, thus achieving a larger process window for a constant plasticity caused by microstructural transformation.

[0125] The effect of Nb in the steel according to the invention can be attributed to one or more of the following mechanisms during continuous annealing as part of continuous hot-dip galvanizing: (i) Suppression of recrystallization and inhibition of grain growth in austenite at temperatures > 750 °C. (ii) Accelerated formation of ferrite through nucleation of fine Nb precipitates in the slow cooling section and corresponding enrichment of carbon in the austenite (see Fig. 3). The accelerated ferrite formation enables processes with higher process speeds or shorter slow cooling sections to be implemented, where the formation of ferrite would otherwise not be possible. The Nb content also allows the proportion of ferrite to be specifically adjusted, thus allowing the technological properties to be controlled without changing the annealing cycle.iii) Reduction of the martensite start temperature by a fine austenite grain, enrichment of carbon via (ii) and inhibition of martensite nucleation and consequently the possibility of the formation of fine (granular) bainite with retained austenite as a second phase at low cooling stop temperatures. iv) Grain refinement by retarding bainite kinetics during overaging and changing the bainite morphology from lath-like to granular. v) Alignment of the technological properties and the retained austenite content for different cooling stop temperatures, resulting in a wider process window in large-scale production. vi) Higher retained austenite content in the hot-dip coated steel strip after final cooling to ambient temperature. Figure 3 illustrates the transformation behavior of the inventive steel F compared to the reference steel H during cooling.Starting from the austenite region Y, the steel F according to the invention already forms ferrite in the slow cooling section X1 due to the increased Nb content from approx. 735 °C (ferrite transformation region Z), so that the remaining austenite becomes enriched in carbon and approx. 28 ol% ferrite is present in the final structure. In reference steel H, ferrite formation is strongly suppressed by the high C content and comparatively low Nb content and the ferrite content in the final structure is too low at < 10 vol.%. Using the results from Table 3, it was deduced that ferrite can preferentially be formed in the slow cooling section if the following relationship applies: (100C + 10Mn) / [4.5 x (Si + 0.9Al + Cr) + 200xNb] < 4.5 with the respective alloy contents in weight%.

[0126] Table 4 shows the proportion of special Z3 grain boundaries of the Nb-containing steels B to E according to the invention in comparison to the reference steel A without targeted Nb addition.

[0127]

[0128] Table 4: Proportions of special grain boundaries relative to the total grain boundary area for large angle grain boundaries with a

[0129] Disorientation angle > 15°. The proportions of Z3 grain boundaries were determined with a maximum tolerance of 10° to the Z3 orientation relationship. Misorientation refers to the orientation relationship or rotation between two grain orientations. The disorientation angle is the smallest rotation angle of all crystallographically equivalent misorientations.

[0130] Electron backscatter diffraction (EBSD) was used to measure the individual grain orientations of the microstructure, which yield the misorientations of the grain boundaries. Misorientation is defined as the orientation relationship between two grains. With increasing content of p = 4.5 x (Si + 0.9Al + Cr) + 200xNb, the proportion of coherent Z3 grain boundaries decreases (relative to proportions of the total grain boundary area for large-angle grain boundaries with a disorientation angle > 15°, whereby the length fractions of grain boundaries measured with EBSD are usually adopted as area fractions of the grain boundaries in the EBSD maps). In this case, a grain boundary is defined as a Z3 grain boundary if its misorientation deviates by a maximum of 10° from the exact misorientation of 60°. <111> The decrease in the proportion of Z3 grain boundaries is also accompanied by a decrease in the proportion of grain boundaries with a disorientation angle in the range 57 - 63 ° (Table4), which are typically observed as a characteristic 60° peak in the disorientation angle distribution of multiphase steels. Such special Z3 grain boundaries arise, among other things, when massive martensite lamellae form before entering the overaging zone or when lower bainite (parallel bainite laths) develops.

[0131] Figure 4 shows a section of a band contrast map (Kikuchi band contrast) measured by electron backscatter diffraction for an area with a) lower bainite from steel A with annealing cycle II and an area with b) granular bainite from steel D with annealing cycle II. Dark areas indicate grain boundaries and grains with higher dislocation density.

[0132] It was observed that in the steel according to the invention, with increasing content of elements corresponding to the value p, the microstructure development shifts from tempered martensite and / or lower bainite (Fig. 4a) toward a fine granular bainite (Fig. 4b). Due to weak variant selection, granular bainite is characterized by a smaller number of Z3 grain boundaries (weak 60° peak in the disorientation angle distribution), see article "S. Zajac, V. Schwinn, K.-H. Tacke; Mater. Sci. Forum 500-501 (2005) 387-394." The advantage of this type of bainite is that, due to its formation mechanism, granular bainite leads to the formation of a carbon-rich second phase. If the carbon concentration of this second phase is sufficiently high, retained austenite is stabilized. For this reason, it is advantageous for the steel according to the invention to set a lower proportion of Z3 grain boundaries in the final microstructure.In order to stabilize the highest possible contents of residual austenite and to be able to use low cooling stop temperatures, a proportion of Z3 grain boundaries of < 30% is advantageous in the steel according to the invention.

Claims

Patent claims 1. Method for producing a hot-dip coated high-strength steel strip with plasticity induced by microstructure transformation, comprising the following steps: - Production of a hot-rolled steel strip, in particular from a slab heated to over 1200 °C, consisting of the following elements in weight %: C: from 0.15 to 0.205, Mn: from 1.9 to 2.6, AI: from 0.2 to 0.7, Si: from 0.5 to 0.9, Cr: from 0.2 to 0.5, Note: from 0.01 to 0.06, Mo: < 0.15, B: < 0.001 , P: < 0.02, S: < 0.005, and optionally one or more of the following elements in weight %: Ti: 0.005 to 0.060, V: 0.001 to 0.060, N: 0.0001 to 0.016, Ni: 0.01 to 0.5 and Cu: 0.01 to 0.3, Residual iron, including usual steel-associated elements, where for a value p = 4.5 x ([Si] + 0.9 x [Al] + [Cr]) + 200 x [Nb], where [Si], [Al], [Cr] and [Nb] are the proportions of the corresponding elements in weight %, 8 < p < 16 holds. - Pickling and optional cold rolling of the hot-rolled steel strip to a cold-rolled steel strip, - Subsequent continuous annealing as part of a continuous hot-dip coating process of the cold- or hot-rolled steel strip at a maximum temperature between 750 °C and 950 °C inclusive, in particular between 800 and 870 °C, for a total duration of 10 s to 1200 s, in particular from 50 s to 650 s, - subsequent cooling of the cold- or hot-rolled steel strip to an intermediate temperature in a temperature range of 620 to 760 °C with an average cooling rate CRi of up to 10 K / s, - subsequent further cooling of the cold- or hot-rolled steel strip from the intermediate temperature to a cooling stop temperature in a temperature range between 200 °C and 450 °C inclusive, in particular between 280 °C and 450 °C inclusive, with an average cooling rate CR2 > CRi and a maximum of 150 K / s, and subsequent holding of the temperature in the temperature range between 200 °C and 450 °C inclusive, in particular between 280 °C and 450 °C inclusive, for 25 to 500 s, - followed by hot-dip coating of the cold- or hot-rolled steel strip at a temperature between 380 and 500 °C and - followed by final cooling of the hot-dip coated cold- or hot-rolled steel strip at an average cooling rate of 1 K / s to 50 K / s to ambient temperature.

2. Method according to claim 1 , characterized in that 10 < p < 16 for the value p and the expression [Si] + 0.9 x [AI] < 1 ,2, in particular [Si] + 0.9 x [AI] < 1 ,0, applies, wherein [Si] and [AI] are the proportions of the corresponding elements in weight % of the hot-rolled steel strip.

3. Method according to claim 1 or 2, characterized in that the Nb content in ppm in the hot-rolled steel strip is > 200, in particular > 300.

4. Method according to one of claims 1 to 3, characterized in that the proportion of Mn in the hot-rolled steel strip is between 1.95 and 2.4 wt%, the proportion of C in the hot-rolled steel strip is at least 0.16 wt% and preferably the expression (100[C] + 10[Mn]) / (4.5 x ([Si] + 0.9[Al] + [Cr]) + 200x[Nb]) < 4.5 applies, where [Si], [Al], [Cr], [C] and [Mn] are the proportions of the corresponding elements in the hot-rolled steel strip in wt%.

5. Method according to one of claims 1 to 4, characterized in that the sum of the proportions of the elements Cr and Mo in the hot-rolled steel strip in weight % is less than 0.5, i.e. [Cr] + [Mo] < 0.

5.

6. Method according to one of claims 1 to 5, characterized in that the intermediate temperature is in a temperature range of 650 to 730 °C and the steel strip has a microstructure with at least 10 volume% when this temperature is reached. It contains ferrite.

7. Method according to one of claims 1 to 6, characterized in that the cooling stop temperature is < 400 °C, preferably < 350 °C, and after final cooling to ambient temperature more than 8 volume % austenite is present in the microstructure, wherein the temperature at which the cold- or hot-rolled steel strip is held before hot-dip coating is < 400 °C, preferably < 350 °C.

8. Method according to any one of claims 1 to 7, characterized in that the hot-dip coated high-strength steel strip has a tensile strength R m of at least 900 MPa and a uniform elongation A g exhibits at least 8%.

9. Method according to any one of claims 1 to 8, characterized in that the hot-dip coated steel strip is tempered with a rolling degree of at most 2%, wherein the tempering process reduces the R p The 0.2 yield strength increases by at least 20 MPa.

10. Method according to any one of claims 1 to 9, characterized in that the hot-rolled steel strip has a Ti content of at least 0.005 wt%, a N content of at most 0.008 wt%, an Al content of at most 0.5 wt%, and TiN and TiAlN particles with a diameter of > 0.96 pm in a total area fraction of at least 1 pm 2 / mm 2 on a measuring surface of at least 100 mm 2in the slab before reheating and on a measuring surface of at least 20 mm 2 are present in hot-dip coated high-strength steel strip.

11. Hot-dip coated high-strength steel strip with plasticity induced by microstructure transformation, in particular produced by a method according to one of claims 1 to 10, consisting of the following elements in weight %: C: from 0.15 to 0.205, Mn: from 1.9 to 2.6, AI: from 0.2 to 0.7, Si: from 0.5 to 0.9, Cr: from 0.2 to 0.5, Note: from 0.01 to 0.06, Mo: < 0.15, B: < 0.001 , P: < 0.02, S: < 0.005, and optionally one or more of the following elements in weight %: Ti: 0.005 to 0.060, V: 0.001 to 0.060, N: 0.0001 to 0.016, Ni: 0.01 to 0.5 and Cu: 0.01 to 0.3, Residual iron, including usual steel-associated elements, where for a value = 4.5 x ([Si] + 0.9 x [Al] + [Cr]) + 200 x [Nb], where [Si], [Al], [Cr] and [Nb] are the proportions of the corresponding elements in weight %, 8 < p < 16 holds, where the steel strip is a product of R m -Tensile strength and uniform elongation A g exhibits a pressure greater than 8000 MPa %, particularly greater than 9000 MPa %, and particularly advantageously between 9900 and 13000 MPa %,.

12. Hot-dip coated high-strength steel strip according to claim 11, characterized in that the area fraction of special Z3 grain boundaries with a maximum deviation of 10° to the Z3 orientation relation of 60° <111 >, based on the total grain boundary area for large-angle grain boundaries with a disorientation angle > 15°, is less than 30%.

13. Hot-dip coated high-strength steel strip according to claim 11 or 12, characterized in that the steel strip has a yield strength ratio R p 0.2 / Rm has a coefficient of < 0.87 and a bake-hardening value BH2 of > 25 MPa.

14. Hot-dip coated high-strength steel strip according to one of claims 11 to 13, characterized in that the microstructure of the hot-dip coated high-strength steel strip comprises at least the following components: 8-16 volume% retained austenite, > 10 and < 40 volume% ferrite, at least a sum of 50 volume% of bainite, tempered martensite and fresh martensite.

15. Hot-dip coated high-strength steel strip according to claim 14, characterized in that the microstructure of the hot-dip coated high-strength steel strip has at least two of the following properties: - the total proportion of bainite and fresh martensite by volume is greater than the Percentage of tempered martensite by volume, - where the proportion of granular bainite in volume % is higher than the proportion of lower bainite in volume % and - There is a proportion of at least 2% by volume of fresh martensite in the entire structure.

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