Method for producing a high-strength flat steel product having a multiphase microstructure, and corresponding high-strength flat steel product

Abstract

The invention relates to a method for producing a high-strength flat steel product having an Rm tensile strength of at least 700 MPa on the basis of a cold-rolled and continuously annealed steel strip having a multiphase microstructure, the method comprising the following steps: (i) producing a hot-rolled steel strip consisting of the following elements in percent by weight: C: from ≥ 0.13 to ≤ 0.20, Mn: from ≥ 1.40 to ≤ 2.40, Al: from ≥ 0.06 to ≤ 0.60, Si: from ≥ 0.03 to ≤ 0.70, Mo: < 0.10, Nb: from ≥ 0.01 to ≤ 0.1, P: < 0.02, S: < 0.05, Ca: ≤ 0.004, B: ≤ 0.001, Cr: ≤ 0.50, Ni: ≤ 0.10, Cu: ≤ 0.10, and optionally one or more elements from: N: ≥ 0.001 to ≤ 0.10, Ti: ≥ 0.01 to ≤ 0.10, V: ≥ 0.01 to ≤ 0.20, the remainder iron, including typical steel-accompanying elements, the sum of the proportions in percent by weight of Si and Al being in a range from ≥ 0.60 to ≤ 1.20, (ii) cold-rolling the hot-rolled steel strip to form a cold-rolled steel strip, (iii) subsequently continuously annealing the cold-rolled steel strip, in particular in the context of a continuous hot-dip coating process, at a maximum temperature between 750°C and, inclusive, 950°C for the total duration of 10 s to 1000 s, (iv) subsequently cooling the continuously annealed, cold-rolled steel strip, in one step or in a plurality of steps, to a cooling stop temperature in a temperature range from 250°C to 450°C at an average cooling rate CR1 of 0.5 to 100 K / s and subsequently holding the temperature in the temperature range from 250°C to 450°C for 30 to 600 s and optionally subsequently performing hot-dip coating of the cold-rolled steel strip at a temperature between 380 and 500°C in the context of the continuous hot-dip coating process, and (v) subsequently performing final cooling of the continuously annealed, cold-rolled steel strip to ambient temperature at an average cooling rate of 2 K / s to 100 K / s in order to create the high-strength flat steel product. The invention also relates to a corresponding high-strength flat steel product.

Description

[0001] Process for producing a high-strength flat steel product with a multi-phase structure and corresponding high-strength flat steel product

[0002] The present invention is based on a process for producing a high-strength flat steel product with an R m -Tensile strength of at least 700 MPa based on a cold-rolled and continuously annealed steel strip.

[0003] The present invention is further based on a corresponding high-strength flat steel product with R m -Tensile strength of at least 700 MPa, which has a steel strip with a multi-phase structure.

[0004] The document EP 3663425 A1 describes a corresponding method for producing a high-strength steel flat product with the following steps: (i) producing a hot-rolled steel strip, (ii) cold-rolling the hot-rolled steel strip to a cold-rolled steel strip, (iii) subsequently continuously annealing the cold-rolled steel strip as part of a continuous hot-dip coating process at a maximum temperature between 750 °C and 950 °C inclusive, (iv) subsequently cooling the continuously annealed, cold-rolled steel strip to a cooling stop temperature in a temperature range from 440 °C to 470 °C and subsequent hot-dip coating of the cold-rolled steel strip at a temperature between 430 °C and 480 °C as part of the continuous hot-dip coating process and (v) subsequently finally cooling the continuously annealed, cold-rolled steel strip to ambient temperature to create the high-strength flat steel product.In one of the examples, the steel strip has a composition of the following elements in weight percent: C: 0.127, Mn: 1.53, Al: 0.011, Si: 0.67, Nb: 0.074, with the sum of the weight percent Si and Al being 0.68. This document further describes the resulting high-strength flat steel product in the form of a hot-dip galvanized steel strip. The tensile strength of this hot-dip galvanized steel strip in this example exceeds 800 MPa, and the steel strip has a multi-phase microstructure, which, in addition to bainite, tempered, and fresh martensite, contains over 80% ferrite by volume and a maximum of 2% retained austenite by volume.

[0005] In high-strength dual-phase and multi-phase steels, the achievable elongation or ductility, and thus the formability, generally 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 through plasticity induced by microstructural transformation. This requires a sufficient content of retained austenite in the microstructure.

[0006] A higher Si content can stabilize a proportion of residual austenite in the microstructure, thereby achieving higher elongations through the well-known TRIP effect (TRIP: Transformation Induced Plasticity). However, if continuous galvanizing is also desired, the presence of Si significantly impairs galvanizability by impairing the galvanizing reaction when the steel strip is immersed in the zinc melt. At high Si contents, strongly adhering scale can also form even in the hot strip, which complicates surface quality and further processing. In some cases, Si can be substituted with higher Al contents to stabilize residual austenite. However, excessively high Al contents can be detrimental to hot ductility and castability during continuous casting. In addition, Si and Al increase the Ac3 transformation temperature, so high annealing temperatures are necessary for continuous galvanizing, which should be avoided for process and cost reasons.The sum of the Si + Al contents should therefore be as low as possible when producing hot-dip galvanized steel strip. Low combined Si and Al contents are also beneficial to prevent liquid metal embrittlement.

[0007] In addition to the effects of Si and Al, patent EP 2 831 299 B2 describes the advantageous use of Cr in combination with Si and Al to produce higher contents of retained austenite with a minimum content of Si + 0.8Al + Cr of 1.4 wt.% for 5 - 20 vol.% retained austenite. In addition to high costs and limitations in weldability, a high Cr content has the disadvantage that, under conventional process control, the hot-rolled strip can become too hard due to increased hardenability to subsequently be cold-rolled into a cold-rolled strip. By increasing hardenability, Cr also has a significant impact on microstructure development during continuous annealing. Therefore, Cr cannot be added to an arbitrarily high 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.

[0008] The document JP 2001 303226 A describes a process for producing a flat steel product in the form of a hot-dip galvanized and subsequently annealed steel strip (“galvannealed steel sheet”) based on a steel strip consisting of the following elements in % by weight: C: 0.05 to 0.20, Si: 0.02 to 1.50, Mn: 0.50 to 3.0, P: 0.10 or less, Al: 0.05 to 2.0, Ti and / or Nb in a range from 0.007 to 0.25, Nb: 0.005 to 0.25 and Ti (%) +2 Nb (%): 0.30 (%) or less, the remainder being iron and usual elements accompanying steel. In one of the examples, the steel strip has a composition of the following elements in weight%: C: 0.20, Mn: 2.00, Al: 0.10, Si: 0.43, Nb: 0.060, Ti: 0.050, P: 0.010, S: 0.003 and N: 0.005, the remainder being iron and usual steel-accompanying elements, the sum of the proportions in weight% of Si and Al being accordingly 0.53.In this example, the resulting flat steel product has a tensile strength below 700 MPa and the microstructure of the steel strip has an austenite content of only 2 volume%.

[0009] The invention is based on the object of providing a method for producing a high-strength flat steel product and a corresponding high-strength flat steel product, in which the high-strength flat steel product has overall high strength in combination with high uniform elongation or formability.

[0010] The problem is solved by the subject matter of the independent patent claims. Advantageous developments of the invention are revealed by the features of the subclaims.

[0011] In the process according to the invention for producing a high-strength flat steel product with an R m-Tensile strength of at least 700 MPa based on a cold-rolled and continuously annealed steel strip having a multi-phase structure, the process is intended to comprise the following steps:

[0012] (i) Producing a hot-rolled steel strip consisting of the following elements in % by weight: C: from > 0.13 to < 0.20, Mn: from > 1.40 to < 2.40, Al: from > 0.06 to < 0.60, Si: from > 0.03 to < 0.70, Mo: < 0.10, Nb: from > 0.01 to < 0.10, P: < 0.02, S: < 0.05, Ca: < 0.004, B: < 0.001, Cr: < 0.50, Ni: < 0.10, Cu: < 0.10, and optionally one or more elements from: N: > 0.001 to < 0.10, Ti: > 0.01 to < 0.10, V: > 0.01 to < 0.20, balance iron, including usual steel-accompanying elements, where the sum of the proportions in weight% of Si and Al is in a range of > 0.60 to < 1.20,

[0013] (ii) cold rolling the hot-rolled steel strip into a cold-rolled steel strip,

[0014] (iii) subsequently continuously annealing the cold-rolled steel strip, in particular in a continuous hot-dip coating process, at a maximum

[0015] Temperature between 750 °C and 950 °C inclusive for a total duration of 10 s to 1000 s,

[0016] (iv) subsequently cooling the continuously annealed, cold-rolled steel strip in one or more stages to a cooling stop temperature in a temperature range of 250 °C to 450 °C with an average cooling rate CRi of 0.5 to 100 K / s and then holding the temperature in the temperature range of 250 °C to 450 °C for 30 to 600 s and optionally subsequent hot-dip coating of the cold-rolled steel strip at a temperature between 380 and 500 °C as part of the continuous hot-dip coating process and

[0017] (v) Subsequently, the continuously annealed, cold-rolled steel strip is finally cooled to ambient temperature at an average cooling rate of 2 K / s to 100 K / s to create the high-strength flat steel product. For the purposes of the present invention, the ambient temperature T is understood to be a temperature of 60 °C or less.

[0018] For the steel strip composition described here, in which the sum of the Si and Al contents is limited to 1.20 wt.%, the stated manufacturing steps and parameters result in a high-strength flat steel product that is nevertheless easy to form. These properties are quantitatively determined within the scope of the present invention using a R m -Tensile strength of at least 700 MPa and a product of R m - Tensile strength and uniform elongation A gwhich lies in the range between 9000 and 14000 MPa%. The fact that these properties are achieved so reliably in the flat steel product is primarily due to the fact that, for the composition in question, the selection of a relatively low cooling stop temperature results in a multiphase structure in which the residual austenite content lies between 5 and 15 volume percent.

[0019] Since the sum of the Si and Al contents is limited to 1.20% by weight in the described composition of the steel strip, the following positive properties also result: relatively good galvanizability, surface quality, machinability in the production of the flat steel product, further processability of the produced flat steel product, etc.

[0020] At this point, it should be noted that steps (iii) - (v) in the context of a continuous hot-dip coating process exclude so-called "galvanealing," since the final cooling process of step (v), with its cooling to the ambient temperature of a maximum of 60 °C, immediately follows the hot-dip coating process of step (iv). Furthermore, it should be noted that the terms cooling stop temperature and cooling temperature TC are used synonymously.

[0021] According to a preferred embodiment of the invention, the cooling stop temperature to which the cold-rolled steel strip is cooled (in step (iv)) is in a temperature range of 250 °C to 400 °C, preferably in a temperature range of 250 °C to 350 °C. A cooling stop temperature in these temperature ranges ensures that a residual austenite content of 5 volume % and 15 volume % in the microstructure of the steel strip of the flat steel product is reliably established.

[0022] According to a further preferred embodiment of the invention, the cooling of the cold-rolled steel strip takes place in several stages over at least one intermediate temperature in a temperature range of 620 to 760 °C. Such multi-stage cooling allows the adjustment of the ferrite content in the microstructure and thus enables a targeted adjustment of the strength.

[0023] With regard to cold rolling of the hot-rolled steel strip, it is advantageous to perform this cold rolling with a cold rolling degree of less than 80%. A higher cold rolling degree requires such high forming forces that there is a risk of the strip breaking during cold rolling.

[0024] According to yet another preferred embodiment of the invention, it is provided that, to achieve a desired roughness of the flat steel product, the cold-rolled steel strip is subjected to skin-passing and / or, to achieve a desired yield strength, the cold-rolled steel strip is subjected to stretch-bend straightening, resulting in a total degree of deformation of up to a maximum of 1%. Higher degrees of deformation would reduce the proportion of residual austenite in the microstructure, since the deformation itself triggers part of the trip effect, which is undesirable.

[0025] The following procedure is used to produce the hot-rolled steel strip:

[0026] (i) heating an ingot or slab having the composition specified for the hot-rolled steel strip to a temperature above 1100 °C,

[0027] (ii) subsequent hot rolling of the ingot or slab above a temperature of 800 °C to the hot-rolled steel strip and

[0028] (iii) Coiling the hot-rolled steel strip at a temperature between 480 °C and 750 °C.

[0029] It is further preferably provided that in the resulting high-strength flat steel product, the multi-phase structure of the steel strip contains, in addition to bainite, tempered, self-tempered and fresh martensite, a proportion of 20 volume % to 80 volume % ferrite and between 5 volume % and 15 volume % residual austenite.

[0030] In the high-strength flat steel product according to the invention with a steel strip consisting of the following elements in weight %: C: from > 0.13 to < 0.20, Mn: from > 1.40 to < 2.40, Al: from > 0.06 to < 0.60, Si: from > 0.03 to < 0.70, Mo: < 0.10, Nb: from > 0.01 to < 0.10, P: < 0.02, S: < 0.05, Ca: < 0.004, B: < 0.001, Cr: < 0.50, Ni: < 0.10, Cu: < 0.10, and optionally one or more of the elements from: N: > 0.001 to < 0.10, Ti: > 0.01 to < 0.10 and V: > 0.01 to < 0.20, the remainder iron, including usual steel-accompanying elements, whereby the sum of the proportions in weight% of Si and Al is in a range of > 0.60 to < 1.20%, it is intended that this steel strip furthermore has a multi-phase structure, which in addition to bainite, tempered and fresh martensite contains a proportion of 20 volume% to 80 volume% ferrite and between 5 volume% and 15 volume% residual austenite, whereby the steel flat product furthermore has the following properties: an R m-Tensile strength of at least 700 MPa and a product of R m -Tensile strength and uniform elongation A g , which is in the range between 9000 and 14000 MPa %.

[0031] This flat steel product is both high-strength and easy to form. These properties are quantitatively tested within the scope of the present invention using an Rm tensile strength of at least 700 MPa and a product of R m -Tensile strength and uniform elongation A g , which ranges between 9000 and 14000 MPa%. Since the sum of the Si and Al contents is limited to 1.2% by weight in the described composition of the steel strip, the following positive properties also result: relatively good galvanizability, surface quality, machinability during the production of the flat steel product, further processability of the resulting flat steel product, etc.

[0032] This high-strength flat steel product is manufactured, in particular, using the aforementioned method. According to a preferred embodiment of the high-strength flat steel product according to the invention, the microstructure of the steel strip has a ferrite content of 40 to 80% by volume. The limited ferrite content allows for reliable adjustment of strength and elongation values.

[0033] According to a further preferred embodiment of the invention, the flat steel product has a zinc-based coating on the steel strip created by a hot-dip coating process. This leads to improved corrosion resistance and is largely standard in automotive construction.

[0034] Finally, it is advantageously provided that the product of R m -Tensile strength and uniform elongation A gbetween 10,000 and 14,000 MPa%. This is achieved by an analysis according to the invention and a procedure with suitable subcooling.

[0035] The following describes the importance of the individual alloying elements in steel. 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 following will discuss the effects of alloying elements in more detail.

[0036] 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, consequently, the strength of the diffusion-free phase. To ensure sufficient strength and retained austenite content, the minimum C content is set at 0.13 wt.%. Excessively high C contents also typically prove detrimental to weldability and liquid metal embrittlement. Furthermore, they lead to higher strengths and reduce the solubility of microalloying elements such as Ti, V, and Nb. For these reasons, the maximum C content in the steel according to the invention is limited to 0.20 wt.%.

[0037] Aluminum (Al) is typically added to steel to bind the oxygen and nitrogen dissolved in iron. This creates aluminum oxides and aluminum nitrides. These precipitates can increase the nucleation sites and thus refine grains, improving toughness and strength. In its dissolved state, aluminum, like silicon, shifts ferrite formation to shorter times, allowing the formation of sufficient amounts of ferrite. It also suppresses carbide formation, thus delaying the 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 and castability during continuous casting.Al also causes an undesirable increase in the Ac3 transformation temperature. The Al content is therefore limited to a range of 0.06 wt% to a maximum of 0.60 wt%. In particular, the Al content can be limited to a maximum of 0.50 wt% to avoid harmful AlN precipitation, which can lead to AlN embrittlement. However, to keep the required silicon content low and to substitute it with aluminum, a content greater than 0.20 wt% is preferred.

[0038] 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 0.03 wt.% and the maximum Si content at 0.70 wt.%.

[0039] 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 coating and zinc adhesion. The Mn content is therefore set at 1.40 to 2.40 wt%.

[0040] 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 0.10 wt.%.

[0041] 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 impair the galvanizing quality. The optional Cr content is therefore set at values ​​up to 0.50 wt.%.

[0042] 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 0.10 wt.%.

[0043] 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 0.10 wt.%.

[0044] 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.

[0045] 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 0.10 wt.%.

[0046] 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 up to 0.05 wt.%.

[0047] Vanadium (V): The formation of carbide and nitride 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. Vanadium therefore has little grain-refining effect due to the small number of precipitates present in the austenite.

[0048] Austenite grain growth is 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 retards transformation. One advantage of vanadium is its high solubility in austenite and the high volume fraction of fine precipitates caused by the low precipitation temperature. The optional V content is therefore limited to values ​​of up to 0.20 wt.%.

[0049] Boron (B) forms nitrides or carbides with nitrogen as well as with carbon; however, this is generally not the desired result. Firstly, due to their low solubility, only a small amount of precipitates forms, and secondly, these are mostly precipitated at the grain boundaries. An increase in surface hardness is not achieved (with the exception of 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 more affine elements. 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: with 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 the boron is predominantly distributed atomically within the grain boundaries 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 because 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, as a companion element, is limited to values ​​of up to 0.001% by weight.

[0050] The features and advantages of the present invention are illustrated below using the following examples with reference to the accompanying drawings and tables.

[0051] Figure 1 is a graphical representation of the residual austenite contents for different steel flat product samples as a function of the cooling temperatures after an annealing treatment during the production of these steel flat product samples,

[0052] Figure 2 is a scanning electron microscope image of a portion of a flat steel product sample according to an embodiment of the invention,

[0053] Figure 3 shows the results of a step test for a steel according to the invention with 5 repeat tests in the form of a graphical representation of the load curve of the load at fracture over the time period until fracture and

[0054] Figure 4 shows the influence of the subcooling temperature on technological properties and retained austenite content as a function of the cooling temperature using the example of a selected melt (melt B3).

[0055] The core of the present invention is the discovery that low cooling temperatures (TC), also called cooling stop temperatures, can have a beneficial effect on the formation of residual austenite. The use of low cooling temperatures makes it possible to avoid higher Si and Al contents and limit their maximum content to 1.20%, which offers advantages in terms of galvanizability, surface quality, machinability during the production of the flat steel product, and further processability of the resulting flat steel product.

[0056] This finding is based on a series of tests with various samples of flat steel products, both from laboratory melts and industrially produced melts. Details of the production and testing of these samples are described below.

[0057] Annealing treatments were performed on cold-rolled strip samples with different chemical analyses (Table 1). For this, the samples were heated to 860 °C at 5 K / s and held for 30 s. They were then cooled at 10 K / s to a cooling temperature TC between 480 °C and 300 °C. The samples were held at this temperature for 300 s and finally cooled to room temperature at 10 K / s. The retained austenite content of the samples was determined using a magnetization yoke. Alternatively, the retained austenite content can also be determined using X-ray diffraction or electron backscatter diffraction (EBSD) on electropolished samples.

[0058] Table 2 summarizes the results of the retained austenite measurements at the various cooling temperatures TC, i.e., the retained austenite contents (RA) for different cooling temperatures after annealing treatment with the following steps: heating to 860 °C, cooling at 10 K / s to the cooling temperature TC (cooling stop temperature); holding for 300 s, and final cooling at 10 K / s. Furthermore, the calculated values ​​("RA(max)-RA(min) difference" of the retained austenite contents in the range between 480 °C and 350 °C and "RA450-RA350 difference" of the retained austenite contents between 450 and 350 °C are shown). Figure 1 shows these results in a diagram in which the retained austenite content RA is plotted in volume % against the cooling temperature TC. The inventive examples are marked with a solid line, while the comparative examples have a dashed / dotted curve.

[0059] As Figure 1 shows, the retained austenite content RA increases—contrary to expectations—with lower cooling temperatures TC. This is surprising, since a trend similar to the counterexamples is actually expected. The maximum retained austenite content should be at temperatures around 450 °C and then decrease with lower temperatures. Thus, highest retained austenite contents can be achieved in the range of 450 °C.

[0060] Characteristic of an inventive course is:

[0061] • At 350 °C, a residual austenite content (RA) of more than 5% must be present

[0062] • RA450 - RA350 < 0%

[0063] • RAmax-RAmin •* 3%

[0064] Melts L3, L6, L7, L9, L10 and L12 to 15 show the one course according to the invention, in which higher contents of residual austenite can be stabilized by low supercooling temperatures.

[0065] Although comparative examples L1 and L2 demonstrate a curve according to the invention, the retained austenite content remains too low at < 5% at 350 °C. Thus, there is insufficient retained austenite available to adequately improve the elongation characteristics. These grades have too low a carbon content to stabilize the retained austenite sufficiently. The carbon content is therefore limited to the range of 0.13 to 0.2. While higher carbon contents increase the possibility of retained austenite formation, the strength also increases with increasing martensite content.

[0066] Silicon and aluminum, individually or in combination, counteract the formation of cementite and thus enable the enrichment of carbon in the austenite. The upper limits for Si and Al were set individually at 0.7 Si and 0.6 Al, respectively, and should not exceed 1.2% in total. Higher Si contents particularly impair galvanizing properties, while increasing Al contents make castability in continuous casting more difficult and also significantly increase the temperatures required for austenitization. Higher contents also no longer exhibit the effect stated in the invention that higher contents of retained austenite can be achieved by lowering the cooling temperature. Melts L5, L8, and L11 show high retained austenite contents (RA) across the entire range (480 °C - 350 °C). m ax-RA m(< 3%). Since the addition of Si and aluminum is required to achieve higher levels of retained austenite, minimum amounts of Si were set at 0.3% and Al at 0.06%. The total amount of Si+Al must be > 0.6%.

[0067] Molybdenum severely limits the diffusion capacity of carbon (L4), thus making low cooling temperatures unusable. The use of molybdenum is ruled out.

[0068] Cr increases hardenability and reduces the diffusivity of carbon, thus shifting the optimal cooling temperature upwards and to longer cooling times. Cr increases stability over a wider cooling temperature range. Cr contents of up to 0.5% can be used optionally.

[0069] Vanadium is used as a dissolved element here and also acts similarly to Cr. Concentrations of up to 0.12% can be used optionally to better adjust correspondingly high residual austenite levels. Niobium was added in all cases; it enables a finer microstructure and thus higher residual austenite levels. The niobium content should be in the range of 0.01 to 0.05%.

[0070] Higher aluminum contents, such as those used in the steel according to the invention, can result in 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 > 0.5 pm in total in an area fraction of at least 1 pm2 / mm2 on a measuring area of ​​at least 100 mm2 in the slab before reheating.Since TiN and TiAIN do not dissolve during subsequent annealing processes, they can still be detected in the hot-dip coated high-strength steel strip. The TiN and TiAIN content can be determined quantitatively using energy-dispersive X-ray spectroscopy (EDX).

[0071] The production of the high-strength strip according to the invention from high-strength multi-phase steel with improved formability due to residual austenite is carried out as follows:

[0072] Production of hot-rolled strip of various thicknesses ranging from 1.5 to 6 mm from slabs with the steel chemistry composition according to the invention. Slabs are heated to a temperature of > 1100 °C, preferably above 1200 °C, to largely dissolve carbide and nitride precipitation. The slabs are formed in several stages within a temperature range of 1250 °C to 860 °C. 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 480 and 750 °C and is typically > 600 °C to achieve a predominantly ferritic-pearlitic hot strip structure.

[0073] 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, where the annealing time includes the time for heating and cooling to below 50 °C. Batch annealing can be performed primarily if the hot strip was coiled at temperatures < 600 °C to reduce resistance during cold rolling.

[0074] Pickling and cold rolling of the hot strip into a cold strip, or sheet, to a thickness of typically 0.5 mm to 3 mm.

[0075] Continuous annealing of the cold-rolled strip at a temperature between 750 and 950 °C, preferably > 800 °C, to achieve a high degree of austenitization. This is followed by single-stage or multi-stage cooling. Typically, multi-stage cooling to a cooling temperature is performed, with cooling rates of 0.5 to 100 K / s being used in the individual zones.

[0076] Cooling temperatures range from 450 °C to 250 °C. Preferably in the range < 400 °C, or more advantageously in the range < 350 °C. Holding times at this temperature are > 60 s.

[0077] In the case of optional hot-dip coating, the steel strip must be reheated and hot-dip coated at a temperature between 400 °C and 500 °C.

[0078] The strip is then cooled to < 60 °C.

[0079] Subsequent skin-passing and straightening processes are possible to improve flatness and surface and to influence the mechanical properties, in particular the Rpo,2 yield strength.

[0080] Reference steels and steels according to the invention: Table 3 contains the properties obtained for different operating modes and melts according to the previously presented production, including the optional hot-dip coating, with detailed information on the final annealing treatment.

[0081] With comparative analysis B2, only low retained austenite contents of < 5% could be achieved. These are not suitable for achieving a correspondingly high level of elongation and tensile strength. This is characterized by the product of Rm*Ag remaining < 9000 MPa% and a retained austenite content of less than 5%.

[0082] With the steel composition according to the invention with melt B1, the desired properties can be adjusted using different operating procedures.

[0083] Cooling temperatures in the range of 450 °C and 340 °C result in residual austenite contents well above 5% and the product of tensile strength (Rm) and uniform elongation (Ag) is above 10000 MPa%.

[0084] Microstructure: Dual-phase steels are generally assumed to have a two-phase structure consisting of ferrite and martensite. However, the other components of the microstructure are also of great importance for the properties. In particular, the presence of retained austenite is crucial for improved formability. Increased formability can be achieved by additionally activating another forming mechanism, the deformation-induced transformation (TRIP effect).

[0085] The microstructure components are determined using longitudinal sections at % strip thickness using a scanning electron microscope. The polished samples are etched with Nital for this purpose. The phase fractions are evaluated using image analysis techniques. A distinction is made between the martensite / austenite phase and ferrite, consisting of polygonal, quasi-polygonal, and bainitic ferrite. The residual austenite contained in the microstructure can be determined using a magnetization yoke. Martensite also includes any hard second phase of bainite, even if it has been tempered by reheating or transformation heat. Martensite therefore includes fresh, self-tempered, and tempered martensite.

[0086] Figure 2 shows the microstructure of a steel melt B1 according to the invention after annealing 4. In the example, 42% is martensite and austenite, of which 8% is austenite according to the magnetic yoke determination. The remaining portion of the microstructure consists of polygonal, quasi-polygonal, and bainitic ferrite, accounting for 58%.

[0087] Hydrogen embrittlement: Material according to the invention from melt B1 was tested for its tendency to hydrogen-induced brittle fracture. The test was conducted using a step test, in which an increasing mechanical load is applied in a corrosive environment. Figure 3 shows the results of the step test for steel from melt B1 in the form of a graphical representation of the load curve L at fracture (in %) over the time period Z (in hours) until fracture, i.e., in the form of a load curve LC.

[0088] The test was conducted in accordance with VDA 238-200 using laser-cut specimens. The specimens were annealed at 170 °C for 20 minutes. The maximum notched tensile force (FmK) was then determined. The step test is performed in an aqueous test solution containing 50 + / - 5 g / l NaCl. The specimen is initially loaded with 50% FmK for 24 hours. The load is then increased by 5% FmK every hour until the specimen fractures. If a specimen reaches 95% FmK for one hour, the load is subsequently increased until fracture occurs.

[0089] All specimens met the 60% load level. This demonstrates that the steel has sufficient resistance to hydrogen-induced brittle fracture. The results are summarized in Table 4 and Figure 3.

[0090] Figure 4 shows the influence of the subcooling temperature on technological properties and retained austenite content for melt B3 (some values ​​are shown in the table) as a function of the cooling temperature TC (here, RCS). By reducing the cooling temperature, higher retained austenite contents can be achieved. In addition, an increase in the yield strength is achieved.

[0091] Table 1: Analyses of laboratory and industrial melts (data in weight %)

[0092] Table 2: Retained austenite contents (RA) for different cooling temperatures after annealing: heating to 860 °C, cooling at 10 K / s to the cooling temperature, holding there for 300 s, and final cooling at 10 K / s. Furthermore, the calculated values ​​(RA(max)-RA(min) difference in retained austenite contents in the range between 480 °C and 350 °C) and RA450-RA350 difference in retained austenite contents between 450 °C and 350 °C)

[0093] 30 November 2023 alzgitter Flachstahl GmbH 122265WO / 53416

[0094] 22

[0095] Table 3: Operation and properties for test melts B1, B2 and B3 during continuous annealing with optional hot-dip coating 1 Heating rate An Annealing temperature An Holding time at temperature R1 Cooling rate to Tm M Intermediate cooling temperature C Cooling temperature R2 Cooling rate between Tm and TK kH Holding time at cooling temperature HD Temperature of zinc bath R3 Cooling rate after zinc bath A Retained austenite content in %

[0096] Table 4: Results of the step test on hydrogen-induced brittle fracture for melt

[0097] AN1 for 5 replicate samples.

Claims

Patent claims 1. Process for producing a high-strength flat steel product with an R m - Tensile strength of at least 700 MPa based on a cold-rolled and continuously annealed steel strip having a multi-phase structure, the process comprising the following steps: Manufacturing a hot-rolled steel strip consisting of the following elements in weight %: C: from > 0.13 to < 0.20, Mn: from > 1.40 to < 2.40, Al: from > 0.06 to < 0.60, preferably > 0.20 to < 0.50 Si: from > 0.03 to < 0.70, Mo: < 0.10, Nb: from > 0.01 to < 0.10, P: < 0.02, S: < 0.05, Ca: < 0.004, B: < 0.001, Cr: < 0.50, Ni: < 0.10, Cu: < 0.10, and optionally one or more elements from: N: > 0.001 to < 0.10, Ti: < 0.05, V: < 0.20, Remainder iron, including usual steel-accompanying elements, where the sum of the proportions in weight% of Si and Al is in a range of > 0.60 to < 1.20 Cold rolling of the hot-rolled steel strip to a cold-rolled steel strip, followed by continuous annealing of the cold-rolled steel strip, in particular as part of a continuous hot-dip coating process, at a maximum temperature between 750 °C and 950 °C inclusive for a total duration of 10 s to 1000 s, followed by cooling of the continuously annealed, cold-rolled steel strip in one or more stages to a cooling stop temperature in a temperature range from 250 °C to 450 °C with an average cooling rate CRi of 0.5 to 100 K / s and subsequent holding of the temperature in the temperature range from 250 °C to 450 °C for 30 to 600 s and optionally subsequent hot-dip coating of the cold-rolled steel strip at a temperature between 380 °C and 500 °C as part of the continuous hot-dip coating process and subsequent final cooling of the continuously annealed, cold-rolled steel strip with an average cooling rate of 2 K / s to 100 K / s to ambient temperature to create the high-strength flat steel product.

2. Method according to claim 1, characterized in that the cooling stop temperature to which the cold-rolled steel strip is cooled is in a temperature range of 250 °C to 400 °C, preferably in a temperature range of 250 °C to 350 °C.

3. A method according to claim 1 or 2, characterized in that the cooling of the cold-rolled steel strip takes place in several stages over at least one intermediate temperature in a temperature range of 620 °C to 760 °C.

4. Method according to one of claims 1 to 3, characterized in that the cold rolling of the hot-rolled steel strip is carried out with a cold rolling degree of less than 80%.

5. Method according to one of claims 1 to 4, characterized in that in order to set a desired roughness of the flat steel product, a skin-passing of the cold-rolled steel strip is carried out and / or in order to set a desired yield point, a stretch-bend straightening of the cold-rolled steel strip is carried out with a respective resulting total degree of deformation of up to a maximum of 1%.

6. Method according to one of claims 1 to 5, characterized in that the hot-rolled steel strip is subjected to an annealing treatment and / or descaling before cold rolling.

7. A method according to any one of claims 1 to 6, characterized in that the hot-rolled steel strip is produced as follows: Heating an ingot or slab having the composition specified in claim 1 for the hot-rolled steel strip to a temperature above of 1100 °C, followed by hot rolling of the ingot or slab above a temperature of 800 °C to the hot-rolled steel strip and Coiling of the hot-rolled steel strip at a temperature between 480 °C and 750 °C.

8. Process according to one of claims 1 to 7, characterized in that in the resulting high-strength flat steel product, the multi-phase structure of the steel strip contains, in addition to bainite, tempered, self-tempered and fresh martensite, a proportion of 20% by volume to 80% by volume of ferrite and between 5% by volume and 15% by volume of residual austenite.

9. High-strength flat steel product, in particular produced by a process according to one of claims 1 to 8, with a steel strip consisting of the following elements in % by weight: C: from > 0.13 to < 0.20, Mn: from > 1.40 to < 2.40, AI: from > 0.06 to < 0.60, preferably > 0.20 to < 0.50 Si: from > 0.03 to < 0.70, Mo: < 0.10, Nb: from > 0.01 to < 0.10, P: < 0.02, S: < 0.05, Ca: < 0.004, B: < 0.001, Cr: < 0.50, Ni: < 0.10, Cu: < 0.10, and optionally one or more elements from: N: > 0.001 to < 0.10, Ti: < 0.10, V: < 0.20, Remainder iron, including usual steel-associated elements, where the sum of the proportions in weight% of Si and Al is in a range from > 0.60 to < 1.20 and with a multi-phase structure, which in addition to bainite, tempered and fresh martensite a proportion of 20 volume% to 80 volume% ferrite and between 5 volume% and 15 Volume% residual austenite, whereby the flat steel product further has the following properties: an R m -Tensile strength of at least 700 MPa and a product of R m -Tensile strength and uniform elongation A g , which is in the range between 9000 and 14000 MPa %.

10. High-strength flat steel product according to claim 9, characterized in that the structure of the steel strip has a ferrite content of 40 to 80% by volume.

11. High-strength flat steel product according to claim 9 or 10, characterized by a zinc-based coating on the steel strip created by a hot-dip coating process.

12. High-strength flat steel product according to one of claims 9 to 11, characterized in that the product of R m -Tensile strength and uniform elongation A g between 10000 and 14000 MPa%.

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