High-strength cold-rolled steel having reduced hydrogen embrittlement sensitivity and method for manufacturing same

Abstract

The high-strength cold-rolled flat steel product with reduced sensitivity to hydrogen embrittlement according to the present invention comprises a steel substrate consisting of, in mass%, C: 0.20% to 0.40%, Mn: 1.50% to 3.00%, Si: 0.90% to 1.50%, Al: 0.005% to 1.00%, V: 0.01% to 0.30%, optionally Cr: 0.01% to 1.00%, optionally Mo: 0.005% to 0.20%, optionally B: 0.00001% to 0.002%, optionally Nb and Ti: a total content of Nb and Ti of 0.005% to 0.2%, P: up to 0.020%, S: up to 0.005%, N: up to 0.008%, and the balance being Fe and unavoidable impurities. The remainder being Fe and unavoidable impurities, the sum of the proportions of the impurities being < = 0.8%, and the microstructure of the steel substrate comprising, in area%, 65% to 92% of primary (tempered) martensite and at least 8% of retained austenite (RA), the remainder is filled with up to 27% secondary (untempered) martensite, up to 10% bainite or bainite ferrite and / or up to < = 5% polygonal ferrite, the sum of the proportions of secondary (untempered) martensite, bainite or bainite ferrite and polygonal ferrite < = 27%. The invention also provides a method for reliably manufacturing the flat steel product.

Description

[0001] The present invention relates to a high-strength cold-rolled flat steel product with reduced hydrogen embrittlement sensitivity, and a method for manufacturing the flat steel product.

[0002] In this context, “flat steel products” refer to rolled products whose length and width are significantly greater than their thickness. Therefore, flat steel products include, in particular, steel strips, steel plates and blanks produced therefrom.

[0003] Unless explicitly stated otherwise, all information about the content of steel alloy components in this application is related to mass. Therefore, all % data about the composition of steel alloys or other alloys mentioned in this article, if no reference unit is mentioned, should be understood as "mass percentage" ("mass %").

[0004] In order to achieve the emission reduction targets of the automotive and transport industries through lightweighting while at the same time improving passenger safety, flat steel products that combine higher strength with improved ductility are required. These two conflicting goals of strength and ductility were brought together in the advanced third generation high strength steel concepts (the so-called "AHSS"), which include steels known as "quenched and partitioned" steels ("Q&P steels").

[0005] Q&P steels are also known to use retained austenite ("RA") as a microstructural component to improve strain hardening and tensile strength of the steel, while also improving elongation through the transformation induced plasticity ("TRIP") effect. The retained austenite is embedded in a matrix of quenched and tempered martensite (primary martensite). Small amounts of bainite (bainitic ferrite), polygonal ferrite, and newly formed martensite (secondary martensite) may also be present in the microstructure of Q&P steels.

[0006] As tensile strength increases, the risk of sensitivity to hydrogen embrittlement, which can lead to an unexpected decrease in ductility and strength, also increases.

[0007] It is known that the susceptibility of steel to hydrogen embrittlement can be minimized by adding microalloying elements such as Ti, Nb or V. Microalloying elements form fine carbide or carbonitride precipitates which are coherent or semi-coherent and have a size of about <10 nm, preferably <5 nm.

[0008] In the specialist literature, precipitates formed by microalloying elements are often referred to as "traps for diffusible hydrogen". This is because hydrogen atoms have a relatively strong binding energy for these fine precipitates. As a result, hydrogen atoms that penetrate into the steel in which these fine precipitates are present in the microstructure bind to the interfaces of the precipitates or to dislocations resulting from the mismatch between the precipitates and the surrounding matrix.

[0009] A large number of such traps in the matrix slows down the diffusion of hydrogen atoms in the microstructure, thus reducing their negative influence on deformation and fracture processes.

[0010] However, it is not easy to combine microalloying elements with Q&P steels. Since the microalloying elements form carbides with the carbon present in the steel, they consume part of the carbon needed to stabilize the retained austenite during the partitioning process.

[0011] Against this background, the present invention will specify a high-strength, easily formable flat steel product having a reduced tendency to hydrogen embrittlement.

[0012] This problem is solved by a flat steel product having at least the features of claim 1 .

[0013] The invention further specifies a method which allows such a flat steel product to be produced reliably.

[0014] The method as claimed in claim 12 solves this problem.

[0015] Advantageous embodiments of the invention are given in the dependent claims and, like the general concept of the invention, are explained in detail hereinafter.

[0016] Thus, according to the present invention, a high-strength cold-rolled flat steel product with reduced hydrogen embrittlement sensitivity is composed, in mass %, of:

[0017] C: 0.20% to 0.40%,

[0018] Mn: 1.50% to 3.00%,

[0019] Si: 0.90% to 1.50%,

[0020] Al: 0.005% to 1.00%,

[0021] V: 0.01% to 0.30%,

[0022] Optional Cr: 0.01% to 1%,

[0023] Optional Mo: 0.005% to 0.2%,

[0024] Optional B: 0.00001% to 0.002%,

[0025] Optional Nb and Ti: the total content of Nb and Ti is 0.005% to 0.2%,

[0026] P: up to 0.020%,

[0027] S: up to 0.005%,

[0028] N: up to 0.008%,

[0029] And, the balance is Fe and unavoidable impurities, the total proportion of said impurities is ≤ 0.8%, and

[0030] The flat steel product exhibits a microstructure consisting, in area %, of:

[0031] 65% to 92% primary (tempered) martensite, and

[0032] At least 8% retained austenite (RA),

[0033] The rest is filled in by:

[0034] Up to 27% secondary (untempered) martensite,

[0035] Up to 10% bainite or bainitic ferrite,

[0036] and / or up to ≤ 5% polygonal ferrite,

[0037] The sum of the proportions of the secondary (untempered) martensite, the bainite or bainitic ferrite and the polygonal ferrite is ≤27%.

[0038] The flat steel product according to the invention has a tensile strength of at least 1300 MPa, typically in the range of 1300 MPa to 1600 MPa, a yield strength of at least 1000 MPa, and a total elongation A80 of at least 10%, the tensile strength, the yield strength and the elongation being determined in accordance with currently valid DIN EN ISO 6892 (sample form 2).

[0039] The flat steel product according to the invention not only has a high strength, but also has good formability, as evidenced by a hole expansion ratio HER of more than 20%, which is determined in accordance with ISO 16630, which is currently valid.

[0040] Furthermore, extensive tests have demonstrated that the flat steel product according to the invention has an increased resistance to hydrogen embrittlement, as shown by the quantitative results of slow strain rate tests (with and without a hydrogen-charged medium) carried out in accordance with the currently valid DIN EN ISO 7539-7.

[0041] The invention thus has significant advantages in applications such as structural and crash-related components in the automotive and transportation industries, including battery housings for electric vehicles. These applications require both high yield strength and a large capacity to absorb plastic deformation. High yield strength ensures minimal intrusion in the event of a crash, while high elongation values ​​allow for high energy absorption.

[0042] The method proposed by the invention for producing the flat steel product according to the invention comprises the following working steps:

[0043] a) providing a steel melt, which consists, in mass%, of C: 0.2% to 0.4%, Mn: 1.5% to 3.0%, Si: 0.9% to 1.5%, Al: 0.005% to 1.0%, V: 0.01% to 0.3%, optionally Cr: 0.01% to 1%, optionally Mo: 0.005% to 0.2%, optionally B: 0.00001% to 0.002%, optionally Nb and Ti: the total content of Nb and Ti is 0.005% to 0.2%, P: up to 0.020%, S: up to 0.005%, N: up to 0.008%, and the balance is Fe and unavoidable impurities, the sum of the proportions of said impurities being ≤0.8%;

[0044] b) casting the steel melt into a slab;

[0045] c) fully heating the slab to a reheating temperature of 1000° C. to 1300° C.;

[0046] d) hot rolling the reheated slab into a hot-rolled strip, wherein the hot rolling is completed at a hot rolling finishing temperature of 850° C. to 980° C.;

[0047] e) cooling the hot rolled strip to a coiling temperature of 400° C. to 600° C., wherein the cooling is completed within a maximum of 25 seconds after the hot rolling is completed, and then coiling the hot rolled strip into a coil;

[0048] f) optionally subjecting the hot rolled strip to a pickling treatment;

[0049] g) cold rolling the hot rolled strip into a cold rolled strip, with a cold rolling reduction ratio of 20% to 80%.

[0050] h) final annealing the cold rolled strip by:

[0051] - heating the cold-rolled strip at a heating rate θS of 2°C / s to 10°C / s to a soaking temperature TS which is at least 50°C higher than the Ac3 temperature of the corresponding steel and is at most 950°C,

[0052] - then keeping the cold-rolled strip at a soaking temperature TS for a soaking time tS of more than 40 seconds and less than 200 seconds;

[0053] - subsequently quenching the cold-rolled strip at a quenching rate θQ of 20°C / s to 100°C / s to a quenching stop temperature TQ, which is lower than the martensite start temperature T_MS of the steel and at least equal to a temperature TQ_min, wherein at said temperature TQ_min, in the microstructure of the cold-rolled strip, 65% to 92% of primary martensite is present in terms of area %; and

[0054] - keeping the annealed cold-rolled strip at the quenching stop temperature TQ for 4 to 20 seconds;

[0055] i) subjecting the final annealed cold-rolled strip to an overaging treatment (also called “partitioning treatment”), which comprises:

[0056] - heating the cold-rolled strip to an overaging temperature TP of 380° C. to 460° C.,

[0057] - maintaining the cold-rolled strip at the overaging temperature for 50 to 200 seconds, and

[0058] - Cooling the cold rolled strip to below 100°C at a cooling rate θC of 0.5°C / s to 20°C / s.

[0059] In order to protect the flat steel product according to the invention from corrosion, an anticorrosion coating can be provided on at least one of its surfaces in a conventional manner. Such a coating can be applied by electrolytic plating, hot-dip galvanizing or galvannealing. The coating usually consists of an alloy whose main constituent is zinc ("Zn") or aluminum ("Al"), and further alloying elements such as silicon ("Si"), magnesium ("Mg") and iron ("Fe") can also be added in a well-known manner to optimize the properties of the coating.

[0060] The steel composition of the steel substrate of the flat steel product according to the invention is determined as follows:

[0061] In the steel of the flat steel product according to the invention, carbon ("C") accounts for 0.20 mass % to 0.40 mass %. C is an essential element for increasing the strength of the steel sheet and for reliably obtaining the required amount of stable retained austenite in the microstructure. During the quenching process carried out in the final annealing (working step h) of the method according to the invention) and the subsequent overaging step (working step i) of the method according to the invention), the retained austenite is stabilized by diffusion and distribution of C until room temperature is reached. For the stability of the retained austenite phase, the C content is at least 0.20 mass %. In addition, the required amount and strength of the martensite formed during the first quenching (working step h) of the method according to the invention) or the final quenching (working step i) of the method according to the invention) to temperatures below 100° C., in particular room temperature, is determined by the C content. Steel sheets with a C content of less than 0.2% will not show a combination of sufficiently high strength and good formability. However, if the C content exceeds 0.4 mass %, the martensite start temperature is significantly reduced, which results in only a small amount of martensite and thus a reduced strength level. Furthermore, C contents exceeding 0.4% reduce the weldability of the steel sheet according to the invention. The positive effect of C on the properties of the flat steel product according to the invention can be achieved particularly reliably with a C content of at least 0.22% by mass. The presence of C in the flat steel product according to the invention is particularly effective with a C content of at most 0.3% by mass.

[0062] According to the flat steel product of the present invention, 1.50 mass % to 3.00 mass % of manganese ("Mn") is added to the steel. The Mn element can effectively increase the hardness of the steel and thus increase the strength of the steel. In addition, the presence of the Mn content within the range specified by the present invention can inhibit the formation of ferrite and pearlite during quenching. By using a quenching rate of <100°C / s, a suitable microstructure containing martensite and residual austenite is obtained after the first quenching (working step i)). Mn is also a solid solution strengthening element, which stabilizes austenite by reducing the Ms temperature. In order to reliably obtain these positive effects of the presence of Mn, the Mn content in the steel of the flat steel product of the present invention needs to reach at least 1.5 mass %, in particular at least 1.9 mass %. However, a high Mn content of more than 3.00 mass % will deteriorate the welding properties and cause segregation, thereby also deteriorating the mechanical properties of the flat steel product. In order to reliably avoid these negative effects of Mn, the Mn content in the steel of the flat steel product of the present invention can be limited to 2.8 mass %.

[0063] In the steel for the flat steel product according to the invention, silicon ("Si") accounts for 0.90% by mass to 1.50% by mass. Si contributes to the strength of the steel by solid solution strengthening. In addition, silicon is insoluble in cementite and therefore inhibits the formation of iron carbides during the distribution process (working step i)) and supports the stabilization of the residual austenite, thereby increasing the ductility. In order to achieve these effects, at least 0.90% by mass of Si is required in the steel for the flat steel product according to the invention. In addition, Al in combination with 0.9% Si also shows a similar effect on the stabilization and amount of residual austenite, thereby increasing the ductility. However, high silicon contents of more than 1.50% by mass have an adverse effect on the coatability of the steel sheet.

[0064] The steel of the flat steel according to the invention contains 0.005% to 1.00% by mass of aluminum ("Al"). Al is usually used as a deoxidizer in steelmaking and forms aluminum nitride, which contributes to the strength of the steel. Like Si, Al is insoluble in cementite and therefore inhibits the formation of iron carbides during the decomposition process (working step i)). However, with increasing aluminum content, the Ac3 temperature of the steel for producing the flat steel product according to the invention rises sharply to values ​​that are too high for conventional industrial annealing lines. If Al is used only as a deoxidizer, the Al content can be limited to a maximum of 0.1% by mass to avoid the formation of AlN. If Al is to be used for deoxidation, the Al content should be at least 0.005% by mass. Typically, in the steel alloy according to the invention, the Al content for this purpose ranges from 0.005 to 0.100% by mass, in particular from 0.005 to 0.070% by mass or from 0.005 to 0.060% by mass. However, if Al is used to promote the retention of austenite, a content of up to 1.0% is permitted. In this case, an Al content of at least 0.060 mass %, particularly at least 0.1 mass % or at least 0.3 mass % is suitable.

[0065] In the steel for producing the flat steel product according to the invention, the vanadium ("V") content ranges from 0.01% by mass to 0.3% by mass. V is a microalloying element, the dissolution temperature of which is much lower than that of Ti or Nb. This allows V-based precipitates to partially or completely dissolve during the final annealing (working step h)) and to precipitate during the subsequent cooling process, or most preferably to precipitate by carbon diffusion and distribution during the distribution phase of the final annealing cycle (working step i)). In this way, particularly fine V-based precipitates are formed, which are particularly effective for capturing diffusible hydrogen. The positive effect of V becomes apparent when the concentration of V exceeds 0.01% by mass and continues to increase to at most 0.15% by mass. Saturation of this effect is observed when the concentration exceeds 0.3% by mass. The presence of V has proven to be most effective in the range of up to 0.25% by mass, in particular in the range of up to 0.20% by mass. The effect is particularly pronounced if the V content is at least 0.07% by mass.

[0066] Chromium ("Cr") can optionally be added to the steel alloy of the steel substrate for producing the flat steel product according to the invention in order to effectively prevent the formation of pearlite and bainite and to increase the strength. This effect can be achieved by adding at least 0.01% by mass of Cr. However, in order to avoid the occurrence of grain boundary oxidation, the Cr content is limited to 1% by mass. It is particularly effective when Cr is present in the steel of the flat steel product according to the invention in a Cr content of at least 0.1% by mass. In order to avoid the negative effects of the presence of Cr, the Cr content can be limited to 0.5% by mass.

[0067] As a further optional element, 0.005% to 0.2% by mass of molybdenum ("Mo") can be added to the steel alloy for producing the flat steel product according to the invention. Mo increases the strength of the steel sheet and inhibits the formation of pearlite. Mo contents of at least 0.02% by mass are particularly effective in this respect, with the positive effect of Mo being particularly pronounced at Mo contents of up to 0.15% by mass.

[0068] Niobium ("Nb") and titanium ("Ti") are microalloying elements which can be added to the steel alloy of the flat steel product according to the invention, optionally in combination or individually, and which effectively increase the strength of the steel sheet by precipitation hardening and refining the microstructure. However, the dissolution temperatures of Nb and Ti are much higher than the dissolution temperature of V. This means that, for example, mixed carbides are formed if Ti is present in combination with V. Such mixed carbides tend not to dissolve during the final annealing process, so that during the final annealing process they remain coarse or grow to an ineffective size. Therefore, the sum of the contents of Nb and Ti is limited to a maximum of 0.2% by mass. On the other hand, the advantageous effects of the presence of Nb and / or Ti on the strength and formability of the flat steel product can be reliably utilized when the total content of Nb and / or Ti is at least 0.005% by mass. The strengthening of the steel is particularly effective when the total content of Nb and / or Ti is in the range of 0.02% by mass to 0.15% by mass.

[0069] 0.00001 to 0.002 mass % boron ("B") may optionally be added to the steel of the flat steel product according to the invention to inhibit the formation of ferrite and segregation along grain boundaries, thereby preventing the movement of ferrite. This results in a fine-grained microstructure that contributes to improved mechanical properties of the steel. In this regard, a B content in the range of 0.0001 to 0.001 mass % is particularly effective.

[0070] The term "impurities" includes all elements that enter the steel during the steelmaking process or cannot be completely removed from the steel. In the steel of the flat steel product according to the invention, impurities may be detectable by measurement, but are not listed here as essential or optional components, or, in the case of optionally added elements, are present in very low amounts below the effective limits specified here. In order to avoid a detrimental effect of the impurities as a whole, the sum of the impurity contents is limited to a maximum of 0.8% by mass. In particular, phosphorus ("P"), sulfur ("S") and nitrogen ("N") are unavoidable impurities. However, it has been shown that a P content of up to 0.020% by mass, a S content of up to 0.005% by mass and a N content of up to 0.008% by mass do not reduce the properties of the steel of the flat steel product according to the invention. Typically, 0.001% to 0.020% by mass of P, 0.0001% to 0.005% by mass of S and 0.0001% to 0.008% by mass of N are present in the steel. The copper ("Cu") content is up to 0.5 mass%, the nickel ("Ni") content is up to 0.5 mass%, and the oxygen ("O") content is up to 0.0080 mass%, which are also unavoidable impurities. Of course, other elements such as W, Co, Sn, Ca, Mg, REM, Zr, Te, As, Bi, etc. are also considered impurities.

[0071] The flat steel product according to the invention exhibits a microstructure comprising 65 area % to 92 area % of primary (tempered) martensite and at least 8 area % of retained austenite (RA). The retained austenite component fills the portion not occupied by primary martensite and optionally other microstructural components permitted according to the invention. Thus, if no other components are present in the microstructure of the steel, the retained austenite accounts for 8 area % to 35 area % of the microstructure. However, if only minimal amounts of primary (tempered) martensite (65 area %) and retained austenite (8 area %) are present in the microstructure, the total amount of components that may be present optionally according to the invention is 27 area %. In theory, the corresponding remainder of the microstructure may be filled solely by secondary (untempered) martensite, solely by bainite or bainitic ferrite and / or solely by polygonal ferrite, wherein these optional components usually occur in combination.

[0072] The microstructure of the flat steel product according to the present invention also includes V-based precipitates with a diameter of less than 10 nm, and the density of the precipitates is ≥ 1000 / μm. 2 The flat steel product according to the invention therefore exhibits an increased resistance to hydrogen embrittlement.

[0073] The precipitate density of the V-based precipitates was determined by transmission electron microscopy images in combination with X-ray microanalysis (TEM and EDX) using carbon extraction replicas. The carbon extraction replicas were obtained from longitudinal sections. The magnification of the measurements was between 10,000x and 200,000x. Based on these images, the diameter of the precipitates within the measurement area can be calculated by computer-assisted image analysis. For this purpose, in each case 5 measurement areas were measured. The results of the 5 measurement areas were then averaged. The size of the measurement area depends on the selected magnification and ranges from 18.5μm x 14.5μm at 10,000x magnification to 0.925μm x 0.725μm at 200,000x magnification. The nature of the precipitates was determined simultaneously by EDX (energy dispersive X-ray spectroscopy). In this process, the electron beam emitted by the transmission electron microscope excites the atoms in the precipitates. The elemental distribution in the sample can be determined by the emitted X-rays, which enables the identification of the vanadium precipitates.

[0074] For example, the density of V precipitates can be determined by the following steps:

[0075] 1. Provide carbon extraction replicas of longitudinal slices of flat steel products;

[0076] 2. Using TEM and computer-assisted image analysis at 100,000x magnification, the diameters of all precipitates were determined over five different measurement areas measuring 1.85 μm x 1.45 μm;

[0077] 3. Identify vanadium precipitates based on detected X-ray quanta;

[0078] 4. Count the number of vanadium precipitates with a diameter less than 10 nm in each of the five measurement regions, and determine the density of the vanadium precipitates in each of the five measurement regions;

[0079] 5. The average value of the density of the five measurement areas is determined as the vanadium precipitate density of the flat steel product.

[0080] If the austenite grains are too large, the nucleation sites for martensite are reduced and the martensite packets may be too large, which in turn reduces the local formability of the final product. The width of a single martensite lath is a function of its length. Thinner laths are advantageous for supporting the microstructural processes that occur during the overaging treatment (working step i) of the method according to the invention). According to the invention, the precipitation of fine V-based precipitates along the phase boundaries of martensite and retained austenite grains results in a thin lath thickness of up to 1000 nm, wherein lath lengths of up to 500 nm are often obtained. The microstructure can be measured using a transverse slice located at 1 / 3t of the layer, i.e., on a slice taken at one third of the thickness of the steel substrate sheet. These slices are prepared for scanning electron microscopy (SEM) and treated with 3% Nital etchant. Due to the fineness of the microstructure, the microstructure is observed by scanning electron microscopy at a magnification of 5000x. The measured lath thickness length corresponds to the average value of five measurements.

[0081] In addition, the precipitates located at the phase boundaries can also serve as hydrogen traps, thereby improving hydrogen embrittlement resistance. Controlling the microstructure by precipitating fine V-based precipitates along the phase boundaries of martensite and / or retained austenite grains can lead to excellent hydrogen embrittlement resistance and mechanical properties. Due to the alloy specified by the invention and the method for manufacturing a flat steel product according to the invention, fine V(C,N) precipitates are formed, thereby making the original austenite grains finer and the martensite / retained austenite lath structure thinner. V(C,N) nanoparticles are located along the lath interfaces, where they capture hydrogen that penetrates the steel substrate, thereby improving hydrogen embrittlement resistance.

[0082] The properties of the flat steel product according to the invention explained above can be reliably achieved by the method proposed according to the invention for producing such a flat steel product.

[0083] To produce the flat steel product according to the invention, a steel melt having a composition corresponding to the composition specified in the invention is prepared by conventional methods. The steel melt is then cast by conventional methods to form at least one slab (steps a) and b) of the method according to the invention).

[0084] In preparation for subsequent hot rolling, the slab is reheated to a temperature in the range of at most 1000° C. to 1300° C. or, if the temperature of the slab after casting is high enough, is kept in this temperature range so that the temperature of the entire volume of the slab is equalized (working step c) of the method according to the invention).

[0085] In working step d) of the method according to the invention, the slab is hot rolled into a hot strip at a sufficiently high temperature so that the final rolling temperature is in the range of 850° C. to 980° C. If the temperature during the hot rolling is lower, sufficient static recrystallization will not occur between the rolling steps. As a result, the resulting microstructure will retain a solid texture and a high dislocation density due to dynamic recrystallization. Therefore, the final rolling temperature must not be lower than 850° C. Final rolling temperatures above 980° C. are technically not feasible.

[0086] According to step e) of the method of the invention, the hot rolled strip obtained should be cooled to the coiling temperature within a maximum of 25 seconds after the end of hot rolling to avoid precipitation or formation of polygonal ferrite on the run-out table passed by the hot rolled strip before coiling. Cooling is preferably completed within a maximum of 18 seconds, even more preferably within a maximum of 15 seconds, to avoid such effects.

[0087] In working step e), cooling to the coiling temperature can be carried out in any manner known in the prior art, and the cooling rate is generally 20° C. / s to 1000° C. / s. A high cooling rate can be achieved, for example, by water quenching.

[0088] The coiling temperature should be at most 600°C to avoid the formation of pearlite. In addition, at higher coiling temperatures, oxidizing elements (such as Si, Cr or Mn) diffuse into the grain boundaries and form oxides, which will lead to a decrease in the surface quality of the hot-rolled material, thereby limiting the surface quality of the flat steel product after hot rolling and optional coating. In addition, limiting the coiling temperature to at most 600°C also prevents the formation of unwanted polygonal ferrite. Coiling temperatures below or at most 580°C increase the content of bainite in the microstructure of the hot-rolled material. By setting the coiling temperature at a lower 400°C to 500°C, the positive effect of effectively avoiding grain boundary oxidation can be achieved. However, low coiling temperatures also promote the formation of a large amount of martensite, resulting in a high hardness of the hot-rolled strip, so batch annealing is required during cold rolling; annealing is usually carried out in a temperature range of 550°C to 600°C for more than 6 hours and less than 24 hours to decompose the bainite / martensite. Therefore, by coiling the hot rolled strip at a coiling temperature of 400°C to 600°C, preferably 400°C to 580°C, and particularly usefully 500°C to 580°C, a uniform microstructure can be reliably obtained without producing a large amount of martensite, thereby enabling smaller thickness and width tolerances to be achieved in the subsequent cold rolling process.

[0089] After the hot-rolled strip has been coiled and cooled to room temperature, it can be descaled in a conventional manner, if necessary, for example by pickling, to remove dirt from the strip surface or to improve the surface quality (optional working step f) of the method according to the invention).

[0090] Cold rolling is usually carried out in one or more rolling steps on a conventional tandem rolling mill, and if necessary, intermediate batch annealing as described above can also be carried out. For cold rolling, reversing rolling stands can be used to achieve higher reductions. In this case, it is advantageous to achieve a cold rolling reduction of up to 80%. The minimum reduction obtained during cold rolling should generally be 20% to ensure sufficient recrystallization in the final annealing step (step h) of the method according to the invention). However, a higher reduction helps to obtain a fine-grained microstructure and thus obtain the described product properties. Therefore, in the method according to the invention, when cold rolling the hot-rolled strip into a cold-rolled strip (working step g) of the method according to the invention), a cold rolling reduction of 20% to 80% is appropriate.

[0091] The heat treatment steps h) and i) of the method according to the invention are preferably carried out in a heat treatment line, wherein the corresponding flat steel products are passed through the line one after the other continuously and uninterruptedly.

[0092] During the final annealing of the cold-rolled strip according to working step h) of the method according to the invention, most of the properties of the flat steel product according to the invention are adjusted. When heating the cold-rolled strip to the soaking temperature TS, the average heating rate "θS" should be 2°C / s to 10°C / s. Heating rates above 10°C / s endanger the proper recrystallization before austenitization, while heating rates below 2°C / s are not economical in a continuous annealing line. The soaking temperature TS should be at least 50°C higher than the Ac3 temperature of the steel of the corresponding composition to ensure proper homogenization of C in the fully austenitic microstructure.

[0093] For a given composition, the Ac3 temperature can be determined experimentally in a known manner using dilatometry or estimated according to the following formula (1):

[0094] Ac3[℃]=910℃+(-203√(%C)–15.2%Ni+

[0095] 44.7%Si+31.5%Mo–21.1%Mn))℃ / mass%

[0096] Where %C = the corresponding C content in the steel alloy (in mass %),

[0097] %Ni = corresponding Ni content in steel alloy (in mass %),

[0098] %Si = corresponding Si content in steel alloy (in mass %),

[0099] %Mo = corresponding Mo content in steel alloy (in mass %),

[0100] %Mn = the corresponding Mn content in the steel alloy (in mass %).

[0101] For the steel alloys covered by the alloy parameters of the present invention, the temperature range of Ac3 is generally about 750 °C to about 920 °C. The Ac3 temperature of the steel alloys covered by the alloy parameters of the present invention is experimentally determined by dilatometry according to SEP 1681-1998-06. When measuring, a dilatometer is used to compare the average thermal expansion coefficient of the steel sample with the thermal expansion coefficient of the quartz tube. The quartz tube expands linearly with temperature over a wide temperature range (far exceeding 1000 °C), while the steel sample undergoes phase changes during heating and cooling. Record the change of the linear expansion rate during uniform heating or cooling, so as to show the phase change temperature. For this purpose, commercially available dilatometers can be used, such as Dilatometer 805.

[0102] The upper limit of the soaking temperature TS is 950 °C in order to fully dissolve any V-carbides in the semi-finished product. The soaking time tS should be long enough to achieve chemical homogenization and carbide dissolution. At the same time, however, the soaking time must be limited to prevent excessive growth of austenite grains.

[0103] In order to promote the precipitation of V(C,N) precipitates, the soaking time tS should be at least 40 seconds, preferably more than 40 seconds, and at most 200 seconds (40 s ≤ tS ≤ 200 s), preferably less than 200 seconds (40 s < tS < 200 s). A soaking time tS exceeding 60 seconds and less than 120 seconds (60 s < tS < 120 s) is particularly useful in the present invention.

[0104] After the soaking stage, the steel strip enters the primary cooling step. The average cooling rate "ΘQ" of the flat steel product according to the present invention in the primary cooling step must be high enough to minimize the formation of polygonal ferrite, bainite, bainite ferrite and any precipitated carbides. Therefore, the lower limit of the cooling rate ΘQ is set to 20 °C / s. The upper limit of ΘQ depends on the process stability and the cooling capacity of the primary cooling step. The present invention does not require the cooling capacity of the primary cooling step to be increased above 100 °C / s. On the contrary, a cooling rate exceeding 100 °C / s during the primary cooling process will increase the production cost and lead to insufficient cooling, which is not conducive to the mechanical properties of the final product. Therefore, the cooling rate ΘQ is set to 20 °C / s to 100 °C / s, preferably 30 °C / s to 70 °C / s.

[0105] The quenching stop temperature TQ at the end of the primary cooling step for the cooling carried out in this way must be lower than the martensite start temperature T_MS.

[0106] The martensite start temperature T_MS can be experimentally determined by dilatometry in a known manner or estimated according to the following formula (2):

[0107] T_MS[℃])=539℃+(-423%C–30.4%Mn–7.5%Si+30%Al)℃ / mass%

[0108] Where %C = the corresponding C content in the steel alloy (in mass %),

[0109] %Mn = corresponding Mn content in steel alloy (in mass %),

[0110] %Si = corresponding Si content in steel alloy (in mass %),

[0111] %Al = the corresponding Al content in the steel alloy (in mass %).

[0112] For the steel alloys covered by the alloy parameters of the present invention, the temperature range of T_MS is generally about 265° C. to about 435° C. The T_MS temperature of the steel alloys covered by the alloy parameters of the present invention is experimentally determined by dilatometric measurement according to SEP 1681-1998-06. This determination method is well known to the skilled person and is carried out using a dilatometer, as described above when determining the Ac3 temperature.

[0113] The quenching stop temperature must not fall below a temperature TQ_min at which 65 to 92 area % of primary martensite is present in the microstructure of the cold-rolled strip (T_MS≥TQ≥TQ_min).

[0114] In order to determine the proportion of primary martensite after the initial quenching and thus obtain the temperature TQ_min, the Koistinnen-Marburger equation can be used. This equation is generally expressed as follows:

[0115] f m =1–exp(-0.011(T_MS–TQ))

[0116] Among them, f m It is the part of austenite that transforms into martensite when quenched to a temperature TQ below T_MS (Speer et al., MS&T 2003, pp. 505-522, 2003). According to the present invention, the amount of primary martensite should not exceed 92 area % and not be less than 65 area %.

[0117] After the primary cooling, the flat steel product according to the invention is then held at the corresponding quenching temperature TQ for at least 4 seconds and no more than 20 seconds. This holding is necessary to homogenize the temperature over the entire sheet thickness and to allow the final stage of the isothermal transformation to take place.

[0118] After the holding step, the flat steel product enters the overaging step i) of the method according to the invention, also referred to in the technical terminology as the "separation step". In step i), C is partitioned from the primary martensite to the adjacent retained austenite. However, the separation of C is a thermally activated process that competes with the precipitation of nanoparticles. This means that if the partitioning is carried out at too high a temperature or for too long a time, iron carbides may be formed, which consume a large amount of carbon required to stabilize the retained austenite and reduce both the tensile strength and elongation of the flat steel product. In order to avoid this effect, the invention sets an upper limit of the overaging temperature TP to 460°C. The lower limit of the overaging temperature TP is set to 380°C to ensure the precipitation of a sufficient amount of V(C,N) nanoparticles, which does not occur at lower temperatures. In this regard, limiting the overaging temperature TP to a maximum of 460°C has also proven to be particularly advantageous, since it also prevents the V(C,N) nanoparticles from coarsening, which would reduce the mechanical properties of the flat steel product. By setting the overaging temperature to 380° C. to 460° C. (380° C. TP 460° C.) and limiting the time period tP for which the corresponding flat steel product is kept at the temperature TP to 50 seconds to 200 seconds, the precipitate density of V-based precipitates with a diameter of less than 10 nm in the microstructure of the product according to the invention can reliably reach ≥ 1000 precipitates / μm. 2 .

[0119] After the corresponding time period tP has expired, the annealed steel strip is cooled to below 100°C at a cooling rate θC of 0.5°C / s to 20°C / s. This allows the dispensing to be stopped immediately and any carbide coarsening or other heat-induced effects that would reduce the mechanical properties of the flat steel product according to the invention to be avoided. After the flat steel product has been cooled to below 100°C, the cooling to room temperature is not critical and can be carried out in any suitable manner.

[0120] Tests carried out in accordance with DIN EN ISO 7539-7 have shown that the flat steel products according to the invention have an increased resistance to hydrogen embrittlement. For this purpose, samples of the flat steel products according to the invention were exposed to natural air and other samples of the same flat steel products were exposed to a water-based solution containing 5% NH4SCN as a hydrogenation medium. Subsequently, the maximum tensile stress Rm_max(air) of the samples exposed to the air environment and the maximum tensile stress Rm_max(NH4SCN) of the samples exposed to the hydrogenation medium were determined in the usual standard manner. 4 SCN). From the tensile stress thus determined, the "hydrogen embrittlement sensitivity factor" ("H_sf") is calculated as follows:

[0121] H_sf=1-Rm_max(NH 4 SCN) / Rm_max(air)

[0122] The results show that the hydrogen embrittlement sensitivity factor H_sf of the flat steel products according to the invention is reliably at most 0.35 (H_sf≤0.35). The tests therefore show that the hydrogen embrittlement sensitivity of the flat steel products according to the invention is significantly lower than that of the comparative steel plate samples, which were subjected to the same test conditions and evaluated in the same way, but which consisted of conventional steel plates, the steel of which was not alloyed in the manner according to the invention, i.e., in particular lacked the V content specified in the invention. In contrast to the samples according to the invention, the comparative samples generally showed a hydrogen embrittlement sensitivity factor H_sf of at least 0.35.

[0123] In each of Tables 1 to 5 shown below, the values ​​that do not meet the specifications of the present invention are underlined.

[0124] In order to carry out actual tests and prove the effects obtained by the present invention, we melted a steel melt AM, the composition of which is shown in Table 1. In Table 1, numerical values ​​and examples that are not within the scope of the present invention are underlined.

[0125] Table 2 shows for each melt AM the Ac3 temperature determined by dilatometry in accordance with SEP 1681-1998-06 and the T_MS temperature determined by dilatometry in accordance with SEP 1681-1998-06.

[0126] The melt AM is cast into slabs which are reheated to a temperature of 1150° C. to 1300° C. for subsequent hot rolling.

[0127] The slab reheated in this way is hot rolled into a hot strip on a conventional hot rolling line, and the hot rolling finishing temperature FT ranges from 850° C. to 980° C. Table 2 also shows the hot rolling finishing temperature FT maintained during the hot rolling of steels A to M.

[0128] After leaving the last hot rolling stand of the hot rolling line, the hot strip obtained in each case is cooled within a maximum time period tC of 25 seconds to the respective coiling temperature CT, which ranges from 400° C. to 600° C. Table 2 also lists the respective time period tC and the respective coiling temperature CT.

[0129] After the hot-rolled strip is coiled and cooled to room temperature, it is pickled in a manner known to those skilled in the art to remove dirt from its surface.

[0130] The hot rolled strip is then rolled into a cold rolled strip, and the rolling method is also the same as the ordinary cold rolling method. The total cold rolling degree CRD achieved during the cold rolling process ranges from 20% to 80%, and the calculation formula for each cold rolling degree CRD is as follows:

[0131] CRD=(D1–D2) / D1*100%

[0132] Where D1: thickness of each hot rolled strip before cold rolling

[0133] D2: Thickness of each cold-rolled strip after cold rolling

[0134] Table 2 also lists the respective cold rolling degrees CRD.

[0135] The cold rolled samples of the obtained cold rolled strips are heated to the soaking temperature TS at a heating rate θS for final annealing, and then each sample is kept at the soaking temperature TS for a soaking time tS. After soaking, the samples are immediately quenched at a quenching rate θQ to a quenching stop temperature TQ, so that in the microstructure of the sample annealed in this way, the proportion (PM%) of primary martensite is between 65 area % and 92 area %. After quenching, the samples of the annealed cold rolled strips are then kept at the quenching stop temperature TQ for a period of time tQ of 4s to 20s. Table 3 lists the respective soaking temperatures TS, heating rates θS, soaking times tS, quenching rates θQ, quenching stop temperatures TQ, the proportion PM% of primary martensite (corresponding to tM in Table 5) and the time period tQ.

[0136] After the final annealing, the samples of the cold-rolled strip were subjected to an aging treatment, which is also known in technical terms as a "partitioning treatment".

[0137] To this end, the sample is first heated to a temperature TP and then held at this temperature for a period of time tP. Finally, the sample is cooled at a cooling rate ΘC to a temperature below 100° C. Table 3 also lists the temperature TP and the period of time tP.

[0138] For each sample treated in this way, Table 4 lists the yield strength YS, tensile strength TS and elongation A80 determined according to DIN EN ISO 6892 (sample form 2), the YS / TS ratio, the hole expansion HER determined according to ISO 16630 and the product of TS x A80 for each sample.

[0139] In addition, the microstructure of each sample was analyzed according to ISO 9042 for optical microscopic analysis. Table 5 lists the retained austenite content RA, the primary (tempered) martensite content tM, the secondary (fresh) martensite content fM, the total bainite content, the bainitic ferrite content bF, and the polygonal ferrite content pF determined in this way.

[0140] In addition, the surface density of V(C,N) precipitates in the sample microstructure was determined by preparing carbon replicas and conventional transmission electron microscopy.

[0141] Finally, the hydrogen embrittlement sensitivity factor H_sf was determined according to the method described above. The results show that the hydrogen embrittlement sensitivity of the samples according to the present invention, which meet all the requirements of the present invention in terms of composition, production and heat treatment, is significantly lower than that of the samples that are not alloyed and / or manufactured and heat treated according to the present invention.

[0142]

[0143]

[0144]

[0145]

[0146]

Claims

1. A high-strength cold-rolled flat steel product with reduced hydrogen embrittlement sensitivity, the flat steel product comprising a steel substrate, the steel substrate consisting of the following in mass %: C： 0.20% to 0.40%, Mn: 1.50% to 3.00%, Si: 0.90% to 1.50%, Al: 0.005% to 1.00%, V: 0.01% to 0.30%, Optional Cr: 0.01% to 1.00%, Optional Mo: 0.005% to 0.20%, Optional B: 0.00001% to 0.002%, Optional Nb and Ti: the total content of Nb and Ti is 0.005% to 0.2%, P: up to 0.020%, S: up to 0.005%, N: up to 0.008%, The balance is Fe and unavoidable impurities, the total proportion of which is ≤ 0.8%, and The steel substrate exhibits a microstructure which, measured in area % according to ISO 9042, comprises: 65% to 92% primary (tempered) martensite, and At least 8% retained austenite (RA), The rest is filled in by: Up to 27% secondary (untempered) martensite, Up to 10% bainite or bainitic ferrite, and / or at most ≤5% polygonal ferrite, the sum of the proportions of the secondary (untempered) martensite, the bainite or bainitic ferrite and the polygonal ferrite being ≤27%.

2. The flat steel product according to claim 1, characterized in that The density of V-based precipitates with a diameter of less than 10 nm in the microstructure of the steel substrate of the flat steel product is ≥ 1000 / μm, measured as described in the specification. 2 .

3. The flat steel product according to claim 1 or 2, characterized in that According to the measurement described in the specification, the primary martensite contained in the microstructure of the steel substrate of the flat steel product has a fine lath structure, and the maximum length of the lath is 1000 nm.

4. The flat steel product according to claim 3, characterized in that The maximum length of the laths was 500 nm, measured as described in the description.

5. Flat steel product according to any one of the preceding claims, characterized in that The C content in the steel of the steel substrate of the flat steel product is 0.22% to 0.3% in terms of mass %.

6. The flat steel product according to any one of the preceding claims, It is characterized in that The Mn content in the steel of the steel substrate of the flat steel product is 1.9% to 2.8% in terms of mass %.

7. Flat steel product according to any one of the preceding claims, characterized in that The V content in the steel of the steel substrate of the flat steel product is 0.07% to 0.20% in terms of mass %.

8. The flat steel product according to any one of the preceding claims, It is characterized in that The flat steel product has a yield strength of at least 1000 MPa, a tensile strength of at least 1300 MPa and an elongation A80 of at least 10%, determined in accordance with DIN EN ISO 6892 (sample form 2).

9. Flat steel product according to any one of the preceding claims, characterized in that The flat steel product has a hole expansion ratio, measured according to ISO 16630, of at least 20%.

10. The flat steel product according to any one of the preceding claims, It is characterized in that An anti-corrosion coating is provided on at least one surface of the flat steel product.

11. The flat steel product according to claim 10, characterized in that The coating is applied by electrolytic plating, hot dip galvanizing or galvannealing.

12. A method for producing a flat steel product according to any one of the preceding claims, comprising the following working steps: a) providing a steel melt, which consists, in mass %, of: C： 0.2% to 0.4%, Mn: 1.5% to 3.0%, Si: 0.9% to 1.5%, Al: 0.005% to 1.0%, V: 0.01% to 0.3%, optional Cr: 0.01% to 1.00%, optional Mo: 0.005% to 0.20%, optional B: 0.00001% to 0.002%, optional Nb and Ti: the total content of Nb and Ti is 0.005% to 0.2%, P: up to 0.020%, S: up to 0.005%, N: up to 0.008%, and the balance is Fe and unavoidable impurities, the sum of the proportions of the impurities is ≤0.8%; b) casting the steel melt into a slab; c) fully heating the slab to a reheating temperature of 1000° C. to 1300° C.; d) hot rolling the reheated slab into a hot-rolled strip, wherein the hot rolling is completed at a hot rolling finishing temperature of 850° C. to 980° C.; e) cooling the hot rolled strip to a coiling temperature of 400° C. to 600° C., wherein the cooling is completed within a maximum of 25 seconds after the hot rolling is completed, and then coiling the hot rolled strip into a coil; f) optionally subjecting the hot rolled strip to a pickling treatment; g) cold rolling the hot rolled strip into a cold rolled strip, with a cold rolling reduction ratio of 20% to 80%. h) final annealing the cold rolled strip by: - heating the cold-rolled strip at a heating rate θS of 2°C / s to 10°C / s to a soaking temperature TS which is at least 50°C higher than the Ac3 temperature of the corresponding steel and is at most 950°C, wherein the Ac3 temperature of the corresponding steel is determined using dilatometry according to SEP 1681-1998-06, - then keeping the cold-rolled strip at the soaking temperature TS for a soaking time tS of more than 40 seconds and less than 200 seconds; - subsequently quenching the cold-rolled strip at a quenching rate θQ of 20°C / s to 100°C / s to a quenching stop temperature TQ which is below the martensite start temperature T_MS of the steel and at least equal to a temperature TQ_min, wherein at said temperature TQ_min, 65% to 92% of primary martensite is present in the microstructure of the cold-rolled strip in terms of area %, wherein the T_MS temperature of the steel is determined using dilatometry according to SEP 1681-1998-06; and - keeping the annealed cold-rolled strip at the quenching stop temperature TQ for 4 to 20 seconds; i) subjecting the final annealed cold-rolled strip to an overaging treatment, the overaging treatment comprising the following steps: - heating the cold-rolled strip to an overaging temperature TP of 380° C. to 460° C., - maintaining the cold-rolled strip at the overaging temperature for 50 to 200 seconds, and - Cooling the cold rolled strip to below 100°C at a cooling rate θC of 0.5°C / s to 20°C / s.

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