High-strength cold-rolled steel sheet product with reduced susceptibility to hydrogen embrittlement and method for producing the same

Abstract

The present invention consists of, by 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: total content of 0.005 to 0.2%, P: at most 0.020%, S: at most 0.005%, N: at most 0.008%, and Fe and inevitable impurities as the balance, the total content of the impurities being 0.8% or less, and contains, by area%, 65 to 92% of primary (annealed) martensite and at least 8% of retained austenite (RA), the balance consisting of at most 27% of secondary (non-annealed) martensite, at most 10% of bainite or bainitic ferrite, and / or at most 5% of polygonal ferrite, and the total of the secondary (non-annealed) martensite, the bainite or the bainitic ferrite, and the polygonal ferrite being 27% or less, and provides a high-strength cold-rolled steel sheet product having a steel substrate and having reduced susceptibility to hydrogen embrittlement. The present invention also provides a method capable of reliably manufacturing such a steel sheet product. The abstract needs to be published without drawings.

Description

Technical Field

[0001] The present invention relates to a high-strength cold-rolled steel sheet product with reduced susceptibility to hydrogen embrittlement and a method for manufacturing such a steel sheet product.

Background Art

[0002] It should be understood that the "steel sheet product" in this specification means a rolled product in which the length and width are each significantly larger than the thickness. Therefore, steel sheet products particularly include steel strips, steel sheets, and blanks obtained therefrom.

[0003] All information regarding the content of the steel alloy composition shown in this application is related to mass unless explicitly stated otherwise. Therefore, all % data referring to the composition of the steel alloy or another alloy mentioned in this specification should be understood as information of "% by mass" ("mass%") without referring to the reference unit.

[0004] In the automotive and transportation industries, in order to achieve emissions reduction through weight reduction and at the same time enhance passenger safety, steel sheet products with improved strength and ductility are required. The conflicting goals of strength and ductility are integrated in the advanced third-generation high-strength steel concept (so-called "AHSS"). This includes steels called "quenching and partitioning" steels ("Q&P steels").

[0005] Q&P steels use retained austenite ("RA") as a component of the microstructure, improve the strain hardening and tensile strength of the steel, and increase elongation by the known transformation-induced plasticity ("TRIP") effect. Retained austenite is embedded in the matrix structure of quenched and tempered martensite (primary martensite). The microstructure of Q&P steels may also contain a small amount of bainite (bainitic ferrite), polygonal ferrite, and fresh martensite (secondary martensite).

[0006] As the tensile strength increases, the susceptibility to hydrogen embrittlement also increases, which can cause unexpected decreases in ductility and strength.

[0007] When adding trace alloying elements such as Ti, Nb, or V, the susceptibility of steel to hydrogen embrittlement can be minimized. The trace alloying elements form fine carbide or carbonitride precipitates that are coherent or semi-coherent on the order of less than 10 nm, preferably less than 5 nm.

[0008] The precipitates formed by the trace alloying elements are often referred to in the specialized literature as "traps for diffusible hydrogen". This is because hydrogen atoms have a relatively strong binding energy to such fine precipitates. Therefore, hydrogen atoms that penetrate into the steel having a structure in which these fine precipitates are present bind to the interfaces of the precipitates or to dislocations generated from the mismatch between the precipitates and the surrounding matrix structure.

[0009] When a large number of such traps exist in the matrix structure, the diffusion of hydrogen atoms through the structure is slowed down, and the adverse effects on the processes occurring during deformation and fracture are reduced.

[0010] However, it is not easy to use trace alloying elements in combination with Q&P steel. Since the trace alloying elements form carbides with carbon present in the steel, they are required during partitioning to stabilize retained austenite, and a part of the carbon is consumed. Summary of the Invention

[0011] Against such a background, the present invention needs to identify a high-strength and easily formable steel sheet product with a reduced tendency of hydrogen embrittlement.

[0012] This object is solved by a steel sheet product having at least the features described in claim 1.

[0013] Furthermore, the present invention should identify a process that enables such a steel sheet product to be reliably manufactured.

[0014] This object is solved by the method described in claim 12.

[0015] Advantageous embodiments of the present invention are described in the dependent claims and will be explained in detail below, similar to the general concept of the present invention.

[0016] According to the present invention, a high-strength cold-rolled steel sheet product with reduced susceptibility to hydrogen embrittlement has 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%, Optionally Mo: 0.005 to 0.2%, Optionally B: 0.00001 to 0.002%, Optionally Nb and Ti: total content of 0.005 to 0.2%, P: at most 0.020%, S: at most 0.005%, N: at most 0.008%, and consists of Fe and inevitable impurities as the balance, with the total content of the impurities being 0.8% or less, in area %, 65 to 92% of primary (annealed) martensite, and at least 8% of retained austenite (RA) including, and the balance being at most 27% of secondary (unannealed) martensite, at most 10% of bainite or bainitic ferrite, and / or at most 5% of polygonal ferrite and shows a structure in which the total of the secondary (unannealed) martensite, the bainite or the bainitic ferrite, and the polygonal ferrite is 27% or less.

[0017] The steel plate product according to the present invention has a tensile strength of at least 1300 MPa, usually in the range of 1300 - 1600 MPa, a yield strength of at least 1000 MPa, and a total elongation A80 of at least 10%. The tensile strength, yield strength, and elongation are each determined in accordance with the currently valid DIN EN ISO 6892 (specimen form 2).

[0018] The good formability of the steel plate product according to the present invention is also reflected in a hole expansion ratio HER of more than 20%, despite its high strength. The hole expansion ratio HER is determined in accordance with the currently valid ISO 16630.

[0019] Furthermore, through extensive testing, it has been proven that the steel plate product according to the present invention has improved resistance to hydrogen embrittlement, as quantified in a low strain rate test carried out in accordance with the currently valid DIN EN ISO 7539-7, regardless of the presence of a hydrogen charging medium.

[0020] Therefore, the present invention has significant advantages for structural and crash-related components in the automotive and transportation industries, including applications such as battery housings for electric vehicles. In these applications, both a high yield strength and a large ability to absorb plastic deformation are required. The high yield strength minimizes intrusion during a collision, and the high elongation value enables large energy absorption.

[0021] The manufacturing method of the steel plate product according to the present invention proposed by the present invention is (a) providing a molten steel consisting of, by mass%, C: 0.2 - 0.4%, Mn: 1.5 - 3.0%, Si: 0.9 - 1.5%, Al: 0.005 - 1.0%, V: 0.01 - 0.3%, optionally Cr: 0.01 - 1%, optionally Mo: 0.005 - 0.2%, optionally B: 0.00001 - 0.002%, optionally Nb and Ti: total content of 0.005 - 0.2%, P: at most 0.020%, S: at most 0.005%, N: at most 0.008%, and Fe and unavoidable impurities as the balance, with the total content of the impurities being 0.8% or less; (b) casting the molten steel into a slab; (c) heating the slab to a reheating temperature of 1000 to 1300 °C; (d) hot rolling the reheated slab to obtain a hot-rolled steel strip, wherein the hot rolling ends at a hot rolling finishing temperature of 850 to 980 °C; (e) cooling the hot-rolled steel strip to a coiling temperature of 400 to 600 °C, and finishing the cooling within at most 25 seconds after the end of the hot rolling, and coiling the hot-rolled steel strip in a coil shape; (f) optionally, pickling the hot-rolled steel strip; (g) cold rolling the hot-rolled steel strip with a cold rolling reduction rate of 20 to 80% to obtain a cold-rolled steel strip; (h) heating the cold-rolled steel strip to a soaking temperature TS which is at least 50 °C higher than the Ac3 temperature and at most 950 °C at a heating rate ΘS of 2 to 10 °C / s; immediately holding the cold-rolled steel strip at the soaking temperature TS for a soaking time tS exceeding 40 seconds and less than 200 seconds; immediately quenching the cold-rolled steel strip at a quenching rate ΘQ of 20 to 100 °C / s to a quenching stop temperature TQ which is lower than the martensite start temperature T_MS and at least equal to a temperature TQ_min at which 65 to 92 area% of primary martensite exists in the structure of the cold-rolled steel strip; holding the cold-rolled steel strip at the quenching stop temperature TQ for 4 to 20 s; thereby, finally annealing the cold-rolled steel strip; (i) heating the cold-rolled steel strip to an overaging temperature TP of 380 to 460 °C; holding the cold-rolled steel strip at the overaging temperature for 50 to 200 s; cooling the cold-rolled steel strip to less than 100 °C at a cooling rate ΘC of 0.5 to 20 °C / s; aging the finally annealed cold-rolled steel strip by an overaging treatment (also referred to as "segmented treatment") including the above steps.

[0022] In order to protect the steel plate product according to the present invention from corrosion, corrosion-resistant plating can be provided on at least one of its surfaces in a general manner. Such plating can be applied by electroplating, hot-dip galvanizing, or alloyed hot-dip galvanizing. Usually, the plating consists of an alloy whose main component is zinc ("Zn") or aluminum ("Al"), and additional alloying elements such as silicon ("Si"), magnesium ("Mg"), and iron ("Fe") can be added in a known manner to optimize the properties of the plating.

Embodiments for Carrying Out the Invention

[0023] The composition of the steel substrate of the steel plate product according to the present invention was determined as follows.

[0024] In the steel of the steel plate product according to the present invention, 0.20 to 0.40% by mass of carbon ("C") is present. C is an essential element for increasing the strength of the steel plate and ensuring a necessary amount of stable retained austenite in the structure. Retained austenite is stabilized to room temperature by carbon diffusion and partitioning during quenching performed during the final annealing (step h) of the method of the present invention and during the subsequent overaging treatment step (step i) of the method of the present invention. For the stability of the retained austenite phase, a C content of at least 0.20% by mass is required. Furthermore, the required amount and strength of martensite formed during the first quench (step h) of the method of the present invention or during the last quench (step i) of the method of the present invention to a temperature below 100°C, particularly room temperature, are determined by the carbon content. Steel plates with a carbon content of less than 0.2% do not exhibit sufficiently high strength in combination with good formability. However, when the carbon content exceeds 0.4% by mass, the martensite start temperature drops significantly, and as a result, only a small amount of martensite is formed, thus reducing the strength level. Furthermore, when the carbon content exceeds 0.4%, the welding properties of the steel plate of the present invention deteriorate. The favorable effect of C on the properties of the steel plate product according to the present invention can be particularly reliably obtained when the C content is at least 0.22% by mass. The presence of C in the steel plate product according to the present invention is particularly effective when the C content is at most 0.3% by mass.

[0025] In the steel of the steel plate product according to the present invention, 1.50 to 3.00% by mass of manganese ("Mn") is added. Mn is an element that effectively increases the hardness and thus the strength of the steel. Further, due to the presence of the Mn content within the range defined in the present invention, the formation of ferrite and pearlite during quenching is suppressed. By using a quenching rate of less than 100 °C / s, an appropriate structure containing martensite and retained austenite can be obtained after the first quenching (step i). Mn is also a solid solution strengthening element that stabilizes austenite by lowering the Ms temperature. In order to reliably obtain these favorable effects due to the presence of Mn, the steel of the steel plate product according to the present invention requires at least 1.5% by mass, particularly at least 1.9% by mass of Mn. However, when the Mn content exceeds 3.00% by mass, the welding properties deteriorate, and segregation occurs that also reduces the mechanical properties of the steel plate product. In order to reliably avoid these adverse effects due to the presence of Mn, the Mn content in the steel of the steel plate product according to the present invention can be limited to 2.8% by mass.

[0026] In the steel of the steel plate product according to the present invention, 0.90 to 1.50% by mass of silicon ("Si") is present. Si contributes to the strength of the steel by solid solution strengthening. Further, since silicon is insoluble in cementite, it suppresses the formation of iron carbide during partitioning (step i), promotes the stabilization of retained austenite, and brings about an improvement in ductility. In order to obtain these effects, at least 0.90% by mass of Si is required in the steel of the steel plate product of the present invention. Further, the combination of Al and 0.9% Si shows a similar effect on the stability and amount of retained austenite, and the ductility is improved. However, a high silicon content exceeding 1.50% by mass has an adverse effect on the plating property of the steel plate.

[0027] The steel for the flat steel of the present invention contains 0.005 to 1.00% by mass of aluminum ("Al"). Al is generally used in steelmaking as a deoxidizer and to form aluminum nitride that contributes to the strength of the steel. Similar to Si, since Al does not dissolve in cementite, it suppresses the formation of iron carbide during decomposition (step i). However, as the aluminum content increases, the Ac3 temperature of the steel for manufacturing the steel plate product of the present invention will rise dramatically to a value that is too high in a normal industrial annealing line. When using Al only as a deoxidizer, the Al content can be limited to a maximum of 0.1% by mass to avoid the formation of AlN. To use Al for deoxidation, an Al content of at least 0.005% by mass is required. Usually, in the steel alloy of the present invention, the Al content for these purposes is in the range of 0.005 to 0.100% by mass, particularly in the range of 0.005 to 0.070% by mass or 0.005 to 0.060% by mass. However, when Al is required to promote the retention of austenite, a concentration of up to 1.0% is allowed. In this case, an Al content of at least 0.060% by mass, particularly at least 0.1% by mass or at least 0.3% by mass is appropriate.

[0028] The vanadium ("V") content of the steel constituting the steel plate product according to the present invention is in the range of 0.01 to 0.3% by mass. V is a microalloying element with a melting temperature much lower than that of Ti or Nb. As a result, V-based precipitates partially or completely dissolve during final annealing (step h) and then precipitate during subsequent cooling, or most preferably, precipitate by carbon diffusion and partitioning during the partitioning stage (step i) of the final annealing cycle. This enables the formation of particularly fine V-based precipitates that are particularly effective in capturing diffusible hydrogen. The favorable effect of V becomes prominent when the concentration exceeds 0.01% by mass and continues to increase steadily up to 0.15% by mass. This effect saturates when the concentration exceeds 0.3% by mass. It can be seen that the presence of V is most effective in the range of up to 0.25% by mass, particularly up to 0.20% by mass. This is particularly true when the V content is at least 0.07% by mass.

[0029] Chromium ("Cr") can be added to the alloy of the steel substrate of the steel plate product according to the present invention as needed to effectively delay the formation of pearlite and bainite and 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. The presence of Cr in the steel of the steel plate product according to the present invention is particularly effective at a Cr content of at least 0.1% by mass. In order to avoid the adverse effects caused by the presence of Cr, the Cr content can be limited to 0.5% by mass.

[0030] As a further optional element, molybdenum ("Mo") can be added to the alloy of the steel constituting the steel plate product of the present invention in an amount of 0.005 to 0.2% by mass. Mo improves the strength of the steel plate and suppresses the formation of pearlite. It is particularly effective in this regard when the Mo content is at least 0.02% by mass, and the positive effect of Mo is particularly prominent when the Mo content is at most 0.15% by mass.

[0031] Niobium ("Nb") and titanium ("Ti") can be added to the steel alloy of the steel plate product according to the present invention either in combination or alone as needed, and are microalloying elements that effectively contribute to the strength of the steel plate by precipitation hardening and refinement of the structure. However, the dissolution temperatures of Nb and Ti are much higher than that of V. That is, for example, when Ti is present in combination with V, the formation of mixed carbides can occur. Since this type of mixed carbide tends not to dissolve during final annealing, it remains coarse or grows to an ineffective size during final annealing. Therefore, the total content 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 steel plate product can be surely obtained when the total content of Nb and / or Ti is at least 0.005% by mass. It is particularly effective for the strengthening of the steel when the total range of the Nb and / or Ti content is 0.02 to 0.15% by mass.

[0032] The steel for the steel plate product according to the present invention may be added with 0.00001 to 0.002% by mass of boron ("B") as required, which can suppress the formation of ferrite and segregation along the grain boundaries and prevent their movement. As a result, a fine grain structure that contributes to the mechanical properties of the steel is formed. Particularly effective in this regard is a B content of 0.0001 to 0.001% by mass.

[0033] The term "impurities" includes all elements that enter the steel during steelmaking or cannot be completely removed from the steel. In the steel for the steel plate product according to the present invention, the impurities can be detected approximately by measurement, but are present in a very low content in the case of elements that are not described here as essential or optional components or are added as required, and thus are below the specified effective limits in this specification. To avoid the harmful effects of the total impurities, the total impurity content is limited to a maximum of 0.8% by mass. In particular, phosphorus ("P"), sulfur ("S"), and nitrogen ("N") are inevitable impurities. However, it has been found that the properties of the steel for the steel plate product according to the present invention are not degraded if the P content is at most 0.020% by mass, the S content is at most 0.005% by mass, and the N content is at most 0.008% by mass, respectively. Usually, in the steel, there are 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. A maximum Cu content of 0.5% by mass, a maximum Ni content of 0.5% by mass, and a maximum O content of 0.0080% by mass are also included in the inevitable impurities. Needless to say, other elements such as W, Co, Sn, Ca, Mg, REM, Zr, Te, As, and Bi are also regarded as impurities.

[0034] The steel plate product according to the present invention exhibits a structure comprising 65 to 92 area% of primary (annealed) martensite and at least 8 area% of retained austenite (RA). The retained austenite content fills the portion of the structure not occupied by the primary martensite and other organizational components that are permitted as necessary according to the present invention. Thus, when no other components are present in the steel structure, the retained austenite occupies 8 to 35 area% of the structure. However, when only a minimum amount of 65 area% of primary (annealed) martensite and 8 area% of retained austenite by area ratio are present in the structure, a total of 27 area% of components may be present, which is permitted as necessary according to the present invention. Theoretically, the remaining proportion of each of the structures may be filled by secondary (unannealed) martensite alone, bainite or bainitic ferrite alone, and / or polygonal ferrite alone, and usually, any combination of these components occurs.

[0035] The structure of the steel plate product according to the present invention also contains V-based precipitates with a precipitation density of 1000 or more per 1 μm having a diameter of less than 10 nm. Thereby, the steel plate product according to the present invention has improved resistance to hydrogen embrittlement. 2

[0036] The precipitation density of V-based precipitates is determined by a combination of transmission electron microscope images using carbon extraction replicas and X-ray microanalysis (TEM and EDX). The carbon extraction replicas are obtained from longitudinal sections. The magnification of the measurement is from 10,000 times to 200,000 times. Based on these images, the diameter of the precipitates within the measurement field of view can be calculated by computer-aided image analysis. In each case, five measurement fields of view are measured for this purpose. Next, the results of the five measurement fields of view are averaged. The size of the measurement field of view varies depending on the magnification selected, ranging from 18.5 μm × 14.5 μm at a magnification of 10,000 times to 0.925 μm × 0.725 μm at a magnification of 200,000 times. The nature of the precipitates is simultaneously determined by EDX (energy-dispersive X-ray spectroscopy). In this process, the atoms within the precipitates are excited by the electron beam from the transmission electron microscope. The elemental distribution within the sample is determined from the emitted X-rays, and as a result, the identification of vanadium precipitates can be achieved.

[0037] For example, the density of V precipitates can be determined using the following procedure. 1. Prepare a carbon extraction replica of the longitudinal section of the steel plate product. 2. Using TEM and computer-aided image analysis, determine the diameter of all precipitates on five different measurement fields of view with a size of 1.85 μm × 1.45 μm at a magnification of 100,000 times. 3. Identify vanadium precipitates based on the detected X-ray quanta. 4. Calculate the number of vanadium precipitates with a diameter of less than 10 nm in each of the five measurement fields of view, and determine the density of vanadium precipitates in each of the five measurement fields of view. 5. Determine the density of vanadium precipitates in the steel plate product as the average value of the densities over the five measurement fields of view.

[0038] If the austenite grain size is too large, the nucleation sites of martensite decrease, the dimensions of the martensite packets can increase, and as a result, the local formability of the final product decreases. The width of the individual martensite laths is a function of their length. Thinner laths are advantageous for supporting the tissue processes that occur during the overaging treatment (step i) of the method of the present invention. The precipitation of the fine V-based precipitates of the present invention along the phase boundaries of the martensite and retained austenite grains results in a fine lath thickness having a length of up to 1000 nm, and usually a length of up to 500 nm is obtained. The structure can be determined using a cross-section of the 1 / 3t layer, i.e., a cross-section taken at 1 / 3 of the thickness of the steel substrate. The cross-section is prepared for a scanning electron microscope (SEM) and treated with 3% nital etching. Since the structure is fine, the structure is examined by SEM observation at a magnification of 5000 times. The length of the determined lath thickness corresponds to the average value of 5 measurements.

[0039] Furthermore, the precipitates located at the phase boundaries function as traps for hydrogen, improving the resistance to hydrogen embrittlement. By controlling the structure by precipitating fine V-based precipitates along the phase boundaries of the martensite and / or retained austenite grains, the hydrogen embrittlement resistance and mechanical properties become excellent. The fine V(C,N) precipitates formed as a result of the alloy defined by the present invention and the manufacturing method of the steel sheet product according to the present invention make the prior austenite grains finer, and as a result, the martensite / retained austenite lath structure also becomes thinner. The V(C,N) nanoparticles are arranged along the lath interfaces and capture the hydrogen that penetrates into the steel substrate, improving the resistance to hydrogen embrittlement.

[0040] The characteristics of the steel sheet product according to the present invention described above can be surely obtained by the manufacturing method of such a steel sheet product proposed by the present invention.

[0041] To manufacture the steel sheet product according to the present invention, molten steel having a composition corresponding to the specifications of the present invention is prepared by a conventional method. Next, this melt is cast by a conventional method to form at least one slab (steps (a) and (b) of the process according to the present invention).

[0042] As a preparation for subsequent hot rolling, the slab is reheated to a temperature in the range of 1000 to 1300 °C, or, if the temperature of the slab after casting is sufficiently high, it is held in this temperature range in order to equalize the temperature over the entire volume of the slab of the method of the present invention (step c).

[0043]

[0042] In step (d) of the method according to the invention, the hot rolling of the slab into a hot-rolled steel strip is carried out at a temperature high enough so that the finishing rolling temperature is in the range of 850 to 980 °C. If the temperature during hot rolling is low, sufficient static recrystallization does not occur between the rolling steps. As a result, the resulting microstructure maintains a strong microstructure and a high dislocation density by dynamic recrystallization. For this reason, the finishing rolling temperature must not be less than 850 °C. A finishing rolling temperature exceeding 980 °C is technically unfeasible.

[0044] According to step (e) of the method according to the invention, the cooling of the obtained hot-rolled steel strip to the coiling temperature is completed within at most 25 seconds after the end of hot rolling, and it is necessary to avoid precipitation or the formation of polygonal ferrite on the run-out table through which the hot-rolled steel strip passes before being coiled. To prevent this effect, preferably, the cooling is completed within at most 18 seconds, and more preferably within at most 15 seconds.

[0045] In step (e), the cooling to the coiling temperature can be carried out by any method known in the prior art, and the cooling rate is usually in the range of 20 °C / s to 1000 °C / s. A high cooling rate can be achieved, for example, by water cooling.

[0046] To avoid the formation of pearlite, the coiling temperature needs to be at most 600°C. Furthermore, when the coiling temperature is high, oxidation elements such as Si, Cr, or Mn diffuse to the grain boundaries to form oxides, resulting in a deterioration of the surface quality of the hot-rolled material, and the surface quality after hot rolling and after plating as required for the steel plate product is limited. Furthermore, by limiting the coiling temperature to at most 600°C, the generation of unnecessary polygonal ferrite is also prevented. When the coiling temperature is 580°C or lower, the amount of bainite in the structure of the hot-rolled material increases. By setting the coiling temperature low to 400°C to 500°C, a favorable effect regarding the avoidance of grain boundary oxidation can be obtained. However, a low coiling temperature promotes the formation of a large amount of martensite and may increase the hardness of the hot-rolled steel strip. Therefore, in cold rolling, in order to decompose bainite / martensite, batch annealing is usually required, which is carried out at a temperature of 550°C to 600°C for more than 6 hours and less than 24 hours. A uniform structure containing no large amount of martensite, which enables narrowing of the thickness and width tolerances during subsequent cold rolling, is surely obtained by coiling the hot-rolled steel strip at a coiling temperature of 400°C to 600°C, preferably 400°C to 580°C, particularly usefully 500°C to 580°C.

[0047] The hot-rolled steel strip is coiled into a coil shape and cooled to room temperature. Then, if necessary, the scale of the hot-rolled steel strip is removed by a normal method such as pickling, etc., and the scale present on the surface of the steel strip can be removed or the surface quality of the method of the present invention can be improved (step f if necessary).

[0048] Cold rolling is usually carried out in one or more rolling steps on a conventional tandem rolling mill, and if necessary, intermediate batch annealing is carried out as described above. In the case of cold rolling, a reversing rolling stand can be used to achieve a higher reduction. Here, for the sake of convenience, it is necessary to achieve a cold reduction rate of up to 80%. The minimum reduction rate usually obtained during cold rolling needs to be 20% in order to ensure sufficient recrystallization during the final annealing step (step h) of the method of the present invention. However, a higher reduction is useful for obtaining a fine-grained structure that brings about the described product characteristics. Therefore, in the method of the present invention, when cold rolling the hot-rolled steel strip of the method of the present invention into a cold-rolled steel strip (step g), a cold reduction of 20% to 80% is appropriate.

[0049] The heat treatment steps (h) and (i) of the method according to the present invention are preferably carried out in a heat treatment line through which each steel plate product passes continuously without interruption.

[0050] During the final annealing of the cold-rolled steel strip carried out as step (h) of the method according to the present invention, most of the characteristics of the steel plate product of the present invention are adjusted. The average heating rate “ΘS” when heating the cold-rolled steel strip to the soaking temperature TS is carried out at 2 to 10 °C / s. At a heating rate exceeding 10 °C / s, proper recrystallization before austenitization may be hindered, and a heating rate less than 2 °C / s is not economical in a continuous annealing line. In order to ensure proper homogenization of C within a complete austenite structure, the soaking temperature TS needs to be at least 50 °C higher than the Ac3 temperature of each steel composition.

[0051] The Ac3 temperature of a given composition can be determined experimentally in a known manner using dilatometry or can be estimated according to the following formula (1). Ac3 [°C] = 910 °C + (-203√(%C) - 15.2%Ni + 44.7%Si + 31.5%Mo - 21.1%Mn)) °C / mass% %C = the respective C content of the steel alloy in mass% %Ni = the respective Ni content of the steel alloy in mass% The Si content of each steel alloy in %Si by mass The Mo content of each steel alloy in %Mo by mass The Mn content of each steel alloy in %Mn by mass

[0052] For steel alloys within the scope of the alloy specifications of the present invention, the Ac3 temperature is usually in the range of about 750 °C to about 920 °C. The Ac3 temperature of the steel alloys within the scope of the alloy specifications of the present invention is determined experimentally by dilatometry in accordance with SEP 1681 - 1998 - 06. A dilatometer is used for the determination, and the average coefficient of thermal expansion of the steel sample is compared with the thermal expansion of the quartz tube. The quartz tube expands proportionally to the temperature over a very wide temperature range (far exceeding 1000 °C), while the steel sample transforms both during heating and cooling. The change in linear expansion when the temperature rises or falls equally is recorded, and the transformation temperature is indicated. For this purpose, a commercially available dilatometer such as a Baehr805 dilatometer can be used, for example.

[0053] The upper limit of the soaking temperature TS is 950 °C, which enables sufficient dissolution of any V carbides in the semi-finished product. The soaking time tS needs to be long enough to enable chemical homogenization and carbide dissolution. However, at the same time, it must be sufficiently restricted to prevent excessive growth of austenite grain size.

[0054] To promote the precipitation of V(C,N) precipitates, the soaking time tS needs to be at least 40 seconds, preferably more than 40 seconds, up to 200 seconds (40 s ≤ tS ≤ 200 s), preferably less than 200 seconds (40 s < tS < 200 s). Here, a soaking time tS of more than 60 seconds and less than 120 seconds (60 s < tS < 120 s) is particularly effective.

[0055] After the soaking stage, the steel strip enters the primary cooling step. The average cooling rate "ΘQ" for cooling the steel plate product according to the present invention in the primary cooling step must be high enough to minimize the formation of polygonal ferrite, bainite, bainitic ferrite, and precipitated carbides. Therefore, the lower limit of the cooling rate ΘQ is set to 20 °C / s. The upper limit of ΘQ is determined by the process stability and the cooling capacity of the primary cooling step. In the proposed invention, there is no need to increase the cooling capacity of the primary cooling step to a rate exceeding 100 °C / s. Instead, if the cooling rate during primary cooling exceeds 100 °C / s, the manufacturing cost increases, leading to insufficient cooling and adversely affecting 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.

[0056] The quenching stop temperature TQ at which the cooling executed by this method stops at the end of the primary cooling step must be lower than the martensite start temperature T_MS.

[0057] The martensite start temperature T_MS can be experimentally determined by dilatometry in a known method or can be estimated according to the following formula (2). T_MS [°C]) = 539 °C + (-423%C - 30.4%Mn - 7.5%Si + 30%Al) °C / mass% %C = The respective C content of the steel alloy in mass% %Mn = The respective Mn content of the steel alloy in mass% %Si = The respective Si content of the steel alloy in mass% %Al = The respective Al content of the steel alloy in mass%

[0058] For steel alloys within the alloy specification range of the present invention, the T_MS temperature is usually in the range of about 265 °C to about 435 °C. The T_MS temperature of the steel alloy within the alloy specification range of the present invention is experimentally determined by dilatometry according to SEP 1681-1998-06. This measurement is known to those skilled in the art and is performed using a dilatometer as described above in connection with the determination of the Ac3 temperature.

[0059] The quenching stop temperature shall not be lower than the temperature TQ_min at which 65 to 92 area% of primary martensite exists in the structure of the cold-rolled steel strip (T_MS ≥ TQ ≥ TQ_min).

[0060] To determine the ratio of primary martensite after initial quenching and thereby obtain the temperature TQ_min, the Koistinnen-Marburger equation can be used. This equation is often expressed as follows. f m = 1 - exp(-0.011(T_MS - TQ)) Here, f m is the ratio of austenite that transforms into martensite when quenched to a temperature TQ lower than T_MS (Speer et al., MS&T 2003, pp. 505-522, 2003). According to the present invention, the amount of primary martensite shall not exceed 92 area% and shall not be less than 65 area%.

[0061] Immediately after primary cooling, the steel plate products according to the present invention are held at each quenching temperature TQ for at least 4 seconds and 20 seconds or less. This holding is necessary to equalize the temperature throughout the plate thickness and cause the final stage of isothermal transformation.

[0062] Following the holding step, the steel plate product enters the overaging step (i) of the method of the present invention, which is also called the "separation step" in technical jargon. In step (i), C is distributed from primary martensite towards adjacent retained austenite. However, the separation of C is a thermally activated process and competes with the precipitation of nanoparticles. That is, if the partitioning process is carried out at too high a temperature or over too long a time, iron carbide can be formed, consuming much of the carbon necessary for the stabilization of retained austenite, meaning that both the tensile strength and elongation of the steel plate product decrease. To avoid this effect, in the present invention, the upper limit of the overaging temperature TP is set at 460 °C. The lower limit of the overaging temperature TP is set at 380 °C, ensuring that a sufficient amount of V(C,N) nanoparticles precipitate, which does not occur at lower temperatures. In this regard, limiting the overaging temperature TP to a maximum of 460 °C has proven to be particularly appropriate because it can also avoid the coarsening of V(C,N) nanoparticles that would degrade the mechanical properties of the steel plate product. By setting the overaging temperature to 380 °C to 460 °C (380 °C < TP < 460 °C) and limiting the time tP for which each steel plate product is held at the temperature TP to 50 seconds to 200 seconds, in the structure of the product of the present invention, the precipitation density of V-based precipitates with a diameter of less than 10 nm is 2 surely achieved to be 1000 or more per μm

[0063] After the elapse of each time period tP, 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 immediately stops the partitioning and avoids the coarsening of carbides or other heat-induced effects that could degrade the mechanical properties of the steel plate product according to the present invention. After the steel plate product is cooled to below 100 °C, the cooling to room temperature is not important and can be carried out by any suitable method.

[0064] The improvement in the resistance of the steel plate product of the present invention to hydrogen embrittlement was demonstrated by tests carried out in accordance with DIN EN ISO 7539-7. For this purpose, samples of the steel plate product of the present invention were exposed to natural air, and other samples of the same steel plate product were exposed to an aqueous solution containing 5% NH4SCN as a hydrogen charging medium. Subsequently, the maximum tensile stress Rm_max(air) of the sample exposed to the air atmosphere and the maximum tensile stress Rm_max(NH4SCN) of the sample exposed to the hydrogen charging medium were determined by ordinary standard methods. From the tensile stresses thus determined, the "hydrogen embrittlement susceptibility coefficient" ("H_sf") was calculated as follows. H_sf = 1 - Rm_max(NH4SCN) / Rm_max(air)

[0065] In the case of the steel plate product according to the present invention, it has been found that the hydrogen embrittlement susceptibility coefficient H_sf is surely at most 0.35 (H_sf ≦ 0.35). Thus, by the test, the hydrogen embrittlement susceptibility of the steel plate product according to the present invention was evaluated under the same test conditions and in the same manner, but the steel was not alloyed by the method according to the present invention, that is, it was demonstrated that it is significantly lower than the hydrogen embrittlement susceptibility of the plate sample used for comparison, which is made of a conventional steel plate lacking the V content specified particularly in the present invention. In contrast to the sample of the present invention, in the comparative example, a hydrogen embrittlement susceptibility coefficient H_sf of at least 0.35 is always shown.

[0066] In each of Tables 1 to 5 described below, values not conforming to the specifications of the present invention are underlined.

[0067] To demonstrate the actual tests and the effects obtained by the present invention, molten steels A - M having the compositions shown in Table 1 were melted. In Table 1, values and examples outside the scope of the present invention are underlined.

[0068] In Table 2, the Ac3 temperature determined using the dilatometry method in accordance with SEP 1681-1998-06 and the T_MS temperature determined using the dilatometry method in accordance with SEP 1681-1998-06 are specified for each of the molten steels A to M.

[0069] The molten steels A to M were cast into slabs and then reheated to a temperature of 1150 to 1300 °C for subsequent hot rolling.

[0070] The slabs reheated in this way were heated on a common hot rolling line to a hot rolling finish temperature FT of 850 to 980 °C to obtain hot rolled steel strips. The hot rolling finish temperature FT maintained during the hot rolling of each of the steels A to M is also shown in Table 2.

[0071] After exiting the last hot rolling stand of the hot rolling line, the obtained hot rolled steel strips were each cooled to their respective coiling temperatures CT in the range of 400 °C to 600 °C within a maximum continuous time tC of 25 seconds. The respective continuous time tC and the respective coiling temperature CT are also shown in Table 2.

[0072] After coiling and cooling to room temperature, the hot rolled steel strips were pickled by a method known to those skilled in the art to remove scale from the surface.

[0073] Next, the hot rolled steel strips were rolled by a method corresponding to a general cold rolling method to obtain cold rolled steel strips. The total cold rolling reduction CRD achieved in the cold rolling process is in the range of 20 to 80% respectively, and each cold rolling reduction CRD is calculated as follows. CRD = (D1 - D2) / D1 * 100% D1: The thickness of each hot rolled steel strip before cold rolling D2: The thickness of each cold rolled steel strip after cold rolling

[0074] Each cold rolling reduction CRD is also shown in Table 2.

[0075] After cold rolling, a sample of the cold-rolled steel strip was finally annealed by heating it at a heating rate ΘS to a soaking temperature TS, and then immediately holding each sample at the soaking temperature TS for a soaking time tS. Immediately after soaking, the sample was quenched at a quenching rate ΘQ to a quenching stop temperature TQ. As a result, in the structure of the sample annealed in this way, the proportion %PM of primary martensite was 65 to 92 area %. After quenching, the annealed sample of the cold-rolled steel strip was held at the quenching stop temperature TQ for 4 to 20 seconds over a period tQ. Each soaking temperature TS, heating rate ΘS, soaking time tS, quenching rate ΘQ, quenching stop temperature TQ, proportion %PM of primary martensite (corresponding to tM in Table 5), and period tQ are shown in Table 3.

[0076] After final annealing, the sample of the cold-rolled steel strip was subjected to an overaging treatment, which is also called "partitioning treatment" in technical terms.

[0077] For this purpose, the sample was heated to a temperature TP and then held for a duration tP. Finally, the sample was cooled at a cooling rate ΘC to a temperature below 100°C. The temperature TP and the duration tP are also listed in Table 3.

[0078] For each sample treated in this way, the yield strength YS, tensile strength TS, and elongation A80 determined in accordance with DIN EN ISO 6892 (specimen form 2), the ratio YS / TS, the hole expansion rate HER determined in accordance with ISO 16630, and the product TS×A80 are shown in Table 4.

[0079] Furthermore, for each sample, the structure was analyzed in accordance with ISO 9042 regarding optical microscopes. The retained austenite content RA, primary (reverted) martensite content tM, secondary (fresh) martensite content fM, total bainite content and bainitic ferrite content bF, and polygonal ferrite content pF thus determined are shown in Table 5.

[0080] Furthermore, the areal density of V(C,N) precipitates present in the structure of the sample was determined by preparing a carbon replica and using a general transmission electron microscopy method.

[0081] Finally, the hydrogen embrittlement susceptibility coefficient H_sf was determined by the above method. The samples according to the present invention satisfy all the requirements of the present invention from the viewpoints of composition, manufacturing method, and heat treatment method, and it has been found that the susceptibility to hydrogen embrittlement is significantly lower than that of samples not alloyed, manufactured, or heat-treated according to the present invention.

[0082] [Table 1]

[0083] [Table 2]

[0084] [Table 3]

[0085] [Table 4]

[0086] [Table 5]

Claims

1. By mass percentage, C: 0.20 - 0.40%, Mn: 1.50 - 3.00%, Si: 0.90 - 1.50%, Al: 0.005 - 1.00%, V: 0.01 - 0.30%, Optionally Cr: 0.01 - 1.00%, Optionally Mo: 0.005 - 0.20%, Optionally B: 0.00001 - 0.002%, Optionally Nb and Ti: Total content of 0.005 - 0.2%, P: Maximum 0.020%, S: Maximum 0.005%, N: Maximum 0.008%, and Consisting of Fe and inevitable impurities as the balance, with the total content of the impurities being 0.8% or less, By area percentage, 65 - 92% of primary (tempered) martensite, and At least 8% of retained austenite (RA) Including, with the balance being Maximum 27% of secondary (untempered) martensite, Maximum 10% of bainite or bainitic ferrite, and / or Maximum 5% of polygonal ferrite Consisting of, with the total of the secondary (untempered) martensite, the bainite or the bainitic ferrite, and the polygonal ferrite being 27% or less, showing a structure determined according to ISO 9042, A high-strength cold-rolled steel sheet product provided with a steel substrate, having reduced susceptibility to hydrogen embrittlement.

2. The structure of the steel substrate is such that the precipitation density of V-based precipitates with a diameter of less than 10 nm, which was determined as described in the specification, is 1000 or more per μm ２ The steel sheet product according to claim 1, wherein the precipitation density is 1000 or more per μm

3. The primary martensite contained in the structure of the steel substrate has a fine lath structure having laths with a maximum length of 1000 nm determined as described in the specification, the steel sheet product according to claim 1 or 2.

4. The maximum length of the laths determined as described in the specification is 500 nm, the steel sheet product according to claim 3.

5. The C content in the steel of the steel substrate is 0.22 - 0.3% by mass, the steel sheet product according to any one of claims 1 to 4.

6. The Mn content in the steel of the steel substrate is 1.9 - 2.8% by mass, the steel sheet product according to any one of claims 1 to 5.

7. The V content in the steel of the steel substrate is 0.07 - 0.20% by mass, the steel sheet product according to any one of claims 1 to 6.

8. Having a yield strength of at least 1000 MPa, a tensile strength of at least 1300 MPa, and an elongation A80 of at least 10%, each determined according to DIN EN ISO 6892 (specimen form 2), the steel sheet product according to any one of claims 1 to 7.

9. The steel plate product according to any one of claims 1 to 8, wherein the hole expansion ratio determined according to ISO 16630 is at least 20%.

10. The steel plate product according to any one of claims 1 to 9, wherein corrosion-resistant plating is provided on at least one of the surfaces.

11. The steel plate product according to claim 10, wherein the plating is applied by electroplating, hot-dip galvanizing, or alloyed hot-dip galvanizing.

12. (a) providing molten steel consisting of, by mass%, 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 1.00%, optionally Mo: 0.005 to 0.20%, optionally B: 0.00001 to 0.002%, optionally Nb and Ti: total content of 0.005 to 0.2%, P: at most 0.020%, S: at most 0.005%, N: at most 0.008%, and Fe and inevitable impurities as the balance, with the total content of the impurities being 0.8% or less; (b) casting the molten steel into a slab; (c) heating the slab to a reheating temperature of 1000 to 1300 °C; (d) hot-rolling the reheated slab into a hot-rolled steel strip, the hot-rolling ending at a hot-rolling finishing temperature of 850 to 980 °C; (e) cooling the hot-rolled steel strip to a coiling temperature of 400 to 600 °C, finishing the cooling within at most 25 seconds after the end of the hot-rolling, and coiling the hot-rolled steel strip in a coil shape; (f) optionally, pickling the hot-rolled steel strip; (g) cold-rolling the hot-rolled steel strip with a cold-rolling reduction ratio of 20 to 80% to obtain a cold-rolled steel strip; (h) heating the cold-rolled steel strip to a soaking temperature TS which is at least 50 °C higher than the Ac3 temperature and at most 950 °C at a heating rate ΘS of 2 to 10 °C / s, the Ac3 temperature being determined using a dilatometry method according to SEP 1681-1998-06; immediately holding the cold-rolled steel strip at the soaking temperature TS for a soaking time tS exceeding 40 seconds and less than 200 seconds; Immediately, heat the cold-rolled steel strip at a quenching rate ΘQ of 20 to 100 °C / s to a quenching stop temperature TQ that is at least equal to a temperature TQ_min at which 65 to 92 area% of primary martensite exists in the structure of the cold-rolled steel strip and is lower than the martensite start temperature T_MS. The T_MS temperature is determined using a dilatometry method in accordance with SEP 1681-1998-06. Hold the cold-rolled steel strip at the quenching stop temperature TQ for 4 to 20 s. Thereby, a step of final annealing the cold-rolled steel strip. (i) Heat the cold-rolled steel strip to an overaging temperature TP of 380 to 460 °C. Hold the cold-rolled steel strip at the overaging temperature for 50 to 200 s. Cool the cold-rolled steel strip at a cooling rate ΘC of 0.5 to 20 °C / s to below 100 °C. A step of overaging the finally annealed cold-rolled steel strip by an overaging treatment including the above. A method for manufacturing a steel sheet product according to any one of claims 1 to 11.

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