High strength cold rolled steel sheet product with reduced susceptibility to hydrogen embrittlement and manufacturing method thereof

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 high strength cold rolled steel sheet products with reduced susceptibility to hydrogen embrittlement and to methods for making such steel sheet products. [Background technology]

[0002] In this specification, the term "steel plate product" is understood to mean a rolled product whose length and width are each significantly greater than its thickness, and therefore includes in particular steel strip, steel plate, and blanks obtained therefrom.

[0003] All information regarding the content of steel alloy compositions given in this application relates to mass unless expressly stated otherwise. Therefore, all % data referring to the composition of a steel alloy or another alloy mentioned herein should be understood as "% by mass" ("wt. %) information without reference to a reference unit.

[0004] To achieve weight-saving emissions reductions in the automotive and transportation industry while simultaneously increasing passenger safety, steel products with improved strength and ductility are required. The conflicting goals of strength and ductility have been united in advanced third-generation high-strength steel concepts (so-called "AHSS"), which include steels known as "quench and part" steels ("Q&P steels").

[0005] Q&P steels use retained austenite ("RA") as a structural component to improve the strain hardening and tensile strength of the steel, as well as increase elongation through the known 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 fresh martensite (secondary martensite) may also be present in the structure of Q&P steels.

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

[0007] The addition of trace alloying elements such as Ti, Nb, or V can minimize the susceptibility of steels to hydrogen embrittlement. Trace alloying elements form fine, coherent or semi-coherent carbide or carbonitride precipitates on the order of less than 10 nm, preferably less than 5 nm.

[0008] Precipitates formed by trace alloying elements are often referred to in the technical literature as "traps for diffusible hydrogen" because hydrogen atoms have a relatively strong binding energy to such fine precipitates. Therefore, hydrogen atoms that penetrate into steels with structures containing these fine precipitates will bind to the precipitate interfaces or to dislocations resulting from the incompatibility of the precipitates with the surrounding matrix structure.

[0009] The presence of a large number of such traps in the matrix structure slows the diffusion of hydrogen atoms through the structure, reducing their adverse effects on processes occurring during deformation and fracture.

[0010] However, using trace alloying elements in combination with Q&P steels is not straightforward: they form carbides with the carbon present in the steel, which is needed during partitioning to stabilize the retained austenite, consuming some of the carbon. Summary of the Invention

[0011] Against this background, the present invention seeks to identify high strength, easily formable steel sheet products that have a reduced tendency to hydrogen embrittlement.

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

[0013] Furthermore, the invention should identify a process that makes it possible to reliably manufacture such steel sheet products.

[0014] This object is solved by a method as defined in claim 12.

[0015] Advantageous embodiments of the invention are set forth in the dependent claims and are explained in detail below, as well as the general idea of ​​the invention.

[0016] According to the present invention, a high strength cold rolled steel sheet product having reduced susceptibility to hydrogen embrittlement is In mass%, C: 0.20~0.40%, Mn: 1.50-3.00%, Si: 0.90 to 1.50% Al: 0.005 to 1.00%, V: 0.01 to 0.30%, Optional Cr: 0.01~1% Optional Mo: 0.005~0.2% Optionally, B: 0.00001~0.002%, Optionally, Nb and Ti: total content 0.005 to 0.2%; P: 0.020% maximum, S: Maximum 0.005%, N: 0.008% maximum, and The balance is Fe and unavoidable impurities, and the total content of the impurities is 0.8% or less; Area % 65-92% primary (tempered) martensite, and At least 8% retained austenite (RA) and the remainder is up to 27% secondary (untempered) martensite, maximum 10% bainite or bainitic ferrite, and / or Maximum 5% polygonal ferrite and the total of the secondary (untempered) martensite, the bainite or bainitic ferrite, and the polygonal ferrite is 27% or less.

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

[0018] The good formability of the steel sheet products according to the invention, despite their high strength, is also reflected in a hole expansion ratio HER of more than 20%, the hole expansion ratio HER being determined in accordance with ISO 16630 currently in force.

[0019] Furthermore, extensive testing has demonstrated that the steel sheet products according to the invention have improved resistance to hydrogen embrittlement, with or without a hydrogen loading medium, as quantified in slow strain rate tests carried out in accordance with the currently valid DIN EN ISO 7539-7.

[0020] Thus, the present invention has significant advantages for structural and crash-related components in the automotive and transportation industries, including applications such as battery housings in electric vehicles. These applications require both high yield strength and a large ability to absorb plastic deformation. High yield strength minimizes intrusion during a crash, while high elongation values ​​allow for high energy absorption.

[0021] The method for producing a steel sheet product according to the invention proposed by the present invention comprises: (a) providing molten steel consisting, in mass%, of 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: a total content of 0.005-0.2%, P: 0.020% at most, S: 0.005% at most, N: 0.008% at most, and the balance being Fe and unavoidable impurities, wherein the total content of said impurities is 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, and the hot rolling is finished at a hot rolling finish temperature of 850 to 980 ° C.; (e) cooling the hot-rolled steel strip to a coiling temperature of 400 to 600°C, completing the cooling within a maximum of 25 seconds after the end of the hot rolling, and winding the hot-rolled steel strip into a coil; (f) optionally pickling the hot rolled steel strip; (g) cold rolling the hot rolled steel strip at a cold rolling reduction of 20 to 80% to obtain a cold rolled steel strip; (h) The cold-rolled steel strip is heated to a soaking temperature TS of at least 50°C higher than the Ac3 temperature and up to 950°C at a heating rate ΘS of 2 to 10°C / s; Immediately, the cold-rolled steel strip is held at the soaking temperature TS for a soaking time tS of more than 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 quench stop temperature TQ that is lower than the martensite start temperature TMS and is at least equal to the temperature TQ_min at which 65 to 92 area % of primary martensite is present in the structure of the cold rolled steel strip; The cold-rolled steel strip is held at the quenching stop temperature TQ for 4 to 20 seconds. final annealing of the cold rolled steel strip by (i) The cold-rolled steel strip is heated to an overaging temperature TP of 380 to 460 ° C., The cold-rolled steel strip is held at the overaging temperature for 50 to 200 seconds, The cold-rolled steel strip is cooled to less than 100°C at a cooling rate ΘC of 0.5 to 20°C / s. and overaging the final annealed cold rolled steel strip by an overaging treatment (also called "splitting treatment") which comprises:

[0022] To protect the steel sheet product according to the invention from corrosion, at least one of its surfaces may be provided with a corrosion-resistant coating in a conventional manner. Such a coating may be applied by electroplating, hot-dip galvanizing, or galvannealed hot-dip galvanizing. Typically, the coating 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") may be added in a known manner to optimize the properties of the coating. DETAILED DESCRIPTION OF THE INVENTION

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

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

[0025] The steel of the steel plate product according to the present invention contains 1.50 to 3.00 mass% manganese ("Mn"). Mn is an element that effectively increases the hardness and thus the strength of the steel. Furthermore, the presence of a Mn content within the range specified in the present invention suppresses the formation of ferrite and pearlite during quenching. By using a quenching rate of less than 100°C / s, a suitable structure containing martensite and retained austenite is obtained after the first quench (step i). Mn is also a solid-solution strengthening element that stabilizes austenite by lowering the Ms temperature. To ensure these favorable effects of the presence of Mn, the steel of the steel plate product according to the present invention needs at least 1.5 mass%, particularly at least 1.9 mass%, of Mn. However, if the Mn content exceeds 3.00 mass%, the welding properties deteriorate and segregation occurs, which also reduces the mechanical properties of the steel plate product. To ensure these adverse effects of the presence of Mn are avoided, the Mn content in the steel of the steel plate product according to the present invention may be limited to 2.8 mass%.

[0026] The steel of the steel sheet product according to the present invention contains 0.90 to 1.50 mass% silicon ("Si"). Si contributes to the strength of the steel by solid solution strengthening. Furthermore, silicon is insoluble in cementite, which inhibits the formation of iron carbides during partitioning (step i), promotes the stabilization of retained austenite, and improves ductility. To achieve these effects, at least 0.90 mass% Si is required in the steel of the steel sheet product according to the present invention. Furthermore, the combination of Al and 0.9% Si has a similar effect on the stability and amount of retained austenite, improving ductility. However, a high silicon content exceeding 1.50 mass% adversely affects the galvanizability of the steel sheet.

[0027] The steel of the present invention contains 0.005 to 1.00% by weight of aluminum ("Al"). Al is commonly used in steelmaking as a deoxidizer and to form aluminum nitrides, which contribute to the strength of the steel. Like Si, Al is insoluble in cementite, thus suppressing the formation of iron carbides during decomposition (step i). However, as the aluminum content increases, the Ac3 temperature of the steel from which the present invention is made would dramatically increase to values ​​too high for conventional industrial annealing lines. If Al is used solely as a deoxidizer, the Al content may be limited to a maximum of 0.1% by weight to avoid the formation of AlN. For deoxidation, an Al content of at least 0.005% by weight is required. Typically, in the steel alloy of the present invention, the Al content for these purposes ranges from 0.005 to 0.100% by weight, particularly 0.005 to 0.070% or 0.005 to 0.060% by weight. However, if Al is needed to promote austenite retention, a concentration of up to 1.0% is permitted. In this case, an Al content of at least 0.060% by weight, in particular at least 0.1% by weight or at least 0.3% by weight, is suitable.

[0028] The vanadium ("V") content of the steel constituting the steel sheet product according to the present invention ranges from 0.01 to 0.3% by weight. V is a trace alloy element with a much lower melting temperature than Ti or Nb. V-based precipitates are thereby partially or completely dissolved during the final annealing (step h) and precipitated during the subsequent cooling, or, most preferably, by carbon diffusion and partitioning during the partitioning stage (step i) of the final annealing cycle. This allows the formation of particularly fine V-based precipitates that are particularly effective in trapping diffusible hydrogen. The favorable effect of V becomes evident at concentrations above 0.01% by weight and continues to increase steadily up to 0.15% by weight. This effect saturates at concentrations above 0.3% by weight. It has been found that the presence of V is most effective in the range of up to 0.25% by weight, especially up to 0.20% by weight. This is particularly true for V contents of at least 0.07% by weight.

[0029] Chromium ("Cr") may optionally be added to the alloy of the steel substrate of the steel plate product according to the present invention to effectively delay the formation of pearlite and bainite and increase strength. This effect may be achieved by adding at least 0.01% by weight of Cr. However, to avoid the occurrence of intergranular oxidation, the Cr content is limited to 1% by weight. 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 weight. To avoid the adverse effects of the presence of Cr, the Cr content may be limited to 0.5% by weight.

[0030] As a further optional element, molybdenum ("Mo") may be added in an amount of 0.005 to 0.2 wt. % to the steel alloy from which the steel sheet product of the present invention is made. Mo improves the strength of the steel sheet and inhibits the formation of pearlite. Particularly effective in this respect are Mo contents of at least 0.02 wt. %; the positive effect of Mo is particularly pronounced at Mo contents of up to 0.15 wt. %.

[0031] Niobium ("Nb") and titanium ("Ti") are trace alloy elements that may be added, in combination or alone, as needed to the steel alloy of the steel sheet product according to the present invention. These elements effectively contribute to the strength of the steel sheet by precipitation hardening and refinement of the structure. However, the melting points of Nb and Ti are much higher than that of V. This means that, for example, the presence of Ti in combination with V can result in the formation of mixed carbides. These mixed carbides tend not to dissolve during final annealing and therefore remain coarse or grow to an ineffective size during final annealing. For this reason, the total content of Nb and Ti is limited to a maximum of 0.2% by weight. However, the beneficial effect of the presence of Nb and / or Ti on the strength and formability of the steel sheet product is reliably achieved when the total content of Nb and / or Ti is at least 0.005% by weight. A total range of 0.02 to 0.15% by weight of Nb and / or Ti is particularly effective in strengthening the steel.

[0032] The steel of the steel plate product according to the present invention may optionally contain boron ("B") in an amount of 0.00001 to 0.002% by weight to inhibit the formation of ferrite and segregation along grain boundaries and to prevent their migration, resulting in the formation of a fine grain structure that contributes to the mechanical properties of the steel. A B content of 0.0001 to 0.001% by weight is particularly effective in this regard.

[0033] The term "impurities" includes all elements that are introduced into or cannot be completely removed from steel during steelmaking. In the steel of the steel plate product according to the present invention, impurities can be detected by measurement, but in the case of elements not listed herein as essential or optional components or added as needed, they are present at very low contents and therefore below the effective limits specified herein. To avoid the detrimental effects of the total impurities, the total impurity content is limited to a maximum of 0.8% by weight. In particular, phosphorus ("P"), sulfur ("S"), and nitrogen ("N") are unavoidable impurities. However, it has been found that a maximum P content of 0.020%, a maximum S content of 0.005%, and a maximum N content of 0.008% by weight, respectively, do not degrade the properties of the steel of the steel plate product according to the present invention. Typically, 0.001-0.020% by weight of P, 0.0001-0.005% by weight of S, and 0.0001-0.008% by weight of N are present in the steel. A maximum copper ("Cu") content of 0.5% by weight, a maximum nickel ("Ni") content of 0.5% by weight, and a maximum oxygen ("O") content of 0.0080% by weight are also included as unavoidable impurities. Of course, other elements such as W, Co, Sn, Ca, Mg, REM, Zr, Te, As, and Bi are also considered as impurities.

[0034] Steel sheet products according to the present invention exhibit a structure comprising 65 to 92 area percent primary (tempered) martensite and at least 8 area percent retained austenite (RA). The retained austenite content fills the portion of the structure not occupied by primary martensite and other structural components permitted as required in accordance with the present invention. Thus, if no other components are present in the steel structure, the retained austenite occupies 8 to 35 area percent of the structure. However, if the structure contains only 65 area percent primary (tempered) martensite and a minimum amount of 8 area percent retained austenite, a total of 27 area percent of the components may be present, which is permitted as required in accordance with the present invention. Theoretically, the remaining proportions of each of the structural components may be filled solely by secondary (untempered) martensite, bainite or bainitic ferrite, and / or polygonal ferrite, with any combination of these components typically occurring.

[0035] The structure of the steel sheet product according to the invention contains 1 μm particles with a diameter of less than 10 nm. 2 The steel sheet product according to the present invention also contains V-based precipitates at a precipitate density of 1000 or more per square inch. This improves the resistance to hydrogen embrittlement of the steel sheet product according to the present invention.

[0036] The precipitate density of V-based precipitates is determined by a combination of transmission electron microscope images and X-ray microanalysis (TEM and EDX) using carbon extraction replicas. Carbon extraction replicas are obtained from longitudinal sections. The magnification ranges from 10,000x to 200,000x. Based on these images, the diameter of the precipitates within the measurement field can be calculated using computer-assisted image analysis. In each case, five measurement fields are measured for this purpose. The results of the five measurement fields are then averaged. The size of the measurement field varies depending on the selected magnification, ranging from 18.5 μm x 14.5 μm at 10,000x magnification to 0.925 μm x 0.725 μm at 200,000x magnification. The nature of the precipitates is simultaneously determined by EDX (energy dispersive X-ray spectroscopy). In this process, atoms within the precipitates are excited by an electron beam from a transmission electron microscope. From the emitted X-rays the elemental distribution within the sample can be determined 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-assisted image analysis, determine the diameter of all precipitates on five different measurement fields with a size of 1.85 μm x 1.45 μm at 100,000x magnification. 3. Identify vanadium precipitates based on the detected X-ray quanta. 4. Calculate the number of vanadium precipitates less than 10 nm in diameter in each of the five measurement fields, and determine the density of vanadium precipitates in each of the five measurement fields. 5. Determine the density of vanadium precipitates in the steel plate product as the average value of the density over five measurement fields.

[0038] If the austenite grains are too large, the number of martensite nucleation sites will decrease and the size of the martensite packets may increase, resulting in reduced local formability of the final product. The width of individual martensite laths is a function of their length. Thinner laths are advantageous for supporting the microstructural processes that occur during the overaging treatment (step i) of the method of the present invention. The precipitation of fine V-based precipitates along the phase boundaries of martensite and retained austenite grains in accordance with the present invention results in fine lath thicknesses with lengths of up to 1000 nm, typically up to 500 nm. The microstructure 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 scanning electron microscopy (SEM) and treated with 3% nital etching. Due to the fine structure, the microstructure is examined by SEM observation at a magnification of 5000x. The determined lath thickness length corresponds to the average of five measurements.

[0039] Furthermore, precipitates located at phase boundaries function as hydrogen traps, improving resistance to hydrogen embrittlement. By controlling the microstructure through the precipitation of fine V-based precipitates along the phase boundaries of martensite and / or retained austenite grains, excellent hydrogen embrittlement resistance and mechanical properties are achieved. The fine V(C,N) precipitates formed as a result of the alloy defined by the present invention and the method for manufacturing a steel sheet product according to the present invention result in finer prior austenite grains, which in turn result in thinner martensite / retained austenite lath structures. The V(C,N) nanoparticles, located along the lath interfaces, trap hydrogen penetrating into the steel substrate, improving resistance to hydrogen embrittlement.

[0040] The properties of the steel sheet product according to the invention as explained above can be reliably obtained by the method for manufacturing such a steel sheet product proposed by the invention.

[0041] To manufacture steel plate products according to the invention, molten steel having a composition corresponding to the specifications of the invention is prepared in the usual way, and this melt is then cast in the usual way to form at least one slab (steps (a) and (b) of the process according to the invention).

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

[0043] In step (d) of the method according to the invention, hot rolling of the slab into a hot rolled steel strip is carried out at a sufficiently high temperature so that the finish rolling temperature is in the range of 850°C to 980°C. If the temperature during hot rolling is too low, sufficient static recrystallization does not occur between the rolling steps. As a result, the resulting structure maintains a strong structure and high dislocation density through dynamic recrystallization. For this reason, the finish rolling temperature must not be lower than 850°C. Finish rolling temperatures above 980°C are technically not feasible.

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

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

[0046] To avoid the formation of pearlite, the coiling temperature must be a maximum of 600°C. Furthermore, high coiling temperatures can cause oxidizing elements such as Si, Cr, or Mn to diffuse into grain boundaries and form oxides, degrading the surface quality of the hot-rolled material and limiting the surface quality of the hot-rolled material after hot rolling and after plating as required. Furthermore, limiting the coiling temperature to a maximum of 600°C also prevents the formation of unwanted polygonal ferrite. A coiling temperature of 580°C or less increases the amount of bainite in the structure of the hot-rolled material. Setting the coiling temperature low, between 400°C and 500°C, can have a favorable effect in preventing intergranular oxidation. However, low coiling temperatures can promote the formation of a large amount of martensite, potentially increasing the hardness of the hot-rolled steel strip. Therefore, in cold rolling, batch annealing, typically performed at temperatures between 550°C and 600°C for more than 6 hours but less than 24 hours, is required to decompose the bainite / martensite. A homogeneous structure free from large amounts of martensite, which allows for narrower thickness and width tolerances during subsequent cold rolling, is reliably obtained by coiling the hot-rolled steel strip at a coiling temperature of 400°C to 600°C, preferably 400°C to 580°C, with a particularly useful coiling temperature of 500°C to 580°C.

[0047] After the hot-rolled steel strip has been wound into a coil and cooled to room temperature, the hot-rolled steel strip can be descaled, if necessary, by conventional methods, for example by pickling, to remove any scale present on the surface of the steel strip and to improve the surface quality of the method of the present invention (step f, if necessary).

[0048] Cold rolling is typically performed in one or more rolling steps in a conventional tandem mill, with intermediate batch anneals, as needed, as described above. Higher reductions may be achieved using reversing rolling stands during cold rolling. For convenience, a maximum cold reduction of 80% is required. The minimum reduction typically achieved during cold rolling is 20% to ensure adequate recrystallization during the final annealing step (step h) of the method of the present invention. However, higher reductions are useful for achieving a fine grain structure that results in the described product characteristics. Therefore, in the method of the present invention, a cold reduction of 20% to 80% is appropriate when cold rolling the hot-rolled steel strip of the method of the present invention into cold-rolled steel strip (step g).

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

[0050] The final annealing of the cold-rolled steel strip, performed as step (h) of the method of the present invention, determines most of the properties of the steel sheet product of the present invention. The average heating rate "ΘS" when heating the cold-rolled steel strip to the soaking temperature TS is 2-10°C / s. Heating rates exceeding 10°C / s may prevent proper recrystallization before austenitization, and heating rates below 2°C / s are not economical in continuous annealing lines. To ensure proper homogenization of C within the complete austenitic structure, the soaking temperature TS must be at least 50°C higher than the Ac3 temperature of the respective 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 equation (1): Ac3 [℃] = 910℃ + (-203√(%C) - 15.2%Ni + 44.7%Si + 31.5%Mo - 21.1%Mn))℃ / mass% %C = C content of each steel alloy in mass % %Ni = Ni content of the respective 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 alloy specifications of the present invention, the Ac3 temperature is typically in the range of about 750°C to about 920°C. The Ac3 temperature of the steel alloys within 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 a quartz tube. The quartz tube expands proportionally to temperature over a very wide temperature range (far exceeding 1000°C), while the steel sample undergoes transformation during both 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.

[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 crystals.

[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 seconds ≤ tS ≤ 200 seconds), preferably less than 200 seconds (40 seconds < tS < 200 seconds). Here, a soaking time tS of more than 60 seconds and less than 120 seconds (60 seconds < tS < 120 seconds) is particularly effective.

[0055] After the soaking stage, the steel strip undergoes the primary cooling step. The average cooling rate "ΘQ" for cooling the steel sheet product of 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, it is not necessary to increase the cooling capacity of the primary cooling step to a rate exceeding 100°C / s. Instead, a cooling rate exceeding 100°C / s during primary cooling increases production costs, leads to insufficient cooling, and adversely affects 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 quench stop temperature TQ, at which the cooling carried out in this manner 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 determined experimentally by dilatometry in a known manner or can be estimated according to the following equation (2): T_MS[℃])=539℃+(-423%C-30.4%Mn-7.5%Si+30%Al)℃ / mass% %C = C content of each steel alloy in mass % %Mn = Mn content of the respective steel alloy in mass % %Si = Si content of each steel alloy in mass% %Al = Al content of the respective steel alloy in mass%

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

[0059] The quenching stop temperature must 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 proportion of primary martensite after the initial quench and thereby obtain the temperature TQ_min, the Koistinnen-Marburger equation can be used, which is often expressed as: f m =1-exp(-0.011(T_MS-TQ)) where f m is the fraction of austenite that transforms to martensite upon quenching 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 should not be less than 65 area %.

[0061] Immediately after the first cooling, the steel sheet products according to the invention are held at their respective quenching temperatures TQ for at least 4 seconds and not more than 20 seconds, which is necessary to equalize the temperature throughout the thickness and to allow the final stages of the isothermal transformation to occur.

[0062] Following the holding step, the steel sheet product undergoes the overaging step (i) of the method of the present invention, which is also referred to in technical terms as the "separation step." In step (i), carbon partitions from the primary martensite toward the adjacent retained austenite. However, the separation of carbon is a thermally activated process that competes with the precipitation of nanoparticles. This means that if the partitioning process is carried out at too high a temperature or for too long a time, iron carbides may form, consuming much of the carbon needed to stabilize the retained austenite, resulting in a decrease in both the tensile strength and elongation of the steel sheet product. To avoid this effect, the upper limit of the overaging temperature Tp is set to 460°C in the present invention. The lower limit of the overaging temperature Tp is set to 380°C, ensuring the precipitation of a sufficient amount of V(C,N) nanoparticles, which would not occur at lower temperatures. In this regard, limiting the overaging temperature Tp to a maximum of 460°C has proven particularly appropriate, since it also avoids the coarsening of V(C,N) nanoparticles, which would reduce the mechanical properties of the steel sheet product. By setting the overaging temperature to 380°C to 460°C (380°C TP 460°C) and limiting the time tP for each steel plate product to be held at the temperature TP to 50 seconds to 200 seconds, the precipitation density of V-based precipitates with a diameter of less than 10 nm can be reduced to 1 μm 2 This will ensure that the number of samples per unit is 1000 or more.

[0063] After 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, thereby immediately stopping splitting and avoiding carbide coarsening or other thermally induced effects that may degrade the mechanical properties of the steel sheet product according to the invention. After the steel sheet product has cooled to below 100°C, cooling to room temperature is not critical and can be carried out in any suitable manner.

[0064] The improved resistance to hydrogen embrittlement of the steel plate product of the present invention 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, while other samples of the same steel plate product were exposed to an aqueous solution containing 5% NHSCN as a hydrogen charging medium. Subsequently, the maximum tensile stress Rm_max(air) of the samples exposed to the air atmosphere and the maximum tensile stress Rm_max(NHSCN) of the samples exposed to the hydrogen charging medium were determined in the usual standard manner. From the tensile stresses thus determined, the "hydrogen embrittlement susceptibility factor" ("H_sf") was calculated as follows: H_sf=1-Rm_max(NH4SCN) / Rm_max(air)

[0065] In the case of the steel plate products according to the invention, the hydrogen embrittlement susceptibility factor H_sf was found to be reliably a maximum of 0.35 (H_sf≦0.35). The tests thus demonstrated that the hydrogen embrittlement susceptibility of the steel plate products according to the invention is significantly lower than that of the plate samples used for comparison, which were subjected to the same test conditions and evaluated in the same way, but which consisted of conventional steel plates not alloyed in the way according to the invention, i.e., lacking in particular the V content specified in the invention. In contrast to the samples according to the invention, the comparative examples always showed hydrogen embrittlement susceptibility factors H_sf of at least 0.35.

[0066] In each of Tables 1 to 5 described below, values ​​that do not comply with the specifications of the present invention are underlined.

[0067] For practical testing and to demonstrate the effects obtained by the present invention, molten steel AM was melted with the composition shown in Table 1. In Table 1, values ​​and examples outside the range of the present invention are underlined.

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

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

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

[0071] After leaving the last hot rolling stand of the hot rolling line, the resulting hot-rolled steel strips were cooled to respective coiling temperatures CT in the range of 400°C to 600°C within a maximum duration tC of 25 seconds. The respective durations tC and respective coiling temperatures CT are also shown in Table 2.

[0072] After being coiled and cooled to room temperature, the hot rolled strip was pickled by methods known to those skilled in the art to remove scale from the surface.

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

[0074] The respective cold rolling reduction ratios CRD are also shown in Table 2.

[0075] After cold rolling, the resulting cold-rolled steel strip samples were final annealed by heating to a soaking temperature TS at a heating rate ΘS. Immediately thereafter, each sample was held at the soaking temperature TS for a soaking time tS. Immediately after soaking, the samples were quenched to a quench stop temperature TQ at a quenching rate ΘQ. As a result, the structure of the annealed samples had a primary martensite fraction %PM of 65-92 area %. After quenching, the annealed cold-rolled steel strip samples were held at the quench stop temperature TQ for a period tQ of 4-20 seconds. The respective soaking temperatures TS, heating rates ΘS, soaking times tS, quenching rates ΘQ, quench stop temperatures TQ, primary martensite fraction %PM (corresponding to tM in Table 5), and period tQ are listed in Table 3.

[0076] After the final annealing, the cold rolled strip samples were subjected to an overaging treatment, also known in technical terms as the "splitting treatment".

[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 temperatures TP and durations 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, each determined according to DIN EN ISO 6892 (specimen type 2), the ratio YS / TS, the hole expansion ratio HER determined according to ISO 16630 and the product TS x A80 are given in Table 4.

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

[0080] Furthermore, the areal density of V(C,N) precipitates in the structure of the samples was determined by carbon replica preparation and conventional transmission electron microscopy.

[0081] Finally, the hydrogen embrittlement susceptibility factor H_sf was determined using the method described above. It was found that the samples according to the present invention fulfilled all the requirements of the present invention in terms of composition, manufacturing method, and heat treatment method, and were significantly less susceptible to hydrogen embrittlement than samples that were not alloyed or manufactured and heat treated according to the present invention.

[0082] [Table 1]

[0083] [Table 2]

[0084] [Table 3] TIFF0007806936000004.tif101170

[0085] [Table 4]

[0086] [Table 5]

Claims

1. In mass%, 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 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: maximum 0.020% S: maximum 0.005%, N: 0.008% maximum, and The balance is Fe and unavoidable impurities, and the total content of the unavoidable impurities is 0.8% or less; In area %, 65-92% primary (tempered) martensite, and At least 8% retained austenite (RA) and the remainder is Maximum 27% secondary (untempered) martensite, a maximum total of 10% bainite and / or bainitic ferrite, and / or Maximum 5% polygonal ferrite and exhibiting a structure determined in accordance with ISO 9042, in which the sum of the secondary (untempered) martensite, the bainite and / or the bainitic ferrite, and the polygonal ferrite is 27% or less. A steel substrate is provided. The structure of the steel substrate has a precipitation density of V-based precipitates with a diameter of less than 10 nm, determined as described in the specification, of 1 μm ２ 1. A high strength cold rolled steel product having reduced susceptibility to hydrogen embrittlement, having a modulus of at least 1000 per square meter.

2. 2. The steel plate product of claim 1, wherein the primary martensite contained in the structure of the steel substrate has a microlath structure with laths having a maximum length of 1000 nm determined as described herein.

3. 3. A steel sheet product according to claim 2, wherein the maximum length of the laths determined as described in the specification is 500 nm.

4. 2. The steel plate product according to claim 1, wherein the C content in the steel of the steel substrate is 0.22 to 0.3 mass%.

5. 2. The steel plate product according to claim 1, wherein the Mn content in the steel of the steel substrate is 1.9 to 2.8 mass%.

6. 2. The steel plate product according to claim 1, wherein the V content in the steel of the steel substrate is 0.07 to 0.20 mass%.

7. 2. The steel sheet product according to claim 1, 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 type 2).

8. 2. The steel plate product of claim 1, having a hole expansion ratio determined in accordance with ISO 16630 of at least 20%.

9. 10. The steel sheet product of claim 1, wherein at least one of the surfaces is provided with a corrosion resistant plating.

10. 10. The steel sheet product according to claim 9, wherein the corrosion resistant coating is applied by electroplating, hot dip galvanizing, or galvannealed hot dip coating.

11. (a) providing a molten steel consisting, in mass%, of 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: a combined content of 0.005-0.2%, P: max 0.020%, S: max 0.005%, N: max 0.008%, and the balance being Fe and unavoidable impurities, the combined content of said unavoidable impurities being 0.8% or less; (b) casting the molten steel into a slab; (c) heating the slab to a reheat temperature of 1000-1300°C; (d) hot rolling the reheated slab into a hot rolled steel strip, the hot rolling being finished at a hot rolling finish temperature of 850 to 980°C; (e) cooling the hot-rolled steel strip to a coiling temperature of 400 to 600°C, completing the cooling within a maximum of 25 seconds after the end of the hot rolling, and winding the hot-rolled steel strip into a coil; (f) optionally pickling the hot rolled steel strip; (g) cold rolling the hot rolled steel strip at a cold rolling reduction of 20 to 80% to obtain a cold rolled steel strip; (h) The cold rolled steel strip is heated at a heating rate ΘS of 2 to 10°C / s to a soaking temperature Ts that is at least 50°C above the Ac3 temperature and up to 950°C, the Ac3 temperature being determined using dilatometry in accordance with SEP 1681-1998-06; Immediately, the cold-rolled steel strip is held at the soaking temperature TS for a soaking time tS of more than 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 quench stop temperature TQ that is lower than the martensite start temperature T_MS, the martensite start temperature T_MS being determined using dilatometry in accordance with SEP 1681-1998-06, The cold-rolled steel strip is held at the quenching stop temperature TQ for 4 to 20 seconds. final annealing of the cold rolled steel strip by (i) The cold-rolled steel strip is heated to an overaging temperature TP of 380 to 460°C, The cold-rolled steel strip is held at the overaging temperature TP for 50 to 200 s, The cold-rolled steel strip is cooled to less than 100°C at a cooling rate ΘC of 0.5 to 20°C / s. overaging the final annealed cold rolled steel strip by an overaging treatment comprising: The method for producing a steel plate product according to any one of claims 1 to 8, comprising:

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