Hot-rolled high-strength steel strip

Abstract

The invention relates to a hot-rolled high-strength steel strip having a composition, in wt%, C: 0.06-0.12 wt%, Mn: 1.20-2.20 wt%, Ti: 0.07-0.15 wt%, Si: up to 0.65 wt%, Al: up to 0.10 wt%, P: up to 0.02 wt%, S: up to 0.01 wt%, Nb: up to 0.03 wt%, N: up to 0.01 wt%, optionally one or more elements selected from the group consisting of Cr: up to 0.25 wt%, Mo: up to 0.10 wt%, V: up to 0.30 wt%, P: up to 0.20 wt%, S: up to 0.01 wt%, N: up to 0.01 wt%, and optionally one or more elements selected from the group consisting of Fe: up to 0.20 wt%, Mn: up to 0.20 wt%, N: up to 0.20 wt% %, Cu: up to 0.20 wt.%, Ni: up to 0.50 wt.%, B: up to 0.005 wt.%), the balance being Fe; the invention relates to a steel plate having a microstructure of precipitation-strengthened acicular / bainite ferrite and / or bainite having an MOD index of at least 0.85 at 1 / 4 thickness, a second phase component of at most 5 vol%, an average grain size (GSN 15) of at most 2.0 [mu] m in the rolling direction and based on a grain tolerance angle of 15 degrees, a total large-angle grain boundary length (sigma GBL 15-65) at 1 / 4 thickness of at least 1060 mm1, and a thickness of at least 10 [mu] m. And a gamma-fiber X-ray random intensity ratio measured by XRD at 1 / 2 thickness of at least 27; and wherein the steel strip has at least the following mechanical properties: a yield strength of 600-820 MPa, an ultimate tensile strength of 730-950 MPa, a tensile elongation of at least 11%, a reaming capacity of at least 45%, a work of fracture of at least 100 J. The invention also relates to a method for producing such a hot-rolled high-strength steel strip.

Description

Technical Field

[0001] The present invention relates to a hot rolled high strength steel strip, which is particularly suitable for, but not limited to, automotive parts. The present invention also relates to a method of manufacturing such a hot rolled high strength steel strip. In addition, the present invention relates to an automotive part incorporating the hot rolled high strength steel strip. Background Art

[0002] As is well known, with the increase of the strength of hot rolled (HR) steel strip, the formability decreases, and the fracture toughness in terms of crack sensitivity and the sensitivity to unstable brittle fracture may be compromised. The main application areas of HR steel in transportation and automotive applications are chassis and suspension (C&S). Other areas include frame rails of trucks, battery boxes of electric vehicles, or bumper beams. The typical thickness of hot rolled steel for these applications is less than 8mm, and is generally less than 6mm or less than 4.5mm. For example, thicker HR steel strips of up to 12mm can be used for engineering applications, such as crane booms, or for transportation applications of heavy truck frames. From the perspective of weight reduction, in order to be able to reduce the thickness of the steel strip, the above-mentioned applications must use higher strength steel. These applications of hot rolled steel require mechanical properties that are difficult to coordinate. In addition to high strength, the steel strip should also have good formability so that parts are manufactured by, for example, cold forming, because this is an energy-efficient manufacturing approach compared to hot forming. In addition, for applications such as bumper beams, battery housings, crane booms or frame rails, good impact and fracture toughness or energy absorption capacity are also required. In order to assemble these components, good weldability is also required. However, as the tensile strength of steel increases, the formability parameters decrease. Formability is a general term for steel sheets, which is considered to be a combination of material behavior during several mechanical operations such as stretching, bending, drawing and flanging. Depending on the component geometry, any property of the material or a combination of two or more properties is important during sheet metal forming. In addition to formability, fracture toughness may also be impaired by the increase in strength. This can lead to increased crack sensitivity and increased sensitivity to unstable brittle crack growth under specific mechanical load conditions, such as local compression in deep drawing operations.

[0003] Another important characteristic of steels for structural components in transport and automotive applications is their response to shearing in terms of cracking or delamination. Some steels may be prone to delamination when shearing, i.e. mechanical cutting or punching, with a typical industry-related clearance, i.e. usually in the range of about 10% to 18%, with a typical average of about 12%. This delamination is visible during shearing in the form of cracks or cracks in the fracture zone of the shear edge perpendicular to the movement of the shearing device. These cracks in the fracture zone of the shear edge can act as notches and thus lead to an exacerbation of shear edge fatigue and thus impair the service performance of structural components in transport and automotive applications that are subject to cyclic loads.

[0004] In order to save component weight, the usual approach is to apply high-strength steel and reduce the thickness of the steel strip used to reduce weight. However, this can lead to a loss of stiffness, which is undesirable for some applications in automotive parts. The inherent loss of stiffness caused by reducing the thickness of the steel strip used to make automotive parts can be recovered by optimizing the component geometry, such as producing a deeper flange or a flange with an increased degree of stretch or bending. In order to increase component stiffness through geometry optimization, high-strength steel strips require excellent formability in terms of tensile elongation and hole expansion capability.

[0005] Single-phase precipitation-hardened ferritic high-strength steels, freed from the conventional constraints between global formability (e.g. tensile elongation) and local formability (e.g. hole expansion capability or HEC) at high levels of both forming modes, may have low fracture toughness values ​​and increased edge crack sensitivity under certain conditions, such as under compression, and tend to have increased sensitivity to unstable brittle fracture behavior and delamination during shearing, which impairs shear edge fatigue.

[0006] Patent document EP1616970-A1 discloses a method for manufacturing high-strength hot-rolled steel sheets, comprising the following steps: reheating a steel slab at a temperature range of 1150-1300° C., wherein the steel slab is composed of 0.04-0.15% C, 1.5% or less Si, 0.5-1.6% Mn, 0.04% or less P, 0.005% or less S, 0.04% or less Al, 0.03-0.15% Ti, 0.03-0.5% Mo in weight %. and the balance of Fe and unavoidable impurities; hot rolling the reheated steel slab into a hot-rolled steel sheet at a finishing temperature above the Ar3 transformation temperature; primary cooling the hot-rolled steel sheet in a temperature range of 700-850°C at an average cooling rate of 20°C / s or more; holding the primary cooled steel sheet at a temperature of 680°C or more for longer than 1 second; and secondary cooling the steel strip at a temperature of 550°C or less at an average cooling rate of 30°C / s or more, and then coiling the steel strip. Preferably, the hot-rolled steel sheet is primarily cooled to a temperature range of not only 700-850°C but also (SRT / 3+300)-(SRT / 8+700)°C, where SRT represents the reheating temperature of the steel slab. In the embodiment of EP1616970-A1, this actually results in a primary holding temperature of 667-860°C. The processing conditions are such as to obtain a microstructure consisting of a second phase of ferrite, bainite and / or martensite containing precipitates and other phases, wherein the percentage of ferrite containing precipitates is 40-95% and the percentage of other phases is less than 5%.

[0007] Patent document EP1338665-A1 discloses a method for manufacturing high-strength hot-rolled single-phase ferritic steel using nano-precipitated steel sheets, comprising the following steps: manufacturing a steel slab, which is basically composed of the following in weight %, C: less than 0.06%, Si: less than 0.5%, Mn: 0.5-2.0%, P: less than 0.06%, S: less than 0.005%, Al: less than 0.1%, N: less than 0.006%, Mo: 0.05-0.6%, Ti: 0.02-0.10%, and the balance is Fe, satisfying the equation 0.8≤(C / 12) / [(Ti / 48)+(Mo / 96)]≤1.3; producing a hot-rolled steel sheet by hot rolling the steel slab at a temperature above the Ar3 transformation point; and coiling the hot-rolled steel sheet at a temperature of 550-700°C. The processing conditions are such that a microstructure consisting essentially of a matrix of a ferrite structure single phase and fine precipitates dispersed in the matrix is ​​obtained, wherein the fine precipitates are complex carbides containing Ti and Mo and have a grain size of less than 10 nm, wherein the fine precipitates are dispersed in a volume of 5×10 4 / mm 3 The above quantities are scattered.

[0008] Patent document WO2013 / 167572-A1 discloses a high-strength hot-rolled single-phase ferritic steel with nano-precipitates steel sheet, which has an excellent combination of tensile strength of at least 550 MPa and formability, and contains, by weight%, up to 0.15% C, up to 0.5% Si, 0.5-2.0% Mn, up to 0.06% P, up to 0.008% S, up to 0.1% Al-Sol, up to 0. 0.02% N, 0.02-0.45% V, 0.05-0.7% Mo, optionally 0.01-0.1% Nb, the balance Fe and inevitable impurities, wherein the steel sheet has a precipitation-strengthened and substantially single-phase ferrite microstructure, wherein the volume fraction of the ferrite phase in the microstructure is not less than 97%, and wherein the precipitates in the microstructure comprise fine precipitates of complex carbides containing Mo and V and optionally Nb. The hot rolled steel strip is manufactured by hot rolling at a hot finishing rolling temperature above the Ar3 transformation point and coiling the hot rolled sheet at a temperature range of 700°C-585°C.

[0009] Patent document WO2018 / 193032-A1 discloses a high-strength steel strip having a cementite-free microstructure, comprising: 0.005-0.08 wt. % C; 1.30-2.30 wt. % Mn; 2-35 ppm B; 5-65 ppm N; 0.005-0.1 wt. % Al_tot; 0.03-0.20 wt. % Ti; 0-1.5 wt. % Cu; 0-0.75 wt. % Cr; 0-0.05 wt. % Mo; 0-0.5 0 wt% Ni; 0-0.30 wt% V; 0-0.6 wt% Si; 0-0.01 wt% P; 0-0.01 wt% S; C / (Ti_sol+V)≤0.25, wherein Ti_sol=Ti-((48 / 14)-N), the remainder being iron and unavoidable impurities, said steel strip having a yield strength of at least 570 MPa, a tensile strength of at least 760 MPa, a total elongation (A50) of at least 10.3% and a hole expansion (λ) value of at least 70%. The steel strip is manufactured by hot rolling, cooling the hot rolled strip to a coiling temperature below 500°C at an average cooling rate of 15-100°C / s on a run-out table within 2 seconds between finishing and the start of cooling, then coiling cooling by natural cooling to ambient temperature, cold rolling at a reduction of 50-90%, preferably followed by solution heat treatment to allow bainite transformation.

[0010] There is a need for a hot rolled high strength good punchable steel which, in addition to having sufficient strength to reduce gauge and weight, provides a suitable combination of global (tensile elongation) and local (hole expansion capability) formability in terms of component geometry optimization and sufficient stiffness, while also providing a sufficient level of fracture toughness to minimize brittle fracture behavior and shear edge quality without cracking or delamination, with an industrially relevant shear gap to optimize shear edge fatigue, and thereby achieves an optimized balance of all these material properties, namely formability in terms of tensile elongation and hole expansion capability, fracture toughness in terms of crack sensitivity and work of fracture, and sensitivity to cracking during shearing. Summary of the invention

[0011] As will be understood herein, with respect to any description of alloy compositions or preferred alloy compositions, all references to percentages are by weight unless otherwise indicated.

[0012] As used herein, the term "about" when used to describe a compositional range or amount of an alloying addition means that the actual amount of the alloying addition may vary from the nominal intended amount due to factors such as standard processing variations as understood by those skilled in the art.

[0013] As used herein, the terms "up to" and "up to about" expressly include, but are not limited to, the possibility of zero weight percent of the particular alloy component to which they refer. For example, up to 0.20% Cu may include a steel that contains no Cu.

[0014] As used herein, the term "at 1 / 4 thickness" when used to describe microstructural features determined by microscopy techniques, in particular by electron backscatter diffraction (EBSD), refers to the area below the surface of the steel strip between 1 / 8 and 3 / 8 of the thickness of the steel strip, and preferably the area below the surface of the steel strip between 3 / 16 and 5 / 16 of the thickness of the steel strip.

[0015] As used herein, the term "at 1 / 2 thickness" when used to describe texture features determined by X-ray diffraction (XRD) refers to the area between 3 / 8 and 5 / 8 of the thickness of the steel strip below the surface of the steel strip, and preferably the area between 7 / 16 and 9 / 16 of the thickness of the steel strip below the surface of the steel strip.

[0016] The object of the present invention is to provide a hot rolled high strength deep drawable steel strip having high edge ductility, sufficient fracture toughness and good shear properties with respect to hole expansion capability, substantially without cracking or delamination, which is particularly suitable for complex components for automotive applications.

[0017] The object of the present invention is to provide a hot rolled high strength deep drawable steel strip having a yield strength of 600 to 820 MPa, an ultimate tensile strength of 730 to 950 MPa, high edge ductility with respect to hole expansion capability, sufficient fracture toughness and good shear properties, which is particularly suitable for complex components for automotive applications.

[0018] Another object of the present invention is to provide a method for producing such hot rolled high strength deep drawable steel strip having an improved balance of yield strength between 600 MPa and 820 MPa, ultimate tensile strength between 730 and 950 MPa, high edge ductility in terms of hole expansion capability, sufficient fracture toughness and good shear properties with substantially no cracking or delamination.

[0019] These and other objects and further advantages are met or exceeded by the present invention, which provides a hot rolled high strength steel strip having, in weight %,

[0020] C: 0.06-0.12 wt%, Mn: 1.20-2.20 wt%, Ti: 0.07-0.15 wt%, Si: up to 0.65 wt%, Al: up to 0.10 wt%, P: up to 0.02 wt%, S: up to 0.01 wt%, Nb: up to 0.03 wt%, N: up to 0.01 wt%, optionally one or more elements selected from the following: (Cr: up to 0.25 wt%, Mo: up to 0.10 wt%, V: up to 0.30 wt%, Cu: up to 0.20 wt%, Ni: up to 0.50 wt%, B: up to 0.005 wt%), the balance is Fe and inevitable impurities generated by ironmaking and steelmaking processes;

[0021] And wherein the steel strip has the following microstructure:

[0022] - precipitation-hardened acicular / bainitic ferrite and / or bainite,

[0023] - a misorientation distribution (MOD) index of at least 0.85 at 1 / 4 thickness,

[0024] - up to 5% by volume of second phase components, including cementite, pearlite, martensite and / or retained austenite,

[0025] - a number average grain size (GS) of at most 2.0 μm based on a grain tolerance angle of 15° at 1 / 4 thickness in the rolling direction N15 ),

[0026] - At least 1060mm at 1 / 4 thickness -1 The total high-angle grain boundary length (∑GBL 15-65 ),and

[0027] - a gamma-fiber X-ray random intensity ratio of at least 27 measured by XRD at 1 / 2 thickness;

[0028] And wherein the steel strip has at least the following mechanical properties: yield strength (Rp0.2) of 600-820 MPa, ultimate tensile strength (Rm) of 730-950 MPa, tensile elongation (A50) of at least 11%, hole expansion capacity (HEC) of at least 45%, and work of fracture (WOF) of at least 100 J.

[0029] According to the present invention, it has been found that hot rolled steel strips having these narrow alloy composition ranges and microstructures provide an improved balance of high edge ductility as indicated by hole expansion capability (i.e. HEC ≥ 45%), adequate fracture toughness (e.g. ATCL ≤ 20 mm), and work of fracture (WOF) ≥ 100 J, and good shear properties with substantially no cracking / delamination of the hot rolled deep drawable steel strip, the hot rolled steel strip having a yield strength (Rp0.2) of 600-820 MPa, a tensile strength (Rm) of 730-950 MPa, and a tensile elongation A50 of 11% or more, preferably a tensile elongation A50 of 14% or more.

[0030] Reduced crack sensitivity, in particular reduced edge crack sensitivity, can be objectively expressed as average total crack length (ATCL) following the measurement method described herein. The steel strip according to the present invention has an ATCL of less than 20 mm, preferably less than 10 mm, more preferably less than 8 mm, most preferably less than 5 mm.

[0031] The work of fracture (WOF) is considered a measure of the fracture toughness of the steel sheet and is ≥ 100 J. In one embodiment, WOF ≥ 125 J, preferably WOF ≥ 140 J. In one embodiment, WOF is at least 180 J, preferably at least 200 J.

[0032] The desired microstructure of such hot rolled high strength steel strip is obtained by a narrow composition range and by careful control of the manufacturing processes, such as the finishing temperature (FRT), controlled cooling of the hot rolled steel strip on the run-out table (ROT) and in particular a narrow operating window of the coiling temperature (CT).

[0033] After intensive research, the inventors discovered an empirical relationship between the hole expansion capacity (λ) and the so-called T factor (R 2 is about 0.86), according to the following formula, the T factor is defined as 1.2 times the power of MOD multiplied by the sum of γ-fiber, θ-fiber and {110} <001> The sum of the X-ray random intensity ratios of the Goss texture components divided by the α*-part, {112} <131> , {332} <113> and {554} <225> The sum of the X-ray random intensity ratios of the texture components:

[0034]

[0035] Here, the α*-part is defined as the Euler space or orientation distribution function (ODF) from {115} <110> To {223} <110> of The portion of the α-fiber in the region (and including the X-ray random intensity ratio of these two external texture components) covers the range of 15° to 45° on the α-fiber.

[0036] The T factor is defined to be at least 1.1. In one embodiment, the T factor is at least 1.15, more preferably at least 1.4. An increased T factor provides an increased HEC. N15 Too high (e.g. 1 / 4 thickness, >2.0 μm) and / or ∑GBL 15-65 Too low (1 / 4 thickness, <1060mm -1 ), the work of fracture is too low and some cracking may still occur in shear, so a T factor of at most 1.8 is preferred.

[0037] In order to obtain a hole expansion capability (HEC) of λ≥45%, preferably λ≥50%, and more preferably λ≥60%, the microstructure needs to be essentially acicular / bainitic ferrite and / or bainite fine-grained microstructure. The amount of second phase components including cementite, pearlite, martensite and / or retained austenite should not exceed 5% by volume in total, preferably not more than 3% by volume, because too high a fraction of hard carbon-rich second phase components will lead to too much damage during shearing, which is at the expense of hole expansion capability.

[0038] In addition, it is required to have a misorientation distribution (MOD) index of at least 0.85 at 1 / 4 thickness, preferably at least 0.90, more preferably at least 0.95, and most preferably at least 1.0. In general, an increase in the MOD index is associated with an increase in hole expansion capability. The MOD index is a measure of microstructural properties that is dominated by the phase components that make up the matrix. High-temperature phase components, such as polygonal ferrite, typically have a low MOD index. As the transformation temperature decreases, the ferrite morphology will transform into quasi-polygonal or irregularly shaped ferrite, through acicular and / or bainitic (feathery) ferrite toward upper bainite and ultimately toward lower bainite before reaching martensite. With this change in matrix morphology, the MOD index will increase, and this change in matrix morphology is mainly related to the ROT cooling trajectory and is mainly related to the coiling temperature. In general, a reduction in the coiling temperature will lead to an increase in the MOD index. It is believed that the increase in the MOD index leads to a more complex crystal arrangement, which is beneficial to the improved hole expansion ability, and the increase in the MOD index with atomically reduced coiling temperature is accompanied by refined cementite particles due to reduced diffusivity, which is beneficial to inhibit damage and void formation during shearing, and to inhibit void coalescence and crack formation on the shear edge of the stretch flange.

[0039] An important feature of the present invention is that the microstructure must be sufficiently grain refined and characterized by a number average grain size (GS) of at most 2.0 μm, preferably at most 1.9 μm, more preferably at most 1.8 μm, most preferably at most 1.7 μm at 1 / 4 thickness in the rolling direction and based on a grain tolerance angle of 15°. N15 ), and at least 1060 mm / mm measured by EBSD based on an adjacent grain orientation difference of 15° to 65° at 1 / 4 thickness in the rolling direction 2 (mm -1 ), preferably at least 1100 mm -1 , more preferably at least 1200 mm -1 , most preferably at least 1300mm -1 The total grain boundary length (∑GBL 15-65 ). High-angle grain boundaries effectively inhibit or prevent crack propagation. Therefore, grain refinement and increased high-angle grain boundary density are conducive to increased fracture toughness and shear performance, leaving a smoother fracture zone after shearing, which in turn is conducive to shear edge fatigue. This grain refinement is achieved by controlling the interaction between alloy composition and hot rolling process parameters. Regarding the hot rolling process, lower finishing temperatures lead to a higher temperature at the non-recrystallization temperature T nr Increased rolling below 100° will be beneficial in promoting grain refinement of the final microstructure. However, specific alloying elements in the steel composition will also play a role as they affect T nr In addition to the microalloying elements niobium (if present) and titanium (used to increase Tnr In addition to increasing the fraction of unrecrystallized austenite at the end of hot rolling and before phase transformation, carbon also increases T nr The increase in carbon content leads to a nr This is beneficial to increase the fraction of uncrystallized austenite at the end of the hot rolling process before the austenite to ferrite phase transformation, which in turn will promote grain refinement of the final microstructure. nr In addition to the effect of the Ar3 phase transformation temperature, the Ar3 phase transformation temperature also plays a role in determining the average grain size of the final microstructure. Carbon inhibits the phase transformation of Ar3 austenite to ferrite, which will promote grain refinement. Therefore, carbon inhibits the Ar3 austenite to ferrite phase transformation, which will promote grain refinement. nr The influence of the austenite to ferrite phase transformation temperature Ar3 is beneficial to promote the grain refinement of the final microstructure, which in turn is beneficial to improve the fracture toughness and shear properties.

[0040] Furthermore, the increased fraction of unrecrystallized austenite at the end of hot rolling and before the phase transformation increases {554} <225> In particular, <113> The X-ray random intensity ratio of the texture components is related to the increased fracture toughness and is beneficial to reduce the average total crack length (ATCL) and increase the work of fracture (WOF). Moreover, the increased non-recrystallized austenite fraction at the end of hot rolling and before the phase transformation reduces the {001} <110> Rotating Cube and {110} <001> Goss's X-ray random intensity ratio, which is related to embrittlement, delamination behavior of steel during fracture, and cracking during shear.

[0041] At the same time, a high fraction of unrecrystallized austenite at the end of hot rolling and before the phase transformation must be avoided, as this increases the {112} <110> In particular, {113} <110> and {112} <131> The X-ray random intensity ratio of the texture components, which is associated with increased (in-plane) anisotropy, potentially leads to reduced punchability and reduced hole expansion capability.

[0042] On the other hand, too low a fraction of unrecrystallized austenite at the end of hot rolling and before the phase transformation can lead to {001} <110> Rotating Cube and {110} <001> The excessively high X-ray random intensity ratio of Goss textured components is associated with reduced punchability, embrittlement of the steel, and delamination behavior during fracture or shear.

[0043] The hot rolled steel of the present invention has a microstructure of precipitation-hardened acicular / bainitic ferrite and / or bainite. This means that at least 95% by volume (volume fraction) of the microstructure is formed by precipitation-hardened acicular / bainitic ferrite and / or bainite, and more preferably at least 97% by volume, and most preferably at least 98% by volume.

[0044] The steel has up to 5 volume % of second phase components, including any cementite, pearlite, martensite and / or retained austenite, and preferably up to 3 volume % of second phase components, and therefore has limited internal microstructural damage when sheared; the hole expansion ability is largely plasticity driven and less damage controlled, and is subsequently related to the r-value, and therefore to the punchability of the steel. In other words, the hole expansion ability is mainly determined by the amount of in-plane plastic strain transverse to the load direction and the amount of plastic strain through the thickness. Therefore, for these steels, i.e. high-strength steels with up to 5 volume % of second phase components, preferably up to 3 volume % of second phase components, more preferably up to 2 volume %, an improvement in the punchability of the steel can be found, and thus the average or normal r-value (r m ):

[0045]

[0046] and reduce the in-plane anisotropy coefficient (r p ):

[0047]

[0048] Or increase the minimum r value (r min ). As the strength of the γ-fibers associated with the grains having the {111} planes parallel to the sheet surface increases, the punchability and r-value increase, but at the expense of the strength of the texture component on the α-fibers associated with the grains having the {100} planes parallel to the sheet surface. Therefore, for the hot-rolled high-strength steel strip having a hole expansion capability of at least 45%, preferably at least 50%, and most preferably at least 60% involved in the present invention, the hot-rolled texture is required to have γ-fibers with an X-ray random intensity ratio of at least 27 in the orientation distribution function (ODF) of Φ2=45° at 1 / 2 the strip thickness. In a preferred embodiment, the X-ray random intensity ratio of the γ-fibers is at least 30, more preferably at least 35.

[0049] By increasing the {001} <110> Rotating Cube and {110} <001> Goss strength, and can also (further) increase hole expansion capacity at the expense of fracture toughness. These are texture components associated with embrittlement and delamination. Although this is detrimental to fracture toughness and the shear edge quality of cracking in shear, the increased delamination sensitivity promotes crack extension in the RD-TD plane (RD = rolling direction; TD = transverse direction) parallel to the surface of the steel strip. These preferential crack paths in the RD-TD plane can be beneficial for deflecting crack extension, for example in hole expansion capacity tests, and in this way delay or prevent crack extension through the thickness, i.e. in the ND direction (ND = normal). Since fracture through the entire thickness of the steel strip is the criterion for determining edge ductility by hole expansion capacity tests, any deflection of the main crack extending through the thickness will delay the development of a fully developed through-thickness fracture. This leads to increased hole expansion capacity, although at the same time it may be associated with reduced fracture toughness and increased sensitivity to cracking in shear.

[0050] Carbon is present in an amount of 0.06-0.12 wt.%. In order to obtain sufficient strength, a suitable minimum C content is 0.060 wt.%. In a preferred embodiment, it is at least 0.07 wt.%, and more preferably the C content is at least 0.075 wt.%. Carbon has an influence on the non-recrystallization temperature T nr The influence of the austenite to ferrite phase transition temperature Ar3 is conducive to promoting the grain refinement of the final microstructure, which in turn is conducive to improved fracture toughness and shear properties. In addition, in combination with carbide-forming microalloying elements such as titanium, niobium (if added) and / or vanadium (if added), C is an essential element to achieve precipitation strengthening, and removes C to suppress the excessive fraction of cementite or other carbon-rich secondary phase components in the final microstructure. By optimizing other alloying elements, including Ti, Nb and / or V, an almost uniform precipitation-strengthened acicular / bainitic ferrite and / or bainite microstructure can be obtained, with substantially no or only very little cementite. In a preferred embodiment, the C content is at most 0.10 weight % to prevent excessive second phase components in the final microstructure, which may deteriorate the hole expansion capability.

[0051] The steel strip has Mn in the range of about 1.20-2.20 wt. % to obtain sufficient hardenability and grain refinement. In addition to carbon, manganese also suppresses the Ar3 transition temperature and thus helps promote grain refinement of the final microstructure. In one embodiment, the Mn content is in the range of about 1.40 wt. % to 2.20 wt. % for an improved balance of strength, corrosion resistance, fracture toughness and edge crack sensitivity. Preferably, the Mn content is at least about 1.50 wt. %, more preferably at least about 1.60 wt. %, and most preferably at least 1.70 wt. % to obtain sufficient grain refinement to improve fracture toughness and reduce crack sensitivity. In one embodiment, the Mn content is a maximum of about 2.10 wt. %, more preferably a maximum of about 2.0 wt. %. Too high a Mn content can lead to segregation during the casting process, which adversely affects the desired balance of properties.

[0052] Titanium is another essential alloying element, which is present in an amount of about 0.07 wt. % to 0.15 wt. % because it provides hardenability and, as a carbide former, inhibits the formation of cementite while providing precipitation strengthening by forming small titanium-based carbides. However, depending on the specific chemical composition of the steel, Ti also combines with N, S and C to form nitrides and carbon sulfides. Therefore, it is preferred that at least about 0.08 wt. % Ti is present to combine substantially all of the N and S in the steel and have a sufficient excess of Ti to combine with the C in the steel. In a preferred embodiment, at least about 0.09 wt. % Ti is present. When the amount of Ti present is greater than about 0.15 wt. %, coarse Ti nitrides, carbonitrides and carbides can be formed, which are difficult to dissolve during the process of reheating the slab before hot rolling. In addition, these coarse Ti nitrides, carbonitrides and carbides lead to a deterioration in the hole expansion ability of the steel. In one embodiment, the Ti content does not exceed 0.13 wt. %, preferably does not exceed 0.12 wt. %.

[0053] The composition of the steel according to the invention has a C / (Ti_sol+V) of >0.25, preferably >0.30, more preferably >0.40, and wherein Ti_sol represents the amount of Ti that can form carbides, and to calculate, and emphasize the role of carbon in promoting grain refinement.

[0054] Silicon can be present in an amount of up to about 0.65 wt %, and preferably in an amount of about 0.10 to 0.65 wt %, to improve the strength of the steel by substitutional solid solution strengthening of the iron lattice. In addition, Si is conducive to suppressing the formation of cementite. However, when using a higher amount of Si, the weldability and paintability of the steel deteriorate, so the amount of Si is preferably up to about 0.55 wt %, and in a preferred embodiment up to about 0.50 wt %. In one embodiment, the amount of Si is at least about 0.20 wt %, more preferably at least about 0.25 wt %.

[0055] In the steel according to the present invention, the performance of aluminum is comparable to that of Si. It slows down the carbide precipitation kinetics and inhibits the formation of cementite. When Al is less than about 0.01% by weight, the effect of inhibiting carbide formation is negligible. It is believed that aluminum values ​​below 0.01% by weight are residues of the deoxidation step in the steelmaking process, so the preferred minimum value is about 0.01% by weight. On the other hand, when Al is higher than about 0.10% by weight, excessive oxides can be formed during the thermomechanical processing of steel (slab reheating, hot rolling, coiling, etc.), which can deteriorate formability and toughness. Moreover, as Al increases, a higher amount of surface oxidation can occur at higher temperatures. These scales are harmful to hot rolling, pickling, coating and overall surface appearance. Moreover, when Al is increased to a level that makes steel very brittle and more difficult to hot roll together with the presence of Si, the rolling force during hot rolling increases. Therefore, the presence of Al in the present invention is up to about 0.10% by weight, preferably about 0.01-0.10% by weight, and more preferably about 0.03-0.08% by weight.

[0056] Nitrogen, sulfur and phosphorus are residual elements present in steel due to the steelmaking and refining processes. Their content is limited to up to about 0.01% by weight of S, up to about 0.02% by weight of P and up to about 0.01% by weight of N. Contents above these are harmful to mechanical properties, formability, toughness, shear edge quality, fatigue and weldability. In one embodiment, P is only present at most 0.015% by weight. In one embodiment, S is present in an amount of only up to 0.005% by weight (50ppm), more preferably only up to 0.0025% by weight (25ppm), and most preferably up to 0.0012% by weight (12ppm). N forms titanium nitride with Ti, which acts as a dispersoid for austenite grain size control during reheating. However, too high N can lead to too many coarse TiN particles, which can impair hole expansion ability. Preferably, the N content is up to 0.0075% by weight (75ppm), more preferably up to 0.0065% by weight (65ppm). A suitable minimum N content is about 0.0010 wt% (10 ppm).

[0057] Niobium may be present in the steel in an amount of up to 0.03 wt %. Nb improves the strength of the steel partially but mainly by grain refinement through precipitation hardening. This grain refinement is also beneficial for improving fracture toughness. However, since Nb is a relatively expensive alloying element and can promote anisotropy, the Nb content is preferably no more than 0.025 wt %, more preferably no more than 0.02 wt %, and even more preferably no more than 0.018 wt %. In one embodiment, the Nb content is at least 0.005 wt %, more preferably at least 0.010 wt %. It has been found that by keeping Ti at a sufficiently high level, the use of Nb can be overcome while still achieving the target balance of formability, mechanical properties and improved fracture or edge cracking characteristics. In addition, Nb can promote anisotropy to increase, and in addition has a high tendency to segregate and form coarse NbC particles due to centerline segregation, so avoiding the use of Nb leads to improved hole expansion ability and increased shear edge quality.

[0058] In another embodiment, Nb is not purposefully added but is present only as an unavoidable impurity generated during the iron and steel making process. In practice, this means that it may be present at up to about 0.005 wt. %, more preferably up to about 0.003 wt. %.

[0059] The steel strip may optionally have one or more elements selected from the group consisting of (Cr: up to 0.25 wt. %, Mo: up to 0.10 wt. %, V: up to 0.30 wt. %, Cu: up to 0.20 wt. %, Ni: up to 0.50 wt. %, B: up to 0.005 wt. %).

[0060] Cr increases the hardenability of the steel and may be added to the steel strip in an amount up to about 0.25 wt. %, preferably up to about 0.10 wt. %, as it may also impair the corrosion properties of the uncoated steel strip substrate. In the most preferred embodiment, Cr is not purposefully added but is present only as an unavoidable impurity generated during the iron and steel making process.

[0061] The steel strip may have up to 0.10 wt. % Mo to obtain hardenability. In addition, the presence of Mo may also improve the weldability of the steel strip and may contribute to precipitation strengthening, since Mo as a carbide former may promote complex carbide precipitation in combination with Nb, Ti and / or V.

[0062] Since Mo is an expensive alloying element, in another preferred embodiment, the steel strip contains Mo as a tolerable impurity element. In practice, this means that it may be present in an amount of up to about 0.050 wt. %, more preferably up to about 0.03 wt. %, and most preferably up to about 0.01 wt. %.

[0063] When present, copper can increase the strength of the steel strip by solid solution strengthening as well as precipitation hardening by copper precipitates. However, the Cu content should not exceed about 0.20 wt % because excessively high levels of Cu can promote hot shortness that can occur during hot working of the steel strip, which leads to embrittlement of the steel. In one embodiment, Cu is not added as an intentional alloying element and can be present at up to about 0.10 wt %, more preferably up to about 0.05 wt %, and most preferably up to about 0.03 wt %.

[0064] Up to about 0.50 wt %, preferably up to about 0.30 wt %, of nickel increases impact toughness and counteracts any hot shortness that may occur during hot working of the steel strip due to the presence of copper. In one embodiment, Ni is not added as an intentional alloying element and may be present up to about 0.10 wt %, more preferably up to about 0.05 wt %, and most preferably up to about 0.03 wt %.

[0065] Boron is not necessary to obtain the desired balance of steel strip properties, but may be present in an amount up to about 0.005 wt. %, thus up to 50 ppm, preferably up to about 0.0030 wt. % (30 ppm), most preferably up to about 0.0025 wt. % (25 ppm). B is very effective in increasing the hardenability of the steel, which means that lower cooling rates can be used on the run-out table. In one embodiment, the B content is at least 0.0010 wt. % (10 ppm) and up to 0.0030 wt. % (30 ppm).

[0066] The amount of vanadium present in the steel can be up to about 0.30 wt %, preferably up to about 0.20 wt %, more preferably up to about 0.17 wt %, and most preferably up to about 0.10 wt %. However, V is a relatively expensive alloying element, which is mainly used to replace Ti to obtain its precipitation strengthening effect, and to avoid cementite formation by forming vanadium carbides or composite carbide precipitates with Nb, Ti and / or Mo. In a preferred embodiment, V is purposefully added in an amount of at least about 0.05 wt %, and more preferably at least 0.08 wt %. The combined addition of Ti and V is advantageous because Ti provides a catalytic effect for V precipitation, making V precipitation more effective. In addition, V has a significant effect on the non-recrystallization temperature T nr There is no significant effect, so unlike Nb and Ti, they can affect T nr Since the reduced degree of anisotropy is beneficial to hole expansion capability, the combined addition of Ti and V can be beneficial.

[0067] Furthermore, it has been found that the formation of cementite can be suppressed and the favorable formation of a small amount of martensite plus retained austenite in the microstructure can be better controlled if the amounts of carbide-forming elements Ti, Nb, V and Mo expressed in weight % satisfy the following formula:

[0068]

[0069] where Ti_sol is defined as the amount of free Ti in the solution and is expressed as:

[0070]

[0071] Wherein the amount of Ti and N is expressed in weight %. Preferably, the lower limit of the formula is 2.0, more preferably 2.2, and most preferably 2.4. In one embodiment, the upper limit is preferably 5.0, more preferably 4.0, and most preferably 3.8, in order to further suppress the presence of cementite and / or control the amount of martensite.

[0072] In one embodiment, the steel strip has a composition consisting of, in wt %, C: 0.06-0.12 wt %, Mn: 1.20-2.20 wt %, Ti: 0.07-0.15 wt %, Si: up to 0.65 wt %, Al: up to 0.10 wt %, P: up to 0.02 wt %, S: up to 0.01 wt %, Nb: up to 0.03 wt %, N: up to 0.01 wt %, and optionally one or more elements selected from the following: (Cr: up to 0.25 wt %, Mo: up to 0.10 wt %, V: up to 0.30 wt %, Cu: up to 0.20 wt %, Ni: up to 0.50 wt %, B: up to 0.005 wt %), the balance being Fe and unavoidable impurities resulting from the ironmaking and steelmaking processes, and more preferred ranges are as described herein and claimed.

[0073] According to EN 10002-1 / 1506892-1, A50 specimen geometry was used, and the tensile test was performed parallel to the rolling direction. The tensile test was performed by quasi-static (strain rate 3×10 -4 s -1 ) Tensile test determines the 0.2% offset proof strength or yield strength (Rp0.2), ultimate tensile strength (Rm), uniform elongation (Ag) and tensile elongation (A50). The geometry of the tensile specimen consists in a gauge length of 50 mm in the rolling direction, a width of 12.5 mm and a thickness depending on the final gauge. The steel strength at 0.2% offset strain is measured as yield strength (Rp0.2 or YS).

[0074] The tensile flanging or hole expansion capacity (HEC) of steel strip is determined by hole expansion test. A specimen with a size of 90 mm × 90 mm × the final thickness of the strip is cut from the coiled steel. A hole with a diameter of 10 mm is punched in the middle of the specimen and the hole expansion test is carried out according to ISO / TS16630:2003 (E) standard. The hole expansion test of the sample is carried out with the burr removed from the top. A 60° conical punch is pushed from bottom to top to measure the hole diameter d when a through-thickness crack is formed. f The hole expansion ratio λ is calculated using the following formula, where d o =10mm:

[0075]

[0076] For all the above mechanical tests, at least three samples were tested for each condition and the average values ​​are reported herein.

[0077] The average total crack length (ATCL) is used to assess the susceptibility to crack formation in situations similar to those of industrial applications. The ATCL parameter is determined in ordinary laboratory cylindrical deep draw tests on steel billets measuring 90 mm × 90 mm, with the four corners of the square cut 10 mm in the direction of the two diagonals, as Figure 1A As shown. In the cylindrical deep drawing test, a punch, a punching die and a blank holder are used. In this process, the punch has a diameter of 50 mm and a punch radius of 7 mm. The die has an inner diameter of 62 mm and a radius of 8 mm. This setup Figure 1B The inner diameter is large enough to allow the edge of the formed cup to move freely. The gap, i.e. the distance between the punch wall and the die wall, is 6 mm. The blank holding force is set to 50 kN. During the cylindrical deep drawing test of the cup and the initial stage of deep drawing ( Figure 1C ), due to the high local compressive stress during punching, four regions at the edge of the blank are plastically deformed ( Figure 1D ). This results in localized wrinkling of the edge. At the end of the test and when the blank holding force is released, the four compression areas are subjected to reverse loading due to springback, as these areas of the formed cylindrical cup begin to lose contact with the blank holder. This reverse loading due to springback can lead to nucleation and growth of cracks in the four compression and wrinkle areas of the punched cup (see, for example, Figure 1E The cracks can pass through the entire thickness of the steel strip and be visible on both sides of the drawn cup, i.e., inside and outside, or only on one of the two sides of the cup. The four compression edges of the deep drawn cup (e.g., Figure 1DThe length of all visible cracks on the inside and outside of the cup wall (shown in ). The sum of the lengths of all observable cracks on the inside and outside of the cup wall was averaged on three of the four punched cups and recorded as the average total crack length (ATCL) in mm.

[0078] In addition, the work of fracture (WOF) is determined by compressing the formed cup between a cone and a flat tool, the normal of the plane being parallel to the axis of the cone. The cone has an angle of 90° and contacts the inner wall of the cup. As the flat tool moves toward the cone, it applies a normal force on the flat bottom of the cup. The cone applies a force on the cup wall, causing cracks to grow from the four compressed edge areas of the cup. The tool is placed in a conventional tensile testing machine. The normal force on the flat plate is recorded as a function of the plate displacement. The test is stopped a few millimeters before the flat bottom of the cup contacts the top of the cone. For the selected geometry, a displacement of up to 12 mm can be recorded. A graph of the area under the force versus displacement is calculated, up to 10 mm displacement. The calculated value in joules (kN mm or J) is an estimate of the material's resistance to crack extension and / or initiation, and the work of fracture (WOF) is defined here as the area under the force-displacement curve up to 10 mm displacement.

[0079] Evidence of cracking was determined by shearing the steel with a gap of about 12% and visually inspecting the shear edge quality. Cracks are the result of delamination during shearing, resulting in one or more cracks or fissures in the rupture zone of the shear edge. When inspecting the shear edge, these cracks manifest themselves as small fractures that extend parallel to the sheet surface (and perpendicular to the movement of the shearing device during shearing) and into the material, extending largely parallel to the deformation lines in the shear affected zone (SAZ) caused by shearing.

[0080] In Euler space The X-ray random intensity ratios of various texture components and fibers in the region or the so-called orientation distribution function (ODF) curves were measured as follows. In the range of 0°<α<90°, the ... a - radiation and primary graphite monochromator) to record three incomplete pole figures (110, 200, 211). Scanning 8mm in reflection mode 2region. The samples were ground / polished to medium thickness or 1 / 2T, where T corresponds to the sample thickness. The surface was etched in a solution of 83 ml H2O, 330 ml H2O2, 8 ml Ethyline Glycolmonobutylether and 578 ml H3PO4 for 15 minutes to remove any deformation during the grinding / polishing step prior to texture measurements. The C values ​​required to generate texture data (e.g., orientation distribution function or ODF for short) were obtained using BrukerTexEval software (version 2.4.0.4) using triclinic sample symmetry. ODFs were generated using orthogonal sample symmetry in conjunction with the Van Houtte routine and an in-house developed software routine (T2F, version 2006-36.1).

[0081] The microstructure of the steel strip is analyzed by a technique known in the art - Electron Backscatter Diffraction (EBSD), which in turn allows the area or volume fractions of the various components to be quantified. The EBSD measurements are carried out on a cross section parallel to the rolling direction (RD-ND plane), mounted in a conductive resin and mechanically polished to 1 μm. In order to obtain a completely deformation-free surface, a final polishing step is carried out with colloidal silica (OPS).

[0082] The scanning electron microscope (SEM) used for EBSD measurements was a Zeiss Ultra 55 machine equipped with a field emission gun (FEG-SEM) and an EDAXPEGASUS XM 4HIKARI EBSD system. EBSD scans were collected on the RD-ND plane of the sheet at 1 / 4 thickness. The sample was placed in the SEM at a 70° angle. The accelerating voltage was 15 kV, and the high current option was turned on. A 120 μm aperture was used, and the typical working distance during scanning was 17 mm. In order to compensate for the high tilt angle of the sample, dynamic focus correction was used during scanning.

[0083] EBSD scans were collected using TexSEM Lab (TSL) software: "Oriented Imaging Microscope (OIM) Data Collection Version 7.2". In general, the following data collection settings were used: Hikari camera (standard mode) with 5 × 5 pixel binning combined with background subtraction. In all cases, the scan area was located at 1 / 4 thickness of the sample and care was taken to avoid including non-metallic inclusions in the scan area as much as possible. The EBSD scan size was 100 × 100 μm in all cases with a step size of 0.1 μm and a scan rate of approximately 100 frames per second. Fe(α) and Fe(γ) were used to mark the Kikuchi patterns. The Hough settings used during data collection were: pixel binning pattern size of approximately 96; theta group size of 1; rho fraction of approximately 90; maximum peak count of 10; minimum peak count of 5; Hough type set to classic; Hough resolution set to low; butterfly convolution mask of 9 × 9; peak symmetry of 0.5; minimum peak amplitude of 10; and maximum peak distance of 20.

[0084] EBSD scans were evaluated using TSL OIM Analysis Software Version "8.0×64 [12-14-16]". Typically, the data set was rotated 90° in the RD axis to orient the scan appropriately relative to the measurement direction. Standard grain expansion cleaning was performed (grain tolerance angle (GTA) of 5°, minimum grain size of 5 pixels, and the criterion used was that the grains must contain multiple rows for a single expansion iteration to clean). Next, pseudo-symmetric cleaning was performed (GTA 5, axis ang 30°@111).

[0085] From the 100×100mm area (0.01mm 2 ) and measured the number average grain size (GS) based on a grain tolerance angle of 15° N15 For these scans, the sum of the total grain boundary lengths (ΣGB) of the grain boundaries with misorientation angles from 15° to 65° was also measured (ΣGB 15-65 ) (in mm). ∑GB 15-65 The value is in mm -1 (or 1 / mm) and is a measure of the density of high-angle grain boundaries.

[0086] The misorientation angle distribution (MOD) index of the Fe(α) fraction was calculated using the following method: The normalized misorientation angle distribution (MOD) was calculated from a separate EBSD data set using the TSL OIM analysis software, including all boundaries, with misorientation angles ranging from 5° to 65° and a pixel binning of 1°. Similarly, the normalized theoretical MOD of randomly recrystallized polygonal ferrite was calculated using the same misorientation angle range and pixel binning as the measured curve. In practice, this is the so-called MOD based on "MacKenzie" included in the TSL OIM analysis software. Normalization of the MOD means that the area under the MOD is defined as 1. Then, the MOD index is defined as Figure 2A and Figure 2B The area between the theoretical curve (dashed line) and the measured curve (solid line) in , and can be written as:

[0087]

[0088] Among them, M MOD，i is the intensity of the measured MOD at angle i (ranging from 5° to 65°), R MOD，i is the theoretical strength at angle i or the "MacKenzie" based on the MOD of randomly recrystallized polygonal ferrite.

[0089] Figure 2A and 2B The solid line in FIG. 5 represents the measured MOD, and the dashed line represents the curve of the theoretical misorientation angle of a randomly recrystallized polygonal ferrite structure. Figure 2A The MOD curve for an exemplary sample having a microstructure that is primarily polygonal ferrite in character is shown. Figure 2B The MOD curves for exemplary samples with a microstructure that is predominantly acicular / bainitic or bainitic in character are shown. The MOD index range is defined as from 0 to almost 2; when the measured curve is equal to the theoretical curve, the area between the two curves is 0 (the MOD index will be 0), while if there is (almost) no intensity overlap between the two distribution curves, the MOD index is (almost) 2. Thus, if Figure 2A and Figure 2BAs shown, MOD contains information about the properties of the microstructure, and the MOD index can be used to evaluate the characteristics of the microstructure on a quantitative basis, and is therefore a clearer method than based on conventional methods such as light-optical microscopy. A completely polygonal ferrite microstructure will have a unimodal MOD with most of its intensity in the range of 20° to 50°, and a peak intensity around 45°. In contrast, a completely acicular / bainitic ferrite or bainite microstructure will have a strong bimodal MOD with peak intensities between 5° to 10° and 50° to 60°, and very little intensity in the range of 20° to 50°. Therefore, in this example, a low MOD index and a high 20° to 50° MOD intensity are clear signs of a predominantly polygonal ferrite microstructure, while a high MOD index and a low 20° to 50° MOD intensity are clear signs of a predominantly acicular / bainitic ferrite or bainite microstructure. The inventors have found that a MOD index greater than 0.65 corresponds to an acicular / bainitic ferrite or bainite microstructure, with alloy compositions and process settings as used in the examples to describe the present invention.

[0090] In one aspect of the invention, it relates to a method of manufacturing a hot rolled high strength steel strip as described and claimed herein, the method comprising the following steps, in this order,

[0091] - casting of a slab followed by a step of reheating the solidified slab to a temperature of 1050-1260° C., preferably keeping the slab at a temperature of 1050-1260° C. for a period of about 20-60 minutes, and then hot rolling said slab, or casting of a slab or strip followed by a step of hot rolling said slab or strip;

[0092] - hot rolling the steel slab or strip and completing the hot rolling at a finishing temperature (FRT) between 960°C and Ar3+20°C, preferably between 960°C and 900°C, most preferably between 960°C and 920°C, wherein Ar3 is the temperature at which the transformation of austenite to ferrite begins during cooling. As known in the art, the Ar3 temperature can be calculated according to the following equation:

[0093] Ar3=910℃-203x[C] 1 / 2 +44.7x[Si]+400[Al]+700[P]-30x[Mn]+31.5x[Mo]-11[Cr]+400x[Ti]+104x[V];

[0094] The equation shows that carbon inhibits the phase transformation of Ar3 austenite to ferrite, which will promote grain refinement. In addition, it is also shown that Mn inhibits the Ar3 transformation temperature.

[0095] - accelerated cooling of the hot rolled steel strip to an intermediate temperature (IT) on the run-out roller between 50 and 150°C / s, preferably between 65 and 100°C / s, most preferably between 70 and 80°C / s, to between 500 and 600°C, preferably between 520 and 590°C, and more preferably between 540 and 580°C ROT );

[0096] - further accelerated cooling of the hot rolled steel strip to between 450-580°C at a run-out roller secondary cooling rate (CR2) between 1 and 30°C / s, or preferably between 5 and 25°C / s,

[0097] Coiling temperature preferably between 475-560°C;

[0098] - coiling the hot rolled and cooled steel strip at a coiling temperature (CT) between 450-580° C., preferably between about 475-560° C.;

[0099] - further cooling the coiled hot rolled steel strip to ambient temperature;

[0100] - optionally pickling the hot rolled steel strip; and

[0101] - Optionally providing the hot rolled steel strip with a metal coating, preferably selected from: a Zn layer, a Zn-based alloy layer, a Mg-based alloy layer or an Al-based alloy layer, to provide improved corrosion resistance in use. The metal coating is preferably applied by hot coating or hot dip coating, and may also be applied by electrolytic deposition.

[0102] The manufacturing method described and claimed herein results in a desired microstructure that provides a targeted improved balance of high edge ductility as indicated by hole expansion capability with adequate toughness and good shear properties with substantially no cracking for hot rolled deep-drawable sheet strip having a yield strength in the range of 600 to 820 MPa, a tensile strength in the range of 730 to 950 MPa, and an A50 tensile elongation of 11% or more. The invention is also embodied in a steel strip manufactured by the method described and claimed herein, having the described composition, microstructure, and improved balance of toughness, shear properties, and mechanical properties.

[0103] The present invention is not limited by the casting method. The steel can be cast in a direct strip plant into conventional thick slabs with a casting thickness of 150-350 mm, typically 225-250 mm, and thin slabs with a casting thickness of 50-150 mm. For conventional thick slab casting, reheating of the slab is necessary to reheat the slab from ambient temperature (usually the thick slab has been cooled from the casting temperature to ambient temperature in the slab yard) and to homogenize the slab in terms of composition, and therefore the reheating temperature should be above about 1050°C to also dissolve any precipitates when microalloying elements are present, and to bring the slab to a temperature at which the final hot rolling in the finishing mill can still be carried out at above 960°C. Typically, this requires a (slab) reheating temperature between about 1050-1260°C. For thin slab casting, the cast slab is subjected to a homogenization treatment in a homogenization furnace immediately after casting the thin slab, wherein the homogenization temperature should be above about 1050° C., and typically about 1100-1160° C. This will also prevent the formation of any precipitates when microalloying elements (if any) are present, and also bring the thin slab to a temperature where the final hot rolling in the finishing mill can still be carried out at FRT>Ar3+20° C. According to the present invention, the reheating or homogenization time of the thin slab continuous casting route is preferably 20 minutes or more.

[0104] Hot rolling of steel must be carried out in the austenite phase to control the final microstructure. On an industrial scale of rolling, the FRT should be kept above the Ar3 temperature. In a preferred embodiment, it is above Ar3+20°C, for example, typically above about 890°C, to avoid local hot rolling below Ar3 at the cooler edge or tail of the strip. However, the FRT should not be too high in the austenite region, as this will lead to insufficient grain refinement (too high GS N15 ) and too few high-angle grain boundaries (too low ∑GB 15-65), which can ultimately lead to too low a fracture toughness and, therefore, too high an average total crack length and too low a fracture work. In addition, too high a FRT will promote a texture that is unfavorable with regard to fracture toughness, embrittlement and delamination (e.g., {001} on the θ-fibers). <110> Rotating Cube and {110} <001> Goss), and subsequently leads to increased edge crack sensitivity and increased cracking sensitivity in shear. Therefore, the FRT should not exceed 960°C. A relatively low FRT is beneficial for promoting sufficient grain refinement and promoting texture components that are beneficial to fracture toughness, such as {554} <225> , especially {332} <113> On the other hand, too low FRT, in addition to avoiding local hot rolling below Ar3 at the cooler edge or strip tail, can also stimulate excessive α*-fractions of α-fibers and {112} <332> The intensity of the texture component is related to the increased anisotropy (especially the {113} <110> and {112} <110> ). Too high anisotropy increases the plane anisotropy (increased variation of the r value), which reduces the punchability of the steel and impairs its hole expansion ability. Therefore, FRT should not be too low and should be higher than (Ar3+20°C), preferably higher than 900°C, and more preferably higher than 920°C. The considerations given above illustrate the relevance and complexity of FRT in controlling the final microstructure and its texture to obtain a proper balance between fracture toughness and hole expansion ability and reduce the risk of cracking during shear.

[0105] After hot rolling, the steel strip is cooled on the run-out roller table (ROT) at an accelerated rate to an intermediate temperature (IT) between 500-600°C, preferably between 520-590°C, and most preferably between 540-580°C on the ROT. ROT ), while the primary cooling rate (CR1) of the run-out roller is between 50-150°C / s, preferably between 65-100°C / s, and more preferably between 70-80°C / s. This can be achieved by air cooling, turbulent cooling and water jet cooling, depending on the thickness of the steel strip. It is desirable to accelerate the cooling rate to suppress the recovery and loss of internal stored energy in austenite to promote grain refinement (small GS) of the final microstructure. N15 ) and the increased ∑GB 15-65 The cooling rate should be high enough to avoid the austenite to ferrite phase transformation at high temperature (which may lead to too low MOD index), and preferably promote the austenite to ferrite phase transformation at a relatively low temperature of about 560-630°C on the run-out table, thereby promoting an MOD of ≥ 0.85. The increased cooling rate will promote grain refinement (reduced GS N15 ), improved ∑GB 15-65, thus increasing fracture toughness and reducing crack sensitivity. The increased cooling rate will also suppress texture randomization and thus inhibit the development of those textures from deformed austenite that promote toughness (e.g. {332} <113> ) and suppress those associated with embrittlement and delamination (e.g. {001} <110> Rotating Cube). Unnecessarily high CR1 may affect the flatness of the strip after cooling and may lead to control problems to stop at the correct cooling stop temperature, so a suitable maximum CR1 is about 150°C / s, or preferably about 100°C / s, or most preferably about 80°C / s. For practical reasons, the run-out roller cooling rate (ROT-CR) is defined as the average cooling rate of the steel strip surface.

[0106] Next, at a temperature between 500-600°C, preferably between 520-590°C, more preferably between 540-580°C (IT ROT ) to a coiling temperature between 450-580°C, or preferably between 475-560°C, with a run-out roller secondary cooling rate (CR2) of about 1-30°C / s. In a preferred embodiment, CR2 is about 5-25°C / s. In addition, the hot rolled steel strip is coiled at a temperature between 450-580°C, preferably between 475-560°C. In addition to the finishing temperature, the coiling temperature of the steel strip is also a key process parameter to obtain the desired hot rolled steel strip microstructure, which provides an improved balance of properties as described herein.

[0107] When the coiling temperature is too low, the precipitation kinetics are insufficient and thus a low strength level will be obtained. When the coiling temperature is too high, the grain refinement is insufficient, resulting in reduced fracture toughness and increased edge crack sensitivity. Too high a coiling temperature will also reduce the MOD index, which can impair hole expansion capability and promote cracking during shearing. During coil cooling, further precipitation may occur, as well as some further phase transformations. Unexpectedly, once formed, the precipitates may coarsen during coil cooling. The alloy composition, together with the required coiling temperature range, suppresses this phenomenon. The coiling temperature will help promote small grains of ferrite formed during cooling on the run-out roller, coiling or coil cooling, and also suppress the coarsening of precipitates that strengthen the ferrite matrix, which can be beneficial to improving shear edge quality and suppressing delamination or cracking during shearing.

[0108] In a preferred embodiment, the fraction transformed from austenite to acicular / bainitic ferrite and / or bainite (CTzm) at the end of the run-out table and before coiling is at least 48%, preferably at least 50%, most preferably at least 52%. The increased transformation fraction from austenite to acicular / bainitic ferrite and / or bainite is attributed to a faster phase transformation rate on the run-out table before coiling and results in increased grain refinement, which is beneficial to improved fracture toughness in terms of reduced average total crack length (ATCL) and increased work of fracture (WOF).

[0109] Patent document EP1616970-A1 discloses that a holding temperature below 680°C results in insufficient driving force for ferrite transformation, and subsequently the fraction of ferrite containing precipitates is too low. For the present invention, the phase transformation from austenite to ferrite is carried out in a temperature range below 680°C in order to achieve an increase in grain refinement and ∑GB 15-65 The increase in the steel strip is used to improve the fracture toughness and reduce the crack sensitivity while still having sufficient precipitation kinetics. In addition, patent document EP1616970-A1 discloses that after the steel strip is kept at a temperature above 680°C for more than 1 second, in order to form a secondary phase of bainite and / or martensite and suppress the formation of other phases to less than 5 volume %, it is necessary to perform secondary cooling at an average cooling rate of more than 30°C / s, preferably more than 50°C / s, to a coiling temperature below 550°C, preferably below 450°C, and more preferably below 350°C, and coiling. For the present invention, the coiling temperature is significantly higher, with a value between 450-580°C, preferably between 475-560°C, so as to allow the phase transformation from austenite to ferrite to continue at a relatively low temperature, thereby promoting fine-grained acicular / bainitic ferrite and / or bainite, which is precipitation strengthened with carbide precipitates containing Ti and optionally Mo, Nb and / or V. Grain refinement and enlarged ∑GB 15-65 Provides improved fracture toughness and reduced (edge) crack sensitivity. Coiling below 450°C will result in insufficient acicular / bainitic ferrite and / or bainite formation and loss of precipitation strengthening. Furthermore, it may lead to too high a martensite fraction.

[0110] Although patent document EP1338665-A1 discloses coiling temperatures in the range of 550-700° C., steels produced have a tensile strength of at least 950 MPa and a hole expansion capacity of at least 40%, all of which are produced at coiling temperatures exceeding 600° C. The finishing temperature of all these steels is in the range of 880-930° C. According to the present invention, it has been found that excessively high coiling temperatures lead in particular to increased crack sensitivity, in particular edge crack sensitivity, and also to reduced fracture toughness.

[0111] After the steel strip has cooled to room temperature, the oxides (scale) on the hot rolled steel strip are removed by pickling in an acid solution (e.g. HCl) at warm temperature (about 80-120°C) or by a combination of pickling and mechanical brushing of the strip surface. This step is necessary to make the steel strip surface suitable for direct use as uncoated hot rolled steel or to make it suitable for coating processes when corrosion resistance is optionally required.

[0112] In one embodiment, the thickness of the high strength hot rolled steel strip is in the range of about 1.5 to 8 mm, more preferably about 1.8 to 6 mm, and most preferably about 1.8 to 4.5 mm.

[0113] An important aspect of the present invention is that the hot rolled steel strip is not subsequently subjected to a cold rolling operation with a thickness reduction greater than 1.6%.

[0114] The high strength hot rolled steel strip product may be a bare or uncoated product, or it may have a thin metal coating on one or both of its major surfaces, typically up to about 100 g / m per side of the steel strip. 2 , preferably up to about 50 g / m per side 2 The metal coating is preferably selected from an aluminum alloy coating (such as an Al-Si alloy or an Al-Zn alloy), a zinc coating and a zinc alloy coating (such as a Zn-Al alloy, a Zn-Mg alloy, a Zn-Fe alloy, a Zn-Al-Mg alloy or a Zn-Mg-Al alloy).

[0115] There is no restriction on the composition of the zinc or zinc alloy coating. Although the coating can be applied in various ways, it is preferably hot-dip galvanizing using a standard Gl coating bath. The Zn-based coating may comprise a Zn alloy containing Al as an alloying element. The preferred zinc bath composition contains about 0.10-0.35 wt. % Al, the balance being zinc and unavoidable impurities.

[0116] Other zinc coatings may also be applied. One example includes a zinc alloy coating according to patent document WO2008 / 102009-A1 (which is incorporated herein by reference), in particular a zinc alloy coating consisting of about 0.3-4.0 wt. % Mg and about 0.05-6.0 wt. % Al, preferably about 0.1-5.0 wt. % Al, and optionally up to about 0.2 wt. % of one or more additional elements and unavoidable impurities, and the balance is zinc. A preferred Zn bath containing Mg and Al as the main alloying elements has a composition: about 0.5-3.8 wt. % Al, about 0.5-3.0 wt. % Mg, optionally up to 0.2 wt. % of one or more additional elements; the balance is zinc and unavoidable impurities. Additional elements, which are usually added in small amounts of less than 0.2 wt. %, may be selected from Pb, Sb, Ti, Ca, Mn, Sn, La, Ce, Cr, Ni, Zr and Bi. Pb, Sn, Bi and Sb are usually added to form spangles. Preferably, the total amount of additional elements in the zinc alloy is at most 0.2 wt.%, and more preferably at most 0.1 wt.%. For typical applications, these small amounts of additional elements do not change the properties of the coating or bath to a significant extent. Preferably, when one or more additional elements are present in the coating, each is present in an amount of at most 0.02 wt.%, preferably each is present in an amount of at most 0.01 wt.%. The additional elements are usually added only to prevent the formation of dross in a bath with molten zinc alloy for hot dip galvanizing, or the formation of spangles in the coating.

[0117] In another embodiment, the metal coating comprises a layer of (industrially pure) aluminum or an aluminum alloy. A typical metal bath for hot-dip coating of such an aluminum layer comprises aluminum alloyed with silicon, for example aluminum alloyed with about 8-11% by weight of silicon and up to about 4% by weight of iron, optionally up to 0.2% by weight of one or more additional elements such as calcium, unavoidable impurities, the remainder being aluminum. Silicon is present to prevent the formation of a thick iron-metallic compound layer that reduces adhesion and formability. Iron is preferably present in an amount of about 1-4% by weight, more preferably at least about 2% by weight.

[0118] In one aspect of the present invention, it relates to a galvanized steel strip obtained by hot-dip galvanizing the high-strength hot-rolled steel strip according to the present invention.

[0119] In another aspect of the present invention, it relates to an automotive component, in particular an automotive chassis part, which is made of or combined with a high-strength hot-rolled steel strip according to the present invention, and in particular benefits from an improved balance of strength, formability and improved fracture toughness, high edge ductility and reduced (edge) crack sensitivity. The steel strip can be formed into automotive components in cold forming operations, warm forming or hot forming operations known in the art. Automotive components include, but are not limited to, suspension arms, reinforcement members, body-in-white frame members, as side members, seat frames, seat rails, bumper beams, battery boxes for electric vehicles, all of which have complex shapes. By using hot-rolled high-strength steel strips, these components can be manufactured in high quality, cost-effectively and at high yields. The high-strength hot-rolled steel sheet products of the present invention can also be used for engineering purposes.

[0120] In another aspect of the present invention, it relates to an automobile chassis part using the hot-rolled high-strength steel strip according to the present invention and / or produced according to the present invention, and wherein the hot-rolled high-strength steel strip used for manufacturing the automobile chassis part:

[0121] - having an ultimate tensile strength of 730-950 MPa, a hole expansion ratio of at least 45%, preferably at least 50%, most preferably at least 60%, and a work of fracture (WOF) of at least 100 J, preferably at least 125 J, more preferably at least 200 J, wherein the ultimate tensile strength (Rm), the total tensile elongation (A50) and the sheet thickness t (mm) satisfy the formula (Rm×A50) / t 0.2 ≥8,000; or

[0122] - having an ultimate tensile strength of 730-950 MPa, a hole expansion ratio of at least 45%, preferably at least 50%, most preferably at least 60%, and a work of fracture (WOF) of at least 100 J, preferably at least 125 J, more preferably at least 200 J, wherein the ultimate tensile strength (Rm), the total tensile elongation (A50) and the sheet thickness t (mm) satisfy the formula (Rm×A50) / t 0.2 ≥9,000; or

[0123] - having an ultimate tensile strength of 730-950 MPa, a hole expansion rate of at least 45%, preferably at least 50%, most preferably at least 60%, a work of fracture (WOF) of at least 100 J, preferably at least 125 J, more preferably at least 200 J, and wherein the ultimate tensile strength (Rm), the total tensile elongation (A50) and the sheet thickness t (mm) satisfy the formula (Rm×A50) / t 0.2 ≥10,000; or

[0124] - having a yield strength (Rp0.2) of 600-820 MPa, a yield ratio of 0.8-0.9, and a hole expansion rate of at least 45%, preferably at least 50%, more preferably at least 60%, a work of fracture (WOF) of at least 100 J, preferably at least 125 J, more preferably at least 200 J, and wherein the ultimate tensile strength (Rm), the total elongation (A50) and the sheet thickness t (mm) satisfy the formula (Rm×A50) / t 0.2 ≥10,000. BRIEF DESCRIPTION OF THE DRAWINGS

[0125] The invention will also be described with reference to the accompanying non-limiting drawings, in which:

[0126] Figures 1A to 1E Several features of a method for determining the average total crack length (ATCL) and work of fracture (WOF) of a steel strip product using a common cylindrical deep draw test for a draw cup are shown.

[0127] Figure 2A and Figure 2B An example of a misorientation distribution (MOD) curve is shown to explain the physical meaning and definition of the MOD index to characterize the matrix of the microstructure in a quantitative manner.

[0128] Figure 3 Schematic diagram showing the Euler space or orientation distribution function (ODF) Cross-sections, characterized by relevant texture fibers, sections and components, to illustrate the relationship between texture and hole expansion capability and the occurrence of cracking that the inventors have discovered.

[0129] Figure 4 The relationship of work of fracture (WOF) versus hole expansion capacity (HEC) is shown for Inventive Examples 1A-4A and 6A and Comparative Examples 5A and 7A-13A.

[0130] Figure 5 The work of fracture (WOF) versus the density of high-angle grain boundaries (ΣGBL) of inventive examples 1A–4A and 6A and comparative examples 5A and 7A–13A are shown. 15-65 ) relationship.

[0131] Figure 6 The work of fracture (WOF) of inventive examples 1A-4A and 6A and comparative examples 5A and 7A-13A is shown relative to the number average grain size (GS) based on a grain tolerance angle of 15°. N15 ) relationship.

[0132] Figure 7 The average total crack length (ATCL) versus the high-angle grain boundary density (ΣGBL) of inventive examples 1A–4A and 6A and comparative examples 5A and 7A–13A is shown. 15-65 ) relationship.

[0133] Figure 8 The average total crack length (ATCL) of inventive examples 1A-4A and 6A and comparative examples 5A and 7A-13A is shown relative to the number average grain size (GS) based on a grain tolerance angle of 15°. N15 ) relationship.

[0134] Fig. 9 The hole expansion capacity (HEC) versus T factor is shown for Inventive Examples 1A-4A and 6A without cracking in shear and Comparative Examples 5A, 7B, 8B, 10B, and 11B, and Comparative Examples 9B, 12B, and 13B with cracking in shear.

[0135] Fig.10 The high-angle grain boundary density (ΣGBL) of Inventive Examples 1A-4A and 6A without cracking during shearing and Comparative Examples 5A, 7B, 8B, 10B and 11B and Comparative Examples 9B, 12B and 13B with cracking during shearing is shown. 15-65 ) relative to the T factor.

[0136] Fig.11 The number average grain size (based on a grain tolerance angle (GS) of 15°) of Inventive Examples 1A-4A and 6A and Comparative Examples 5A, 7B, 8B, 10B and 11B without cracking during shearing and Comparative Examples 9B, 12B and 13B with cracking during shearing is shown. N15 ))relative to the T factor. DETAILED DESCRIPTION

[0137] The invention will now be illustrated with reference to non-limiting comparative examples and examples of the invention.

[0138] Example

[0139] Example 1

[0140] Three steel slabs of chemical compositions A, B and C were produced industrially by basic oxygen steelmaking and continuous casting, where A and B are according to the present invention and C is a comparative example. The chemical compositions of these steels are listed in Table 1. For each of steels AC, there is no intentional addition of Cr, Mo, V and Nb. α is also given in Table 1, where α is calculated using the following formula:

[0141]

[0142] in

[0143] Table 1. Chemical composition of steels in wt%. The balance is Fe and unavoidable impurities in the iron and steel making process. Alloy A is an example of the present invention, and alloys B and C are comparative examples.

[0144]

[0145] The slabs have been processed on an industrial scale and the relevant processing parameters used are given in Table 2. The fraction transformed from austenite to acicular / bainitic ferrite and / or bainite (CTzm) is measured at the end of the run-out table and before coiling by real-time, on-line measurement using a detector that generates a primary magnetic field that interacts with the passing hot rolled steel strip, resulting in a secondary magnetic field. This magnetic response of the hot rolled steel strip and the essential difference in magnetic permeability between austenite and ferrite are the essential principles for monitoring the transformation of iron from austenite to ferrite as the steel cools at a specific position on the run-out table. The fraction transformed from austenite to acicular / bainitic ferrite and / or bainite (CTzm) at the end of the run-out table and before coiling is given in Table 2.

[0146] The mechanical properties (Rp, Rm, yield ratio and A50), hole expansion capacity (HEC), average total crack length (ATCL), work of fracture (WOF) and cracking (crack) evaluation according to the methods described in this article are listed in Table 2 as a function of the applied processing parameters.

[0147] According to the methods described herein, the microstructural characteristics of the hot rolled steel were determined by EBSD and the X-ray random intensity ratios of specific texture components, fibers and sections were measured by XRD. The microstructural characteristics and X-ray random intensity ratios are listed in Table 3. Table 3 also shows the T factor, which is defined as the MOD index multiplied to 1.2 times the power of the γ-fiber, θ-fiber and {110} <001> The sum of the X-ray random intensity ratios of the Goss texture components divided by the α*-segment, {112} <131> , {332} <113> and {554} <225> The sum of the X-ray random intensity ratios of the texture components:

[0148]

[0149] Here, the α*-part is defined as the Euler space or orientation distribution function (ODF) from {115} <110> To {223} <110> Extended The portion of the α-fiber in the region (and including the X-ray random intensity ratio of these two external texture components) covers the range of 15° to 45° on the α-fiber. Figure 3 The relevant textured fibers, parts and components are shown in FIG. 1 , which shows a schematic representation of the Euler space or orientation distribution function (ODF) part.

[0150] From the results of Tables 2 and 3, it can be seen that in order to produce a steel having a good balance between hole expansion capability and fracture toughness in terms of average total crack length (ATCL) and work of fracture (WOF) on the one hand, and without cracking in shear on the other hand, the process must be controlled in terms of finishing temperature, hot runout roller cooling trajectory and coiling temperature to promote a fine grained / acicular bainitic ferrite and / or bainite microstructure with a sufficiently high misorientation distribution (MOD) index and a sufficiently high angle grain boundary density and a suitable texture. The inventors have found that the use of an appropriate combination of alloy composition and hot rolling method with the following:

[0151] The finishing temperature (FRT) is at least 900°C and at most 960°C,

[0152] Output roller table (IT ROT ) is at least 500° C. and at most 600° C.,

[0153] The coiling temperature (CT) is at least 475°C and at most 580°C, resulting in a fraction (CTzm) of at least 48% transformed from austenite to acicular / bainitic ferrite and / or bainite at the end of the run-out table and before coiling. According to inventive examples 1A-4A and 6A, a steel exhibiting an acicular / bainitic ferrite and / or bainite microstructure can be obtained, having:

[0154] A MOD index of at least 0.85,

[0155] · Number average grain size (GS) of up to 2.0 μm based on a 15° grain tolerance angle N15 ),

[0156] At least 1060mm -1 The high-angle grain boundary density (ΣGB 15-65 ),

[0157] · a gamma-fiber X-ray random intensity ratio of at least 27, and

[0158] A T factor of at least 1.1 and at most 1.8,

[0159] It produces a high strength steel with a suitable hole expansion capability of at least 45% and sufficient fracture toughness indicated by an average total crack length (ATCL) of at most 20 mm, a work of fracture (WOF) of at least 100 J and no cracking in shear. Figure 4 Shown is a graph of work of fracture (WOF) versus hole expansion capacity (HEC) for inventive examples (no cracking) and comparative examples with and without cracking.

[0160] Figures 5 to 8 It is shown that for all the examples listed in Table 2 (inventive examples and comparative examples), relative to ΣGBL 15-65 and GSN15 The data show a strong correlation between the work of fracture (WOF) and the average total crack length (ATCL) and the average grain size and density of high-angle grain boundaries. The inventors found that the work of fracture (WOF) is a function of ΣGBL. 15-65 and GS N15 The empirical relationship of the function is:

[0161] WOF=3×10 6 ×(∑G(L 15-16 ) 2.47

[0162] as well as:

[0163] WOF=1619.3×(GS N15 ) -4.38

[0164] The inventors found that according to Fig. 9 From the experience of the linear fit of the data shown, the hole expansion capacity is mainly controlled by the texture and the MOD index according to the following equation, in which the hole expansion capacity is plotted against the T factor:

[0165] HEC=28.13×T+15.78

[0166] in

[0167]

[0168] This relationship shows that in order to have a hole expansion capacity of at least 45%, the T factor should be at least 1.1, such as Fig. 9 As shown in FIG. 1 , it is preferably at least 1.15, and more preferably at least 1.4. However, too high a T factor may promote cracking during shearing, as shown in FIG. Fig. 9 Therefore, when GSN15 is too high and ∑GBL 15-65When it is too low, the T factor should be at most 1.8. The high X-ray random intensity ratio of γ-fibers is beneficial to promote the deep drawability of steel, and according to the above description of the T factor, it is also beneficial to the hole expansion ability. The inventors found that the X-ray random intensity ratio of γ-fibers should be at least 27. Although it is beneficial to high hole expansion ability, too high a T factor can lead to cracking during shearing because the inclusion of {001} <110> Rotating Cube Texture Components and {110} <001> The X-ray random intensity ratio of the θ-fibers of the Goss texture components (all related to delamination and embrittlement) becomes too high, and the α*-fraction and {112} <131> The texture components (both associated with increased anisotropy) and {332} <113> and {554} <225> The X-ray random intensity ratio (both associated with increased toughness) becomes too low. Cracking in shear is undesirable because it reduces shear edge fatigue and may impair the serviceability of the structural component.

[0169] Fig.10 and 11 The high-angle grain boundary density (ΣGBL) of Inventive Examples 1A-4A and 6A that did not crack during shearing and Comparative Examples 5A, 7B, 8B, 10B, and 11B and Comparative Examples 9B, 12B, and 13B that cracked during shearing are shown, respectively. 15-65 ) and the number average grain size (GS) based on a grain tolerance angle of 15° N15 ) versus T factor. These figures show the relationship between the grain refinement (GS N15 ≥2.0μm) and promote sufficient high-angle grain boundary density (ΣGBL 15-65 ≥1060mm -1 ), on the one hand, for sufficient fracture toughness in terms of average total crack length (ATCL ≤ 10 mm) and work of fracture (WOF ≥ 100 J), and obtaining an acicular / bainitic ferrite and / or bainite microstructure with a sufficient MOD index (≥ 0.85) and X-ray random intensity ratio of specific texture fibers, parts and components, as shown in Table 3, resulting in a sufficient T factor (at least 1.1, preferably not more than 1.8) to obtain a sufficiently high hole expansion capability (HEC ≥ 45%) without the risk of cracking in shear.

[0170] Examples 5A and 7A-13A are comparative examples because the work of fracture (WOF) is lower than 100 J due to having an average grain size (GS) higher than 2.0 μm based on a 15° grain tolerance angle. N15 ) and less than 1060mm -1 The high-angle grain boundary density (ΣGB 15-65). For all these comparative examples, the fraction (CTzm) transformed from austenite to acicular / bainitic ferrite and / or bainite at the end of the run-out table and before coiling is much lower than for the inventive examples. For comparative examples 5A and 7A-13A, the transformed fraction (CTzm) is less than 48%, in most cases even below 25%, while for inventive examples 1A–4A and 6A this is above 50%. Due to an unfavorable combination of process and alloy composition, the lower transformation fractions of comparative examples 5A and 7A-13A are closely related to the slower transformation rates (compared to the inventive examples) and the subsequent reduced grain refinement of these comparative examples, resulting in a significantly higher crack sensitivity of the comparative examples in terms of average total crack length (ATCL), and above all a significantly lower work of fracture (WOF < 100 J). Comparative example 5A has a too low hole expansion capacity (HEC < 45%), because IT ROT Too high (IT ROT >600℃), resulting in a final microstructure with too low MOD index (<0.85), too low X-ray random intensity ratio of γ-fiber (<27), and too low T factor (<1.1). Comparative Examples 7B, 8B, 10B, and 11B have too low hole expansion ability (HEC<45%) due to too low T factor (<1.1), and too high number average grain size (GS) based on grain tolerance angle of 15°. N15 >2.0μm) and the density of large-angle grain boundaries is too low (ΣGBL 15-65 <1060mm -1 ) and have too low a work of fracture (WOF < 100 J). Comparative Examples 9B, 12B and 13B also have too low a hole expansion capacity (HEC < 45%) and too low a work of fracture (WOF < 100 J), and in addition show cracking in shear (cracking: yes). These comparative examples also have too high a number average grain size (GS based on a grain tolerance angle of 15°) N15 >2.0μm) and too low high-angle grain boundary density (ΣGBL 15-65 <1060mm -1 ), resulting in poor fracture toughness based on too high average total crack length (ATCL>10 mm) and too low fracture work (WOF<100 J). In addition, their T factor is too high (>1.8), resulting in high hole expansion ability (HEC≥45%), but unfortunately cracking in the fracture zone of the shear edge during shearing.

[0171]

[0172]

[0173] Example 2

[0174] Steel slabs of chemical composition D according to the invention have been processed on an industrial scale, the composition of which is given in Table 4 and the relevant processing parameters used are given in Table 5. Nb was added purposefully to the steel composition, while Cr, Mo and V were not added purposefully. In Table 4, α was also calculated using the formula given in Example 1.

[0175] The properties were measured and the results are also listed in Table 5, and the microstructural and textural characteristics are given in Table 6; all in the same manner as described in Example 1.

[0176] This example illustrates that small additions of Nb aid grain refinement in steel when processed according to the invention and provide the desired microstructural features to provide a favourable balance of properties, in particular high strength with suitable hole expansion capability and sufficient fracture toughness, work of fracture and lack of cracking in shear.

[0177] Table 4. Chemical composition of Steel D, in wt %. The balance is Fe and unavoidable impurities in the iron and steel making process. Alloy D is an example of the present invention.

[0178]

[0179] Having now fully described the invention, it will be apparent to those skilled in the art that numerous changes and modifications can be made without departing from the spirit or scope of the invention as described herein.

[0180]

Claims

1. A hot-rolled high-strength steel strip having the following composition, in weight %, C: 0.06-0.12% by weight Mn: 1.20-2.20% by weight Ti: 0.07-0.15 wt%, Si: up to 0.65 wt %, Al: up to 0.10 wt %, P: up to 0.02% by weight S: up to 0.01% by weight Nb: up to 0.03 wt %, N: up to 0.01% by weight Optionally, one or more elements selected from the group consisting of: (Cr: up to 0.25 wt %, Mo: up to 0.10 wt %, V: up to 0.30 wt %, Cu: up to 0.20 wt %, Ni: up to 0.50 wt %, B: up to 0.005 wt %), The balance is Fe and unavoidable impurities; And wherein the steel strip has the following microstructure: - precipitation-hardened acicular / bainitic ferrite and / or bainite, - a misorientation distribution (MOD) index of at least 0.85 at 1 / 4 thickness, - up to 5% by volume of second phase components, including any cementite, pearlite, martensite and / or retained austenite, - a number average grain size (GS) of at most 2.0 μm at 1 / 4 thickness in the rolling direction and based on a 15° grain tolerance angle N15 ), - At least 1060mm at 1 / 4 thickness -1 The total length of high-angle grain boundaries (∑GBL 15-65 ),and - a gamma-fiber X-ray random intensity ratio of at least 27 measured by XRD at 1 / 2 thickness; And wherein the steel strip has at least the following mechanical properties: 600-820MPa yield strength (Rp0.2), Ultimate tensile strength (Rm) between 730-950MPa, Tensile elongation (A50) of at least 11%, At least 45% hole expansion capacity (HEC), and Work of fracture (WOF) of at least 100J.

2. The hot rolled high strength steel strip according to claim 1, wherein the composition has a Mn content of 1.40-2.20 wt. %, preferably 1.40-2.10 wt. %.

3. The hot rolled high strength steel strip according to claim 1 or 2, wherein the composition has a C content of 0.070 to 0.12 wt. %, preferably 0.07 to 0.10 wt. %.

4. The hot rolled high strength steel strip according to any one of claims 1 to 3, wherein the composition has a Ti content of at least 0.08 wt. %, and preferably at least 0.09 wt. %, and more preferably 0.09 to 0.13 wt. %.

5. The hot rolled high strength steel strip according to any one of claims 1 to 4, wherein the composition has a Si content of 0.10-0.65 wt. %, preferably 0.20-0.55 wt. %, more preferably 0.25-0.50 wt. %.

6. The hot rolled high strength steel strip according to any one of claims 1 to 5, wherein the composition has a Cr content of at most 0.10 wt.%.

7. Hot rolled high strength steel strip according to any one of claims 1 to 6, wherein the composition has a Nb content of at most 0.025 wt.-%, preferably at most 0.02 wt.-%.

8. The hot rolled high strength steel strip according to any one of claims 1 to 7, wherein the composition has: in And a is equal to 1.8, and b is equal to 5.8, preferably a is equal to 2.0, and b is equal to 5.0, more preferably a is equal to 2.2, and b is equal to 4.0, and most preferably a is equal to 2.4, and b is equal to 3.

8.

9. The hot rolled high strength steel strip according to any one of claims 1 to 8, wherein the microstructure has a T factor of at least 1.1, preferably at least 1.15, more preferably at least 1.4, and wherein T is defined as: where MOD corresponds to the misorientation distribution index at 1 / 4 thickness, and the characteristic texture fibers, parts and components correspond to the associated X-ray random intensity ratios at 1 / 2 thickness.

10. The hot rolled high strength steel strip according to any one of claims 1 to 9, wherein the microstructure has at most 3 volume % of a second phase component including any of cementite, pearlite, martensite and / or retained austenite.

11. The hot rolled high strength steel strip according to any one of claims 1 to 10, wherein the microstructure has a misorientation distribution (MOD) index at ¼ thickness of at least 0.90, preferably at least 0.95, more preferably at least 1.

0.

12. The hot rolled high strength steel strip according to any one of claims 1 to 11, wherein the microstructure has a microstructure of at least 1100 mm at 1 / 4 thickness in the rolling direction. -1 , preferably at least 1200mm -1 , more preferably at least 1300 mm -1 The total length of high-angle grain boundaries (∑GBL 15-65 ).

13. The hot rolled high strength steel strip according to any one of claims 1 to 12, wherein the steel strip has a metal coating, the metal coating is preferably selected from: a Zn layer, a Zn-based alloy layer, a Mg-based alloy layer, an Al-based alloy layer.

14. A method for manufacturing a hot rolled high strength steel strip according to any one of claims 1 to 13, the method comprising the following steps: - casting of a slab followed by a step of reheating the solidified slab to a temperature between 1050° C. and 1260° C. and hot rolling said slab, or casting of a slab or strip followed directly by a step of hot rolling said slab or strip; - hot rolling the steel slab or steel strip, and completing the hot rolling at a finishing temperature between 960°C and Ar3+20°C, preferably between 960°C and 900°C, more preferably between 960°C and 920°C; - accelerating the hot rolled steel strip to an intermediate temperature (IT) on the run-out roller between 50°C and 600°C, preferably between 520°C and 590°C, more preferably between 540°C and 580°C, at a primary cooling rate on the run-out roller between 50°C and 150°C / s, preferably between 65°C and 100°C / s ROT ); - accelerating the hot rolled steel strip to a coiling temperature between 450°C and 580°C, preferably between 475°C and 560°C, at a secondary cooling rate of the run-out roller between 1-30°C / s, preferably between 5-25°C / s; - coiling the hot rolled and cooled strip at a coiling temperature between 450° C. and 580° C.; - further cooling the coiled hot rolled steel strip to ambient temperature; - optionally pickling the hot rolled steel strip; and - Optionally providing the hot rolled steel strip with a metal coating, preferably selected from: a Zn layer, a Zn-based alloy layer, a Mg alloy layer, an Al-based alloy layer, and preferably applied by hot coating or hot dip coating.

15. An automobile component made of the high-strength hot-rolled steel strip according to any one of claims 1 to 13 or made of the high-strength hot-rolled steel strip obtainable by the method according to claim 14.

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