High-embossable ultra-high strength ductile hot-rolled steel, method of manufacturing the hot-rolled steel and use thereof

Abstract

The present invention relates to hot-rolled (HR) steel strips having high flangeability and high total elongation, bendability and toughness values at ultra-high strength levels, to a method of manufacturing said hot-rolled steel and to its use.

Description

Invention Field

[0001] This invention relates to hot-rolled (HR) strip steel with high rollability, high total elongation, bendability and toughness at ultra-high strength levels, and to methods for manufacturing said hot-rolled steel and its uses. Background of the Invention

[0003] It is well known that as the strength of hot-rolled (HR) steel increases, its formability decreases. The primary applications of HR steel in transportation and automotive applications are chassis and suspension (C&S), such as in lower control arms. Other applications include truck frame longitudinal beams, bumpers, or battery boxes in electric vehicles. Typical thicknesses of HR steel used in these applications are less than 4.5 mm. Thicker HR strips, such as up to 12 mm, can be used in engineering applications such as crane mounts or in transportation applications for heavy-duty truck frames.

[0004] From a weight reduction perspective, it is necessary to use higher strength steel for the above applications in order to reduce the size of the steel strip. Therefore, ultra-high strength steel (UHSS) with an ultimate tensile strength (Rm) typically greater than 1000 MPa can be used for this purpose.

[0005] These applications of HR steel require complex mechanical properties. In addition to high strength, the steel should also possess good formability for cold forming of components, as this is an energy-efficient manufacturing route compared to hot forming. Furthermore, for applications such as bumpers, battery casings, crane mounts, or vehicle frame longitudinal beams, good impact toughness or energy absorption capacity is also required. For assembly components, good weldability is also necessary, typically characterized by the steel's low carbon equivalent.

[0006] However, as the tensile strength of steel increases, its formability parameters decrease. Formability is a general term for steel sheets, which is considered as a combination of material behavior during several mechanical operations such as stretching, bending, drawing, and crimping. Depending on the part geometry, any combination of two or more of the material's properties is important during sheet metal forming. For typical automotive C&S parts, stretchable crimping is also important. This type of formability requires high hardness expansion (HEC) and good total elongation. For manufacturing frame longitudinals, bumpers, or battery casings typically produced by rolling, bendability is important. Good HEC, bendability, and elongation are also required for manufacturing crane hangers.

[0007] Achieving high formability and high impact toughness in steel at ultra-high strength levels is challenging.

[0008] Purpose of the invention

[0009] The purpose of this invention is to provide hot-rolled steel strip with ultra-high strength, high rollability, good elongation, bendability and impact toughness.

[0010] The purpose of this invention is to provide hot-rolled steel strip with excellent weldability.

[0011] The object of this invention is to provide a method for producing such steel.

[0012] Invention Description

[0013] Based on the following hot-rolled steel, achieve one or more objectives:

[0014] Hot-rolled steel strip with ultra-high strength, excellent ductility and rollability, comprising (by weight%):

[0015] ●C: 0.10–0.30;

[0016] ●Si: 0.50–1.50;

[0017] ●Al: 0.010–1.00;

[0018] ●Mn: 1.00–3.00;

[0019] ●(Si+Al): 0.80-2.50;

[0020] and optionally any one or more of the following alloying elements:

[0021] ●V: less than 0.10;

[0022] ●Nb: less than 0.10;

[0023] ●Ti: less than 0.10;

[0024] ●Mo: less than 0.50;

[0025] ●Cr: less than 1.50;

[0026] ●Cu: less than 1.00;

[0027] ●Ni: less than 0.50;

[0028] ●B: Less than 0.0030 (30ppm);

[0029] Inevitably, it also includes

[0030] ●N: Less than 0.0100 (100ppm)

[0031] ●S: Less than 0.005;

[0032] ●P: Less than 0.020;

[0033] The balance consists of Fe and other unavoidable impurities produced during the ironmaking and steelmaking processes;

[0034] It has a yield strength of at least 1100 MPa, an ultimate tensile strength of at least 1200 MPa, a yield ratio of at least 0.85, a total elongation of at least 6.0%, a hole expansion ratio of at least 30%, and a bending angle of at least 70° with a thickness of 1 mm.

[0035] It has a microstructure consisting of 40 to 85 vol% tempered martensite, 60 to 15 vol% primary martensite, less than 1 vol% retained austenite, and essentially no cementite or other carbides.

[0036] Preferred embodiments are disclosed in any dependent claim.

[0037] According to the second aspect, the invention is also embodied in the following method for producing steel according to the invention:

[0038] A method for preparing hot-rolled steel strip with ultra-high strength, excellent ductility, and rollability includes the following steps:

[0039] ● Cast molten steel into thick or thin slabs having the following composition (in weight percent):

[0040] ●C: 0.10–0.30;

[0041] ●Si: 0.50–1.50;

[0042] ●Al: 0.030–1.00;

[0043] ●Mn: 1.00–3.00;

[0044] ●(Si+Al): 0.80-2.50;

[0045] and optionally any one or more of the following alloying elements:

[0046] ●Nb: less than 0.10;

[0047] ●Ti: less than 0.10;

[0048] ●Mo: less than 0.50;

[0049] ●Cr: less than 1.50;

[0050] ●Cu: less than 1.00;

[0051] ●Ni: less than 0.50;

[0052] ●B: Less than 0.0030 (30ppm);

[0053] Inevitably, it also includes

[0054] ●N: Less than 0.0100 (100ppm);

[0055] ●S: Less than 0.005;

[0056] ●P: Less than 0.020;

[0057] The balance consists of Fe and other unavoidable impurities produced during the ironmaking and steelmaking processes;

[0058] ●Heat or reheat the slab;

[0059] ● The slab is hot-rolled into hot-rolled strip, and then...

[0060] The thick slab is rough-rolled to intermediate specifications to disrupt the microstructure of the cast state, and then hot-rolled into hot-rolled strip, or

[0061] o Hot rolling thin slabs into hot-rolled strips via direct rolling

[0062] ●The hot finishing temperature (FRT) is greater than the Ar3 temperature of the steel, where Ar3 is the temperature at which the austenite-to-ferrite transformation begins during cooling.

[0063] ● The hot-rolled strip is cooled more quickly on the output roller conveyor at a cooling rate of over 20°C / s;

[0064] ●Then, the hot-rolled and cooled steel strip is wound at a temperature in the range of (Ms-50)℃ to (Ms-160)℃, where Ms is the martensite initiation temperature of the steel.

[0065] ● Allow the wound hot-rolled strip to be further cooled to ambient temperature;

[0066] ● Pickled hot-rolled steel strip.

[0067] According to a third aspect, the invention is also reflected in the use of hot-rolled steel in the production of parts for transportation or engineering applications.

[0068] The steel according to the present invention contains carbon, silicon, aluminum, and manganese as essential elements. The content ranges (in weight percent) of these alloying elements in the steel are as follows:

[0069] C: 0.10–0.30;

[0070] Si: 0.50–1.50;

[0071] Al: 0.010–1.00;

[0072] Mn: 1.00–3.00;

[0073] Where (Si+Al)≥0.80;

[0074] and optionally any one or more of the following alloying elements:

[0075] V: less than 0.10;

[0076] Nb: less than 0.10;

[0077] Ti: less than 0.10;

[0078] Mo: less than 0.50;

[0079] Cr: less than 1.50;

[0080] Cu: less than 1.00;

[0081] Ni: less than 0.50;

[0082] B: Less than 0.0030 (30ppm);

[0083] Inevitably, it also includes

[0084] N: less than 0.0100 (100ppm)

[0085] S: less than 0.005;

[0086] P: less than 0.020;

[0087] The balance is Fe and other unavoidable impurities produced by the ironmaking and steelmaking processes. Note that all composition percentages are given as weight % unless otherwise stated.

[0088] Carbon is present in the steel in an amount of 0.10-0.30%, preferably 0.10-0.26%, more preferably 0.10-0.23%. Carbon is added primarily for strength and hardenability, as it causes strong solid solution hardening in iron. Carbon ensures that during cooling at the output roller table after hot rolling, austenite does not transform into ferrite and / or pearlite at a rate greater than the critical cooling rate (20°C / s). Less than 0.10% C will not produce the desired Rm level of 1000 MPa or greater, preferably 1200 MPa or greater, and if C is above 0.30%, the weldability of the formed parts may deteriorate. Weldability is also improved by low carbon equivalent values. A suitable minimum carbon content is 0.16%.

[0089] Adding silicon in amounts less than 1.50% is used to improve strength through substitutional solid solution strengthening in the iron lattice. Another important role of Si in steel is that it slows down carbide precipitation (cementite and other carbides). As a result, the martensitic phase will not form harmful iron carbides in its matrix when tempering is performed. When Si is less than 0.50%, the strengthening and carbide formation suppression effects are insufficient to achieve the desired benefits. On the other hand, when Si is greater than 1.50%, excessive oxide formation can occur during the thermomechanical processing of steel (slab reheating, hot rolling, winding, etc.). These oxide scales are detrimental to hot rolling, pickling, coating, and overall surface appearance. Furthermore, the rolling force increases during hot rolling, and when the Si content exceeds 1.50%, the steel becomes brittle to the point that it is very difficult to hot roll. Therefore, according to the invention, the Si content is typically greater than 0.50% and less than 1.50%, preferably in the range of 0.60% to 1.30%, and more preferably in the range of 0.70% to 1.10%.

[0090] In the steel according to the invention, aluminum behaves similarly to Si. When intentionally added, it acts as a solid solution strengthening element in the steel. It also slows down carbide precipitation kinetics during martensitic tempering. When Al is less than 0.030%, the strengthening and carbide-inhibiting effects are negligible. Aluminum values ​​less than 0.030% are considered to be residues from the deoxidation step in the steelmaking process, and therefore a minimum of 0.030% is preferred. On the other hand, when Al is greater than 1.00%, excessive oxide formation can occur during the thermomechanical processing of the steel (slab reheating, hot rolling, winding, etc.). In addition, Al raises the ferrite-to-austenite transformation temperature, thus requiring hot rolling of the steel at higher temperatures to complete the hot rolling in the austenitic phase, since critical zone ferrite appears at lower temperatures. Higher amounts of oxide can occur at higher temperatures. These oxide scales are detrimental to hot rolling, pickling, coating, and overall surface appearance. Furthermore, when Al exceeds 1.00% along with the presence of Si, the rolling force during hot rolling increases to the point that the steel becomes very brittle and difficult to hot roll. Furthermore, an Al content greater than 1.00% can promote ferrite formation during cooling on the output roller conveyor by reducing the incubation time for ferrite formation during continuous cooling. Ferrite is a detrimental phase to this invention because it introduces a brittle interface between primary martensite and tempered martensite. These interfaces act as nucleation sites for damage initiation due to the reduced formability, elongation, and impact toughness of the steel caused by deformation. Therefore, Al is present in this invention in an amount ranging from 0.010-1.00%, preferably 0.030-1.00%, preferably 0.20-0.80%, and more preferably 0.30-0.80%.

[0091] While Si or Al alone can produce solid solution strengthening and inhibit carbide precipitation during martensitic tempering, the synergistic effect of these elements in the presence of both is similar to their individual effects. Therefore, the total content of (Si+Al) in this invention should be at least 0.80%, preferably at least 1.00%, to achieve the desired carbide inhibition effect and the desired strength level. The presence of both Al and Si offers several advantages in facilitating steel processing, particularly during hot rolling, acid treatment, and coating. The presence of a certain amount of Al and Si alters the oxide properties in the skin during high-temperature processing. This makes pickling of the skin after hot rolling easier.

[0092] As will be disclosed later, the initial martensite formed in the hot-rolled steel during the winding stage is tempered during the coil cooling process in this invention. Suppressing carbide formation during this self-tempering (coil cooling) process of the steel is important for this invention due to the individual or synergistic effects of Si and Al. As a result, the martensite reduces its dislocation density without forming carbides. Carbides are detrimental to the elongation, formability, and impact toughness of steel because they are inherently brittle and act as nucleation sites where damage begins during deformation.

[0093] Manganese is present in an amount of 1.00-3.00%. The main role of Mn is to improve strength and toughness. At levels less than 1.00% by weight, the desired effect is not achieved, while amounts greater than 3.00% will cause casting and segregation problems. Furthermore, the deformation mechanism in the steel can be altered to transformation-induced plasticity (TRIP) because of the stabilization to room temperature by Mn austenite, which is detrimental to achieving a good combination of all the target mechanical properties (i.e., impact toughness, formability, and strength) in the product. Preferably, the Mn content is in the range of 1.20-2.70%. In embodiments, the total Mn content is 1.40-2.60%, preferably 1.50-2.50%, more preferably 1.60-2.50%. In embodiments, a suitable minimum amount of Mn is 1.65%, and a suitable maximum amount of Mn will be 1.95%.

[0094] In addition to the aforementioned effects of the essential alloying elements C, Si, Al, and Mn in the steel of this invention, another collective effect of these alloying elements is improved hardenability. They help prevent the formation of pearlite or ferrite phases during cooling after austenitization. This characteristic, exceeding a certain cooling rate during the output roll cooling period after hot rolling and before winding, prevents the steel from having these phases. The presence of these softer phases (ferrite) and inhomogeneous phases (pearlite) is detrimental to obtaining good mechanical and formable properties in the final product because they promote brittle and discontinuous mesophases in the microstructure.

[0095] One or more microalloying elements, selected from V, Nb, Ti, and Mo, are optionally present. These microalloying elements increase strength through precipitation hardening of their carbides, nitrides, or carbonitrides. They also improve the weldability of the steel.

[0096] Chromium, another optional element in this invention, also improves the hardenability of steel.

[0097] Copper (when present) enhances the strength of steel through both solid solution strengthening and precipitation hardening of copper precipitates. Nickel improves impact toughness and counteracts any hot brittleness that may occur in steel during hot working due to the presence of copper.

[0098] If present as alloying elements, the preferred additions of these optional alloying elements (in weight percent) are:

[0099] V: 0.010-0.10

[0100] Nb: 0.010-0.10

[0101] Ti: 0.010-0.10

[0102] Mo: 0.050-0.50

[0103] Cr: 0.10-1.50

[0104] Cu: 0.030-1.00

[0105] Ni: 0.020-0.50

[0106] Nitrogen, sulfur, and phosphorus are residual elements present in steel due to the steel production and refining processes. Their amounts are limited to S < 0.005%, P < 0.020%, and N < 0.0100%. Amounts exceeding these limits are detrimental to mechanical properties, formability, and weldability. Preferably, S < 0.002%, and N between 0.0005% and 0.0100%. Nitrogen within the specified range produces a similar effect to carbon and contributes to strength by forming carbonitrides, which are microalloying elements.

[0107] The optional alloying elements, including nitrogen, sulfur, and phosphorus, can be changed independently within a specified range. They were found to have an additive rather than a synergistic effect in the steel according to the invention.

[0108] According to a second aspect, the invention is also embodied in a method for manufacturing hot-rolled strip to achieve a desired microstructure in the final product. Therefore, the method according to the invention is a method for producing hot-rolled steel having the chemical composition discussed above.

[0109] The method for producing steel includes the following steps:

[0110] - Casting molten steel into slabs;

[0111] - Reheat the slab, preferably at a temperature of 1100°C or higher and preferably for 30 minutes or more;

[0112] - Roughly roll the slab to intermediate dimensions, typically in the range of 35-45mm, to disrupt the microstructure of the cast state;

[0113] - Hot rolling of steel into strip, preferably using a hot finishing rolling temperature (FRT) greater than the Ar3 temperature of the steel, where Ar3 is the temperature at which the austenite-to-ferrite transformation begins during cooling;

[0114] -Use a cooling rate of over 20°C / s on the output roller conveyor to accelerate the cooling of hot-rolled strip;

[0115] - Hot-rolled and cooled steel strip is wound at a temperature ranging from (Ms-50)℃ to (Ms-160)℃, where Ms is the martensite initiation temperature of the steel (in °C).

[0116] - Cool the steel coil to room temperature;

[0117] - Pickled hot-rolled steel strip;

[0118] -Optionally, the hot-rolled strip may be coated with Zn or a Zn-based alloy or an Al-based alloy or any other coating;

[0119] To avoid misunderstanding, Ms is expressed in °C. Preferably, FRT is greater than Ar3+50 °C. Figure 1 A schematic representation of the hot rolling and cooling to room temperature process superimposed on a continuous cooling transformation (CCT) diagram. Room temperature is defined as approximately 20°C. Reheating is preferably performed for 60 minutes or longer, particularly when the hot rolling process according to the invention is carried out in a conventional hot strip mill based on thick slabs.

[0120] This invention is not limited to casting methods. Steel can be cast into conventional thick slabs with a casting thickness between 150 and 350 mm, typically 225 to 250 mm, and into thin slabs with a casting thickness between 50 and 150 mm in a strip direct casting and rolling mill. Schematic examples of methods including conventional hot strip mills and thin slab casting / direct rolling mills are shown in Figures 2a and 2b, respectively. For conventional thick slab casting, reheating of the slab is necessary to reheat the slab from ambient temperature (typically thick cast slabs have been cooled from casting temperature to ambient temperature in the slab storage area) and to homogenize the slab in terms of composition. Therefore, the reheating temperature should be greater than 1100°C, also used to dissolve any precipitates when alloying elements are present, and to bring the slab to such a temperature that final hot rolling in a finishing mill is still possible at FRT>Ar3. Often this requires a (slab) reheating temperature between 1150 and about 1250°C. For thin slab casting, the cast slab is immediately homogenized in a homogenizing furnace after casting, wherein the homogenization temperature should be greater than 1100°C and typically about 1125 to 1150°C. This also prevents the formation of any precipitates when alloying elements are present and also ensures that the thin slab reaches a temperature such that final hot rolling in a finishing mill is still possible at FRT>Ar3. The reheating or homogenization time for the thin slab casting path according to the invention is preferably 30 minutes or more.

[0121] Hot rolling of the steel must be performed in the austenitic phase to ensure the absence of ferrite in the final microstructure. Another purpose of hot rolling in the austenitic phase is to reduce the hot rolling stress, and therefore the finishing rolling temperature (FRT) is preferably maintained at a temperature at least 50°C higher than that of the steel's Ar3.

[0122] After hot rolling, the steel strip is cooled on the output roller conveyor. It is required that the steel be cooled at a rate above the critical cooling rate to avoid any undesirable phase transformation from austenite. In particular, ferrite and pearlite must not form, as these are detrimental to the mechanical and formability properties of the final product. Therefore, the ROT-CR must exceed the critical cooling rate to avoid the formation of ferrite and pearlite. There is no critical maximum ROT-CR, because as long as the critical cooling rate mentioned above is exceeded throughout the thickness of the strip, the transformation from austenite is ensured. Unnecessarily high ROT-CR can affect the flatness of the strip after cooling and cause problems in controlling the stopping at the correct cooling stop temperature; therefore, a suitable maximum ROT-CR is about 300 °C / s, preferably about 200 °C / s, and more preferably about 150 °C / s. Depending on the thickness of the strip, the practical ROT-CR range is 20 to 100 °C / s, as this can be achieved through air cooling, laminar flow cooling, or water spray cooling. For practical reasons, the output roller cooling rate (ROT-CR) is defined as the average cooling rate of the strip surface.

[0123] Next, the hot-rolled steel strip is wound at a temperature lower than Ms, within a temperature range of (Ms-50)°C to (Ms-160)°C. Winding below Ms ensures that the subsequent coil cooling begins with a mixture of martensite and austenite phases, with an initial martensite content in the range of 40-85% by volume. If the initial martensite content is higher than this amount, or in other words, if the winding temperature (CT) is lower than Ms-160°C, the required tempering effect for initial martensite is not achieved, and as a result of less available time during coil cooling and excessively low temperatures at which effective tempering occurs, high ductility, formability, and impact toughness are not achieved in the steel. If the initial martensite content is less than 40% by volume, excessive tempering of martensite can occur, where the product is not the ultra-high strength steel in the context of this invention.

[0124] During the coil cooling process, the tempering of the initial martensite occurs continuously. Simultaneously, as the steel cools in the coil, new primary martensite forms. Due to the presence of Si and Al in the steel, carbides do not form in the tempered martensite. Furthermore, due to a certain distribution of carbon from martensite to austenite, a very small amount of austenite can remain untransformed at room temperature (also known as retained austenite), but its amount is preferably limited to a maximum of 1% by volume, including 0% by volume.

[0125] After the steel has cooled to room temperature, the oxides (scale) on the hot-rolled steel are removed by pickling in an acid solution (e.g., HCl) at a warm temperature (80-120°C) or by a combination of surface mechanical brushing and pickling. This step is necessary to make the steel surface suitable for direct use as uncoated HR steel or, when corrosion resistance is optionally required, to make it suitable for coating methods.

[0126] Optionally, HR steel strips may be coated, for example by hot-dip coating or electrophoretic coating, using Zn or Zn-based alloys, or Al-based alloys or any other coating techniques to produce good corrosion resistance during service.

[0127] The above methods result in a desired microstructure with the target mechanical properties. This invention is also reflected in the steel articles manufactured according to the above methods and the chemical composition (in volume %) of steel containing the following microstructure:

[0128] - Tempered martensite (initial martensite during winding): 40-85%, preferably at least 50%, more preferably at least 60%;

[0129] - Primary martensite (martensite formed during cooling after winding): 15-60%, preferably up to 50%, more preferably up to 40%;

[0130] - Residual austenite: up to 1% by volume, including 0% by volume.

[0131] - Cementite or any other metal carbides: 0% by volume

[0132] The steel chemical composition, method, and microstructure according to the present invention result in the following mechanical and formability properties.

[0133] - Yield strength (Rp): at least 1100 MPa

[0134] - Ultimate tensile strength (Rm): at least 1200 MPa

[0135] - Yield ratio (Rp / Rm): at least 0.85

[0136] -Total elongation: at least 6.0% JIS5

[0137] - Hole enlargement capability: at least 30%

[0138] Bending angle at -1mm thickness: at least 70°

[0139] Preferably, the Charpy impact toughness is at least 40 joules at -40°C and at least 100 joules at room temperature.

[0140] The steel chemical composition, method, and microstructure according to the present invention preferably result in the following mechanical and formability properties.

[0141] - Yield strength (Rp): at least 1100 MPa

[0142] - Ultimate tensile strength (Rm): at least 1200 MPa

[0143] - Yield ratio (Rp / Rm): at least 0.85

[0144] -Total elongation: at least 8.5% JIS5

[0145] - Hole enlargement capability: at least 50%

[0146] Bending angle at -1mm thickness: at least 80°

[0147] -Charpy impact toughness: at least 40 joules at -40°C and at least 100 joules at room temperature.

[0148] The strength of steel is primarily determined by the presence of its hard components in the microstructure. Martensite is a strong phase in steel, and due to the low-temperature tempering (less than Ms) during coiling cooling, it does not lose much strength. Therefore, both primary and tempered martensite contribute to the ultra-high strength values ​​achieved in this invention. Furthermore, the absence of carbides due to the presence of Si and Al in the steel reduces the initiation of damage during deformation, resulting in a high total elongation. Retained austenite is minimized to less than 1% by volume because it is detrimental to the low stability of impact toughness.

[0149] Retained austenite is generated by carbon transforming from martensite to distributed austenite during the coiling cooling process. Carbon improves the stability of austenite by lowering the Ms temperature. However, retained austenite is intentionally avoided in this invention because its mechanical stability is difficult to control during different deformation and forming processes. Retained austenite should possess very high mechanical stability for its beneficial effects of increased elongation (i.e., stretchability) and impact toughness. It requires very high carbon supersaturation along with a fine film morphology to improve these properties. High carbon supersaturation is very difficult to achieve in low-temperature continuous cooling processes, such as during coiling cooling. When the mechanical stability of austenite is low, it rapidly transforms into martensite and forms a brittle interface with the matrix phase that affects the total elongation. Retained austenite with low mechanical stability transforms even faster during dynamic loading processes, such as impact, and reduces impact toughness. Therefore, in this invention, a more uniform microstructure is produced by tempering martensite and primary martensite, without a large amount of retained austenite. In other words, the presence of retained austenite is intentionally avoided, and its maximum amount is limited to 1% by volume.

[0150] Another motivation for avoiding or minimizing the retained austenite phase in this invention is to reduce the tendency for liquid metal embrittlement (LME) during the welding of Zn or Zn-alloy coated steels. Zn or Zn-alloy coated steels with retained austenite phases in their microstructure are known to be more susceptible to LME during welding.

[0151] This was achieved by using low-temperature winding in the range of (Ms-50)℃ to (Ms-160)℃, which is a temperature range for stabilizing a large amount of austenite without the expected large amount of carbon distribution.

[0152] Rp, Rm, and total elongation were determined from quasi-static strain at room temperature (strain rate 3 × 10⁻⁶). -4 s -1 Tensile testing is performed according to EN10002-1 / 150 6892-1 using JIS No. 5 specimen geometry parallel to the rolling direction. The geometry of the tensile specimen is a length of 50 mm, a width of 25 mm, and a thickness of 3.2 mm along the rolling direction. The strength of the steel at 0.2% permanent deformation is measured as the yield strength (Rp or YS). The ratio of yield strength to ultimate tensile strength (Rp / Rm) is expressed as the yield ratio.

[0153] Flexibility was determined in both the longitudinal and transverse directions on a 3.2 mm thick, 40 mm × 30 mm specimen using a three-point bend test according to VDA 238-100. The bend axis was along the 30 mm dimension and the bend radius was 0.4 mm. The bend angles obtained from strips with different thicknesses (2.8, 3.2, and 3.5 mm, respectively) were converted to the angle corresponding to a 1.0 mm thickness using the following formula: Bending angle at 1.0 mm thickness = Measured angle × Square root of actual thickness in mm. From these converted bend angles, the lowest values ​​of the longitudinal and transverse specimens for a specific heat treatment condition were taken to claim the scope of the invention.

[0154] The flammability or enlargement capacity (HEC) of steel is determined by an enlargement test. A 90mm × 90mm × 3.2mm specimen is cut from the coiled steel. A 10mm diameter hole is punched in the center of the specimen, and an enlargement test is performed according to ISO / TS16630:2003(E). The HEC value is determined by the following formula: HEC = (Enlargement of initial hole diameter / Initial diameter) × 100%.

[0155] Charpy impact toughness was measured according to ASTM A370 using full-size Charpy V-notch (CVN) specimens (55 mm × 10 mm × 10 mm). Testing was conducted in two sheet directions, with the V-notch machined parallel and perpendicular to the rolling direction.

[0156] For all of the above mechanical tests, at least three specimens should be tested for each condition and the average value should be reported.

[0157] Microstructure was analyzed using a combination of techniques including optical microscopy, X-ray diffraction (XRD), scanning electron microscopy (SEM), and dilatation. Dilatation was performed on a (l×w×t) 10mm×5mm×3.2mm sample by heating the sample to 950℃ at a rate of 10℃ / s, holding for 2 minutes, and then cooling to room temperature at a rate of 100℃ / s (for Ms quenching) or 0.3℃ / s (for Ar3 slow cooling). The temperatures of Ms and Ar3 were determined from the dilatation data. The amount of initial martensite (i.e., tempered martensite after cooling) after coiling steel was determined by using the Koistinen-Marburger formula given in the following literature: D.P. Koistinen, E.R. Marburger, “A general equation prescribing the extent of the austenite-martensite transformation in pureiron-carbon alloys and plain carbon steels”, Acta Metallurgica, Vol. 7, 1959, pp. 59-60.

[0158] f = 100 · {1 - exp(-(1.10 x 10)} -2· (M s -CT))}

[0159] Where Ms is the martensite initiation temperature (in °C) and CT is the winding temperature (in °C), so (Ms-CT) reflects the supercooling that is less than Ms at the start of winding cooling and is therefore a measure of the initial amount of martensite.

[0160] The amount of retained austenite was determined by XRD at a position of 1 / 4 thickness of the sample. The sample was analyzed using a Panalytical Xpert PRO standard powder diffractometer (Co K). α XRD patterns were recorded on the surface of the radiometer from 45° to 165° (2Θ). Phase proportions were quantitatively determined using the Bruker Topas software package for Rietveld refinement, followed by Rietveld analysis. The amounts of carbides, ferrite, pearlite, and bainite in the microstructure were determined by analyzing high-resolution SEM images. The primary martensite fraction was obtained by subtracting the initial martensite and other phase fractions (if applicable, retained austenite, carbides, and other definite phases) from the total amount.

[0161] There are no restrictions on the composition of the zinc or zinc alloy coating. Although the coating can be applied in various ways, hot-dip galvanizing using a standard GI coating bath is preferred. Zn-based coatings may contain Zn alloys with Al as an alloying element. Preferred zinc bath compositions contain 0.10-0.35% by weight Al, with the balance being zinc and unavoidable impurities.

[0162] Other Zn coatings may also be applied. Examples include zinc alloy coatings according to WO 2008 / 102009, particularly zinc alloy coatings consisting of: 0.3-4.0 wt% Mg and 0.05-6.0 wt% Al, preferably 0.1-5.0 wt% Al, and optionally up to 0.2 wt% of one or more additional elements, as well as unavoidable impurities and the balance being zinc. Preferred Zn baths containing Mg and Al as the main alloying elements have a composition of: 0.5-3.8 wt% Al, 0.5-3.0 wt% Mg, optionally up to 0.2 wt% of one or more additional elements; the balance being zinc and unavoidable impurities. Additional elements typically added in small amounts (less than 0.2 wt%) may be selected from the group consisting of: Pb, Sb, Ti, Ca, Mn, Sn, La, Ce, Cr, Ni, Zr, and Bi. Pb, Sn, Bi, and Sb are typically added to form zinc flowers. Preferably, the total amount of additional elements in the zinc alloy is at most 0.2% by weight. These small amounts of additional elements do not significantly alter the properties of the coating or bath for typical applications. Preferably, when one or more additional elements are present in the coating, each is present in an amount of <0.02% by weight, more preferably in an amount of <0.01% by weight. Additional elements are typically added only to prevent the formation of dross in a bath containing molten zinc alloy for hot-dip galvanizing, or the formation of zinc flowers in the coating.

[0163] In another embodiment, the metallic coating comprises a (commercially pure) aluminum layer or aluminum alloy layer. A typical metal bath for hot-dip coating such an aluminum layer comprises aluminum alloyed with silicon, for example, aluminum alloyed with 8 to 11% by weight of silicon and up to 4% iron, optionally up to 0.2% of one or more additional elements such as calcium, unavoidable impurities, with the balance being aluminum. Silicon is present to prevent the formation of a thick iron-metal intermetallic compound layer that reduces adhesion and formability. Iron is preferably present in an amount between 1% and 4%, more preferably at least 2%. Example

[0164] Seven steel ingots with chemical compositions of the present invention AB and DH, and comparative steel C, with dimensions of 200 mm × 100 mm × 100 mm, were cast by melting the feed in a vacuum induction furnace. The chemical compositions of these steels are given in Table 1. Steels AB and DH contain C, Si, and Al within the boundaries defined by the present invention, while the comparative steel contains Al and Si outside the boundaries defined by the present invention. All ingots were reheated at 1200 °C for 1 hour and rough-rolled to a thickness of 25 mm. The strips were then reheated again at 1200 °C for 30 minutes and hot-rolled using FRT (which is in the austenitic phase region for all steels) at temperatures greater than 900 °C to their final thicknesses of 2.8, 3.2, 3.5, and 12 mm. The Ar3 and Ms of the steels, measured by expansion determination, are also given in Table 1.

[0165] Immediately after hot rolling, the steel is subjected to output roller cooling at various cooling rates and then coil cooling simulation is completed in a muffle furnace by cooling to room temperature from different initial CTs. The strip is then pickled in a conventional manner to remove oxides.

[0166] Table 2 summarizes the various processing conditions for the steels. A, B, and C have similar Ms and Ar3. For steel A, the FRT temperature is 953°C, for steel B it is 939°C, and for steel C it is 945°C, all of which are at least 50°C above Ar3. For steels A and B, a slow output roller cooling rate (ROT-CR) of 3°C / s is used, which is outside the lower boundary defined by this invention. Additionally, for steels A and B using FRT and CT within the boundaries defined by this invention, two additional winding temperatures (200°C and 480°C) are used, which are outside the boundaries defined by this invention. The CT at 200°C is much lower than (Ms-160)°C, and 480°C is higher than the Ms for these steels. These conditions are used for comparative purposes. For steel C, which has a chemical composition outside the scope of this invention, all processing conditions (FRT, ROT-CR, and CT) are selected within the boundaries defined by this invention. For steel DF, one set of process parameters (FRT, ROT-CR, and CT) is within the scope claimed by this invention; however, for another set, FRT and ROT-CR remain the same, but only CT is kept higher for comparison purposes. In this case, a CT of 375°C is used, which is higher than that of steel DF (Ms-50°C). In fact, this CT is slightly higher than their Ms temperatures. For steel GH, all processing parameters are kept within the boundaries claimed by this invention.

[0167] For the various mechanical and microstructural characterizations as described, specimens were selected from the final steel strip. 12 mm thick hot-rolled steel was used to prepare Charpy impact specimens, while 2.8, 3.2, and 3.5 mm thick steel strips were used for all other characterizations.

[0168] For various processing conditions, the phase content in the microstructure is presented in Table 3, the tensile properties are given in Table 4, the bending and HEC test results are given in Table 5, and the Charpy impact toughness is given in Table 6. The following are the abbreviations and symbols used in the tables presenting the tensile and bending test results: Rp = yield strength, Rm = ultimate tensile strength, A JIS5 = Total elongation using JIS5 specimens, BA = bending angle, L = longitudinal specimens where the bending axis is parallel to the rolling direction, T = transverse specimens where the bending axis is perpendicular to the rolling direction.

[0169] Table 3 shows that steels A and B, processed using FRT and ROT-CR within a defined boundary of a CT range of 275-375°C, achieve microstructures consisting of less than 85 vol% and at least 15 vol% tempered martensite and primary martensite, respectively, which fall within the range required by this invention. Furthermore, the steels are free of carbides in their microstructures and have a retained austenite content of less than 1 vol% for these processing conditions. Their microstructures contain no other phases such as ferrite, bainite, or pearlite.

[0170] When steel AB undergoes ROT-CR > 20°C / s and FRT at temperatures above 50°C above Ar3, it produces a microstructure with abundant bainite and retained austenite, and when wound at 480°C (greater than Ms for these steels), a certain amount of primary martensite forms below Ms. For these winding conditions, carbides are absent in these steels due to their Al and Si content. The high retained austenite content is caused by carbon enrichment in the austenite during the bainite transformation (greater than Ms) and the formation of primary martensite during the cooling of the coil (less than Ms). This bainitic microstructure with abundant retained austenite differs from the microstructure anticipated in this invention.

[0171] Similarly, when steels A and B are wound at 200°C, which is much lower than the Ms-160°C required for these steels, the microstructure of the steels also becomes different from that required by the present invention. This CT condition has more than 85% by volume tempered martensite and less than 15% by volume primary martensite. On the other hand, steels A and B, when having FRT and CT within the range required by the present invention but with a slower ROT-CR of 3°C / s (less than 20°C / s), exhibit significant amounts of ferrite and pearlite in their microstructure, in addition to a bainitic matrix and a large amount of retained austenite. Ferrite, pearlite, and bainite are formed due to the slow ROT-CR prior to winding.

[0172] During processing of steel C with all parameters within the range required by this invention, a significant amount of carbides (2.3 vol%) are formed due to the small amounts of Si and Al in steel C.

[0173] As a result of the microstructure described above, the properties given in Tables 4 to 6 are obtained. Steels A and B achieve Rp greater than 1100 MPa and Rm greater than 1200 MPa, with a yield ratio higher than 0.85 along with total elongation (A). JIS5 The yield strength is above 8.5%. When the CT is too high (480°C), the minimum levels of Rp and Rm required in this invention are not achieved in steels A and B due to the presence of softer phases (bainite and retained austenite) – although the total elongation is high. The low Rp value at CT of 480°C also results in a yield ratio of less than 0.85. On the other hand, when the CT is too low (CT = 200°C), Rp and Rm are greater than the target values ​​with a high yield ratio, but the total elongation is too low (<8.5%). The low total elongation is caused by the excessively high amount of initial martensite present in the microstructure (= tempered martensite > 85% by volume) and the lack of tempering effect during coil cooling due to the low availability of time and the low temperature for effective tempering.

[0174] Using a slow ROT-CR process at 3℃ / s, the Rp and Rm values ​​in steels A and B are below 1100 MPa and 1200 MPa, respectively, due to the formation of softer phases including bainite, ferrite, pearlite, and retained austenite. Furthermore, the yield ratio is less than 0.85, although the total elongation is high.

[0175] Steel C exhibits low Rp, Rm, yield ratio, and total elongation values ​​due to the presence of carbides in its microstructure, resulting from the absence of carbide-inhibiting elements Si and Al. Carbides are detrimental to mechanical properties and promote damage during deformation. Therefore, steel C yields low levels of tensile properties.

[0176] Similar to tensile properties, steels A and B also exhibit high bendability and HEC when processed within the defined processing variables (FRT, ROT-CR, and CT) of this invention (Table 5). A minimum bend angle of 80° and a minimum HEC value of 50% were achieved at a thickness of 1.0 mm. However, when CT is high and greater than the steel's Ms (i.e., 480°), both the minimum bend angle and HEC are low and less than the target values ​​of 80° and 50%, respectively, at a thickness of 1.0 mm. This is due to the multiphase properties of the microstructure of these steels, which contain primary martensite, bainite, and retained austenite (Table 3), resulting in multiple damage initiation sites at the interfaces of these phases during deformation. Martensite, present as primary martensite and formed from retained austenite due to the TRIP effect, is a stronger phase than bainite and any untransformed retained austenite. On the other hand, for ideal processing conditions with optimized amounts of tempered martensite and primary martensite in steels A and B, there is a low difference in hardness or strength between these phases, resulting in uniform deformation during bending and reaming. This leads to high HEC and bending values ​​when using ideal processing conditions.

[0177] Furthermore, in the case of steels A and B with excessively low CT (less than Ms (200°C), bendability and HEC values ​​are also low because of the excessive amount of tempered martensite (>85% by volume) in their microstructure (Table 3). This primary martensite, present only at the start of winding due to the lack of effective tempering, causes low ductility (which is also reflected in their total elongation values ​​in Table 4), resulting in the low formability of these steels as measured by bendability and HEC.

[0178] When using slow ROT-CR (3℃ / s), the presence of the softer phases of ferrite and pearlite also deteriorates the bendability and HEC value of steels A and B, as can be seen from Table 5. This is due to the brittle interface between these softer phases and the harder phases of bainite and martensite, which is formed after the transformation of retained austenite during loading.

[0179] Steel C exhibits very poor bendability and HEC values, significantly lower than the minimum values ​​of 80° and 50%, respectively, at a thickness of 1.0 mm. These poor formability parameters in steel C are caused by the very low Al and Si contents in the steel, which promote carbide formation (Table 3), even when processing variables are within the ranges specified in this invention.

[0180] The Charpy impact toughness of steel A processed according to the present invention is above 100 J and 40 J in transverse specimens (which show lower values ​​than longitudinal specimens) when tested at room temperature and -40°C, respectively. The same values ​​obtained using a high CT of 480°C and a lower CT of 200°C are significantly lower than the minimum values ​​obtained using the ideal processing path described in this invention. As previously explained, these low toughness values ​​are caused by brittle fracture resulting from the spontaneous transformation of retained austenite to martensite, the presence of heterogeneous harder and softer phases, and a low tempering effect. Furthermore, the presence of carbides in steel C also contributes to the poor Charpy impact toughness values ​​at room temperature and low temperatures.

[0181] Therefore, as discussed above, these examples illustrate that when the composition is designed according to the invention and steel is processed according to the invention, the steel achieves, as expected, high tensile, formable, and tough properties due to their microstructure. When working outside the boundaries defined by the invention, the same favorable combination of all properties is not achieved.

[0182] Table 1: Chemical composition of steel (wt%) (I: This invention, C: Comparative)

[0183]

[0184] The Nb content in steel AG is at a residual level. No niobium is added as an alloying element in these steels. Niobium is added as an alloying element in steel H.

[0185] Steel AB has a Mn content of approximately 1.8%, while steel DH has a Mn content of approximately 2.35%, along with varying amounts of Si and Nb. The role of Mn is to reduce Ms, and it causes a shift in the ratio of tempered martensite to primary martensite. Although steel DH shows a lower Mn content than steels A and B... JIS5 HEC and bending angle, however, steel DH meets the requirements.

[0186] As shown in Table 3, steel DH achieves a microstructure consisting of less than 85 vol% and at least 15 vol% tempered martensite and primary martensite, respectively, within the defined boundaries of CT at 275°C (steel DF) and 300°C (steel GH), using FRT and ROT-CR, which fall within the range required by this invention. Furthermore, the steels are free of carbides in their microstructures and have a retained austenite content of less than 1 vol% for these processing conditions. Their microstructures do not contain any other phases such as ferrite, bainite, or pearlite.

[0187] Steel DH also achieves a yield strength (Rp) greater than 1100 MPa and a yield strength (Rm) greater than 1200 MPa, with a yield ratio of 0.85 or higher, along with a total elongation (A). JIS5 The bending flexibility and HEC are also high when the steel DH is machined within the defined processing variables (FRT, ROT-CR, and CT) of this invention (Table 5). A minimum bending angle of 70° and a minimum HEC value of 30% were achieved at a thickness of 1.0 mm in these steels.

[0188] However, when the CT of steel DF is 375°C (slightly higher than the Ms of steel), the microstructure of the steel contains a certain amount of bainite and more than 1% by volume of retained austenite (Table 3), which is not intended to achieve the microstructure of the present invention, and the contents of primary martensite and tempered martensite also fall outside the scope defined by the present invention. For a CT of 375°C, these unexpected microstructures do not result in exceptionally high strength in these steels. Rp values ​​are less than 1100 MPa and Rm values ​​are less than 1200 MPa, with yield ratios below 0.85, due to the presence of softer phases of bainite and retained austenite (Table 4), although good bendability and HEC values ​​are achieved (Table 5). Therefore, steels with chemical compositions that can be produced within the defined scope of the present invention may not achieve all the desired mechanical properties without processing within the defined window of the present invention.

[0189] Table 2: Processing Variables Applied to Steel

[0190]

[0191] *A 12 mm thickness is required for Charpy testing; other thicknesses are used to determine microstructure, tensile properties, bending angle, and HEC.

[0192] Table 3: Microstructure of steel (B: bainite, P: pearlite, F: ferrite)

[0193]

[0194]

[0195] Table 4: Tensile Properties of Steel

[0196]

[0197]

[0198] Table 5: Bending capacity and HEC of steel

[0199]

[0200]

[0201] AC bending angle is measured at a thickness of 3.2mm; *DF bending angle is measured at a thickness of 2.8mm; **G and H bending angles are measured at a thickness of 3.5mm.

[0202] Table 6: Charpy impact toughness of steel (L = longitudinal specimen, T = transverse specimen)

[0203]

[0204] Brief description of the attached diagram

[0205] The invention will now be explained by way of the following non-limiting drawings.

[0206] Figure 1 The thermomechanical processing of the present invention is illustrated in the figure.

[0207] Figure 2a shows a schematic diagram of a hot rolling mill for processing thick cast slabs, and Figure 2b shows a thin slab casting equipment with a direct rolling mill.

[0208] Figure 3 and 4 Shows the geometry and definition of the curved and Charpy samples.

Claims

1. Hot-rolled steel strip with ultra-high strength, excellent ductility and rollability, comprising (by weight%): ●C:0.10–0.30； ●Si: 0.50–1.50; ●Al:0.010–1.00； ●Mn: 1.00–3.00; ●(Si+Al): 0.80-2.50; and optionally any one or more of the following alloying elements: ●V: less than 0.10; ●Nb: less than 0.10; ●Ti: less than 0.10; ●Mo: less than 0.50; ●Cr: less than 1.50; ●Cu: less than 1.00; ●Ni: less than 0.50; ●B: Less than 0.0030 (30ppm); Inevitably, it also includes ●N: Less than 0.0100 (100ppm) ●S: Less than 0.005; ●P: Less than 0.020; The balance consists of Fe and other unavoidable impurities produced during the ironmaking and steelmaking processes; It has a yield strength of at least 1100 MPa, an ultimate tensile strength of at least 1200 MPa, a yield ratio of at least 0.85, a total elongation of at least 6.0%, a hole expansion ratio of at least 30%, and a bending angle of at least 70° with a thickness of 1 mm. It has a microstructure consisting of 40 to 85 vol% tempered martensite, 60 to 15 vol% primary martensite, less than 1 vol% retained austenite, and essentially no cementite or other carbides.

2. The hot-rolled steel strip according to claim 1, comprising (in weight percent): ●Al:0.030–1.00。 3. The hot-rolled steel strip according to claim 1 or 2, having a yield strength of at least 1100 MPa, an ultimate tensile strength of at least 1200 MPa, a yield ratio of at least 0.85, a total elongation of at least 8.5%, a hole expansion ratio of at least 50%, a bending angle of at least 80° at a thickness of 1 mm, and a Charpy impact toughness of at least 40 J at -40°C and at least 100 J at room temperature; It has a microstructure consisting of 40 to 85 vol% tempered martensite, 60 to 15 vol% primary martensite, less than 1 vol% retained austenite, and essentially no cementite or other carbides.

4. The hot-rolled steel strip according to claim 1 or 2, comprising one or more of the following elements in amounts (by weight%): ●V:0.010-0.10； ●Nb: 0.010-0.10; ●Ti: 0.010-0.10; ●Mo: 0.050-0.50; ●Cr:0.10-1.50； ●Cu: 0.030-1.00; ●Ni: 0.020-0.50; ●N:0.0005-0.0100； ●S: At most 0.

002.

5. The hot-rolled steel strip according to claim 1 or 2, wherein the microstructure consists of at least 55% by volume tempered martensite and at most 45% by volume primary martensite.

6. The hot-rolled steel strip according to claim 1 or 2, wherein the steel contains at least 1.65% by weight of Mn and at most 2.50% by weight of Mn.

7. The hot-rolled steel strip according to claim 1 or 2, wherein the sum of Al and Si is at least 1.00% by weight.

8. The hot-rolled steel strip according to claim 1 or 2 is provided with a metallic coating, which can be obtained by hot-dip coating.

9. The hot-rolled steel strip according to claim 1 or 2, provided with a Zn-layer or a Zn-based alloy layer or an Al-based alloy layer, which can be obtained by hot-dip coating.

10. The hot-rolled steel strip according to claim 8, wherein the zinc alloy coating comprises the following: 0.3-4.0 wt% Mg and 0.05-6.0 wt% Al, and optionally up to 0.2 wt% of one or more additional elements, as well as unavoidable impurities and the balance being zinc.

11. The hot-rolled steel strip according to claim 8, wherein the zinc alloy coating comprises the following: 0.3-4.0% Mg and 0.1-5.0% Al, and optionally up to 0.2% wt% of one or more additional elements, as well as unavoidable impurities and the balance being zinc.

12. A method for preparing hot-rolled steel strip with ultra-high strength, excellent ductility and rollability, comprising the following steps: ● Cast molten steel into thick or thin slabs having the following composition (in weight percent): ●C:0.10–0.30； ●Si: 0.50–1.50; ●Al:0.030–1.00； ●Mn: 1.00–3.00; ●(Si+Al): 0.80-2.50; and optionally any one or more of the following alloying elements: ●Nb: less than 0.10; ●Ti: less than 0.10; ●Mo: less than 0.50; ●Cr: less than 1.50; ●Cu: less than 1.00; ●Ni: less than 0.50; ●B: Less than 0.0030 (30ppm); Inevitably, it also includes ●N: Less than 0.0100 (100ppm); ●S: Less than 0.005; ●P: Less than 0.020; The balance consists of Fe and other unavoidable impurities produced during the ironmaking and steelmaking processes; ●Heat or reheat the slab; ● The slab is hot-rolled into hot-rolled strip, and then... The thick slab is rough-rolled to intermediate specifications to disrupt the microstructure of the cast state, and then hot-rolled into hot-rolled strip, or o Hot rolling thin slabs into hot-rolled strips via direct rolling ●The hot finishing temperature (FRT) is greater than the Ar3 temperature of the steel, where Ar3 is the temperature at which the austenite-to-ferrite transformation begins during cooling; ● The hot-rolled strip is cooled more quickly on the output roller conveyor at a cooling rate of over 20°C / s; ●Then, the hot-rolled and cooled steel strip is wound at a temperature in the range of (Ms-50)℃ to (Ms-160)℃, where Ms is the martensite initiation temperature of the steel. ● Allow the wound hot-rolled strip to be further cooled to ambient temperature; ● Pickled hot-rolled steel strip.

13. The method of claim 12, wherein the slab is heated or reheated at a temperature of 1100°C or higher for 30 minutes or more.

14. The method of claim 12, wherein the intermediate dimension is in the range of 35-45 mm.

15. The method of claim 12, wherein the slab comprises one or more of the following elements in weight percent: ●Nb: 0.010-0.10; ●Ti: 0.010-0.10; ●Mo: 0.050-0.50; ●Cr:0.10-1.50； ●Cu: 0.030-1.00; ●Ni: 0.020-0.50; ●N:0.0005-0.0100。 16. The method according to any one of claims 12 to 15, wherein the microstructure consists of at least 55 vol% tempered martensite and at most 45 vol% primary martensite.

17. The method according to any one of claims 12 to 15, wherein - The steel contains at least 1.65% by weight of Mn and at most 2.50% by weight of Mn, and / or wherein - The sum of Al and Si is at least 1.00% by weight.

18. The method according to any one of claims 12 to 15, wherein a metallic coating is provided, which can be obtained by hot-dip coating.

19. The method according to any one of claims 12 to 15, wherein a Zn-layer or a Zn-based alloy layer or an Al-based alloy layer is provided, which can be obtained by hot-dip coating.

20. The method of claim 19, wherein the zinc alloy coating comprises: 0.3-4.0 wt% Mg and 0.05-6.0 wt% Al, and optionally up to 0.2 wt% of one or more additional elements, as well as unavoidable impurities and the balance being zinc.

21. The method of claim 19, wherein the zinc alloy coating comprises: 0.3-4.0% Mg and 0.1-5.0% Al, and optionally up to 0.2% wt% of one or more additional elements, as well as unavoidable impurities and the balance being zinc.

22. Use of hot-rolled steel according to any one of claims 1 to 11, for parts used in transportation or engineering applications.

23. The use of the hot-rolled steel according to claim 22, for vehicle chassis or suspension parts.

24. The use of the hot-rolled steel according to claim 22, for use in lower control arms, frame longitudinal beams, bumpers or battery housings or frames of heavy-duty trucks, or crane gantry.

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