HOT-ROLLED ULTRA-HIGH STRENGTH DUCTILE STEEL WITH HIGH FLANGING CAPACITY, METHOD OF MANUFACTURING SAID HOT-ROLLED STEEL AND USE THEREOF

Abstract

The present invention relates to a hot-rolled (HR) steel strip with high flangeability at ultra-high strength levels with high overall values ​​of elongation, pliability or bending capacity and toughness, to a method for manufacturing said hot-rolled steel and its use.

Description

HOT-ROLLED ULTRA-HIGH STRENGTH DUCTILE STEEL WITH HIGH FLANGING CAPACITY, METHOD OF MANUFACTURING SAID HOT-ROLLED STEEL AND USE THEREOF FIELD OF INVENTION The present invention relates to a hot-rolled (HR) steel strip with high flangeability at ultra-high strength levels with high values ​​of total elongation, pliability or bending capacity and toughness, to a method of manufacturing said hot-rolled steel and use thereof. BACKGROUND OF THE INVENTION It is well known that as the strength of hot-rolled (HR) steel increases, its formability decreases. A major application area for HR steels in transportation and automotive applications is chassis and suspension (C&S), such as in lower control arms. Other areas include truck frame rails, bumper beams, and battery boxes for electric vehicles. The typical thickness of HR steels used for these applications is less than 4.5 mm. Thicker gauge HR steel strips, up to 12 mm, can be used in engineering applications, such as crane booms, or in transportation applications for heavy-duty truck structures. From a weight reduction perspective, it is imperative to use higher-strength steels for the aforementioned applications in order to reduce the gauge of the steel strip. Therefore, ultra-high-strength steels (UHSS) with a maximum tensile strength (Rm) typically exceeding 1000 MPa would be useful for this purpose. These applications of high-strength steels demand mechanical properties that are difficult to reconcile. In addition to high strength, the steel must also have good formability to manufacture the component by cold forming, as this is an energy-efficient manufacturing method compared to hot forming. Furthermore, good impact resistance or energy absorption capacity is also required for applications such as bumper beams, battery housings, crane booms, or frame rails. Good weldability, typically characterized by a low-carbon equivalent of steel, is also necessary for assembling the components. However, as the tensile strength of steels increases, formability parameters decrease. Formability is a generic term for sheet steel that encompasses a combination of the material's behavior during various mechanical operations, such as drawing, bending, drawing, and flanging. Depending on the component's geometry, any one or a combination of two or more material attributes is important during sheet metal forming. For typical automotive C&S parts, the stretchability of flanges is also important. This type of formability requires high hole expansion capacity (HEC) and good total elongation. For manufacturing frame rails, bumper beams, or battery housings, which are typically produced by rolling, bendability or pliability is important.The manufacture of crane arms also requires good HEC, foldability or bending capacity and elongation or extension. Achieving high formability and high impact hardness values ​​in steels at ultra-high strength levels is a challenge. BRIEF DESCRIPTION OF THE INVENTION OBJECTIVE OF THE INVENTION It is an objective of the invention to provide a hot-rolled steel strip that has ultra-high strength combined with high flanging capacity, good elongation, foldability or bending capacity, and impact resistance. It is also an objective of the invention to provide a hot-rolled steel strip that has excellent weldability. It is also an objective of the invention to provide a method for producing such steels. BRIEF DESCRIPTION OF THE FIGURES The invention will now be explained using the following non-limiting figures. Figure 1 shows a schematic of the thermo-mechanical processing of this invention. Figure 2a shows a schematic representation of a hot rolling mill for processing thick cast slabs, and Figure 2b shows a thin slab casting installation with a direct rolling mill. Figures 3 and 4 show the geometry and definitions for bending samples and Charpy samples. DETAILED DESCRIPTION OF THE INVENTION One or more of the objectives are achieved with hot-rolled steel according to claim 1. Preferred embodiments are described in any of the dependent claims. According to a second aspect, the invention also materializes in the method according to claim 10 for producing the steels according to the invention. In accordance with a third aspect, the invention also materializes in the use of hot-rolled steel for the production of a part for transport or engineering applications. The steel according to the invention contains carbon, silicon, aluminum, and manganese as essential elements. The ranges of the contents of these alloying elements (in % by weight) in the steel are as follows: C: 0.10 to 0.30; Yes: 0.50 to 1.50; Al: 0.010 to 1.00; Mn: 1.00 to 3.00; where (Si + Al) > 0.80; and optionally one or more of the following alloying elements: V: less than 0.10; Nb: less than 0.10; zcopnn / zznz / E / YiAi 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 (30 ppm); inevitably also includes N: less than 0.0100 (100 ppm). S: less than 0.005; P: less than 0.020; The remainder is iron and other unavoidable impurities resulting from the iron and steel manufacturing process. Note that all composition percentages are given as % by weight, unless otherwise stated. Carbon is present in steel in an amount of 0.10 to 0.30%, preferably 0.10 to 0.26%, and more preferably 0.10 to 0.23%. Carbon, which causes strong solid solution hardening in iron, is added primarily for strength and hardenability. Carbon ensures that during cooling of the draw-off table after hot rolling, austenite does not transform into ferrite and / or pearlite above a critical cooling rate (20°C / s). Less than 0.10% carbon will not yield the desired tensile strength (Rm) of 1000 MPa or more, preferably 1200 MPa or more, and if carbon exceeds 0.30%, the weldability of the formed parts may be impaired. Weldability is also improved by a low carbon equivalent. A suitable minimum carbon content is 0.16%. Silicon is added in amounts less than 1.507° to increase strength by strengthening the substitutional solid solution in the iron lattice. Another important effect of silicon in steels is that it inhibits carbide precipitation (cementite and other carbides). As a result, the martensite phase, when quenched, will not form harmful iron carbides in its matrix. When silicon content is below 0.507°, the strengthening and carbide-suppressing effects are insufficient to achieve the intended benefits. On the other hand, when silicon content is above 1.507°, excessive oxide formation can occur during thermomechanical processing (slab reheating, hot rolling, coiling, etc.) of the steel. This oxide inclusion is detrimental to hot rolling, pickling, coating, and the overall surface appearance.Furthermore, rolling forces during hot rolling increase and the steel becomes brittle when the Si content exceeds 1.50 7° to a level that makes the steel very difficult to hot roll. Therefore, the amount of Si according to the invention is normally greater than 0.50 7° and less than 1.50 7°, preferably in the range of 0.60 7° to 1.30 7°, and more preferably in the range of 0.70 to 1.10 7°. Aluminum behaves similarly to silicon in steel according to the invention. It acts as a solid solution strengthening element in steel when deliberately added. It also slows the kinetics of carbide precipitation during martensite tempering. When the aluminum content (Al) is less than 0.030%, the strengthening effects and the suppression of carbide formation are negligible. Aluminum concentrations (µCpnn / µZn / µE / µYiAi) below 0.030% are considered residues from the deoxidation stage during steelmaking, and therefore a minimum value of 0.030% is preferred. On the other hand, when the aluminum content is above 1.00%, excessive oxide formation can occur during thermomechanical processing (reheating of slabs, hot rolling, winding, etc.) of the steel.Furthermore, aluminum (Al) increases the ferrite-to-austenite transformation temperature, necessitating higher hot rolling temperatures to complete the process in the austenitic phase, as intercritical ferrite appears at lower temperatures. Increased oxidation can occur at these higher temperatures. This oxide inclusion is detrimental to hot rolling, pickling, coating, and overall surface appearance. Additionally, rolling forces during hot rolling increase when Al exceeds 1.00%, especially when combined with silicon (Si) levels that make the steel very brittle and difficult to hot roll. Moreover, Al content above 1.00% can also promote ferrite formation during cooling on the drawtable by reducing the incubation time required for ferrite formation during continuous cooling.Ferrite is a detrimental phase to this invention because it introduces brittle interfaces with fresh and quenched martensite. These interfaces act as nucleation sites for the initiation of deformation damage, reducing the formability, elongation, and impact strength of the steel. Therefore, the aluminum in the present invention is present in an amount of 0.010 to 1.00%, preferably 0.030 to 1.00%, preferably 0.20 to 0.80%, and most preferably in the range of 0.30 to 0.80%. While either Si or Al individually can produce the effects of solid solution strengthening and carbide precipitation inhibition during martensite tempering, when both elements are present, their synergistic effect is similar to their individual effects. Therefore, the total (Si + Al) content in this invention should be at least 0.80%, preferably at least 1.00%, to achieve the desired carbide suppression effects and strength levels. When both Al and Si are present, there are several advantages that facilitate steel processing, particularly during hot rolling, pickling, and coating. The presence of some Al with Si alters the oxide characteristics in the flakes during high-temperature processing. This facilitates the pickling of the flakes after hot rolling. As will be described later, the initial martensite that forms in hot-rolled steel during the winding stage is quenched during coil cooling in this invention. The suppression of carbide formation during this self-quenching (coil cooling) of the steel, due to the individual or synergistic effects of Si and Al, is important to the invention. As a result, the martensite reduces its dislocation density without forming carbides. Carbides are detrimental to the elongation, formability, and impact strength of steel because they are brittle and act as nucleation sites for the initiation of damage during deformation. Manganese is present in an amount of 1.00 to 3.00%. The main effect of Mn is to increase strength and toughness. At levels below 1.00% by weight, the desired effects are not achieved, while at amounts above 3.00%, casting problems and segregation will occur. Furthermore, the deformation mechanism in steel can change to transformation-induced plasticity (TRIP) due to the stabilization of austenite by Mn at room temperature, which does not lead to achieving a good combination of all the desired mechanical properties in the product (i.e., impact resistance, formability, strength). Preferably, the Mn content is in the range of 1.20 to 2.70%. In one form, Mn ranges from 1.40 to 2.60%, preferably from 1.50 to 2.50%, and more preferably from 1.60 to 2.50%. In one form, a minimum suitable amount of Mn is 1.65% and a maximum suitable amount of Mn would be 1.95%. Apart from the effects of the essential alloying elements in inventive steel—namely, C, Si, Al, and Mn—described above, another collective effect of these alloying elements is to increase the steel's hardenability. They help prevent the formation of pearlite or ferrite phases during cooling after austenitizing. This characteristic allows the steel to avoid these phases above a certain cooling rate during the draw-out cooling process after hot rolling and before winding. The presence of these softer (ferrite) and non-homogeneous (pearlite) phases is detrimental to obtaining good mechanical and formability properties in the final product because they promote brittle and inconsistent interfaces in the microstructure. Optionally, one or more microalloying elements, selected from group V, Nb, Ti, and Mo, are present. These microalloying elements increase strength through precipitation hardening of their carbides, nitrides, or carbonitrides. They also improve the weldability of the steel. Chromium, another optional element of this invention, also increases the hardenability of steel. Copper, when present, increases the strength of steel through both solid solution strengthening and precipitation hardening via copper precipitates. Nickel increases impact strength and counteracts any hot shortening that might occur during hot working of the steel due to the presence of copper. If present as alloying elements, the preferred additions of these optional alloying elements (in %) are: V: 0.010 to 0.10 Nb: 0.010 aO.10 Ti: 0.010 to 0.10 Mo: 0.050 to 0.50 Cr: 0.10 to 1.50 Cu: 0.030 to 1.00 Ni: 0.020 to 0.50 Nitrogen, sulfur, and phosphorus are residual elements present in steel as a result of the steelmaking and refining processes. Their quantities 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 has a carbon-like effect and contributes to strength through the formation of carbonitrides of the microalloying elements. zcopnn / zznz / E / YiAi The optional alloying elements and the elements nitrogen, sulfur, and phosphorus can vary independently of each other within the specified ranges. They were found to have an additive, not a synergistic, effect on the steels according to the invention. According to a second aspect, the invention also embodies a manufacturing process for a hot-rolled strip that achieves the desired microstructure in the final product. Consequently, the method according to the invention is a method for producing hot-rolled steel with the chemistry discussed above. The steelmaking process comprises the following steps: mold molten steel into slabs; reheat the slabs, preferably to a temperature of 1100°C or more and preferably for a time of 30 minutes or more; rough rolling of the slab to an intermediate gauge, typically in the range of 35 to 45 mm, to break the freshly cast structure; hot rolling of the steel into a strip, preferably with a final hot rolling temperature (FRT) above the Ar3 temperature of the steel, wherein Ar3 is the temperature at which the transformation of austenite to ferrite begins during cooling; accelerated cooling of the hot-rolled strip on the output table with a cooling rate greater than 20°C / s; roll the hot-rolled steel strip cooled to a temperature in the range of (Ms-50)°C to (Ms-160)°C, where Ms is the initial martensite temperature (in °C) of the steel; Cooling of steel coils at room temperature; pickling of hot-rolled steel strip; Optionally, coat the hot-rolled strip with Zn or a Zn-based alloy or an Al-based alloy or any other coating; To avoid misunderstandings, Ms is expressed in °C. Preferably, the FRT is above Ar3 + 50°C. Figure 1 shows a schematic representation of the hot rolling and room-temperature cooling process superimposed on a schematic diagram of the continuous cooling transformation (CCT). The room temperature is defined as approximately 20°C. Reheating is preferably carried out for a period of 60 minutes or more, particularly when the hot rolling process according to the invention is performed on a conventional slab-based hot strip mill. The invention is not limited by the casting method. The steel can be cast as a conventional thick slab with a casting thickness of between 150 and 350 mm, and typically from 225 to 250 mm, as well as a thin slab with a casting thickness of between 50 and 150 mm in a direct strip mill. Schematic examples of a process involving a conventional hot strip mill and a thin slab direct / casting mill are shown in Figures 2a and 2b, respectively.For conventional casting of thick slabs, it is necessary to reheat the slab from room temperature (typically, thin-cast slabs have been cooled from casting temperature to room temperature in a slab storage area) to homogenize the slab composition. Therefore, the reheating temperature must be above 1100°C to dissolve any precipitates when microalloying elements are present and to bring the slab to a temperature such that the final hot rolling on the finishing mill can still be performed at FRT>Ar3. This often requires a reheating temperature (slab) of between 1150 and approximately 1250°C.For thin slab casting, the cast slab undergoes a homogenization treatment in a homogenization furnace immediately after casting, where the homogenization temperature must be above 1100°C, and is typically approximately 1125 to 1150°C. This also prevents the formation of precipitates when microalloying elements are present and brings the thin slab to a temperature such that the final hot rolling on the finishing mill can still be performed at FRT>Ar3. According to the invention, the reheating or homogenization time for the thin slab casting route is preferably 30 minutes or more. Hot rolling of steel must be carried out in the austenitic phase to ensure that no ferrite is present in the final microstructure. Another purpose of hot rolling in the austenitic phase is to reduce hot-rolled strength; therefore, the final rolling temperature (FRT) is preferably maintained at a temperature that is at least 50°C higher than the Ar3 of the steel. After hot rolling, the steel strip is cooled on a drawtable. Here, the steel must be cooled at a rate exceeding the critical cooling rate to prevent any unwanted phase transformations from austenite. In particular, ferrite and pearlite must not form because they are detrimental to the mechanical and formability properties of the final product. Therefore, the critical cooling rate (CCR) must exceed the critical cooling rate to prevent ferrite and pearlite formation. There is no single critical maximum CCR that guarantees transformation from austenite, provided the aforementioned critical cooling rate is exceeded throughout the strip's thickness.An unnecessarily high ROT-CR can affect the flatness of the strip after cooling and cause control problems in stopping at the correct cooling stop temperature. Therefore, a suitable maximum ROT-CR is approximately 300°C / s, preferably approximately 200°C / s, and more preferably around 150°C / s. A practical ROT-CR range is 20 to 100°C / s, as this can be achieved by air cooling, laminar cooling, or waterjet cooling, depending on the strip thickness. For practical reasons, the exit table cooling rate (ROT-CR) is defined as the average cooling rate of the strip surface. Next, the hot-rolled steel strip is coiled at a temperature below the steel's melting point (Ms), in the range of (Ms-50)°C to (Ms-160)°C. Coiling below Ms ensures that subsequent cooling of the coil begins with a mixture of martensite and austenite phases, with an initial martensite content in the range of 40 to 85% by volume. If the initial martensite content exceeds this amount, or in other words, if the cooling temperature (CT) is below Ms-160°C, the required tempering effect on the initial martensite is not achieved, resulting in insufficient ductility, formability, and impact toughness in the steel. This is due to the reduced cooling time and the temperature being too low for effective tempering.If the initial martensite content zcopnn / zznz / E / YiAi is less than 40% by volume, then excessive tempering of the martensite may occur, and the product is not an ultra-high-strength steel in the context of this invention. During coil cooling, the initial martensite undergoes continuous tempering. Simultaneously, as the steel cools in the coil, new fresh martensite forms. Due to the presence of silicon and aluminum in the steel, carbides do not form in the quenched martensite. Furthermore, due to carbon partitioning from martensite to austenite, very small amounts of austenite may 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. Once the steel has cooled to room temperature, the oxides (scale) on hot-rolled steels are removed by pickling in an acidic solution (e.g., HCl) at tempered temperatures (80 to 120°C) or by a combination of pickling and mechanical surface brushing. This step is necessary to make the steel surface suitable for direct use as uncoated HR steel or to prepare it for the coating process, when optionally required for corrosion resistance. Optionally, the HR steel strip can be coated, for example, by hot-dip coating or electroplating, with Zn or a Zn-based alloy, or an Al-based alloy, or any other coating technique to provide good corrosion resistance in service. The above process results in the desired microstructure to obtain the desired mechanical properties. The invention also embodies a steel article manufactured according to the above process and steel chemistry containing the following microstructure (in % by volume): quenched martensite (initial martensite during winding): 40 to 85%, preferably at least 50%, more preferably at least 60%; fresh martensite (martensite that forms during the cooling of the coil after winding): 15 to 60%, preferably no more than 50%, more preferably no more than 40%; Retained austenite: maximum 1% by volume including 0% by volume Cementite or any other metallic carbide: 0% by volume The chemistry, process, and microstructure of the steel result, according to the invention, in the following mechanical and formability properties. Elastic limit (Rp): at least 1100 MPa Maximum tensile strength (Rm): at least 1200 MPa Performance ratio (Rp / Rm): at least 0.85 Total elongation: at least 6.0% JIS5 Hole expansion capacity: at least 30% Bending angle @ 1 mm thickness: at least 70° Preferably, the Charpy impact resistance is at least 40 joules at -40°C and at least 100 joules at room temperature. zcopnn / zznz / E / YiAi The chemistry, process, and microstructure of the steel result, according to the invention, preferably in the following mechanical and formability properties. Elastic limit (Rp): at least 1100 MPa Maximum tensile strength (Rm): at least 1200 MPa Performance ratio (Rp / Rm): at least 0.85 Total elongation: at least 8.5% JIS5 Hole expansion capacity: at least 50% Bend angle @ 1 mm thickness: at least 80° Charpy impact toughness: at least 40 joules at -40°C and at least 100 joules at room temperature. The strength values ​​of steel result primarily from the presence of its hard constituents in the microstructure. Martensite is a strong phase in steel, and due to the low-temperature tempering below the Ms during coil cooling, the martensite retains much of its strength. Therefore, both fresh and quenched martensite in this invention contribute to achieving the ultra-high strength values. 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 below 1% by volume, as it is detrimental to impact toughness due to its low stability. Retained austenite results from the partitioning of carbon from martensite to austenite during coil cooling. Carbon increases the stability of austenite by lowering the Ms temperature. However, in the present invention, retained austenite is deliberately avoided because it is difficult to control its mechanical stability during different deformation and forming processes. Retained austenite should have very high mechanical stability due to its beneficial effect of increasing elongation (i.e., stretchability) and impact strength. It requires very high carbon saturation along with a thin-film morphology to enhance these properties. Achieving high carbon supersaturation during a continuous, low-temperature cooling process, such as coil cooling, is very difficult.When the mechanical stability of austenite is low, it rapidly transforms into martensite, creating brittle interfaces with the matrix phase that affect total elongation. Retained austenite with low mechanical stability transforms even more rapidly under dynamic loading processes such as impact, reducing impact toughness. Therefore, in this invention, a more homogeneous microstructure has been created using quenched and fresh martensite, without the presence of a large amount of retained austenite. In other words, the presence of retained austenite is deliberately avoided, and its maximum amount has been limited to 1% by volume. Another motivation for avoiding or minimizing the retained austenite phase in this invention is to reduce the susceptibility to liquid metal embrittlement (LME) during welding of Zn or Zn-alloyed steel. It is known that Zn-coated steels or Zn alloys with a retained austenite phase in their microstructures are more prone to LME during welding. zcopnn / zznz / E / YiAi This has been achieved by using low temperature winding in the range of (Ms-50)°C to (Ms160)°C, which is a temperature range in which substantial carbon partitioning is not expected to stabilize large amounts of austenite. The yield strength (Rp), ultimate tensile strength (Rm), and total elongation were determined from quasi-static tensile tests (strain rate 3 x 10⁴ s⁻¹) at room temperature using JIS No. 5 specimen geometry. Tensile tests were performed parallel to the rolling direction according to EN 10002-1 / 150 6892-1. The tensile specimen geometry consisted of a gauge length of 50 mm in the rolling direction, a width of 25 mm, and a thickness of 3.2 mm. The steel's resistance to a compensated strain of 0.2% is measured as the yield strength (Rp or YS). The ratio between the yield strength and the ultimate tensile strength (Rp / Rm) is expressed as the yield ratio. The bending capacity was determined by three-point bending tests following the VDA 238-100 standard on specimens 3.2 mm thick, measuring 40 mm x 30 mm in both the longitudinal and transverse directions. The bending axis was along the 30 mm dimension, and the bending radius was 0.4 mm. The bending angles obtained from strips of different thicknesses (2.8, 3.2, and 3.5 mm thick, respectively) were converted to the corresponding angles at 1.0 mm thickness using the following formula: bending angle at 1.0 mm thickness = measured angle x square root of the actual thickness in mm. From these converted bending angles, for a specific heat treatment condition, the lowest value from the longitudinal and transverse samples was taken to claim the ranges of this invention. The flangeability of the steel, or the hole expansion capacity (HEC), was determined by hole expansion tests. Samples measuring 90 mm x 90 mm x 3.2 mm were cut from the coiled steel. A 10 mm diameter hole was drilled in the center of the samples, and hole expansion tests were performed according to ISO / TS 16630:2003(E). The HEC value was determined using the formula: HEC = (expansion of the initial hole diameter / initial diameter) x 100%. Charpy impact toughness was measured using full-size Charpy V-notch (CVN) samples (55 mm x 10 mm x 10 mm) according to ASTM A370. Tests were performed in both directions of the sheet by machining the V-notch parallel and perpendicular to the rolling direction. For all the mechanical tests above, at least three specimens were tested for each condition and the average values ​​are reported. The microstructure was analyzed using a combination of techniques: optical microscopy, X-ray diffraction (XRD), scanning electron microscopy (SEM), and dilatometry. Dilatometry tests on (I x w x t) 10 mm x 5 mm x 3.2 mm specimens were performed by heating the specimens at a rate of 10°C / s to 950°C, holding them for 2 minutes, and then cooling them to room temperature at a rate of 100°C / s (turn-off for Ms) or 0.3°C / s (slow cooling for Ar3). The temperatures of Ms and Ar3 were determined from the dilatometry data.The amount of initial martensite, (i.e. also the quenched martensite that is quenched after cooling of the coil), after rolling the steel was determined using the Koistinen-Marburger formula given in the following bibliography: “A general equation prescribing the extent of the austenite-martensite transformation in pure iron-carbon alloys and plain carbon steels” by DP Koistinen, RE Marburger, Acta Metallurgica, vol. 7, 1959, pp. 59 and 60. / = 100 {1 - exp(—(1.10 x 10“2(Ms- CT))} Where Ms is the initial martensite temperature (in °C) and CT is the winding temperature (in °C), so (Ms-CT) reflects the undercooling below Ms at the beginning of coil cooling and is therefore a measure of the amount of initial martensite. The amount of retained austenite was determined by XRD at a percentage thickness of the samples. XRD patterns were recorded in the 45–165° (2Θ) range using a standard Panalytical Xpert PRO powder diffractometer (Co₂Kα radiation). Quantitative determination of phase proportions was performed using Rietveld analysis with the Bruker Topas software package for Rietveld refinement. The amounts of carbides, ferrite, pearlite, and bainite in the microstructure were determined by analyzing high-resolution SEM images. Subtracting the initial martensite and other phase fractions (retained austenite, carbide, and other determined phases) from the total amount yielded the fresh martensite fractions. The composition of the zinc or zinc alloy coating is not limited. Although the coating can be applied in various ways, hot-dip galvanizing using a standard Gl coating bath is preferred. The Zn-based coating may comprise a Zn alloy containing Al as an alloying element. A preferred zinc bath composition contains 0.10 to 0.35 wt% Al, the remainder being zinc and unavoidable impurities. Other Zn coatings may also be applied. One example comprises a zinc alloy coating according to WO 2008 / 102009, in particular a zinc alloy coating layer consisting of 0.3 to 4.0 wt% Mg and 0.05 to 6.0 wt% Al, preferably 0.1 to 5.0 wt% Al, and optionally in most cases 0.2 wt% of one or more additional elements along with unavoidable impurities, the remainder being zinc. A preferred Zn bath comprising Mg and Al as the principal alloying elements has the composition: 0.5 to 3.8 wt% Al, 0.5 to 3.0 wt% Mg, optionally a maximum of 0.2 wt% of one or more additional elements; the remainder being zinc and unavoidable impurities. An additional element typically added in a small amount of less than 0.2% by weight could be selected from the group comprising Pb, Sb, Ti, Ca, Mn, Sn, La, Ce, Cr, Ni, Zr and B1. Pb, Sn, B1 and Sb are usually added to form flake.Preferably, the total amount of additional elements in the zinc alloy is at most 0.2% by weight. These small amounts of an additional element do not significantly alter the properties of the coating or the bath for typical applications. Preferably, when one or more additional elements are present in the coating, each is present in an amount < 0.02% by weight, and preferably each in an amount < 0.01% by weight. Typically, the additional elements are added only to prevent slag formation in the molten zinc alloy bath for hot-dip galvanizing, or to prevent the formation of flake in the coating layer. In another embodiment, the metallic coating comprises a layer of (commercially pure) aluminum or an aluminum alloy layer. A typical metal bath for hot-dip coating of such an aluminum layer comprises aluminum alloyed with silicon, for example, aluminum alloyed with 8 to 11 wt% silicon and at most 4 wt% iron, optionally at most 0.2 wt% of one or more additional elements such as calcium, unavoidable impurities, the remainder being aluminum. The silicon is present to prevent the formation of a thick iron-metal intermetallic layer that reduces adhesion and formability. The iron is preferably present in amounts between 1 and 4%, more preferably at least 2%. EXAMPLES Steel ingots of seven inventive chemistries, A, B, D, and H, and a comparative steel C measuring 200 mm x 100 mm x 100 mm, were cast by melting the charges in a vacuum induction furnace. The chemical compositions of these steels are given in Table 1. Steels A and B and D through H contain C, Si, and Al within the limits defined by the invention, while the comparative steel has Al and Si outside the limits defined by the invention. All ingots were reheated for 1 hour at 1200°C and as-rolled to a thickness of 25 mm. The strips were then reheated to 1200°C for 30 minutes and hot-rolled to their final thicknesses of 2.8, 3.2 mm, 3.5, and 12 mm with FRTs above 900°C, which are in the austenitic phase range for all steels. Ar3 and Ms values ​​for the steels, measured by dilatometry, are also given in Table 1. After hot rolling, the steels were immediately quenched on the test bench or exit at various cooling rates, and then coil quenching simulations were performed in a muffle furnace, cooling to room temperature from different starting CTs. The strips were then pickled to remove oxides using conventional methods. The various processing conditions for the steels are summarized in Table 2. Steels A, B, and C have similar Ms and Ar3 values. The FRT temperature for steel A was 953°C, for steel B 939°C, and for steel C 945°C, all of which are at least 50°C above Ar3. For steels A and B, a slow exit table cooling rate (ROT-CR) of 3°C / s was used, which is outside the lower limit defined by the invention. In addition, two further winding temperatures (200°C and 480°C), also outside the limits defined by this invention, were used for steels A and B with FRT and CT within the limits required by the invention. The CT at 200°C is much lower than (Ms-160)°C, and 480°C is above the Ms of these steels. These conditions were used for comparative purposes.For steel C, which has a chemistry outside the scope of this invention, all processing conditions (FRT, ROT-CR, and CT) were selected within the defined limits of this invention. For steels D through F, one set of processing parameters was within the ranges claimed for this invention (FRT, ROT-CR, and CT); however, for the other set, FRT and ROT-CR were kept equal, but only CT was kept higher for comparison purposes. In this case, a CT of 375°C was used, which was higher than the (Ms-50°C) for steels D through F. In fact, this CT was slightly higher than their Ms temperatures. For steels G and H, all processing parameters were kept within the limits required by the present invention. Specimens were extracted from the final steel strips for various mechanical and microstructural characterizations as described. Hot-rolled steels 12 mm thick were used to prepare Charpy impact samples, while steel strips 2.8, 3.2, and 3.5 mm thick were used for all other characterizations. For various processing conditions, the phase contents in the microstructures are presented in Table 3, the tensile properties are given in Table 4, the results of the bending and HEC tests in Table 5, and the Charpy impact strength is given in Table 6. The following are the abbreviations and symbols that have been used in the tables to present the results of the tensile and bending tests: Rp = yield strength, Rm = ultimate tensile strength, Ajiss = total elongation using JIS5 specimens or test pieces, BA = bend angle, L = longitudinal specimen or test piece where the bending axis is parallel to the rolling direction, T = transverse specimen or test piece where the bending axis is perpendicular to the rolling direction. Table 3 shows that steels A and B achieved microstructures consisting of quenched martensite and fresh martensite of less than 85% by volume and at least 15% by volume, respectively, through the use of FRT and ROT-CR within the defined limits for a CT ranging from 275 to 375°C, which falls within the range required for this invention. Furthermore, the steels do not contain carbide in their microstructures, and the retained austenite content was less than 1% by volume under these processing conditions. Their microstructures did not contain any other phases such as ferrite, bainite, or pearlite. Steels A and B, when subjected to a ROT-CR >20°C / s with FRT at a temperature above 50°C above Ar3, produced microstructures with substantial amounts of bainite and retained austenite, with some fresh martensite forming below Ms when wound at 480°C, which is above the Ms of these steels. No carbides were present in these steels under this winding condition due to their Al and Si content. The high retained austenite content was caused by carbon enrichment in the austenite during the bainitic transformation above Ms and by fresh martensite forming during coil cooling below Ms. These bainitic microstructures with high amounts of retained austenite differ from the microstructures anticipated in this invention. Similarly, when steels A and B are rolled at 200°C, which is much lower than their Ms160°C, the microstructure of the steels also becomes different from that required for this invention. This CT condition has more than 85% by volume of quenched martensite and less than 15% by volume of fresh martensite. On the other hand, steels A and B, when their FRT and CT were within the range required by this invention, but with a slower ROT-CR of 3°C / s (less than 20°C / s), showed considerable amounts of ferrite and pearlite in their microstructures, in addition to a bainite matrix and substantial amounts of retained austenite. The ferrite, pearlite, and bainite formed due to the slow ROT-CR before rolling. zcopnn / zznz / E / YiAi C steel formed a considerable amount of carbides (2.3% by volume) during processing with all parameters within the range required for this invention due to the low amounts of Si and Al in C steel. As a result of the microstructures described above, the properties were obtained as shown in Tables 4 to 6. Steels A and B achieved a tensile strength (Rp) above 1100 MPa and a yield strength (Rm) above 1200 MPa with a yield ratio (RQ) greater than 0.85 and a total elongation (AJIS5) greater than 8.5%. When the heat transfer temperature (CT) is too high (480°C), the minimum Rp and Rm levels sought 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 values ​​at the 480°C temperature also resulted in a yield ratio (RQ) below 0.85. On the other hand, when CT is too low (CT = 200°C), Rp and Rm are above the target values ​​with a high performance ratio but the total elongation is too low (<8.5%).The low total elongation is caused by too high an amount of initial martensite (= tempered martensite > 85% by volume) present in the microstructure and the lack of an annealing effect during coil cooling due to the lower availability of time and a temperature that is too low for effective annealing to take place. With a slow ROT-CR of 3°C / s, the Rp and Rm values ​​for steels A and B are less than 1100 MPa and 1200 MPa, respectively, due to the formation of the softer phases of bainite, ferrite, pearlite, and retained austenite. Furthermore, the yield ratio is less than 0.85, although the total elongation is high. Due to the presence of carbides in its microstructure, resulting from the absence of carbide-suppressing elements Si and Al, C steel exhibited low values ​​for Rp, Rm, yield ratio, and total elongation. Carbides are detrimental to mechanical properties and promote damage during deformation. Consequently, C steel has shown low tensile strength. Similar to the tensile properties, the bending capacity and HEC are also high in steels A and B when processed within the defined processing variables (FRT, ROT-CR, and CT) of this invention (Table 5). A minimum bending angle of 80° at 1.0 mm thickness was achieved, and a minimum HEC value of 50% was also obtained. However, when CT is high and above the Ms of the steels (i.e., 480°C), the minimum bending angle and HEC are low and below the target values ​​of 80° at 1.0 mm thickness and 50%, respectively. This is because, due to the multiphase nature of the microstructures of these steels, which contain fresh martensite, bainite, and retained austenite (Table 3), numerous damage initiation sites were present at the interfaces of these phases when deformation was carried out.Martensite, present as fresh martensite and formed due to the TRIP effect of retained austenite, is a stronger phase than bainite and any untransformed retained austenite. Furthermore, under ideal processing conditions with optimal amounts of quenched and fresh martensite in steels A and B, there was a low difference in hardness or strength between these phases, resulting in homogeneous deformation during bending and hole expansion. This led to high HEC and bending values ​​under ideal processing conditions. Furthermore, with a CT value that is too low for steels A and B, below Ms (200°C), the bendability and HEC values ​​are also low due to the excessively high amount of quenched martensite (>85% by volume) present in their microstructures (Table 3). The lack of effective tempering of this initial martensite present at the very beginning of winding resulted in low ductility, which is also reflected in their total elongation values ​​in Table 4, causing the low formability of these steels as measured by bendability and HEC. When a slow ROT-CR (3°C / s) is used, the presence of the softer ferrite and pearlite phases also deteriorated the bending capacity and HEC values ​​of steels A and B, as shown in Table 5. This is due to the brittle interfaces of these softer phases and the harder bainite and martensite phases that are obtained after the transformation of the retained austenite during loading. Steel C achieved very low values ​​for bending capacity and HEC, much lower than the minimum values ​​of 80° at 1.0 mm thickness and 50%, respectively. These poor formability parameters in steel C were caused by the very low Al and Si contents in the steel, which promoted carbide formation (Table 3), even though the processing variable was within the range specified for this invention. The Charpy impact toughness of steel A, processed according to the invention, in cross-sectional specimens (which showed lower values ​​than longitudinal specimens) is greater than 100 J and 40 J when tested at room temperature and -40°C, respectively. The same values ​​with the high CT of 480°C and the low CT of 200°C are significantly lower than the minimum values ​​obtained with the ideal processing route described in this invention. As explained above, these low toughness values ​​are caused by brittle fracture resulting from the spontaneous transformation of retained austenite to martensite, the presence of heterogeneously harder and softer phases, and the effects of insufficient tempering. Furthermore, the presence of carbides in steel C resulted in poor Charpy impact toughness values ​​at both room and cryogenic temperatures. Therefore, as discussed above, these examples illustrate that when steel is designed according to the composition of the invention and processed according to the invention, the steel achieves high tensile, formability, and toughness properties, as intended, due to its microstructural effects. The same good combination of all properties is not achieved when working outside the limits defined in the invention. zcopnn / zznz / E / YiAi Table 1: Chemical compositions of steels in % by weight (I: Inventive, C: Comparison). Steel C Si Al Mn P s N Nb Ms (°C) Ar3 (°C) A 0.20 1.00 0.036 1.82 0.010 0.0010 0.0020 - 428 827 IB 0.20 0.81 0.31 1.82 0.010 0.0012 0.0012 - 435 835 CI 0.21 0.10 0.02 1.85 0.010 0.0012 0.0006 - 430 829 CD 0.20 1.0 0.03 2.31 0.001 0.0001 0.0012 0.001 366 834 IE 0.21 0.81 0.31 2.30 0.001 0.0001 0.0020 0.001 374 854 IF 0.21 0.80 0.03 2.30 0.001 0.0001 0.0008 0.001 374 831 IG 0.18 0.81 0.033 2.39 0.001 0.0001 0.0030 0.001 390 840 IH 0.18 0.80 0.032 2.40 0.001 0.0001 0.0040 0.019 377 835 I zcopnn / zznz / E / YiAi Niobium content in steels A through G is at a residual level. No niobium was added as an alloying element in these steels. Niobium was added as an alloying element in steel H. Steels A and B have a Mn content of approximately 1.8%, and steels D through H have a Mn content of approximately 2.35%, with varying amounts of Si and Nb. The effect of Mn is a reduction in Ms, causing a change in the ratio of quenched to fresh martensite. Although steels D through H exhibit lower values ​​for A JIS5, HEC, and bend angle than steels A and B, they are suitable for their intended purpose. As shown in Table 3, steels D through H achieved microstructures consisting of quenched martensite and fresh martensite of less than 85% by volume and at least 15% by volume, respectively, through the use of FRT and ROT-CR within the defined limits for a CT of 275°C (steels D through F) and 300°C (steels G and H), which fall within the range required for this invention. Furthermore, the steels do not contain carbide in their microstructures, and the retained austenite content was less than 1% by volume under these processing conditions. Their microstructures did not contain any other phases such as ferrite, bainite, or pearlite. Steels D through H also achieved a partial tensile strength (Rp) above 1100 MPa and a maximum tensile strength (Rm) above 1200 MPa with a yield strength (QS) of 0.85 or higher, along with a total elongation (AJIS5) greater than 6% (Table 4). The bending capacity and HEC are also high in steels D through H when processed within the defined processing parameters (FRT, ROTCR, and CT) of this invention (Table 5). A minimum bending angle of 70° at a thickness of 1.0 mm and a minimum HEC value of 30% have been achieved in these steels. However, when the tempering temperature (CT) of steels D to F is high, i.e., 375°C, which is slightly above the melting point (Ms) of the steels, the microstructures of the steels contain some bainite and more than 1% by volume of retained austenite (Table 3), which is not the desired microstructure of this invention. The fresh martensite and quenched martensite contents also fall outside the range defined in this invention. These undesirable microstructures do not lead to the desired ultra-high strengths in these steels for a CT of 375°C. The compressive strength (Rp) values ​​are below 1100 MPa and the ultimate tensile strength (Rm) values ​​are below 1200 MPa, with a yield ratio of less than 0.85 due to the presence of softer bainite and retained austenite phases (Table 4), although good pliability and high coefficient of friction (HEC) values ​​are achieved (Table 5).Therefore, steels that may have a chemical composition within the defined limits of this invention may not achieve all the desired mechanical properties if the processing is not carried out within the defined windows of this invention. Table 2: Processing variables applied to steels. zcopnn / zznz / E / YiAi Steel Final Thickness (mm) FRT (°C) ROT-CR (°C / s) Winding Temperature (°C) A 3.2 / 12* 953 31 275, 300, 325, 350, 375 1 3.2 / 12* 953 31 480, 200 c 3.2 / 12* 953 3 325 c B 3.2 939 34 275, 300, 325, 350, 375 1 3.2 939 34 480, 200 c 3.2 939 3 325 c C 3.2 / 12 945 34 350 c D 2.8 962 45 275 1 2.8 45 375 c E 2.8 965 48 275 1 2.8 48 375 c F 2.8 968 53 275 1 2.8 53 375 c G 3.5 973 47 300 1 H 3.5 976 46 300 1 *A thickness of 12 mm is required for Charpy tests; other thicknesses were used to determine microstructure, tensile properties, bending angle, and HEC. Table 3: Microstructure of steels (B: bainite, P: pearlite, F: ferrite) Maple ROTCR (°C / s) Tde coiled (°C) Martensita Templada (Vol.%) Martensita Fresca (Vol.%) Austenita Retenida (Vol.%) Carbide (Vol.%) Other (Vol.%) A 1 31 275 81.4 18.2 0.4 0 .... 1 2 31 300 75.5 23.8 0.7 0 .... 1 3 31 325 67.8 31.9 0.3 0 .... 1 4 31 350 57.6 42.3 0.1 0 .... 1 5 31 375 44.2 55.7 0.1 0 .... 1 6 31 480 ____ 13.8 11.2 0 B: 75.0 c 7 31 200 91.6 8.3 0.1 0 — c 8 3 325 ____ ---- 5.3 0 B=63.5 F+P=31.2 c B 1 34 275 82.8 16.7 0.5 0 .... 1 2 34 300 77.3 22.4 0.3 0 .... 1 3 34 325 70.2 29.8 0.0 0 .... 1 4 34 350 60.7 39.3 0.0 0 .... 1 5 34 375 48.3 51.7 0.0 0 .... 1 6 34 480 ____ 12.1 10.5 0 B: 77.4 c 7 34 200 92.5 7.5 0.0 0 — c 8 3 325 ____ ---- 6.1 0 B=65.3 F+P=28.6 c C 34 350 39.2 58.5 0.0 2.3 — c D 45 275 63.2 36.2 0.6 0 — 1 45 375 10.2 63.0 5.3 0 B=21.5 c E 48 275 66.3 33.0 0.7 0 ... 1 48 375 11.4 74.0 4.3 0 B=10.3 c F 53 275 66.3 33.2 0.5 0 ... 1 53 375 10.4 75.0 4.5 0 B=10.1 c G 47 300 62.8 36.8 0.4 0 — 1 H 46 300 57.1 42.6 0.3 0 ... 1 zcopnn / zznz / E / YiAi Table 4: Tensile properties of steels. ROT-CR Steel (°C / s) Winding Temperature (°C) Rp (MPa) Rm (MPa) AJIS5 (%) Yield Ratio (-) A 1 31 275 1188 1379 10.1 0.86 I 2 31 300 1176 1351 8.8 0.87 I 3 31 325 1178 1333 9.3 0.88 I 4 31 350 1143 1292 9.3 0.89 I 5 31 375 1110 1217 9.2 0.91 I 6 31 480 693 983 16.8 0.70 c 7 31 200 1215 1435 6.3 0.85 c 8 3 325 670 932 11.7 0.72 c B 1 34 275 1148 1327 9.0 0.87 I 2 34 300 1150 1321 9.1 0.87 I 3 34 325 1148 1300 10.0 0.88 I 4 34 350 1137 1263 10.4 0.90 I 5 34 375 1127 1228 9.6 0.92 I 6 34 480 705 979 15.3 0.72 c 7 34 200 1220 1428 6.4 0.85 I 8 3 325 663 927 12.1 0.72 c C 34 350 781 1079 6.1 0.73 c D 1 45 275 1140 1331 6.9 0.86 I 2 45 375 755 944 13.6 0.80 c E 1 48 275 1132 1334 7.0 0.85 I 2 48 375 767 913 11.5 0.84 c F 1 53 275 1122 1314 6.8 0.85 I 2 53 375 835 1044 7.0 0.80 c G 47 300 1123 1300 8.9 0.86 IH 46 300 1212 1374 8.7 0.88 I Table 5: Foldability or bending capacity and HEC of steels.Acero ROTOR (°C / s) Tde Bobinado (°C) BA-L(°) Medido BA-T @ Medido (°) BA-L @ 1.0 mm (°) BA-T @ 1.0 mm (°) HEC (%) A 1 31 275 45.0 56.7 80.5 101.4 67 I 2 31 300 56.1 61.4 100.4 109.8 72 I 3 31 325 68.5 67.3 122.5 120.4 79 I 4 31 350 77.8 72.9 139.2 130.4 94 I 5 31 375 62.2 79.0 111.3 141.3 97 I 6 31 480 35.5 37.2 63.5 66.5 27 C 7 31 200 37.2 38.3 66.5 68.5 19 c 8 3 325 39.2 49.1 70.1 71.7 21 c B 1 34 275 61.4 59.5 109.8 106.4 52 I 2 34 300 59.3 61.5 106.1 110.0 83 I 3 34 325 48.5 69.0 86.8 123.4 77 I 4 34 350 69.1 69.3 123.6 124.0 76 I 5 34 375 55.7 75.0 99.6 134.2 98 I 6 34 480 33.5 35.1 59.9 62.8 21 c 7 34 200 34.8 36.1 62.3 64.6 18 c 8 3 325 37.2 39.7 66.5 71.0 23 c C 34 350 35.8 37.2 64.0 66.5 15 c D* 1 45 275 93.6 76.2 156.6 127.5 41 I 2 45 375 75.8 107.6 126.8 180.0 57 c E* 1 48 275 53.3 43.4 89.2 72.6 42 I 2 48 375 77.0 93.6 128.8 156.6 38 c F* 1 53 275 57.7 68.0 96.6 113.8 32 I 2 53 375 90.6 105.8 151.6 177.0 35 c G** 47 300 43.9 61.4 82.2 114.8 76 I H** 46 300 45.3 44.1 84.8 82.4 42 I. zcopnn / zznz / E / YiAi Bending angle from A to C measured at 3.2 mm; Bending angle from D to F measured with a thickness of 2.8 mm; Bending angle from G, H measured with a thickness of 3.5 mm. Table 6: Charpy Impact Strength or Toughness of Steels (L = longitudinal specimen, T = transverse specimen) ROT-CR Steel (°C / s) Winding Temperature (°C) Charpy Impact Strength (J) Ambient Temperature - 40°C CLTLTA 1 31 275 136.9 109.6 90.8 42.1 2 31 300 135.8 111.2 93.2 45.7 3 31 325 140.1 115.8 95.1 46.3 4 31 350 143.7 120.0 99.8 49.1 5 31 375 144.0 120.2 100.3 49.5 6 31 480 90.1 75.1 55.1 30.2 7 31 200 76.1 69.2 30.5 26.1 8 3 325 57.8 53.2 23.5 19.1 C 34 350 70.2 75.8 28.3 15.6

Claims

1. A hot-rolled steel strip having ultra-high strength, excellent ductility and flangeability, characterized in that it comprises (in % by weight): C: 0.10 to 0.30; Si: 0.50 to 1.50; Al: 0.010 to 1.00; Mn: 1.00 to 3.00; (Si + Al): 0.80 to 2.50; inevitably also comprising N: less than 0.0100 (100 ppm); S: less than 0.005; P: less than 0.020; the remainder being Fe and other unavoidable impurities resulting from the iron and steelmaking process, having a yield strength of at least 1100 MPa, an ultimate tensile strength of at least 1200 MPa, a yield or creep ratio of at least 0.85, and a total elongation of at least 6.0%, a hole expansion ratio of at least 30% and a bend angle at 1 mm thickness of at least 70°; having a microstructure consisting of 40 to 85% by volume of quenched martensite, 60 to 15% by volume of fresh martensite, less than 1% by volume of retained austenite and substantially no cementite or other carbides.

2. The hot-rolled steel strip according to claim 1, further characterized in that it comprises one or more of the following alloying elements (in % by weight): 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 (30 ppm).

3. The hot-rolled steel strip according to claim 1, further characterized in that it comprises (in % by weight): Al: 0.030 to 1.00; 4. The hot-rolled steel strip according to any of claims 1 to 3, further characterized in that it has a yield strength of at least 1100 MPa, a maximum tensile strength of at least 1200 MPa, a spring ratio of at least 0.85, a total elongation of at least 8.5%, a hole expansion ratio of at least 50%, a bend angle at 1 mm thickness of at least 80°, and a Charpy impact strength of at least 40 J at -40°C and at least 100 J at room temperature; having a microstructure consisting of 40 to 85% by volume of quenched martensite, 60 to 15% by volume of fresh martensite, less than 1% by volume of retained austenite, and substantially no cementite or other carbides.

5. The hot-rolled steel strip according to any of claims 1 to 4, further characterized in that it comprises one or more of the following elements in the following quantities (% by weight): V: 0.010 to 0.10; Nb: 0.010 to 0.10; Ti: 0.010 to 0.10; Mo: 0.050 to 0.50; Cr: 0.10 to 1.50; Cu: 0.030 to 1.00; Ni: 0.020 to 0.50; N: 0.0005 to 0.0100; S: at most 0.

002.

6. The hot-rolled steel strip according to any of the preceding claims, further characterized in that the microstructure consists of at least 55% by volume of quenched martensite and at most 45% by volume of fresh martensite.

7. The hot-rolled steel strip according to any of the preceding claims, further characterized in that the steel comprises at least 1.65% by weight of Mn and at most 2.50% by weight of Mn.

8. The hot-rolled steel strip according to any of the preceding claims, further characterized in that the sum of Al and Si is at least 1.00% by weight.

9. The hot-rolled steel strip according to any of the preceding claims, further characterized in that it is provided with a metallic coating layer, such as 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 9, further characterized in that the zinc alloy coating layer consists of 0.3 to 4.0 wt% of Mg and 0.05 to 6.0 wt% of Al, preferably 0.1 to 5.0 wt% of Al, and optionally a maximum of 0.2 wt% of one or more additional elements together with unavoidable impurities, the remainder being zinc.

11. A method for manufacturing a hot-rolled steel strip having ultra-high strength, excellent ductility, and flangeability, characterized in that it comprises the steps of: casting molten steel into thick or thin slabs having a composition (in wt%) C: 0.10 to 0.30; Si: 0.50 to 1.50; Al: 0.030 to 1.00; (Si + Al): 1.00 to 3.00; (Si + Al): 0.80 to 2.50; inevitably also comprising N: less than 0.0100 (100 ppm); S: less than 0.005; P: less than 0.020; the remainder being Fe and other unavoidable impurities resulting from the iron and steelmaking process; Heat or reheat the slabs, preferably to a temperature of 1100°C or more, and preferably for a time of 30 minutes or more;hot rolling the slab into a hot-rolled strip by or rough rolling a thick slab to an intermediate gauge, generally in the range of 35 to 45 mm, to break the freshly cast structure, followed by finish hot rolling into a hot-rolled strip, or by or hot rolling the thin slab to a hot-rolled strip by direct rolling where the final hot rolling (FRT) temperature is above the Ar3 temperature of the steel, where Ar3 is the temperature at which the transformation of austenite to ferrite begins during cooling; accelerate cooling the hot-rolled strip on the exit table with a cooling rate exceeding 20 °C / s; followed by winding the hot-rolled and cooled steel strip to a temperature in the range of (Ms-50)°C to (Ms-160)°C, where Ms is the initial martensitic temperature of the steel;allow the coiled hot-rolled strip to cool further to room temperature; pickle the hot-rolled steel strip; 12. The method according to claim 11, further characterized in that the slab comprises one or more of the following alloying elements (in % by weight): 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 (30 ppm).

13. The method according to claim 11, further characterized in that the slab comprises one or more of the following elements in the following quantities (in % by weight): V: 0.010 to 0.10; Nb: 0.010 to 0.10; Ti: 0.010 to 0.10; Mo: 0.050 to 0.50; Cr: 0.10 to 1.50; Cu: 0.030 to 1.00; Ni: 0.020 to 0.50; N: 0.0005 to 0.0100.

14. The method according to any of claims 11 to 13, further characterized in that the microstructure consists of at least 55% by volume of quenched martensite and at most 45% by volume of fresh martensite.

15. The method according to any of claims 11 to 14, further characterized in that: the steel comprises 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.

16. The method according to any of claims 11 to 15, further characterized in that it is provided with a metallic coating layer, such as a Zn layer or a Zn-based alloy layer, or an Al-based alloy layer, which can be obtained by hot-dip coating.

17. The method according to claim 16, further characterized in that the zinc alloy coating layer consists of 0.3 to 4.0 wt% of Mg and 0.05 to 6.0 wt% of Al, preferably 0.1 to 5.0 wt% of Al, and optionally a maximum of 0.2 wt% of one or more additional elements together with unavoidable impurities, the remainder being zinc.

18. Use of a hot-rolled steel as claimed in any of claims 1 to 10 for a part for a transport or engineering application.

19. Use of a hot-rolled steel as claimed in claim 18 for a chassis or suspension part of a vehicle, such as a lower control arm, frame rail, bumper beam, battery box, heavy truck frame, or crane arm.