High-flange ultra-high-strength ductile hot-rolled steel, method for manufacturing hot-rolled steel and uses thereof

Abstract

The present invention relates to hot-rolled (HR) strip steel having high flangeability at an ultra-solid level having high total elongation, bendability, and toughness values, and to a method for manufacturing such hot-rolled steel and its use.

Description

Technology Field

[0001] The present invention relates to a hot-rolled strip steel having high flangeability at an ultra-high strength level, along with high total elongation, bendability, and toughness values, a method for manufacturing said hot-rolled steel, and uses thereof. Background Technology

[0002] It is well known that as the strength of hot-rolled (HR) steel increases, its formability decreases. In transportation and automotive applications, the primary applications for hot-rolled steel are chassis and suspension (C&S) components, such as lower control arms. Other areas include truck frame rails, bumper beams, or battery boxes for electric vehicles. The typical thickness of HR steel used in these applications is less than 4.5 mm. Thicker gauge HR strip steel, such as up to 12 mm, can be used in engineering applications like crane booms or in transportation applications for heavy truck frames.

[0003] In terms of weight reduction, it is essential to use higher strength steel for the above application so that the gauge of the steel strip can be reduced. Therefore, ultra-high strength steels (UHSS) with an ultimate tensile strength (Rm) of generally 1000 MPa or more would be useful for this purpose.

[0004] These applications of hot-rolled steel require mechanical properties that are difficult to harmonize. In addition to high strength, the steel must possess good formability for making parts through cold forming, as cold forming is an energy-efficient manufacturing path compared to hot forming. Furthermore, applications such as bumper beams, battery housings, crane booms, or frame rails require excellent impact toughness or energy absorption capabilities. To assemble these parts, excellent weldability, generally characterized by the steel's low carbon equivalent, is also required.

[0005] However, as the tensile strength of steel increases, formability parameters decrease. Formability is a general term for steel sheets considered as a combination of material behaviors during various mechanical operations such as stretching, bending, drawing, and flanging. Depending on the part geometry, some or a combination of two or more material properties are important during sheet metal forming. For typical automotive C&S parts, stretch-flangeability is also important. This type of formability requires high hole expansion capacity (HEC) and excellent total elongation. Bending is important for manufacturing frame rails, bumper beams, or battery housings, which are generally produced by roll forming. The manufacture of crane booms also requires excellent HEC, bendability, and elongation.

[0006] It is difficult to achieve high formability and high impact toughness values ​​for steel at ultra-high strength levels. The problem to be solved

[0007] The object of the present invention is to provide a hot-rolled steel strip having ultra-high strength combined with high flanging ability, excellent elongation, bendability, and impact toughness.

[0008] In addition, the objective of the present invention is to provide a hot-rolled steel strip with excellent weldability.

[0009] In addition, the objective of the present invention is to provide a method for manufacturing such steel. means of solving the problem

[0010] One or more of the above objectives are achieved by the hot-rolled steel according to claim 1. A preferred embodiment is disclosed in any one of the dependent claims.

[0011] According to a second aspect, the present invention is also implemented by a method according to claim 10 for manufacturing steel according to the present invention.

[0012] According to a third aspect, the present invention is also embodied in the use of hot-rolled steel for the production of parts for transportation or engineering applications.

[0013] The steel according to the present invention contains carbon, silicon, aluminum, and manganese as essential elements. The content range (in weight%) of these alloying elements in the steel is as follows:

[0014] C: 0.10 ~ 0.30;

[0015] Si: 0.50 ~ 1.50;

[0016] Al: 0.010 ~ 1.00;

[0017] Mn: 1.00 ~ 3.00;

[0018] Here (Si + Al) ≥ 0.80;

[0019] And optionally, any one or more of the following alloy elements:

[0020] V: Less than 0.10;

[0021] Nb: Less than 0.10;

[0022] Ti: Less than 0.10;

[0023] Mo: Less than 0.50;

[0024] Cr: Less than 1.50;

[0025] Cu: Less than 1.00;

[0026] Ni: Less than 0.50;

[0027] B: Less than 0.0030 (30 ppm);

[0028] Inevitably, it also includes the following:

[0029] N: Less than 0.0100 (100 ppm);

[0030] S: Less than 0.005;

[0031] P: Less than 0.020;

[0032] The remainder consists of Fe and other unavoidable impurities arising from the iron and steelmaking processes. Unless otherwise specified, all composition percentages are provided in weight percent.

[0033] Carbon (C) is present in the 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 prevents the austenite from deforming into ferrite and / or pearlite by exceeding the critical cooling rate (20°C / s) during runout table cooling after hot rolling. C less than 0.10% does not provide the desired Rm level of 1000 MPa or more, preferably 1200 MPa or more, and if C is higher than 0.30%, the weldability of the formed part may be poor. Weldability is also improved by a lower carbon equivalent value. An appropriate minimum value for carbon is 0.16%.

[0034] Silicon (Si) is added in amounts of less than 1.50% to increase strength through substitutional solid solution strengthening of the iron lattice. Another important effect of Si in steel is that it slows down carbide precipitation (cementite and other carbides). Consequently, the martensite phase does not form harmful iron carbides in its matrix when subjected to tempering. If the Si content is less than 0.50%, the strengthening and inhibition effects on carbide formation are insufficient to obtain the intended benefits. On the other hand, if the Si content exceeds 1.50%, excessive oxide formation may occur during the thermomechanical treatment of the steel (slab reheating, hot rolling, coiling, etc.). These oxide scales are detrimental to hot rolling, pickling, coating, and the overall surface appearance. Furthermore, if the Si content exceeds 1.50%, the rolling force increases during hot rolling, and the steel becomes brittle to a level that makes hot rolling very difficult. Accordingly, the amount of Si according to the present invention 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%.

[0035] Aluminum (Al) exhibits behavior comparable to Si in the steel according to the present invention. When intentionally added, it acts as a solid solution strengthening element in the steel. Additionally, it slows down the carbide precipitation rate during the tempering of martensite. When Al is less than 0.030%, the strengthening and suppressing effects on carbide formation are negligible. Since an aluminum value of less than 0.030% is considered a residue of the deoxidation step during steelmaking, a minimum value of 0.030% is desirable. On the other hand, if Al exceeds 1.00%, excessive oxide formation may occur during the thermomechanical treatment of the steel (slab reheating, hot rolling, coiling, etc.). Furthermore, Al increases the ferrite-austenite transformation temperature, requiring the steel to be hot-rolled at higher temperatures to finish the hot rolling process in the austenite phase, as intercritical ferrite appears at lower temperatures. A greater amount of oxidation may occur at higher temperatures. These oxide scales are detrimental to hot rolling, pickling, coating, and the overall surface appearance. Furthermore, when Al exceeds 1.00% in the presence of Si, the rolling force during hot rolling increases to a level that makes the steel very brittle and makes hot rolling difficult. Additionally, an Al content exceeding 1.00% can promote ferrite formation during cooling on a runout table by reducing the incubation time for ferrite formation during continuous cooling. Ferrite is a detrimental phase to the present invention because it introduces brittle interfaces with fresh martensite and tempered martensite. These interfaces act as nucleation sites for the initiation of damage due to deformation, thereby reducing the formability, elongation, and impact toughness of the steel. Accordingly, in the present invention, Al is present in the range of 0.010 to 1.00%, preferably 0.030 to 1.00%, preferably 0.20 to 0.80%, and more preferably 0.30 to 0.80%.

[0036] While Si or Al individually provide solid solution strengthening effects and can suppress carbide precipitation during martensite tempering, the synergistic effect of these elements when both are present is similar to the individual effects. Therefore, in the present invention, the total content of (Si + Al) must be at least 0.80%, preferably at least 1.00%, to achieve the desired carbide suppression effect and desired strength level. The presence of both Al and Si may offer several advantages, particularly in facilitating the treatment of steel during hot rolling, pickling, and coating. The presence of certain amounts of Al and Si alters the oxide properties of the scale during high-temperature treatment. This makes the pickling of the scale easier after hot rolling.

[0037] As described below, the initial martensite formed in the hot-rolled steel during the coiling stage is tempered during coil cooling in the present invention. It is important in the present invention to suppress the formation of carbides during the auto-tempering (coil cooling) of the steel due to the individual or synergistic effects of Si and Al. Consequently, the martensite does not form carbides and only reduces the dislocation density. Since carbides are inherently brittle and act as nucleation sites for damage initiation during deformation, they are detrimental to the elongation, formability, and impact toughness of the steel.

[0038] Manganese (Mn) is present in an amount of 1.00 to 3.00%. The primary effect of Mn is to increase strength and toughness. At levels below 1.00 wt%, the desired effect is not achieved, and if it exceeds 3.00 wt%, casting and segregation problems may occur. Additionally, the deformation mechanism of the steel may be altered to TRIP (transformation-induced plasticity) due to the stabilization of austenite to room temperature by Mn, which does not help achieve an excellent combination of all desired mechanical properties (i.e., impact toughness, formability, and strength) in the product. Preferably, the Mn content is in the range of 1.20 to 2.70%. In one embodiment, Mn reaches 1.40 to 2.60%, preferably 1.50 to 2.50%, and more preferably 1.60 to 2.50%. In one embodiment, the appropriate minimum amount of Mn will be 1.65% and the appropriate maximum amount of Mn will be 1.95%.

[0039] In addition to the aforementioned effects of the essential alloying elements of the present invention, namely C, Si, Al, and Mn, another collective effect of these alloying elements is to increase the hardenability of the steel. They help prevent the formation of pearlite or ferrite phases during cooling after austenitization. This function enables the steel to avoid these phases at a cooling rate above a certain level during runout table cooling after hot rolling and before coiling. The presence of these soft (ferrite) and non-uniform (pearrite) phases is detrimental to obtaining excellent mechanical and formability properties in the final product, as they promote brittle and incoherent interfaces in the microstructure.

[0040] One or more micro-alloying elements selected from the group consisting of V, Nb, Ti, and Mo are optionally present. These micro-alloying elements increase strength through precipitation hardening by their respective carbides, nitrides, or carbonitrides. They also improve the weldability of the steel.

[0041] Chromium, another optional element of the present invention, also increases the hardenability of steel.

[0042] In the presence of copper, it increases the strength of steel through solid solution strengthening and precipitation hardening via copper deposits. Nickel increases impact toughness and prevents hot shortness that can occur during the hot treatment of steel due to the presence of copper.

[0043] When present as alloying elements, the preferred addition amounts (in weight%) of these optional alloying elements are as follows:

[0044] V: 0.010 ~ 0.10

[0045] Nb: 0.010 ~ 0.10

[0046] Ti: 0.010 ~ 0.10

[0047] Mo: 0.050 ~ 0.50

[0048] Cr: 0.10 ~ 1.50

[0049] Cu: 0.030 ~ 1.00

[0050] Ni: 0.020 ~ 0.50

[0051] Nitrogen (N), sulfur (S), and phosphorus (P) are residual elements present in steel as a result of the steelmaking and smelting processes. Their amounts are limited to S < 0.005%, P < 0.020%, and N < 0.0100%. Amounts greater than these are detrimental to mechanical properties, formability, and weldability. Preferably, S < 0.002% and N is 0.0005 to 0.0100%. Nitrogen within the specified range provides an effect similar to C and contributes to strength through the formation of carbonitrides of microalloying elements.

[0052] The above-mentioned selective alloying elements and the elements of nitrogen, sulfur, and phosphorus may vary independently of each other within a specified range. It has been found that they have an additive effect rather than a synergistic effect in the steel according to the present invention.

[0053] According to a second aspect, the present invention is also embodied in a manufacturing process of a hot-rolled strip that achieves a desired microstructure in the final product. Thus, the method according to the present invention is a method for manufacturing hot-rolled steel having the chemical properties discussed above.

[0054] The method for making the above steel includes the following steps:

[0055] - The step of casting molten steel into slabs;

[0056] - A step of reheating the above slab preferably at a temperature of 1100°C or higher, preferably for a time of 30 minutes or more;

[0057] - A step of rough-rolling the slab to an intermediate gauge, generally in the range of 35 to 45 mm, to break down the cast structure;

[0058] - Preferably, a step of hot rolling the steel into a strip at a finishing hot rolling temperature (FRT) higher than the Ar3 temperature of the steel, wherein Ar3 is the temperature at which the transformation from austenite to ferrite begins during cooling;

[0059] - A step of performing accelerated cooling of the hot-rolled strip on a runout table at a cooling rate exceeding 20℃ / s;

[0060] - A step of winding the hot-rolled and cooled steel strip at a temperature in the range of (Ms-50)℃ to (Ms-160)℃, wherein Ms is the martensite start temperature of the steel (in ℃);

[0061] - Step of coil cooling the above steel to room temperature;

[0062] - A step of pickling the above hot-rolled steel strip;

[0063] - Optionally, a step of coating the hot-rolled strip with Zn or a Zn-based alloy or an Al-based alloy or any other coating.

[0064] To avoid misunderstanding, Ms is denoted in °C. Preferably, FRT is higher than Ar3 + 50°C. FIG. 1 shows a schematic representation of the hot rolling and cooling to room temperature processes superimposed on a schematic continuous cooling conversion (CCT) diagram. Room temperature is defined as approximately 20°C. Reheating is preferably performed for at least 60 minutes when carrying out the hot rolling process according to the present invention in a conventional hot strip mill, particularly when based on thick slabs.

[0065] The present invention is not limited by the casting method. Steel can be cast as conventional thick slabs with a casting thickness of 150 to 350 mm, typically 225 to 250 mm, as well as as thin slabs with a casting thickness of 50 to 150 mm in a direct strip casting mill. Schematic examples of a process including a conventional hot strip mill and a thin slab casting / direct rolling mill are shown in FIGS. 2a and 2b, respectively. In the case of conventional thick slab casting, slab reheating is required to reheat the slab to ambient temperature (usually thick cast slabs are cooled from the casting temperature to the ambient temperature of the slab yard) and to homogenize the slab in relation to its composition. Therefore, the reheating temperature must be higher than 1100°C to dissolve precipitates in the presence of microalloying elements and to heat the slab to a temperature at which the final hot rolling in the finishing mill can still be performed at FRT > Ar3. Often, this requires a (slab) reheating temperature between 1150 and about 1250°C. In the case of thin-slab casting, the cast slab undergoes homogenization treatment in a homogenization furnace immediately after casting the thin slab, where the homogenization temperature must be 1100°C or higher, generally about 1125 to 1150°C. This also prevents the formation of precipitates when microalloying elements are present and heats the thin slab to a temperature at which final hot rolling in the finishing mill can still be performed at FRT > Ar3. According to the present invention, the reheating or homogenization time for the thin slab casting path is preferably 30 minutes or more.

[0066] Hot rolling of steel must be performed in the austenite phase so that ferrite does not exist in the final microstructure. Since another purpose of hot rolling in the austenite phase is to reduce the hot rolling force, it is desirable to maintain the finishing rolling temperature (FRT) at a temperature at least 50°C higher than the Ar3 of the steel.

[0067] After hot rolling, the steel strip is cooled on a runout table. The requirement here is that the steel must be cooled at a rate higher than the critical cooling rate to avoid unwanted phase transformations from austenite. In particular, ferrite and pearlite must not be formed, as they are detrimental to the mechanical and formability of the final product. Therefore, the runout table cooling rate (ROT-CR) must exceed the critical cooling rate to avoid the formation of ferrite and pearlite. There is no critical maximum ROT-CR, as transformation from austenite is guaranteed as long as the aforementioned critical cooling rate is exceeded across the thickness of the strip. Since an unnecessarily high ROT-CR affects the flatness of the strip after cooling and causes control issues that may result in stopping at the correct cooling stop temperature, the appropriate maximum ROT-CR is about 300°C / s, preferably about 200°C / s, and more preferably about 150°C / s. The practical ROT-CR range is 20 to 100°C / s, as this can be achieved through air cooling, laminar cooling, or water jet cooling depending on the thickness of the strip. For practical reasons, ROT-CR is defined as the average cooling rate of the strip surface.

[0068] Next, the hot-rolled steel strip is coiled at a temperature below Ms of the steel within the temperature range of (Ms-50)°C to (Ms-160)°C. Coiling below Ms is intended to allow subsequent coil cooling to begin with a phase mixture of martensite and austenite, with an initial martensite content in the range of 40 to 85 volume%. If the initial martensite content is higher than this amount, or in other words, if the coiling temperature (CT) is below Ms-160°C, the required tempering effect of the initial martensite is not obtained, and the high ductility, formability, and impact toughness of the steel are not achieved due to the short available time during coil cooling and the temperature being too low for effective tempering to occur. If the initial martensite content is less than 40 volume%, excessive tempering of the martensite may occur, and the product is not an ultra-high strength steel in the context of the present invention.

[0069] During coil cooling, the tempering of the initial martensite occurs continuously. At the same time, as the steel cools from the coil, new fresh martensite is formed. Because Si and Al are present in the steel, carbides are not formed in the tempered martensite. Additionally, due to some of the carbon in the transition from martensite to austenite, a very small amount of austenite may remain undeformed at room temperature (also known as retained austenite), but the amount is preferably limited to a maximum of 1 volume%, including 0 volume%.

[0070] After the steel has cooled to room temperature, the oxides (scale) on the hot-rolled steel are removed by pickling in an acidic solution (e.g., HCl) at a warm temperature (80–120°C) or by a combination of pickling and mechanical brushing of the surface. This step is necessary to make the steel surface suitable for direct use as uncoated HR steel or suitable for the coating process, if optionally required for corrosion resistance.

[0071] Optionally, HR steel strips may be coated with Zn or Zn-based alloys, or Al-based alloys, for example by hot-dip coating, electro-coating, or other coating technologies that provide excellent corrosion resistance during use.

[0072] The above process generates a desired microstructure to obtain intended mechanical properties. The present invention is also embodied in a steel product having steel chemical properties manufactured according to the above process and comprising the following microstructure (volume%):

[0073] - Tempered martensite (initial martensite during winding): 40–85%, preferably at least 50%, more preferably at least 60%;

[0074] - Fresh martensite (martensite formed during coil cooling after winding): 15–60%, preferably up to 50%, more preferably up to 40%;

[0075] - Retained austenite: Max. 1 volume%, including 0 volume%.

[0076] - Cementite or other metal carbides: 0 volume%

[0077] The chemical properties, process, and microstructure of the steel according to the present invention result in the following mechanical and formability properties.

[0078] - Yield strength (Rp): Minimum 1100 MPa

[0079] - Ultimate tensile strength (Rm): Minimum 1200 MPa

[0080] - Yield Ratio (Rp / Rm): Minimum 0.85

[0081] - Total elongation: Minimum 6.0% JIS5

[0082] - Hall Expansion Ability: Minimum 30%

[0083] - Bending angle @ 1mm thickness: Minimum 70˚

[0084] Preferably, the Charpy impact toughness is at least 40 Joules at -40°C and at least 100 Joules at room temperature.

[0085] According to the present invention, the chemical properties, process, and microstructure of the steel preferably result in the following mechanical and formability properties.

[0086] - Yield strength (Rp): Minimum 1100 MPa

[0087] - Ultimate tensile strength (Rm): Minimum 1200 MPa

[0088] - Yield Ratio (Rp / Rm): Minimum 0.85

[0089] - Total elongation: Minimum 8.5% JIS5

[0090] - Hall Expansion Ability: Minimum 50%

[0091] - Bending angle @ 1mm thickness: Minimum 80˚

[0092] - Charpy impact toughness: Minimum 40 Joules at -40℃, minimum 100 Joules at room temperature.

[0093] The strength value of the steel is determined primarily by the presence of hard components within the microstructure. Martensite is a strong phase in steel, and due to low-temperature tempering of less than Ms during coil cooling, martensite does not significantly loosen the strength. Therefore, in the present invention, both fresh martensite and tempered martensite are responsible for achieving 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 to the steel during deformation, thereby providing high total elongation values. Residual austenite is minimized to less than 1 volume% because it is detrimental due to low stability regarding impact toughness.

[0094] Retained austenite is generated due to the splitting of carbon from martensite to austenite during coil cooling. Carbon lowers the Ms-temperature, thereby increasing the stability of the austenite. However, in the present invention, retained austenite is intentionally avoided because it is difficult to control the mechanical stability of the retained austenite during different deformation and forming processes. Retained austenite must possess very high mechanical stability for the beneficial effects of increasing elongation (i.e., ductility) and impact toughness. To improve these properties, very high carbon saturation is required along with a fine film-like morphology. It is very difficult to achieve high carbon supersaturation in low-temperature continuous cooling processes, such as coil cooling. If the mechanical stability of the austenite is low, it rapidly transforms into martensite and creates a brittle interface with the matrix phase, which affects the total elongation. Retained austenite with low mechanical stability transforms much faster during dynamic loading processes, such as impact, and reduces impact toughness. Accordingly, in the present invention, a more homogeneous microstructure was produced by tempered martensite and fresh martensite without the presence of a large amount of retained austenite. That is, the presence of retained austenite was intentionally avoided and its maximum amount was limited to 1 volume%.

[0095] 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 containing a retained austenite phase in their microstructure are known to be more susceptible to LME during welding.

[0096] This was achieved using low-temperature coiling in the range of (Ms-50)°C to (Ms-160)°C, which is a temperature range where substantial carbon splitting is not expected to stabilize a large amount of austenite.

[0097] Rp, Rm, and total elongation were determined by quasistatic testing at room temperature (strain 3 x 10⁻⁶) using a JIS No. 5 specimen geometry through tensile testing parallel to the rolling direction in accordance with EN 10002-1 / 150 6892-1. -4 s -1 ) was determined in a tensile test. The shape of the tensile specimen was configured with a gauge length of 50 mm, a width of 25 mm, and a thickness of 3.2 mm in the rolling direction. The strength of the steel at a 0.2% offset strain 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.

[0098] Flexibility was determined by a three-point bending test according to the VDA 238-100 standard on 40 mm x 30 mm specimens with a thickness of 3.2 mm in both longitudinal and transverse directions. The bending axis followed 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, respectively) were converted to the angle corresponding to a thickness of 1.0 mm using the following formula: Bending angle at 1.0 mm thickness = Measured angle x Square root of actual thickness (mm). From these converted bending angles, the lowest values ​​of the longitudinal and transverse specimens for specific heat treatment conditions were taken to claim the scope of the present invention.

[0099] The flangeability or hole expansion capacity (HEC) of the steel was determined by a hole expansion test. Specimens with dimensions of 90 mm x 90 mm x 3.2 mm were cut from coiled steel. A 10 mm diameter hole was drilled in the center of the specimen, and the hole expansion test was performed according to the ISO / TS 16630:2003(E) standard. The HEC value was determined by the formula: HEC = (Expansion of initial hole diameter / Expansion of initial diameter) × 100%.

[0100] Charpy impact toughness was measured according to ASTM A370 standards using full-size Charpy V-notch (CVN) specimens (55 mm x 10 mm x 10 mm). The test was performed in both sheet directions by creating a V-notch parallel and perpendicular to the rolling direction (RD).

[0101] For all of the above mechanical tests, at least three specimens are tested for each condition and the average value is reported.

[0102] The microstructure was analyzed using a combination of techniques such as optical microscopy, X-ray diffraction (XRD), scanning electron microscopy (SEM), and dilatometrics. (lxwxt) Dilatometric tests on 10 mm x 5 mm x 3.2 mm specimens were performed by heating the specimens to 950°C at a rate of 10°C / s, holding for 2 minutes, and cooling to room temperature at a rate of 100°C / s (quenching for Ms) or 0.3°C / slow cooling for Ar3. From the dilatometric data, the Ms and Ar3 temperatures were determined. The amount of initial martensite (i.e., also tempered martensite after coil cooling) after coiling the steel was determined using the Koistinen-Marburger formula provided in the following literature: “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, and Acta Metallurgica, Vol. 7, 1959, pp. 59-60.

[0103]

[0104] Here, Ms is the martensite start temperature (°C) and CT is the winding temperature (°C), so (Ms-CT) is a measure of the amount of initial martensite, as it reflects the supercooling of less than Ms when coil cooling begins.

[0105] The amount of retained austenite was determined by XRD at the 1 / 4 thickness position of the sample. The XRD pattern was measured using a Panalytical Xpert PRO standard powder diffractometer (Co K a-방사선 It was recorded in the range of 45–165° (2θ). Quantitative determination of phase ratios was performed by Rietveld analysis using 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. The fresh martensite fraction was obtained by subtracting the initial martensite and other phase fractions (retained austenite, carbides, and other determined phases) from the total amount.

[0106] The composition of the zinc or zinc alloy coating is not limited. The coating can be applied in various ways, but hot-dip galvanizing using a standard GI coating bath is preferred. Zn-based coatings may include Zn alloys containing Al as an alloying element. The preferred composition of the zinc bath contains 0.10 to 0.35 wt% Al, and the remainder is zinc and unavoidable impurities.

[0107] Other Zn coatings may also be applied. An example includes a zinc alloy coating according to WO 2008 / 102009, which specifically comprises 0.3 to 4.0 wt% Mg and 0.05 to 6.0 wt% Al, preferably 0.1 to 5.0% Al, and optionally up to 0.2 wt% of unavoidable impurities and additional elements, with the remainder being zinc. A preferred composition of a Zn bath comprising Mg and Al as the main alloying elements is as follows: 0.5 to 3.8 wt% Al, 0.5 to 3.0 wt% Mg, optionally up to 0.2% of one or more additional elements; with the remainder being zinc and unavoidable impurities. The additional element, typically added in small amounts of less than 0.2 wt%, may be selected from the group comprising Pb, Sb, Ti, Ca, Mn, Sn, La, Ce, Cr, Ni, Zr, and Bi. Pb, Sn, Bi, and Sb are generally added to form spangles. Preferably, the total amount of additional elements in the zinc alloy is at most 0.2 wt%. Such small amounts of additional elements do not significantly alter the properties of the coating or bath in general applications. Preferably, when one or more additional elements are present in the coating, each is present in an amount < 0.02 wt%, and preferably each is present in an amount < 0.01 wt%. The additional elements are generally added only to prevent the formation of dross in the molten zinc alloy bath for hot-dip galvanizing or to form spangles in the coating layer.

[0108] In other embodiments, the metal coating comprises a (commercially pure) aluminum layer or an aluminum alloy layer. A typical metal bath for hot-dip coating, such as an aluminum layer, comprises aluminum alloyed with silicon, for example, 8 to 11 weight percent silicon, up to 4 percent iron, optionally up to 0.2 percent calcium, one or more additional elements, unavoidable impurities, and the remainder being aluminum. Silicon is present to prevent the formation of a thick iron-metal intermetallic layer that reduces adhesion and formability. Iron is preferably present in an amount of 1 to 4 percent, more preferably at least 2 percent. Meanwhile, the hot-rolled steel strip may be used in chassis or suspension parts of a vehicle, such as a lower control arm, frame rail, bumper beam or battery box, or frame of a heavy truck, or crane boom. Brief explanation of the drawing

[0109] The present invention will now be described by the following non-limiting drawings. FIG. 1 is a schematic diagram of the thermomechanical treatment of the present invention. FIG. 2a is a schematic diagram of a hot rolling mill for processing thick cast slabs. Figure 2b illustrates a thin slab casting facility with a direct rolling mill. Figures 3 and 4 show the shape and definition of the bending sample and Charpy sample. Specific details for implementing the invention

[0110] Seven steel ingots of the chemical substances AB and DH of the present invention, with dimensions of 200 mm x 100 mm x 100 mm, and a comparative example steel C were cast by melting the charge 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 defined boundaries of the present invention, whereas the comparative example steel contains Al and Si outside the defined boundaries of 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 at 1200°C for 30 minutes and hot rolled to final thicknesses of 2.8, 3.2 mm, 3.5, and 12 mm for all steels at a FRT greater than 900°C in the austenite phase field. Ar3 and Ms for the steels, measured by the expansion measurement method, are also shown in Table 1.

[0111] After hot rolling, the steels were immediately cooled on a runout table at various cooling rates, then cooled to room temperature from a different starting CT, and coil cooling simulations were performed in a muffle furnace. The strips were then pickled to remove oxides using the conventional method.

[0112] The various treatment conditions of the steels are summarized in Table 2. A, B, and C have similar Ms and Ar3. The FRT temperatures for steel A were 953°C, for steel B 939°C, and for steel C 945°C, all 50°C higher than Ar3. For steels A and B, a slow runout table cooling rate (ROT-CR) of 3°C / s was used, which falls outside the defined lower limit of the present invention. Additionally, two additional coil temperatures (200°C and 480°C), which fall outside the defined boundaries of the present invention, were used for steels A and B, having FRT and CT within the boundaries required by the present invention. A CT of 200°C is significantly lower than (Ms-160)°C, and 480°C is higher than Ms. These conditions are for comparison purposes. For steel C, which has chemical properties outside the scope of the present invention, all processing conditions (FRT, ROT-CR, and CT) were selected within the defined boundaries of the present invention. For steel DF, one set of process parameters was within the range (FRT, ROT-CR, and CT) claimed for the present invention, but for the other set, FRT and ROT-CR were kept the same, while CT was kept higher for comparison. In this case, a CT of 375°C was used, which is higher than (Ms-50°C) for steel DF. In fact, this CT was slightly higher than their Ms temperature. For steel GH, all processing parameters were maintained within the boundaries required by the present invention.

[0113] The specimens were extracted from the final steel strips for various mechanical and microstructural characterizations as described. 12 mm thick hot-rolled steel was used to prepare the Charpy impact specimens, while 2.8, 3.2, and 3.5 mm thick steel strips were used for all other characterizations.

[0114] For various treatment conditions, the phase content of the microstructure is presented in Table 3, tensile properties are given in Table 4, bending and HEC test results are in Table 5, and Charpy impact toughness is given in Table 6. The symbols and abbreviations used in the tables to denote tensile and bending test results are as follows: Rp = yield strength, Rm = ultimate tensile strength, A JIS5 = Total elongation using JIS5 specimen, BA = bending angle, L = longitudinal specimen with the bending axis parallel to the rolling direction, T = transverse specimen with the bending axis perpendicular to the rolling direction.

[0115] Table 3 shows that steels A and B achieved microstructures consisting of less than 85 volume% and at least 15 volume% of tempered martensite and fresh martensite, respectively, using FRT and ROT-CR within the defined boundaries for CT in the range of 275°C to 375°C, which falls within the range required by the present invention. Furthermore, the steels were free of carbides in their respective microstructures, and the residual austenite content was less than 1 volume% under these processing conditions. Their microstructures did not contain other phases such as ferrite, bainite, or pearlite.

[0116] Steels A and B produced a microstructure with a significant amount of bainite and retained austenite when subjected to ROT-CR exceeding 20°C / s at an FRT 50°C higher than Ar3, and a small amount of fresh martensite formed below Ms when coiled at 480°C, which is higher than Ms of these steels. Due to the Al and Si content, no carbides were present in these steels under these coiling conditions. A high content of retained austenite was caused by carbon enrichment of austenite during bainite transformation above Ms and fresh martensite formed during coil cooling below Ms. These bainite microstructures with a large amount of retained austenite differ from the microstructure intended in the present invention.

[0117] Similarly, when steels A and B are coiled at 200°C, which is much lower than the Ms-160°C of these steels, the microstructures of these steels also differ from those required for the present invention. These CT conditions contain more than 85 volume% tempered martensite and less than 15 volume% fresh martensite. On the other hand, when steels A and B have FRT and CT within the range required for the present invention but ROT-CR is 3°C / s (less than 20°C / s), they exhibited significant amounts of ferrite and pearlite in their respective microstructures in addition to a bainite matrix and a significant amount of retained austenite. Ferrite, pearlite, and bainite were formed due to a slow ROT-CR prior to coiling.

[0118] Due to the small amount of Si and Al in steel C, steel C formed a significant amount of carbide (2.3 volume%) while being processed under all parameters within the range required for the present invention.

[0119] As a result of the microstructure described above, properties as shown in Tables 4 to 6 were obtained. Steels A and B achieved a yield ratio greater than 0.85, Rp greater than 1100 MPa, and Rm greater than 1200 MPa, along with a total elongation (AJIS5) greater than 8.5%. When the CT is too high (480°C), the minimum levels of Rp and Rm targeted in the present invention are not achieved in steels A and B due to the presence of softer phases (bainite and retained austenite), even if the total elongation is high. The low Rp value for the CT at 480°C also resulted 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 higher than the target values ​​along with a high yield ratio, but the total elongation is too low (<8.5%). The low total elongation is due to the excessive amount of initial martensite present in the microstructure (= tempered martensite > 85 volume%), insufficient available time resulting in a lack of tempering effect during coil cooling, and the temperature being too low for effective tempering to occur.

[0120] In slow ROT-CR of 3℃ / s, the Rp and Rm values ​​of steels A and B are lower than 1100 MPa and 1200 MPa, respectively, because soft phases of bainite, ferrite, pearlite, and retained austenite are formed. In addition, the total elongation is high, but the yield ratio is lower than 0.85.

[0121] Steel C achieved low Rp, Rm, yield ratio, and total elongation values ​​due to the presence of carbides in the microstructure resulting from the absence of carbide-inhibiting elements Si and Al. Carbides are detrimental to mechanical properties and promote damage during deformation. Consequently, low tensile properties were obtained in steel C.

[0122] Similar to tensile properties, bendability and HEC are also high in steels A and B when processed within the defined processing parameters (FRT, ROT-CR, and CT) of the present invention (see Table 5). A minimum bending angle of 80° was achieved at a thickness of 1.0 mm, and a minimum HEC value of 50% was also obtained. However, when CT is high and the Ms of the steel is exceeded (i.e., 480°), the minimum bending angle and HEC are low, falling below the target values ​​of 80° and 50%, respectively, at a thickness of 1.0 mm. This is because, due to the multiphase nature of the microstructure of these steels, which includes fresh martensite, bainite, and retained austenite (see Table 3), multiple damage initiation sites exist at the interfaces of these phases when deformation is performed. Both the martensite existing as martensite and the martensite formed due to the TRIP effect of retained austenite are stronger phases than bainite and untransformed retained austenite. Meanwhile, under ideal processing conditions in which steels A and B have optimal amounts of tempered martensite and fresh martensite, there existed differences in low hardness or strength between these phases, resulting in homogeneous deformation during bending and hole expansion. Consequently, high HEC and bending values ​​were obtained under ideal processing conditions.

[0123] In addition, when the CT of steels A and B is too low (less than Ms (200℃)), the bendability and HEC values ​​are also low because too much (>85 volume%) of tempered martensite is present in the microstructure (see Table 3). The lack of effective tempering of this initial martensite present at the beginning of coiling caused low ductility (which is also reflected in the total elongation values ​​in Table 4), and caused low formability of these steels as measured by bendability and HEC.

[0124] When using slow ROT-CR (3℃ / s), as can be seen in Table 5, the bendability and HEC values ​​of steels A and B were also reduced due to the presence of softer phases of ferrite and pearlite. This is due to the brittle interface between the softer and harder phases of bainite and martensite obtained after the transformation of residual austenite during loading.

[0125] Steel C achieved 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 of Steel C were caused by the very low Al and Si content in the steel, which promoted the formation of carbides, even though the processing variables were within the range specified for the present invention (see Table 3).

[0126] In the transverse specimen (showing lower values ​​than the longitudinal specimen), the Charpy impact toughness of steel A treated according to the present invention is higher than 100J and 40J when tested at room temperature and -40°C, respectively. At the high CT of 480°C and the low CT of 200°C, the same values ​​are significantly lower than the minimum values ​​in the ideal treatment path described in the present invention. As previously explained, these low toughness values ​​are caused by the spontaneous transformation of retained austenite into martensite, the coexistence of harder and softer phases, and brittle fracture resulting from the low tempering effect. Additionally, due to the presence of carbides in steel C, the Charpy impact toughness values ​​are poor at both room temperature and cryogenic temperatures.

[0127] Accordingly, as discussed above, these embodiments exemplify that when steel is designed according to the composition of the present invention and processed according to the present invention, the steel achieves high tensile, formability, and toughness properties as intended due to microstructural effects. The same excellent combination of all properties is not achieved when working outside the boundaries defined in the present invention.

[0128] The Nb content of steel AG is at a residual level. Niobium was not added as an alloying element to this steel. Niobium was added as an alloying element to steel H.

[0129]

[0130] The Nb content of steel AG is at a residual level. Niobium was not added as an alloying element to these steels. Niobium was added as an alloying element to steel H.

[0131] Steels A and B have 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 effect of Mn is a reduction in Ms, which shifts the ratio of tempered martensite to fresh martensite. Although steel DH exhibits lower values ​​than steels A and B in terms of A JIS5, HEC, and bending angle, steel DH is nevertheless suitable for the purpose.

[0132] As shown in Table 3, steel DH achieved a microstructure consisting of less than 85 volume% and at least 15 volume% of tempered martensite and fresh martensite, respectively, using FRT and ROT-CR within the defined boundaries for CT of 275°C (steel DF) and CT of 300°C (steel GH), which fall within the range required for the present invention. Furthermore, these steels were free of carbides in their microstructure, and the residual austenite content was less than 1 volume% under these treatment conditions. The microstructures of the steels do not contain other phases such as ferrite, bainite, or pearlite.

[0133] Steel DH also achieved a yield ratio of greater than 0.85, Rp greater than 1100 MPa, and Rm greater than 1200 MPa, along with a total elongation (AJIS5) of greater than 6% (Table 4). Flexibility and HEC are also high in steel DH when processed within the defined processing variables (FRT, ROT-CR, and CT) of the present invention (see Table 5). A minimum bending angle of 70˚ and a minimum HEC value of 30% were achieved in these steels at a thickness of 1.0 mm.

[0134] However, when the CT of the steel DF is high, that is, 375°C, which is slightly higher than the Ms of these steels, the microstructure of these steels contains a certain amount of bainite and more than 1 volume% of retained austenite, which is not intended to be achieved in the microstructure of the present invention (see Table 3), and the content of fresh martensite and tempered martensite also falls outside the range defined in the present invention. This unintended microstructure does not provide the desired ultra-high strength of these steels at a CT of 375°C. Because softer phases of bainite and retained austenite are present, the Rp value is less than 1100 MPa, the Rm value is less than 1200 MPa, and the yield ratio is less than 0.85 (see Table 4), but excellent bendability and HEC values ​​are achieved (see Table 5). Therefore, steels that can have a chemical composition within the defined range of the present invention may not achieve all desired mechanical properties if the treatment is not performed within the defined range of the present invention.

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Claims

Claim 1 In a hot-rolled steel strip, the composition of the steel is in weight percent as follows: - C: 0.10 to 0.30; - Si: 0.50 to 1.50; - Al: 0.010 to 1.00; - Mn: 1.00 to 3.00; where (Si + Al): 0.80 to 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; - Cu: less than 1.00; - Ni: less than 0.50; - B: less than 0.0030 (30 ppm); and also, inevitably, comprises: - N: less than 0.0100 (100 ppm); - S: less than 0.005; - P: less than 0.020; and the remainder is A hot-rolled steel strip 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 6.0%, a hole expansion capacity of at least 50%, a bending angle of at least 80˚ at 1 mm, and Charpy impact toughness of 40J and at least 100J at -40℃ and room temperature, respectively, and having a microstructure of 40 to 85 volume% tempered martensite, 60 to 15 volume% fresh martensite, less than 1 volume% retained austenite, and no cementite or other carbides. Claim 2 A hot-rolled steel strip according to claim 1, comprising, in weight% units, - Al: 0.030 to 1.

00. Claim 3 delete Claim 4 A hot-rolled steel strip according to claim 1 or 2, comprising one or more of the following elements in the following amounts (in weight %): - V: 0.010 ~ 0.10 - Nb: 0.010 ~ 0.10 - Ti: 0.010 ~ 0.10 - Mo: 0.050 ~ 0.50 - Cu: 0.030 ~ 1.00 - Ni: 0.020 ~ 0.50 - N: 0.0005 ~ 0.0100; - S: max. 0.

002. Claim 5 A hot-rolled steel strip according to claim 1 or 2, wherein the microstructure consists of at least 55 volume% tempered martensite and up to 45 volume% fresh martensite. Claim 6 A hot-rolled steel strip according to claim 1 or 2, wherein the steel comprises 1.65 weight% to 2.50 weight% of Mn. Claim 7 A hot-rolled steel strip according to claim 1 or 2, wherein the total amount of Al and Si (Al + Si) is at least 1.00 weight%. Claim 8 A hot-rolled steel strip according to claim 1 or 2, provided with a metal coating layer which is a Zn layer, a Zn-based alloy layer, or an Al-based alloy layer obtainable by hot-dip coating. Claim 9 A hot-rolled steel strip according to claim 8, wherein the metal coating layer comprises 0.3 to 4.0 wt% Mg and 0.05 to 6.0 wt% Al, optionally up to 0.2 wt% one or more additional elements and unavoidable impurities, and the remainder being zinc. Claim 10 A method for manufacturing a hot-rolled steel strip, wherein the molten steel comprises the following composition (in weight%): C: 0.10 to 0.30; Si: 0.50 to 1.50; Al: 0.030 to 1.00; Mn: 1.00 to 3.00; (Si + Al): 0.80 to 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; Cu: less than 1.00; Ni: less than 0.50; B: less than 0.0030 (30 ppm); and inevitably also comprises: N: less than 0.0100 (100 ppm); S: less than 0.005; P: less than 0.020; the remainder being Fe and other elements generated in iron and steelmaking processes A step of casting into a slab having unavoidable impurities; - a step of heating or reheating the slab at a temperature of 1100°C or higher for a time of 30 minutes or more; - a step of (a) rough-rolling the slab to an intermediate gauge in the range of 35 to 45 mm to break down the cast structure, and then finishing hot-rolling it into a hot-rolled steel strip, or (b) hot-rolling the slab by direct rolling to hot-roll the slab into a hot-rolled steel strip, wherein the finishing hot-rolling temperature (FRT) is higher than the Ar3 of the steel, and Ar3 is the temperature at which austenite begins to transform into ferrite during cooling; - a step of performing accelerated cooling of the hot-rolled steel strip on a runout table at a cooling rate of more than 20°C / s; - subsequently, coiling the hot-rolled and cooled steel strip at a temperature in the range of (Ms-50)°C to (Ms-160)°C. The method comprises a step, wherein Ms is the martensite start temperature of the steel; a step of further cooling the coiled hot-rolled steel strip to room temperature; and a step of pickling the hot-rolled steel strip, wherein the hot-rolled steel strip 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, and at least 6.A method for manufacturing a hot-rolled steel strip having a total elongation of 0%, a hole expansion capacity of at least 50% and a bending angle of at least 80˚ at 1 mm, and Charpy impact toughness of 40J and at least 100J at -40℃ and room temperature, respectively, and a microstructure having 40 to 85 volume% tempered martensite, 60 to 15 volume% fresh martensite, less than 1 volume% retained austenite, and no cementite or other carbides. Claim 11 In claim 10, a method for manufacturing a hot-rolled steel strip, wherein the slab comprises one or more of the following elements in the following amounts (in weight %): - V: 0.010 ~ 0.10 - Nb: 0.010 ~ 0.10 - Ti: 0.010 ~ 0.10 - Mo: 0.050 ~ 0.50 - Cu: 0.030 ~ 1.00 - Ni: 0.020 ~ 0.50 - N: 0.0005 ~ 0.0100. Claim 12 A method for manufacturing a hot-rolled steel strip according to claim 10 or 11, wherein the microstructure consists of at least 55 volume% of tempered martensite and up to 45 volume% of fresh martensite. Claim 13 A method for manufacturing a hot-rolled steel strip according to claim 10 or 11, wherein the steel comprises at least 1.65 weight% to 2.50 weight% of Mn, and / or the total amount of Al and Si (Al + Si) is at least 1.00 weight%. Claim 14 A method for manufacturing a hot-rolled steel strip according to claim 10 or 11, wherein a metal coating layer, which is a Zn layer or a Zn-based alloy layer or an Al-based alloy layer obtainable by hot-dip coating, is provided. Claim 15 A method for manufacturing a hot-rolled steel strip according to claim 14, wherein the metal coating layer comprises 0.3 to 4.0 wt% Mg and 0.05 to 6.0 wt% Al, optionally up to 0.2 wt% of one or more additional elements and unavoidable impurities, and the remainder being zinc. Claim 16 Hot-rolled steel strip according to claim 1 or 2 used in parts for transportation or engineering applications. Claim 17 Hot-rolled steel strip according to Article 16 used in lower control arms, frame rails, bumper beams or battery boxes or frames of heavy trucks, or chassis or suspension parts of vehicles such as crane booms.

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