Hot-rolled high-strength steel strip

Abstract

The invention relates to a hot-rolled high-strength steel strip. The invention also relates to a method of manufacturing such hot-rolled high-strength steel strip.

Description

[0001] HOT-ROLLED HIGH-STRENGTH STEEL STRIP

[0002] FIELD OF THE INVENTION

[0003] The invention relates to a hot-rolled high-strength steel strip which is in particular suitable for, but not limited to, use in automotive components. The invention also relates to a method of manufacturing such a hot-rolled high-strength steel strip. Furthermore, the invention relates to an automotive part incorporating the hot-rolled high-strength steel strip.

[0004] BACKGROUND TO THE INVENTION

[0005] It is well-known in the art that as the strength of hot-rolled (HR) steel strip increases, the formability decreases, whereas the fracture toughness, crack susceptibility, and susceptibility to unstable brittle fracture are all compromised. A major area of application for HR steels in transport and automotive applications is the chassis and suspension (C&S) domain. Other areas include frame rails of trucks, bumper beams or battery boxes for electrical vehicles. The typical thickness of hot-rolled steels used for these applications is less than 8 mm, and is commonly less than 6 mm or less than 4.5 mm. Thicker gauge HR steel strip such as up to 12 mm can be used in engineering applications such as crane booms or in transport applications for frames of heavy trucks. From the weight reduction perspective it is imperative that higher strength steels should be employed for the above applications in order to be able to reduce the gauge of the steel strip. These applications of the hot-rolled steels demand mechanical properties that are difficult to reconcile. Besides high strength, the steel strip should also have good formability for making the component via e.g. cold-forming because this is an energy efficient manufacturing route in comparison with hot-forming. Furthermore, good impact and fracture toughness or energy absorption capacity is also required for applications like bumper beams, battery housings, crane booms or frame rails. For assembling the components, also a good weldability is required. However, as the tensile strength of the steels increases, the formability parameters decrease. Formability is a generic term for steel sheets which is viewed as a combination of material behaviour during several mechanical operations such as stretching, bending, drawing and flanging. Depending on the component geometry any or a combination of two or more attributes of the material is of importance during sheet metal forming. Next to formability, also the fracture toughness may be compromised with an increase of the strength. This can lead to increased crack susceptibility and increased susceptibility to unstable brittle crack propagation under specific mechanical loading conditions, such as local compression in deep-drawing operations. Another important characteristic for steels used for structural components in transport and automotive chassis applications is their corrosion resistance. For many of these applications the steels are not protected against corrosion with a metallic coating layer, such as e.g. a Zn- layer or Zn-based alloy layer, a Mg-alloy layer, applied to the steel by means of heat-to-coat or hot-dip coating. Instead, these steels are only protected by a phosphate and / or corrosion- preventive paint layer which provide a much less effective protection against corrosion compared to metallic coating layers. However, aforementioned structural components in transport and automotive chassis applications are in many cases exposed to the outdoor elements and are typically subject to abrasive or abusive conditions which may lead to local removal of a corrosion-preventive layer. This in turn impairs local corrosive protection and leads to a situation in which the exposed steel substrate is directly exposed to moisture. Therefor it is desirable for the steel substrate itself to have an improved corrosion resistance.

[0006] To spare component weight, the common approach is to apply high-strength steel and to reduce the thickness of the steel strip. However, this may lead to a loss in stiffness, which for some applications in automotive parts is undesirable. The intrinsic loss in stiffness by reducing the thickness of the steel strip used to manufacture automotive components, can be regained by optimisation of the component geometry, e.g. creating deeper flanges or flanges with an increased degree of stretching or bending. To allow for increased component stiffness via geometry optimisation, the high-strength steel strip requires an excellent formability in terms of tensile elongation and hole expansion capacity.

[0007] Single-phase precipitation-strengthened ferritic high-strength steels that break free from the conventional constraints between global formability (e.g. tensile elongation) and local formability (e.g., hole-expansion capacity or HEC) with both formability modes at a high level, may have low fracture toughness values and increased edge crack sensitivity in certain conditions, e.g., under compression, and are prone to have increased susceptibility to unstable brittle fracture behaviour as well as delamination upon shearing, which impairs sheared-edge fatigue.

[0008] Patent document EP1616970-A1 discloses a method for manufacturing a high-strength hot-rolled steel sheet comprising the steps of: reheating a steel slab consisting of, in wt.%, 0.04-0.15% C, 1.5% or less Si, 0.5-1.6% Mn, 0.04% or less P, 0.005% or less S, 0.04% or less Al, 0.03-0.15% Ti, 0.03-0.5% Mo, by mass, and balance of Fe and inevitable impurities, in a temperature range from 1150-1300°C; hot rolling the reheated steel slab at a finishing temperature of the Ar3transformation temperature or above into a hot rolled steel sheet; primarily cooling the hot rolled steel sheet in a temperature range from 700-850°C at an average cooling rate of 20°C / s or more; holding the primarily cooled steel sheet at a temperature of 680°C or above for more than 1 sec; and secondarily cooling the steel sheet at a temperature of 550°C or below at an average cooling rate of 30°C / s or more, followed by coiling the steel sheet. Preferably, the hot-rolled steel sheet is primarily cooled to a temperature range not only from 700-850°C but also from (SRT / 3+300) to (SRT / 8+700)°C, where the SRT designates the reheating temperature of the steel slab. In the examples of EP1616970-A1 this leads in practice to primary holding temperatures ranging from 667-860°C. The processing conditions are such that a microstructure is obtained which consists of ferrite containing precipitates, second phase of bainite and / or martensite, and other phase, wherein the percentage of the ferrite containing precipitates is 40-95%, and the percentage of the other phase being 5% or less. No details are disclosed about the corrosion resistance of the steels.

[0009] Patent document EP1338665-A1 discloses a method for manufacturing a high-strength hot-rolled single-phase ferritic steel with nano-precipitates steel sheet, comprising the steps of: producing a steel slab which consists essentially of, in wt.%, 0.06% or less C, 0.5% or less Si, 0.5-2.0% Mn, 0.06% or less P, 0.005% or less S, 0.1% or less Al, 0.006% or less N, 0.05-0.6% Mo, 0.02-0.10% Ti, and the balance being Fe, and satisfies the equation of 0.8 ≤ (C / 12) / [(Ti / 48)+(Mo / 96)] ≤ 1.3; producing a hot rolled steel sheet by hot rolling said steel slab at a temperature of Ar3 transformation point or higher; and coiling said hot rolled steel sheet at a temperature of 550-700°C. The processing conditions are such that a microstructure is obtained that consists essentially of a matrix of ferrite structure single phase and fine precipitates which are composite carbides containing Ti and Mo, with a grain size of smaller than 10 nm dispersed in said matrix wherein said fine precipitates are dispersed at a number per unit volume of 5 x 104 / mm3or higher. No details are disclosed about the corrosion resistance of the steels.

[0010] Patent document W02013 / 167572-A1 discloses a high-strength hot-rolled single-phase ferritic steel with nano-precipitates steel sheet with an excellent combination of tensile strength of at least 550 MPa and formability, comprising in wt.%, at most 0.15% of C, at most 0.5% of Si, 0.5-2.0% of Mn, at most 0.06% of P, at most 0.008% of S, at most 0.1% of Al_sol, at most 0.02% of N, 0.02-0.45% of V, 0.05-0.7% of Mo, optionally between 0.01-0.1% of Nb, balance of Fe and inevitable impurities, wherein the steel sheet has a precipitation strengthened and substantially single-phase ferritic microstructure, wherein the volume fraction of the ferrite phase in said microstructure is not lower than 97%, and wherein the precipitates in said microstructure comprise fine precipitates of composite carbides containing Mo and V and optionally Nb. The hot-rolled steel strip has been manufactured by means of hot rolling at a finish hot rolling temperature of Ar3 transformation point or higher, and coiling the hot-rolled sheet in the temperature range of between 700-585°C. No details are disclosed about the corrosion resistance of the steels.

[0011] Patent document W02018 / 193032-A1 discloses a high-strength steel strip having a cementite-free microstructure comprising: 0.005-0.08 wt.% C; 1.30-2.30 wt.% Mn; 2-35 ppm B; 5-65 ppm N; 0.005-0.1 wt.% Al Jot; 0.03-0.20 wt.% Ti; 0-1.5 wt.% Cu; 0-0.75 wt.% Cr; 0-0.05 wt.% Mo; 0-0.50 wt.% Ni; 0-0.30 wt.% V; 0-0.6 wt.% Si; 0-0.01 wt.% P; 0-0.01 wt.% S; C / (Ti_sol+V) ≤ 0.25 with Ti_sol=Ti-((48 / 14)-N), remainder iron and inevitable impurities, the steel strip having a yield strength of at least 570 MPa, a tensile strength of at least 760 MPa, a total elongation (A50) of at least 10.3 % and a hole expansion ratio (A) value of at least 70%. The steel strip has been manufactured by hot rolling, cooling the hot-rolled strip with an average cooling rate of 15-100°C / s on the run-out table within a time period of 2 seconds between the finish rolling and start of cooling to a coiling temperature below 500°C, and then coil cooling by natural cooling to ambient temperature, cold rolling with a reduction between 50-90%, and preferably followed by a solution heat-treatment to allow the bainitic transformation. No details are disclosed about the corrosion resistance of the steels.

[0012] There is a demand for hot-rolled high-strength well-drawable steels that, next to sufficient strength for down gauging and weight saving, offer a suitable combination of global (tensile elongation) and local (hole-expansion capacity) formability for component geometry optimisation and adequate stiffness while also providing an adequate level of fracture toughness to minimise brittle fracture behaviour as well as improved corrosion resistance, and thereby realising an optimum balance for all these material characteristics, i.e., formability in terms of tensile elongation and hole-expansion capacity, fracture toughness in terms of crack susceptibility and work of fracture, and corrosion resistance.

[0013] DESCRIPTION OF THE INVENTION

[0014] As will be appreciated herein, for any description of steel compositions or preferred steel compositions, all references to percentages are by weight percent unless otherwise indicated.

[0015] As used herein, the term "about" when used to describe a compositional range or amount of an alloying addition means that the actual amount of the alloying addition may vary from the nominal intended amount due to factors such as standard processing variations as understood by those skilled in the art.

[0016] The term “up to” and “up to about”, as employed herein, explicitly includes, but is not limited to, the possibility of zero weight-percent of the particular alloying component to which it refers. For example, up to 0.20% Cu may include a steel having no Cu. As used herein, the term “at % thickness” when used to describe microstructural features refers to an area between 1 / 8 and 3 / 8 of the steel strip thickness below the surface of the steel strip, and preferably between 3 / 16 and 5 / 16 of the steel strip thickness below the surface of the steel strip.

[0017] As used herein, the term “at1 / 2 thickness” when used to describe textural features refers to an area between 3 / 8 and 5 / 8 of the steel strip thickness below the surface of the steel strip, and preferably between 7 / 16 and 9 / 16 of the steel strip thickness below the surface of the steel strip.

[0018] It is an object of the invention to provide a hot-rolled high-strength deep-drawable steel strip with high edge ductility in terms of hole-expansion capacity, adequate fracture toughness and good corrosion resistance, that is in particular suitable for complex components for automotive applications.

[0019] It is an object of the invention to provide a hot-rolled high-strength deep-drawable steel strip having a yield strength ranging between 660 MPa to 820 MPa, an ultimate tensile strength ranging between 760 and 950 MPa, and with a high edge ductility in terms of hole-expansion capacity, adequate fracture toughness and good corrosion resistance that is in particular suitable for complex components for automotive applications.

[0020] It is another object of the invention to provide a hot-rolled high-strength deep-drawable steel strip having an improved balance of having a yield strength ranging between 660 MPa to 820 MPa, an ultimate tensile strength ranging between 760 and 950 MPa, and with high edge ductility in terms of hole-expansion capacity of at least 50%, a work of fracture of at least 185 J, and a good corrosion resistance ML-90 of less than 20 g after 90 cycles in a SAE J2334-April 2016 corrosion test.

[0021] It is yet another object of the invention to provide a method of manufacturing such a hot- rolled high-strength deep-drawable steel strip having an improved balance of having a yield strength ranging between 660 MPa to 820 MPa, an ultimate tensile strength ranging between 760 and 950 MPa, and with high edge ductility in terms of hole-expansion capacity, adequate fracture toughness and good corrosion resistance.

[0022] It is another object of the invention to provide a method of manufacturing such a hot- rolled high-strength deep-drawable steel strip having an improved balance of having a yield strength ranging between 660 MPa to 820 MPa, an ultimate tensile strength ranging between 760 and 950 MPa, and with high edge ductility in terms of hole-expansion capacity of at least 50%, a work of fracture of at least 185 J, and a good corrosion resistance ML-90 of less than 20 g after 90 cycles in a SAE J2334-April 2016 corrosion test.

[0023] These and other objects and further advantages are met or exceeded by the present invention providing a hot-rolled high-strength steel strip having a composition, in wt.%, C: 0.035 to 0.120 wt.%, Mn: 1.00 to 2.20 wt.%, Ti: up to 0.150 wt.%, Si: up to 0.65 wt.%, Mo: up to 0.50 wt.%, Nb: up to 0.100 wt.%, Cr: up to 0.090 wt.%, Al: up to 0.10 wt.%, P: up to 0.02 wt.%, S: up to 0.01 wt.%, N: up to 0.01 wt.%, optionally one or more elements selected from the group consisting of: (V: up to 0.20 wt.%, Cu: up to 0.20 wt.%, Ni: up to 0.50 wt.%, and B: up to 0.005 wt.%), and balance Fe and inevitable impurities resulting from the ironmaking and steelmaking process; and wherein the steel strip has a microstructure of: precipitation-strengthened acicular / bainitic ferrite and / or bainite, and at most 5 vol.% of second-phase constituents, including cementite, pearlite, martensite, and / or retained-austenite; and wherein the steel strip has at least the following mechanical and corrosion resistance properties: a yield strength (Rp0.2) ranging from 660 to 820 MPa, an ultimate tensile strength (Rm) ranging between 760 and 950 MPa, a tensile elongation (A50) of at least 11%, a hole- expansion capacity (HEC) of at least 50%, and a work of fracture (WOF) of at least 185 J, and a corrosion resistance ML-90 of less than 20 g mass loss after 90 cycles (ML-90) in a SAE J2334-April 2016 corrosion test.

[0024] The microstructure, the Rp0.2, the Rm, the A50, the HEC, and the WOF are measured as herein described.

[0025] In accordance with the invention it has been found that the hot-rolled steel strip having these narrow alloy compositional ranges in combination with the microstructure provides for an improved balance of a combined high edge ductility as expressed by hole-expansion capacity (viz., HEC ≥50 %) with adequate fracture toughness (e.g., ATCL ≤20 mm), and work of fracture (WOF) >185 J, and a very good corrosion resistance when tested in accordance with the SAE J2334-April 2016 corrosion test, viz. a mass loss of less than 20 g after 90 cycles (ML-90), the hot-rolled steel strip having a yield strength (Rp0.2) ranging between 660- 820 MPa, a tensile strength (Rm) ranging between 760-950 MPa, and a tensile elongation A50 of 11% or higher, and preferably tensile elongation A50 of 14% or higher. A reduced crack sensitivity, in particular a reduced edge-crack sensitivity, can be objectively expressed in an average total crack length (ATCL) following the measuring method as herein described. The steel strip according to this invention has an ATCL of less than 20 mm, preferably of less than 15 mm, and most preferably of less than 10 mm.

[0026] The work of fracture (WOF) is considered as a measure for the fracture toughness of the steel sheet and is at least 185 J. In an embodiment the WOF is at least 200 J, and preferably at least 225 J.

[0027] Based on extensive research, the inventors have found that by careful control of the steel composition and of the hot rolling process, the finish rolling temperature (FRT) and the coiling temperature (CT) to below 560°C in particular, and preferably to below 550°C, a hot-rolled steel strip is provided having a significantly improved corrosion resistance and still providing a complex phase microstructure such that a CP-type of steel is created having a desirable set of tensile properties (viz. in the range of CP800 steels), excellent hole-expansion capacity, low crack susceptibility, and good fracture toughness. The hot-rolled steel is in particular an ideal candidate for manufacturing exposed complex automotive chassis parts.

[0028] To achieve a hole-expansion capacity (HEC) A of at least 50%, and preferably of at least 60%, the microstructure requires a predominantly fine-grained acicular / bainitic ferrite and / or bainite microstructure. The amount of second-phase constituents, including cementite, pearlite, martensite, and / or retained-austenite should in total not exceed 5 vol.%, and preferably does not exceed 3 vol.%, as too high a fraction of hard carbon-rich second-phase constituents will lead to too much damage upon shearing, which is at the expense of the hole- expansion capacity.

[0029] In a preferred embodiment of the invention the microstructure has to be sufficiently grain-refined and is further being characterized by an average grain size by number at Ci- thickness in the rolling direction and based on a grain tolerance angle of 15° (GSNIS) of at most 2.5 pm, preferably of at most 2.2 pm, more preferably of at most 2.0 pm, and most preferably of at most 1.8 pm. Also grain refinement is beneficial for increased fracture toughness. This grain refinement can be realised by controlling an adequate interplay between steel composition and hot-rolling processing parameters. In relation to the hot-rolling process, a lower finish rolling temperature, leading to increased rolling below the non-recrystallisation temperature Tnrwill be beneficial to promote grain refinement of the final microstructure. However, also specific alloying elements in the steel composition will play a role as they have an influence on the Tnr. Apart from micro-alloying elements niobium, if present, and titanium to increase the Tnrto increase the fraction non-recrystallised austenite at the end of hot- rolling and prior to phase transformation, also carbon has a beneficial influence to increase Tnr. An increase of the carbon content leads to an increase of Tnr, which is beneficial to increase the fraction unrecrystallised austenite at the end of the hot-rolling process prior to austenite-to- ferrite phase transformation, which in turn will promote grain refinement of the final microstructure. Next to the effect of hot rolling and Tnr, also the Ar3 phase transformation temperature plays a role to determine the average grain size of the final microstructure. Carbon suppresses the Ar3 austenite-to-ferrite phase transformation, which will promote grain refinement. Thus carbon is beneficial to promote grain refinement of the final microstructure through its effect on the non-recrystallisation temperature Tnrand the austenite-to-ferrite phase transformation temperature Ar3, which in turn is beneficial for improved fracture toughness.

[0030] In an embodiment, in addition a MisOrientation Distribution (MOD) index of at least 0.85 at1 / 4-thickness is required, preferably of at least 0.95, more preferably of at least 1.0, and most preferably of at least 1.2. In general, an increase of the MOD index is associated with an increase of the hole-expansion capacity. The MOD index is a measure for the character of the microstructure, which is dominated by the phase constituents making up the matrix. High- temperature phase constituents, for instance polygonal ferrite will typically have a low MOD index. As the transformation temperature drops, the ferrite morphology will shift to quasi- polygonal or irregular-shaped ferrite, through acicular and / or bainitic (feathered) ferrite towards upper bainite and eventually lower bainite, before reaching martensite. In general, a decrease of the coiling temperature will lead to an increase of the MOD index. It is believed that an increased MOD index leads to a more intricate crystallographic arrangement that is beneficial for improved hole-expansion capacity and that the increased MOD index from a reduced coiling temperature comes with refined cementite particles due to reduced diffusivity.

[0031] The hot-rolled steels of this invention have a microstructure being precipitation- strengthened acicular / bainitic ferrite and / or bainite. This means that at least 95 vol.% (volume fraction) of the microstructure is formed by the precipitation-strengthened acicular / bainitic ferrite and / or bainite, and more preferably by at least 97 vol.%, and most preferably by at least 98 vol.%.

[0032] The steels have at most 5 vol.% second-phase constituents, including any cementite, pearlite, martensite, and / or retained-austenite, and preferably at most 3 vol.% second-phase constituents, and hence limited internal microstructural damage upon shearing.

[0033] Carbon is present in an amount between 0.035 and 0.120 wt.%. To achieve sufficient strength, a suitable minimum C content is 0.045 wt.%. In a preferred embodiment the minimum C content is at least 0.060 wt.%, and more preferably at least 0.075 wt.%. Carbon is beneficial to promote grain refinement of the final microstructure through its effect on the non- recrystallisation temperature Tnrand the austenite-to-ferrite phase transformation temperature Ar3, which in turn is beneficial for improved fracture toughness. In addition, C is an essential element to achieve precipitation strengthening in combination with carbide-forming micro- alloying elements like titanium, niobium (if added) and / or vanadium (if added), and to scavenge C to suppress a too high fraction of cementite or other carbon-rich secondary phase constituents in the final microstructure. By optimizing other alloying elements, including Ti, Nb, and / or V, it is possible to obtain an almost uniform precipitation-strengthened acicular / bainitic ferrite and / or bainite microstructure with substantially no or very little cementite. In a preferred embodiment the C content is at most 0.100 wt.% to prevent too much second-phase constituents in the final microstructure, which may deteriorate in particular hole-expansion capacity.

[0034] The steel strip has Mn in a range of 1.00 wt.% to 2.20 wt.% to achieve sufficient hardenability and grain refinement. Next to carbon also Mn increases Tnr and suppresses the Ar3 transformation temperature and is thus beneficial to promote grain refinement of the final microstructure. In an embodiment the Mn content is in a range of 1.40 wt.% to 2.20 wt.% for an improved balance in strength, corrosion resistance, fracture toughness, and edge crack sensitivity. Preferably the Mn content is at least 1.50 wt.%, more preferably at least 1.60 wt.%, and most preferably at least 1.70 wt.%, to obtain sufficient grain refinement improving the fracture toughness and reducing the crack susceptibility. In an embodiment the Mn content is maximum 2.10 wt.%, and more preferably maximum 2.00 wt.%. Too high a Mn content may lead to segregation during casting which adversely affects the required balance in properties and has an adverse effect on the corrosion resistance.

[0035] It is an important feature of the invention that the chromium (Cr) content in the steel according to this invention is kept very low. The Cr content should not exceed 0.090 wt.%, and preferably should not exceed 0.080 wt.%. In an embodiment the Cr content does not exceed 0.050 wt.%, and more preferably does not exceed 0.030 wt.%, and most preferably does not exceed 0.020 wt.%. Most preferably, Cr is present as an inevitable impurity resulting from the ironmaking and steelmaking process and whereby sufficiently care is taken that any scrap materials used in the steelmaking process have a low presence of Cr.

[0036] The inventors have found that each of the elements Mn, Cr and Mo have a notable adverse effect on the corrosion resistance, in particular when tested in a SAE J2334-April 2016 corrosion test. Based on extensive research, an empirical relationship has been found that may reasonable predict the corrosion resistance, viz the amount of weight loss in grams after 90 cycles (ML-90), according to SAE J2334-April 2016. When the effect of the coiling temperature is not taken into account, the empirical relationship is:

[0037] ML-90 = 8.59 Mn + 15.63 Cr + 8.24 Mo, with Mn, Cr and Mo contents in wt.%.

[0038] This empirical relationship shows that the adverse effect of the presence of Cr on the mass loss and hence the corrosion resistance is about twice the adverse effect of each of Mn and Mo. As mentioned herein above, it has been found that also the coiling temperature (CT), and thus the microstructure of the hot-rolled steel strip, plays a role. In has been found that by controlling the coiling temperature within the defined range the ML-90 can be further reduced by about 2 to 6 grams.

[0039] It has been found that the hot-rolled steel strip in accordance with this invention provides a good corrosion resistance both in marine and in urban conditions, as the final usage can be in each of these environments. This was evaluated using both the VDA 621-415 and the SAE J2334 test schemes.

[0040] Molybdenum (Mo) is an alloying element that can be present up to 0.50 wt.% and it may improve also the weldability of the steel strip and may contribute to precipitation strengthening as Mo as carbide-forming element can promote composite carbide precipitates in combination with Nb, Ti, and / or V. In an embodiment Mo is present up to 0.35 wt.%.

[0041] The inventors have found that too high a Mo content may adversely affect the corrosion resistance, in particular when present in combination with a high Mn content. In an embodiment the Mo-content is only up to 0.10 wt.%, and preferably up to 0.050 wt.%. In a more preferred embodiment Mo is not purposively added and only present as an inevitable impurity resulting from the ironmaking and steelmaking process.

[0042] Thus broadly defined the steel composition in accordance with this invention may comprise:

[0043] Mn: 1.00 to 2.20 wt.%, preferably 1.00 to 2.10 wt.%;

[0044] Mo: up to 0.50 wt.%, preferably up to 0.35 wt.%;

[0045] Cr: up to 0.090 wt.%, preferably up to 0.050 wt.%;

[0046] Si: up to 0.65 wt.%, preferably up to 0.50 wt.%, more preferably 0.10 to 0.30 wt.%.

[0047] In an embodiment of the steel composition the Mn content is kept at the lower end of the defined range to increase in particular the corrosion resistance together with a purposive addition of Mo to achieve the required level of mechanical properties and fracture toughness through sufficient hardenability and obtaining the required microstructure. The steel may comprise:

[0048] Mn: 1 .00 to 1 .65 wt.%, preferably 1 .00 to 1 .50 wt.%;

[0049] Mo: 0.050 to 0.50 wt.%, preferably 0.150 to 0.350 wt.%;

[0050] Cr: up to 0.090 wt.%, preferably up to 0.050 wt.%;

[0051] Si: up to 0.65 wt.%, preferably up to 0.50 wt.%, more preferably 0.10 to 0.30 wt.%.

[0052] In another embodiment of the steel composition the Mn is kept at the higher end of the defined range to achieve the required level of mechanical properties and fracture toughness while also achieving a very good corrosion resistance by controlling the amounts of both Cr and Mo at low levels. The steel composition may comprise:

[0053] Mn: 1.50 to 2.20 wt.%, preferably 1.60 to 2.10 wt.%; most preferably 1.70 to 2.10 wt.%;

[0054] Mo: up to 0.10 wt.%, preferably up to 0.050 wt.%;

[0055] Cr: up to 0.090 wt.%, preferably up to 0.050 wt.%;

[0056] Si: up to 0.30 wt.%, preferably 0.10 to 0.30 wt.%.

[0057] Titanium is an alloying element that can be present up to 0.150 wt.% and it provides precipitation strengthening as it acts as a carbide forming element suppressing the formation of cementite while providing precipitation strengthening via the formation of small Ti-based carbides. However, Ti also combines with N, S and C to form nitrides, and carbo-sulphides, depending on the specific chemical composition of the steel. Therefore, preferably at least 0.050 wt.% Ti is present to bind substantially all the N and S in the steel and to have sufficient excess Ti to combine with C in the steel. In a preferred embodiment at least about 0.080 wt.% Ti is present. When more than 0.150 wt.% Ti is present, coarse Ti nitrides, carbo-nitrides, and carbides may form which are difficult to dissolve during reheating of the slab prior to hot rolling. Furthermore, these coarse Ti nitrides, carbo-nitrides, and carbides can lead to a deterioration of the hole expansion capacity and fracture toughness of the steel. In an embodiment the Ti content does not exceed 0.120 wt.%.

[0058] Silicon may be present up to 0.65 wt.%, and is preferably present in an amount of 0.10 to 0.65 wt.% to improve the strength of the steel by substitutional solid solution strengthening of the iron lattice. Furthermore, Si is beneficial to suppress cementite formation. However, when using higher amounts of Si the weldability and coatability of the steel deteriorates, hence the amount of Si is preferably at most 0.50 wt.%, and in a preferred embodiment at most 0.40 wt.%, and most preferably up to 0.30 wt.%. In an embodiment the amount of Si is at least 0.10 wt.%. Aluminium behaves comparable to Si in the steel according to the invention. It slows down the carbide precipitation kinetics and suppresses the formation of cementite. When Al is less than about 0.01 wt.%, the effects of suppression of carbide formation are negligible. Values of aluminium lower than 0.01 wt.% are deemed to be residuals from the deoxidation step during steelmaking, and therefore a minimum value of about 0.01 wt.% is preferred. On the other hand, when Al is above about 0.10 wt.% there can be excessive oxide formation during thermo- mechanical processing (slab reheating, hot rolling, coiling, etc.) of the steel, which can deteriorate formability and toughness. Also, with increasing Al higher amounts of surface oxidation can occur at higher temperatures. These oxide scales are detrimental for hot rolling, pickling, coating and overall surface appearance. Also, the rolling forces during hot rolling increase when the Al increases in combination with the presence of Si to such a level making the steel very brittle and more difficult to hot roll. Therefore, Al in the present invention is present in an amount of up to 0.10 wt.%, preferably in a range of 0.01-0.10 wt.%, and more preferably in the range of 0.030-0.080 wt.%.

[0059] Nitrogen, sulphur and phosphorus are residual elements present in the steel as a result of steel making and refining process. Their amounts are limited to up to 0.01 wt.% S, up to 0.02 wt.% P, and up to 0.01 wt.% N. Amounts higher than these are detrimental for mechanical properties, formability, toughness, sheared-edge quality, fatigue, and weldability. In an embodiment P is present only up to 0.015 wt.%. In an embodiment S is present only up to 0.005 wt.% (50 ppm) and more preferably only up to 0.0025 wt.% (25 ppm), and most preferably up to 0.0012 wt.% (12 ppm). N forms titanium nitrides with Ti which act as dispersoids for austenite grain size control during reheating. However, too high N can lead to too much coarse TiN particles that can impair hole expansion capacity. Preferably the N content is up to 0.0075 wt.% (75 ppm), and more preferably up to 0.0065 wt.% (65 ppm). A suitable minimum N content is about 0.0010 wt.% (10 ppm).

[0060] Niobium can be present in the steel up to 0.100 wt.%. Nb improves the strength of the steel partly by precipitation hardening but foremost by grain refinement. This grain refinement is also beneficial for increased fracture toughness. However, as Nb is a rather expensive alloying element and can promote anisotropy, the Nb content preferably does not exceed 0.080 wt.%, more preferably it does not exceed 0.070 wt.%. In an embodiment the Nb content is at least 0.003 wt.%, and preferably at least 0.005 wt.%. Also, Nb can promote increased anisotropy and additionally has a high tendency to segregate and to form coarse NbC particles due to centre- line segregation, thus the avoidance of the use of Nb results in an improved hole-expansion capacity. In another embodiment Nb is not purposively added and only present as an inevitable impurity resulting from the ironmaking and steelmaking process. In practice this means it can be present up to 0.0030 wt.%.

[0061] The steel strip may have optionally one or more elements selected from the group consisting of: up to 0.20 wt.% V, up to 0.20 wt.% Cu, up to 0.50 wt.% Ni, and up to 0.005 wt.% B).

[0062] Vanadium can be present in the steel up to 0.20 wt.%, more preferably up to 0.170 wt.%, and most preferably up to 0.100 wt.%. However, V is a relatively costly alloying element that is mostly used to replace Ti for its precipitation strengthening effect and to avoid cementite formation by forming vanadium carbides or composite carbide precipitates with Nb, Ti, and / or Mo.

[0063] Copper, when present, may increase the strength of the steel strip by both solid solution strengthening as well as precipitation hardening through copper precipitates. However, the Cu content should not exceed 0.20 wt.% as a too high level of Cu can promote hot shortness that can occur during hot working of the steel strip, which leads to embrittlement of the steel. In an embodiment the Cu is not added as a purposive alloying element and can be present up to 0.10 wt.%, more preferably up to 0.050 wt.%, and most preferably up to 0.030 wt.%.

[0064] Nickel up to 0.50 wt.%, and preferably up to 0.30 wt.%, increases the impact toughness and counters any hot shortness that can occur during hot working of the steel strip due to the presence of copper. In an embodiment Ni is not added as a purposive alloying element and can be present up to 0.10 wt.%, more preferably up to 0.050 wt.%, and most preferably only up to 0.030 wt.%.

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

[0066] In an embodiment the steel strip has a composition consisting of, in wt.%, 0.035 to 0.120 wt.%, Mn: 1.00 to 2.20 wt.%, Ti: up to 0.150 wt.%, Si: up to 0.65 wt.%, Mo: up to 0.50 wt.%, Nb: up to 0.100 wt.%, Cr: up to 0.090 wt.%, Al: up to 0.10 wt.%, P: up to 0.02 wt.%, S: up to 0.01 wt.%, N: up to 0.01 wt.%, optionally one or more elements selected from the group consisting of: (V: up to 0.20 wt.%, Cu: up to 0.20 wt.%, Ni: up to 0.50 wt.%, and B: up to 0.005 wt.%), and the balance Fe and inevitable impurities resulting from the ironmaking and steelmaking process, and with more preferred ranges as herein described and claimed.

[0067] The 0.2% offset proof strength or yield strength (Rp0.2), ultimate tensile strength (Rm), uniform elongation (Ag) and tensile elongation (A50) were determined from quasistatic (strain rate 3 x 10-4s-1) tensile tests at room temperature with A50 specimen geometry with tensile testing parallel to the rolling direction according to EN 10002-1 / 150 6892-1. The geometry of the tensile specimens consisted in 50 mm gauge length in the rolling direction, 12.5 mm in width and a thickness depending on the final gauge. The strength of the steel at 0.2% offset strain is measured as the yield strength (Rp0.2 or YS).

[0068] The stretch-flangeability of the steel strip or the hole expansion capacity (HEC) was determined by hole expansion tests. Specimens of dimension 90 mm by 90 mm by final thickness of the strip were cut from the as-coiled steel. A hole of 10 mm diameter was punched in the middle of the specimens, and the hole expansion tests were carried out according to ISO / TS 16630:2003(E) standard. Hole expansion testing of the samples was done with upper burring. A conical punch of 60° was pushed up from below and the hole diameter dfwas measured when a through-thickness crack formed. The hole expansion ratio A was calculated using the formula below with d0= 10 mm:

[0069] For all the above mechanical tests, at least three specimens were tested for each condition and the average values are reported herein.

[0070] The average total crack length (ATCL) is used to assess the susceptibility of crack formation in situations similar to industrial applications and for the steel according to this invention should be <20 mm. The ATCL parameter is determined in common laboratory cylindrical deep drawing tests on steel blanks, measuring 90 mm x 90 mm, with the four corners of the square 10 mm cut in the direction of the two diagonals as shown in Fig. 1A. In the cylindrical deep drawing tests a punch, a draw die, and a blank holder is used. In this procedure the punch has a 50 mm diameter with a punch radius of 7 mm. The die has an inner diameter of 62 mm and a radius of 8 mm. This set up is shown schematically in Fig. 1 B. The inner diameter is large enough to allow free movement of the edge of the formed cup. The clearance, i.e., the distance between the punch wall and die wall is 6 mm. The blank holder force is set at 50 kN. During the initial stage of the cylindrical deep drawing test and the deep drawing of a cup (Fig. 1 C), four regions at the edge of the blank are plastically deformed due to high local compressive stresses during drawing (Fig. 1 D). This results in local wrinkling of the edge. At the end of the test and upon release of the blank holder force, the four compressed regions are subjected to reverse loading due to spring back as these regions of the formed cylindrical cup start to lose contact with the blank holder. This reverse loading due to spring back may lead to the nucleation and growth of cracks in the four compressed and wrinkled regions of the drawn cup (see for example Fig. 1 E). Cracks may either be through the full thickness of the steel strip and visible at both sides, i.e., inner and outside, of the drawn cup or only visible at one of both sides of the cup. The length of all visible cracks on the inside and outside of the four compressed edges (as illustrated in Fig. 1 D) of the deep-drawn cup is measured using a magnifying glass of 10x equipped with a scale grid. The sum of the length of all observable cracks on the inside and outside of the cup wall is averaged over three of four drawn cups and reported as the average total crack length (ATCL) and expressed in mm.

[0071] Further, the work of fracture (WOF) is determined by compressing the formed cups between a cone and flat tool, the normal of the flat plane and the axis of the cone being parallel. The cone has an angle of 90° and touches the inner wall of the cup. As the flat tool moves towards the cone it exerts a normal force on the flat bottom of the cup. The cone exerts forces on the wall of the cup resulting in crack growing from the four regions of compressed edges of the cup. The tools are placed in a conventional tensile testing machine. The normal force on the flat plate is recorded as a function of the plate displacement. The test stops a few mm before the flat bottom of the cup touches the top of the cone. For the selected geometry displacement up to 12 mm can be recorded. The area under the force vs. displacement graph was calculated up to 10 mm displacement. The calculated value in Joules (kN mm or J) is an estimate of the material resistance to crack propagation and / or initiation and the work of fracture (WOF) is defined here as the area below the force-displacement curve up to a displacement of 10 mm.

[0072] The microstructure of the steel strip was analysed by means of Electron Back Scatter Diffraction (EBSD), a technique well known in the art, which in turn also allows the quantification of the area or volume fraction of the various components. The EBSD measurements were conducted on cross sections parallel to the rolling direction (RD-ND plane) mounted in a conductive resin and mechanically polished to 1 pm. To obtain a fully deformation free surface, the final polishing step was conducted with colloidal silica (OPS).

[0073] The Scanning Electron Microscope (SEM) used for the EBSD measurements is a Zeiss Ultra 55 machine equipped with a Field Emission Gun (FEG-SEM) and an EDAX PEGASUS XM 4 HIKARI EBSD system. EBSD scans were collected on the RD-ND plane of the sheets at Ci- thickness. The samples were placed under a 70° angle in the SEM. The acceleration voltage was 15 kV with the high current option switched on. A 120 pm aperture was used and the typically working distance was 17 mm during scanning. To compensate for the high tilt angle of the sample, the dynamic focus correction was used during scanning.

[0074] The Scanning Electron Microscope (SEM) used for the EBSD measurements is a Zeiss Ultra plus machine equipped with a Field Emission Gun (FEG-SEM) and an Oxford instruments Symmetry S2 EBSD system. EBSD scans were collected on the RD-ND plane of the sheets at quarter-thickness. The samples were placed under a 70° angle in the SEM. The acceleration voltage was 20 kV with the high current option switched on. A 120 pm aperture was used and the typically working distance was 17 mm during scanning. To compensate for the high tilt angle of the sample, the dynamic focus correction was used during scanning.

[0075] The EBSD scans were captured using the Oxford instruments AZtec® 6.1 software. Typically, the following data collection settings were used: The Symmetry S2 camera in “speed 2” mode combined with auto background substraction. The scan area was in all cases located at a position of % the sample thickness and care was taken to avoid as much as possible to include non-metallic inclusions in the scan area. The EBSD scan size was in all cases 200x150 p.m, with a step size of 0.1 p.m, and a scan rate of approximately 1000 frames per second. Fe(a) and Fe(y) were used to index the Kikuchi patterns. The indexing settings used during data collections were: Indexing mode = Optimized-BD; number of bands = 11 ; Hough resolution is 50; solution refinement = on. After collection the scans were exported to h5oina data format. This resulted in a typically hit rate >97.5% for the scans.

[0076] The EBSD scans (in h5oina format) were evaluated with Oxford instruments AZtecCrystal 3.1 software. Cleanup procedure: step 1) wild spike removal; step 2) zero solution Removal (level 4) step 3) zero solution removal (level 6).

[0077] The MisOrientation angle Distribution (MOD) index of the Fe(a) partition is calculated using the following method: the normalised misorientation angle distribution (MOD), including all boundaries, ranging from misorientation angles of 5° to 65° with a binning of 1°, was calculated from the partitioned EBSD data set using the TSL OIM Analysis software. Similarly, the normalised theoretically MOD of randomly recrystallized polygonal ferrite was calculated with the same misorientation angle range and binning as the measured curve. In practice this is the so-called “MacKenzie” based MOD included in the TSL OIM Analysis software. Normalisation of the MOD means that the area below the MOD is defined as 1. The MOD index is then defined as the area between the theoretical curve (the dashed line) and the measured curve (the solid line) in Fig. 2A and Fig. 2B and can be written as: with MMODJ as the intensity at angle i (ranging from 5° to 65°) of the measured MOD and RMODJ as the intensity at angle i of the theoretical or “MacKenzie” based MOD of randomly recrystallized polygonal ferrite.

[0078] The solid line in Fig. 2A and 2B represents the measured MOD and the dashed curve represents the theoretical misorientation angle curve for a randomly recrystallized polygonal ferrite structure. Fig. 2A shows a MOD curve for an exemplary sample with a microstructure having a predominantly polygonal ferrite character. Fig. 2B shows a MOD curve of an exemplary sample with a microstructure having a predominantly acicular / bainitic or bainitic character. The MOD index ranges by definition from 0 to almost 2; when the measured curve is equal to the theoretical curve, the areas between the two curves is 0 (MOD index will be 0), whereas if there is (almost) no intensity overlap between the two distribution curves, the MOD index is (almost) 2. So, as illustrated in Fig. 2A and Fig. 2B, the MOD contains information on the nature of the microstructure and the MOD index can be used to assess the character of a microstructure based on a quantitative and hence more unambiguous approach than based on conventional methods such as light-optical microscopy. A fully polygonal ferrite microstructure will have a unimodal MOD with most of the intensity in the 20° to 50° range and a peak intensity around 45°. In contrast, a fully acicular / bainitic ferrite or bainitic microstructure will have a strong bimodal MOD with peak intensities in between 5° to 10° and 50° to 60° and little intensity in the range of 20° to 50°. Hence, a low MOD index and a high 20° to 50° MOD intensity in the present example is a clear signature of a predominantly polygonal or quasi-polygonal ferrite microstructure, whereas a high MOD index and a low 20° to 50° MOD intensity is a clear signature of a predominantly acicular / bainitic ferrite or bainitic microstructure. Inventors found that a MOD index above 0.85 corresponds with a predominantly acicular / bainitic ferrite or bainitic microstructure with the steel compositions and process settings as used in the example to describe the present invention.

[0079] In an aspect of the invention it relates to a method of manufacturing a hot-rolled high- strength steel strip as herein described and claimed, the method comprising the steps of, in that order, casting a slab, followed by the step of reheating the solidified slab to a temperature between 1050-1260°C, preferably holding the slab at a temperature between 1050°C and 1260°C for a time of about 20-60 minutes, and hot rolling said slab, or casting a slab or strip followed by the step of hot rolling said slab or strip; hot rolling the steel slab or strip and finishing said hot rolling at a finish rolling temperature (FRT) between 960°C and Ar3+20°C, preferably between 960°C and 900°C, and most preferably between 960°C and 920°C, and where Ar3 is the temperature at which transformation of austenite to ferrite starts during cooling. As known in the art the Ar3 temperature can be calculated according to the following equation, elements concentrations are in wt.%:

[0080] Ar3 = 910°C - 203x[C]1 / 2+ 44.7x[Si] + 400x[AI] + 700x[P]- 30x[Mn] +31.5x[Mo] - 11 [Cr] + 400x[Ti] + 104x[V], cooling the hot-rolled strip to the coiling temperature with an average cooling rate of at least 35°C / s, and preferably in a range of 35 to 60°C / s; coiling the hot-rolled and cooled steel strip at a coiling temperature (CT) between 475°C and 560°C, and preferably between 500°C and 550°C; allowing the coiled hot-rolled steel strip to further cool to ambient temperature; optionally pickling the hot-rolled steel strip; and optionally providing the hot-rolled steel strip with a metallic coating layer, preferably selected from the group comprising: a Zn-layer, Zn-based alloy layer, Mg-based alloy layer, or an Al-based alloy layer, to provide improved corrosion resistance in service. The metallic coating layer is preferably applied by means of heat-to-coat or hot-dip coating, and may also be applied via electrolytic deposition.

[0081] The method of manufacturing herein described and claimed result in the desired microstructure providing for the aimed improved balance of high edge ductility as expressed by hole-expansion capacity with adequate toughness and good corrosion resistance for a hot-rolled deep-drawable sheet strip with a yield strength ranging from 660 to 820 MPa, a tensile strength ranging between 760 to 950 MPa, and a A50 tensile elongation of 11% or higher. The invention is also embodied in a steel strip manufactured by the method described herein and claimed having said composition, microstructure and improved balance of toughness, shearing performance and mechanical properties.

[0082] The invention is not limited by the casting method. The steel can be cast as a conventional thick-slab having a cast thickness of between 150-350 mm, and typically of 225- 250 mm, as well as a thin-slab having a cast thickness of between 50-150 mm in a direct strip plant. For conventional thick-slab casting, reheating of the slab is necessary to reheat the slab from ambient temperatures (usually the thick cast slabs have cooled down from the casting temperature to ambient temperatures in a slab yard) and to homogenise the slab with respect to composition, and therefore the reheating temperature should be above about 1050°C also to dissolve any precipitates when micro-alloying elements are present and to bring the slab to such a temperature that the final hot rolling in the finishing mill can still be performed above 960°C. Often this requires a (slab) reheating temperature of between about 1050-1260°C. For thin-slab casting the cast slab is subjected to a homogenisation treatment in a homogenising furnace immediately after casting the thin slab wherein the homogenisation temperature should be above about 1050°C, and is typically about 1100-1160°C. This would also prevent any precipitates from forming when micro-alloying elements, if any, are present and also bring the thin slab to such a temperature that the final hot rolling in the finishing mill can still be performed at FRT>Ar3+20°C. According to the invention the reheating or the homogenisation time for the thin slab casting route is preferably 20 minutes or more.

[0083] The hot rolling of the steel must be carried out in the austenitic phase to control the final microstructure. On an industrial scale of rolling the FRT should be kept above the Ar3 temperature. In a preferred embodiment the above Ar3+20°C, e.g. typically above about 890°C, to avoid hot rolling locally below Ar3 at colder edges or the tail of the strip. However, the FRT should not be too high in the austenite region as that may lead to insufficient grain refinement (e.g. too high GSNIS), which can ultimately lead to too low fracture toughness and hence a too high average total crack length and too low work of fracture. On the other hand, a too low FRT, next to avoiding hot rolling locally below Ar3at colder edges or the tail of the strip, a too low FRT can stimulate increased anisotropy. A too high degree of anisotropy increases the planar anisotropy (increased variation in r-value), which can reduce the drawability of the steel and impair its hole-expansion capacity. Hence the FRT should not be too low and be above (Ar3 + 20°C), preferably is above 900°C, and more preferably is above 920°C.

[0084] In an embodiment, after hot rolling, the steel strip is accelerated cooled on a run-out table (ROT) to an intermediate temperature on the ROT (ITROT) between 480-580°C, preferably between 500-560°C, and with a run-out table primary cooling rate (CR1) between 50-150°C / s, preferably between 65-100°C / s. This can be achieved through air cooling, laminar cooling or water jet cooling depending on the thickness of the steel strip. An accelerated cooling rate is desired to suppress recovery and loss of internal stored energy in the austenite in order to promote grain refinement (small GSNIS) of the final microstructure. Preferably the cooling rate should be high enough to avoid austenite-to-ferrite phase transformation at elevated temperatures (potentially leading to a too low MOD index) and to preferably promote austenite- to-ferrite phase transformation at relatively low temperatures on the run-out-table promoting a high MOD index ≥0.85, and preferably of at least 0.95, more preferably of at least 1.0, and most preferably of at least 1.2. Increased cooling rate will promote grain refinement and hence increased fracture toughness and reduced crack susceptibility. Increased cooling rate will also suppress texture randomisation and hence suppress the loss in intensity of those textures developed from deformed austenite that promote toughness (e.g., {332}<113>) and suppress those that are associated with embrittlement and delamination (e.g., {001}<110> rotated Cube). An unnecessarily high CR1 may affect the flatness of the strip after cooling and can cause control problems to stop at the correct cooling stop temperature and therefore a suitable maximum CR1 is 150°C / s, or preferably 100°C / s. For practical reasons the run-out table cooling rate (CR-ROT) is defined as the average cooling rate of the surface of the steel strip.

[0085] Next in an embodiment, the hot-rolled steel strip at a temperature (ITROT) between 480- 580°C, preferably between 500-560°C, is actively cooled to a coiling temperature between 475- 560°C, with a run-out table secondary cooling rate (CR2) of about 1 to 30°C / s. In a preferred embodiment the CR2 is in a range of about 5-25°C / s.

[0086] Further, the hot-rolled steel strip is coiled at a temperature between 475-560°C, and preferably between 500-550°C. Next to the finish rolling temperature, the coiling temperature of the steel strip is a key process parameter to arrive at the required microstructure of the hot- rolled steel strip providing for the improved balance in properties as herein described.

[0087] When the coiling temperature is too low, e.g. below 475°C, there is insufficient kinetics for precipitation and consequently low strength levels will be achieved. When the coiling temperature is too high there is insufficient grain refinement leading to reduced fracture toughness and increased edge-crack susceptibility. A too high coiling temperature will also reduce the MOD index, which can impair hole-expansion capacity. During coil cooling further precipitation may take place, as well as some further phase transformation. Undesirably, precipitates once formed may coarsen during coil cooling. The steel composition in combination with the claimed coiling temperature range supresses this phenomenon. This coiling temperature will help to promote small grains of ferrite formed during cooling on the run-out- table, coiling, or coil cooling, and furthermore suppress coarsening of precipitates that strengthen the ferrite matrix.

[0088] In patent document EP1616970-A1 it is disclosed that a holding temperature below 680°C leads to insufficient driving force for ferrite transformation and subsequent a too low fraction ferrite containing precipitates. For the present invention, the austenite-to-ferrite phase transformation is enforced at a temperature range below 680°C in order to achieve increased grain refinement and increased ZGB15-65 (high-angle grain boundary density) for improved fracture toughness and reduced crack susceptibility, while still having sufficient kinetics for precipitation. Furthermore, patent document EP1616970-A1 discloses that after holding the steel strip longer than 1 sec at a temperature of 680°C or above, it is necessary to apply secondary cooling to a coiling temperature of 550°C or below, preferably 450°C or below, and more preferably 350°C or below at an average cooling rate of 30°C / s or more, preferably 50°C / s or more, and coiling in order to form the secondary phase of bainite and / or martensite and to suppress the formation of other phase at 5 vol.% or less. For the present invention, the coiling temperature is considerably higher with values between 475-560°C, in order to allow the austenite-to-ferrite phase transformation to continue at relatively low temperatures to promote fine-grained acicular / bainitic ferrite and / or bainite, which is precipitation strengthened with carbide precipitates comprising Ti, Mo, Nb, and / or V. The grain refinement and increased ZGB15. 65 provide for further improved fracture toughness. Coiling below 475°C will lead to insufficient acicular / bainitic ferrite and / or bainite formation and loss in precipitation strengthening. Furthermore, it may lead to too high a martensite fraction.

[0089] Though patent document EP1338665-A1 discloses a coiling temperature in the range of 550-700°C, the steels produced having a tensile strength of at least 950 MPa and a hole- expansion capacity of at least 40%, were all produced with coiling temperatures exceeding 600°C. The finish rolling temperatures for all these steels were in the range of 880-930°C. In accordance with the present invention it has been found that too high a coiling temperature leads amongst others to reduced fracture toughness.

[0090] After the steel strip has cooled to room temperature, the oxides (scale) on the hot-rolled steel strips may be removed for example either by pickling in an acid solution (e.g., HCI) at warm temperatures (about 80-120°C) or by a combination of pickling and mechanical brushing of the strip surface. This step may be necessary for rendering the steel strip surface suitable for direct use as uncoated hot-rolled steel or making it amenable to the coating process, when optionally needed for corrosion resistance.

[0091] In an embodiment the thickness of the high-strength hot-rolled steel strip is in a range of about 1.5 to 8 mm, and more preferably of about 1.8 to 6 mm, and most preferably of about 1.8 to 4.5 mm.

[0092] It is an aspect of the invention that the hot-rolled steel strip is subsequently not subjected to a cold rolling operation having a thickness reduction of more than 1.6%. The high-strength hot-rolled steel strip product can be a bare product or uncoated product.

[0093] In an embodiment high-strength hot-rolled steel strip product is provided on one or both of its main surfaces with a thin metallic coating layer, typically up to about 100 g / m2per side of the steel strip, and preferably up to about 50 g / m2per side, preferably being applied by means of heat-to-coat, hot-dip coating, or electrolytic deposition. The metallic coating is preferably selected from the group comprising an aluminium alloy coating (e.g., an Al-Si alloy or Al-Zn alloy), a zinc coating, and a zinc alloy coating (e.g., a Zn-AI alloy, Zn-Mg alloy, Zn-Fe alloy, Zn-AI- Mg alloy, or Zn-Mg-AI alloy).

[0094] The composition of the zinc or zinc alloy coating layer is not limited. Although the coating layer can be applied in various ways, hot-dip galvanising is preferred using a standard Gl coating bath. The Zn based coating layer may comprise a Zn alloy containing Al as an alloying element. A preferred zinc bath composition contains about 0.10-0.35 wt.% Al, the remainder being zinc and unavoidable impurities.

[0095] Other zinc coating layers may also be applied. An example comprises a zinc alloy coating according to patent document W02008 / 102009-A1 and incorporated herein by reference, in particular a zinc alloy coating layer consisting of about 0.3-4.0 wt.% Mg and about 0.05-6.0 wt.% Al, preferably about 0.1-5.0 % Al, and optionally at most about 0.2 wt.% of one or more additional elements along with unavoidable impurities and the remainder being zinc. A preferred Zn bath comprising Mg and Al as main alloying elements has the composition: about 0.5-3.8 wt.% Al, about 0.5-3.0 wt.% Mg, optionally at most 0.2 wt.% of one or more additional elements; the balance being zinc and unavoidable impurities. An additional element typically added in a small amount of less than 0.2 wt.%, could 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 usually added to form spangles. Preferably, the total amount of additional elements in the zinc alloy is at most 0.2 wt.%, and more preferably at most 0.1 wt.%. These small amounts of an additional element do not alter the properties of the coating nor the bath to any significant extent for the usual applications. Preferably, when one or more additional elements are present in the coating, each is present in an amount up to 0.02 wt.%, preferably each is present in an amount up to 0.01 wt.%. Additional elements are usually only added to prevent dross forming in the bath with molten zinc alloy for the hot-dip galvanising, or to form spangles in the coating layer.

[0096] In another embodiment the metallic coating comprises a (commercially pure) aluminium layer or an aluminium alloy layer. A typical metal bath for hot-dip coating such an aluminium layer comprises aluminium alloyed with silicon e.g. aluminium alloyed with about 8-11 wt.% of silicon and at most about 4 wt.% of iron, optionally at most 0.2 wt.% of one or more additional elements such as calcium, unavoidable impurities, the remainder being aluminium. Silicon is present in order to prevent the formation of a thick iron-metallic intermetallic layer which reduces adherence and formability. Iron is preferably present in amounts between about 1 and 4 wt.%, more preferably at about least 2 wt.%.

[0097] In an aspect of the invention it relates to a galvanized steel strip obtained by hot dip galvanizing the high-strength hot-rolled steel strip according to this invention.

[0098] In another aspect of the invention it relates to an automotive component, in particular an automotive chassis part, made from or incorporating the high-strength hot-rolled steel strip according to this invention and taking benefit from amongst others improved balance of strength, formability and the improved fracture toughness, high edge ductility, and good corrosion resistance. The steel strip can be shaped into an automotive component in a cold forming operation, warm forming or hot forming operation as are known in the art. The automotive component includes, but is not limited to, a suspension arm, a reinforcement member, body-in-white frame member, as side member, a seat frame, a seat rail, bumper beam, battery boxes for electrical vehicles, all having an intricate shape. By using the hot-rolled high-strength steel strip, these components can be fabricated with high quality, cost efficiently and with high yields. The high-strength hot-rolled steel product according to this invention can be used also for engineering applications.

[0099] In yet another aspect of the invention it relates to an automotive chassis part using the hot-rolled high-strength steel strip according to and / or produced according to the invention, and wherein the hot-rolled high-strength steel strip for making the automotive chassis part either: has an ultimate tensile strength ranging from 760-950 MPa, a hole-expansion ratio of at least 50%, and preferably of at least 60%, a work of fracture (WOF) of at least 185 J, preferably of at least 200 J, in which ultimate tensile strength (Rm), total tensile elongation (A50), and sheet thickness t (mm) satisfy the equation of (Rm x A50) 1 102≥ 7000, and a corrosion resistance ML-90 of less than 20 g mass loss after 90 cycles in a SAE J2334-April 2016 corrosion test; or has an ultimate tensile strength ranging from 760-950 MPa, a hole-expansion ratio of at least 50%, and preferably of at least 60%, a work of fracture (WOF) of at least 185 J, and preferably of at least 200J, in which ultimate tensile strength (Rm), total tensile elongation (A50), and sheet thickness t (mm) satisfy the equation of (Rm x A50) / 102≥ 8000, and a corrosion resistance ML-90 of less than 20 g mass loss after 90 cycles in a SAE J2334-April 2016 corrosion test; or has an ultimate tensile strength ranging from 760-950 MPa, a hole-expansion ratio of at least 50%, and preferably of at least 60%, a work of fracture (WOF) of at least 185 J, preferably of at least 200 J, and in which ultimate tensile strength (Rm), total tensile elongation (A50), and sheet thickness t (mm) satisfy the equation of (Rm x A50) 1 102≥ 9000, and a corrosion resistance ML-90 of less than 20 g mass loss after 90 cycles in a SAE J2334-April 2016 corrosion test; or has a yield strength (Rp0.2) ranging from 660-820 MPa, a hole-expansion ratio of at least 50%, and preferably of at least 60%, a work of fracture (WOF) of at least 185 J, preferably of at least 200 J, more preferably of at least 225 J, and in which ultimate tensile strength (Rm), total elongation (A50), and sheet thickness t (mm) satisfy the equation of (Rm x A50) 1 102≥ 9000 and a corrosion resistance ML-90 of less than 20 g mass loss after 90 cycles in a SAE J2334-April 2016 corrosion test.

[0100] DESCRIPTION OF THE FIGURES

[0101] The invention shall also be illustrated with reference to the appended non-limiting figures, in which:

[0102] Fig. 1A to Fig 1 E show several features of the method to determine to average total crack length (ATCL) and work of fracture (WOF) of a steel strip product using common cylindrical deep drawing tests for drawing cups.

[0103] Fig. 2A and Fig. 2B show examples of MisOrientation Distribution (MOD) profiles used to explain the physical meaning and definition of the MOD index to characterise the matrix of the microstructure in a quantitative manner.

[0104] The invention will now be illustrated with reference to non-limiting comparative and examples according to the invention.

[0105] EXAMPLE

[0106] Steel ingots of six of dimensions 320 x 100 x 100 mm were cast by melting charges in a vacuum induction furnace. All the ingots were reheated for 1 hour at 1240°C and rough-rolled to 35 mm thickness and left to air cool to ambient temperature before sawing into smaller blocks of 75x105x35 mm. Then, these blocks were reheated again to 1240°C for 40 minutes, and hot rolled to their final thickness of about 3.0 to 3.5 mm in 4 rolling passes. The chemical compositions of these steels having a thickness (t) of about 3 mm are given in Table 1. Steel composition A, D and F are according to the invention. After the final rolling pass, the hot-rolled steel plates, measuring in length and width about 750x105 mm, were directly transferred to the run-out table (ROT) and actively cooled with a mixture of water and pressurized air (CRROTis the average cooling rate of the surface of the steel strip at the ROT) before being placed in a furnace to replicate slow coil cooling. This was done with furnace temperatures (CT - coiling temperature) of 540°C and 610°C (see Table 2).

[0107] The mechanical properties (Rp, Rm, and A50), hole expansion capacity (HEC), average total crack length (ATCL), work of fracture (WOF), the corrosion resistance in accordance with the methods herein described, and the microstructure features GSNIS, MOD index, the vol.% of acicular ferrite / bainitic ferrite (AF / BF / B), the vol.% of (quasi-)polygonal ferrite (Q)FP, and the vol.% of second-phase constituents, all measured by the methods as herein described, are listed in Table 2 as function of the processing parameters applied.

[0108] The corrosion resistance has been measured in accordance with the SAE J2334-April 2016 corrosion test using a commercial available Q-FOG Controlled Relative Humidity (CRH) Cyclic Corrosion Tester (model CRH1100). In Table 2 the mass loss in gram after 37 cycles (ML-37), after 60 cycles (ML-60), and after 90 cycles (ML-90) are presented, and the results of ML-90 are decisive for the steels. One cycle takes 24 hours.

[0109] From Tables 1 and 2 at least the following can be seen:

[0110] Steel compositions B, C, and E are outside the claimed steel composition at least due to too high a Cr-content.

[0111] Although compositions A, D, and F are according to the invention, steel A-2, D-8 and F-12 having been processed at too high a coiling temperature and do not have the required microstructure and provide too low fracture toughness properties. A high coiling temperature results in a high amount of quasi-polygonal ferrite and consequently in a low MOD index. Whereas a low coiling temperature in accordance with the invention results in a predominantly acicular / bainitic ferrite and / or bainite microstructure and very low amount of quasi-polygonal ferrite, if any, and consequently in a higher MOD index.

[0112] Composition B having too high a Cr content does not provide the required corrosion resistance irrespective of the coiling temperature. Steel B-4 also falls short at least on the hole- expansion capacity (HEC) related to the higher coiling temperature. Composition C has also too high a Cr content, but about half the amount of composition B. Steel C-5 has a better corrosion resistance compared to steel B-3, but still does not provide the required balance in properties as also the work of fracture (WOF) is too low. Steel C-6 falls short on the tensile properties and the fracture toughness related to the high coiling temperature.

[0113] Composition D is according to the invention, but only meets the further requirements in microstructure and mechanical properties when manufactured at a low coiling temperature (steel D-7). When manufactured at high coiling temperature (steel D-8) the microstructure, the tensile properties and the fracture toughness fall short.

[0114] Composition E having too high a Cr content does not provide the required corrosion resistance irrespective of the coiling temperature. Steel E-10 also falls short at least on the fracture toughness related to the higher coiling temperature.

[0115] Composition F is according to the invention, but only meets the further requirements in properties when manufactured at a low coiling temperature (steel F-11). When manufactured at high coiling temperature (steel F-12) at least the microstructure and the fracture toughness fall short.

[0116] From these results it can be seen that in accordance with this invention a hot-rolled steel strip can be provided when having the narrow alloy compositional ranges avoiding the presence of substantial amounts of Cr in combination with being manufactured using a low coiling temperature of less than 560°C, and preferably of less than 550°C, having the required microstructure, and it is this combination that results in an desirable improved balance of a combined high edge ductility as expressed by hole-expansion capacity (HEC) of at least 50% with adequate fracture toughness and work of fracture (WOF) of at least 185 J, and a very good corrosion resistance when tested in accordance with the SAE J2334-April 2016 corrosion test, viz. a mass loss of less than 20 g after 90 cycles, and the hot-rolled steel strip has a yield strength (Rp0.2) ranging between 660-820 MPa, a tensile strength (Rm) ranging between 760- 950 MPa, and a tensile elongation A50 of 11% or higher. Table 1. Chemical rnmpositiur of the steels, in wt.%.The bdk'iue is he arc mevtab e impurities.

[0117] Table 2. Processing condition, microstructure, tensile properties, edge ductility, fracture toughness and corrosion resistance of the steels tested.

Claims

CLAIMS1. Hot-rolled high-strength steel strip having a composition, in wt.%,C: 0.035 - 0.120 wt.%,Mn: 1.00 - 2.20 wt.%,Ti: up to 0.150 wt.%,Si: up to 0.65 wt.%,Mo: up to 0.50 wt.%,Nb: up to 0.100 wt.%,Cr: up to 0.090 wt.%,Al: up to 0.10 wt.%,P: up to 0.02 wt.%,S: up to 0.01 wt.%,N: up to 0.01 wt.%, optionally one or more elements selected from the group consisting of: up to 0.30 wt.% V, up to 0.20 wt.% Cu, up to 0.50 wt.%Ni , up to 0.005 wt.% B, and balance Fe and inevitable impurities; and wherein the steel strip has a microstructure of: precipitation-strengthened acicular / bainitic ferrite and / or bainite, and at most 5 vol.% of second-phase constituents, including any cementite, pearlite, martensite, and / or retained-austenite; and wherein the steel strip has at least the following mechanical and corrosion resistance properties: a yield strength (Rp0.2) ranging from 660 to 820 MPa, an ultimate tensile strength (Rm) ranging between 760 and 950 MPa, a tensile elongation (A50) of at least 11%, a hole-expansion capacity (HEC) of at least 50%, a work of fracture (WOF) of at least 185 J, a corrosion resistance ML-90 of less than 20 g mass loss after 90 cycles in a SAE J2334 corrosion test-April 2016, and wherein the microstructure, the yield strength, the ultimate tensile strength, the tensile elongation, the hole-expansion capacity, and the work of fracture are measured as described in the description.

2. Hot-rolled high-strength steel strip according to claim 1 , wherein the composition has a Cr-content of up to 0.080 wt.%, preferably of up to 0.050 wt.%, and more preferably of up to 0.030 wt.%.

3. Hot-rolled high-strength steel strip according to claim 1 or 2, wherein the composition has a C content in a range of 0.045 to 0.120 wt.%, and preferably in a range of 0.060 to 0.100 wt.%.

4. Hot-rolled high-strength steel strip according to any one of claims 1 to 3, wherein the composition has a Ti-content of at least 0.050 wt.%, and preferably at least 0.080 wt.%.

5. Hot-rolled high-strength steel strip according to any one of claims 1 to 4, wherein the composition has a Si-content in a range of 0.10 to 0.65 wt.%, preferably of 0.10 to 0.50 wt.%, and more preferably of 0.10 to 0.40 wt.%, and most preferably of 0.10 to 0.30 wt.%.

6. Hot-rolled high-strength steel strip according to any one of claims 1 to 5, wherein the composition hasMn: 1.00 to 2.10 wt.%; Mo: up to 0.35 wt.%;Cr: up to 0.050 wt.%; andSi: up to 0.50 wt.%, preferably of 0.10 to 0.30 wt.%.

7. Hot-rolled high-strength steel strip according to any one of 1 to 6, wherein the compositionMn: 1 .00 to 1 .65 wt.%, preferably 1 .00 to 1 .50 wt.%;Mo: 0.050 to 0.50 wt.%, preferably 0.150 to 0.350 wt.%;Cr: up to 0.090 wt.%, preferably up to 0.050 wt.%; andSi: up to 0.65 wt.%, preferably up to 0.50 wt.%, more preferably 0.10 to 0.30 wt.%.

8. Hot rolled high-strength steel strip according to any one of claims 1 to 6, wherein the composition has:Mn: 1.50 to 2.20 wt.%, preferably 1.60 to 2.10 wt.%; most preferably 1.70 to 2.10 wt.%;Mo: up to 0.10 wt.%, preferably up to 0.050 wt.%;Cr: up to 0.090 wt.%, preferably up to 0.050 wt.%; andSi: up to 0.30 wt.%, preferably 0.10 to 0.30 wt.%.

9. Hot-rolled high-strength steel strip according to any one of claims 1 to 8, wherein the microstructure has a MisOrientation Distribution (MOD) index of at least 0.85 at Ci- thickness, preferably of at least 0.95, more preferably of at least 1.0, as measured as described in the description.

10. Hot-rolled high-strength steel strip according to any one of claims 1 to 9, wherein the microstructure has at most 3 vol.% second-phase constituents, including any cementite, pearlite, martensite, and / or retained-austenite.

11. Hot-rolled high-strength steel strip according to any one of claims 1 to 10, wherein the steel strip has an average total crack length (ATCL) of less than 20 mm, preferably of less than 15 mm, and more preferably of less than 10 mm, as measured as described in the description12. Hot-rolled high-strength steel strip according to any one of claims 1 to 11 , wherein the steel strip is provided with a metallic coating layer, preferably selected from the group comprising: a Zn-layer, a Zn-based alloy layer, a Mg-based alloy layer, an Al-based alloy layer.

13. Method of manufacturing a hot-rolled high-strength steel strip according to any one of claims 1 to 12, the method comprising the steps of: casting a slab, followed by the step of reheating the solidified slab to a temperature between 1050°C and 1260°C and hot rolling said slab, or casting a slab or strip followed directly by the step of hot rolling said slab or strip; hot rolling the steel slab or strip and finishing said hot rolling at a finish rolling temperature between 960°C and Ar3+20°C, and preferably between 960°C and 900°C, and more preferably between 960°C and 920°C, and wherein Ar3 is the temperature at which transformation of austenite to ferrite starts during cooling;Cooling the hot-rolled strip to the coiling temperature with an average cooling rate of at least 35 °C / s, and preferably in a range of 35 to 60 °C / s;coiling the hot-rolled and cooled strip at a coiling temperature between 475°C and 560°C, and preferably between 500°C and 550°C; allowing the coiled hot-rolled steel strip to further cool to ambient temperature; optionally pickling of the hot-rolled steel strip; and - optionally providing of the hot-rolled steel strip with a metallic coating layer, preferably selected from the group comprising: a Zn-layer, Zn-based alloy layer, a Mg- alloy layer, an Al-based alloy layer, and preferably being applied by means of heat-to-coat, hot-dip coating, or electrolytic deposition.

14. An automotive component made from the high-strength hot-rolled steel strip according to any one of claims 1 to 12 or made from the high-strength hot-rolled steel strip obtained by the method according to claim 13.

15. An automotive component according the claim 14 wherein the component is an automotive chassis part.

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