FIELD OF THE INVENTION

The invention relates to a hot-rolled high-strength steel strip and which is in particular suitable for 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.

BACKGROUND TO THE INVENTION

It is well-known that as the strength of hot-rolled (HR) steel strip increases, the formability decreases. A major area of application for HR steels in transport and automotive applications is the chassis and suspension (C&S). Other areas include frame rails of trucks, bumper beams or battery boxes for electrical vehicles. The typical thickness of HR steels used for these applications is 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. Therefore, ultra-high strength steels with ultimate tensile strength (Rm) of typically over 950 MPa would be useful for this purpose. These applications of the HR steels demand mechanical properties that are difficult to reconcile. Beside a 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.

To save component weight, the common approach is to apply high-strength steel and to reduce the thickness of the steel strip used to save weight. 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.

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 and delamination during shearing, which impairs sheared- edge fatigue.

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 to 0.15 % C, 1.5 % or less Si, 0.5 to 1.6 % Mn, 0.04 % or less P, 0.005 % or less S, 0.04 % or less Al, 0.03 to 0.15 % Ti, 0.03 to 0.5 % Mo, by mass, and balance of Fe and inevitable impurities in a temperature range from 1150 to 1300 °C; hot rolling the reheated steel slab at a finishing temperature of the Ar3 transformation temperature or above into a hot rolled steel sheet; primarily cooling the hot rolled steel sheet in a temperature range from 700 to 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 to 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 to 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 to 95%, and the percentage of the other phase being 5% or less.

Patent document EP1338665-A1 discloses a method for manufacturing a high strength hot rolled 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 to 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 to 0.6 % Mo, 0.02 to 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 to 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 10 ⁴ /μm ³ or higher.

There is a demand for hot-rolled high-strength steel strips having high formability characteristics and having reduced crack sensitivity.

DESCRIPTION OF THE INVENTION

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

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.

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 0.03% Cr may include a steel strip composition having no Cr.

It is an object of the invention to provide a hot-rolled high-strength steel strip having a high total elongation (A50 or A80) combined with a high hole-expansion capacity.

It is an object of the invention to provide a hot-rolled high-strength steel strip having a high total elongation (A50 or A80) combined with a high hole-expansion capacity and having reduced crack sensitivity, in particular reduced edge-crack sensitivity.

It is another object of the invention to provide a method of manufacturing such a hot- rolled high-strength steel strip having an improved balance of total elongation, hole-expansion capacity and reduced crack sensitivity, in particular reduced edge-crack sensitivity.

These and other objects and further advantages are met by the present invention providing a hot-rolled high-strength steel strip having a composition consisting of, in wt.%, 0.02 - 0.13 wt.% C, 1.20 - 3.50 wt.% Mn,

0.10 - 1.0 wt.% Si;

0.05 - 1.0 wt.% Mo;

0.02 - 0.70 wt.% Al,

0.04 - 0.25 wt.% Ti, up to 0.010 wt.% N (100 ppm), preferably up to 0.0065 wt.% N (65 ppm), up to 0.02 wt.% P, preferably up to 0.015 wt.% P, up to 0.01 wt.% S, preferably up to 0.0025 wt.% S (25 ppm), up to 0.0050 wt.% B (50 ppm), preferably up to 0.0030 wt.% B (30 ppm), optionally one or more elements selected from the group consisting of:

(up to 1.5 wt.% Cu, preferably up to 0.6 wt.% Cu, and more preferably up to 0.10 wt.% Cu, up to 0.70 wt.% Cr, preferably up to 0.40 wt.% Cr, and more preferably up to 0.25 wt.% Cr, up to 0.50 wt.% Ni, preferably up to 0.3 wt.% Ni, and more preferably up to 0.10 wt.% Ni, up to 0.30 wt.% V, preferably up to 0.20 wt.% V, up to 0.10 wt.% Nb, preferably up to 0.03 wt.% Nb), and balance being Fe and inevitable impurities resulting from the ironmaking and steelmaking process; and wherein the steel strip has a microstructure of at least 95 vol.% ferrite and at most 5 vol.% martensite, and preferably at least 0.2 vol.% martensite; has a MisOrientation Distribution (MOD) index of at least 0.65 at %-thickness; an area fraction of at most 55% with a Kernel Average Misorientation (KAM) of 0-1; an area fraction of at least 45% with a Kernel Average Misorientation (KAM) of 1-5; and a total grain boundary length (ΣGB[5°-65 ⁰ ]) of at least 1400 1/mm; and wherein the steel strip has at least the following mechanical properties: an ultimate tensile strength (Rm) of at least 950 MPa, preferably of at least 960 MPa, and more preferably at least 980 MPa; a total elongation (A50) of at least 10%, and preferably of at least 14%; and a hole expansion ratio (λ) value of at least 40%.

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 high strength (Rm), total elongation (A50) and hole expansion ratio. It is also an important finding of this invention that the microstructure of the steel strip results in a reduced edge crack sensitivity in a forming operation. A reduced crack sensitivity, in particular reduced edge-crack sensitivity, can be objectively expressed in an average total crack length (ATCL) following the measuring method as herein described. In a preferred embodiment the steel strip according to this invention has an ATCL of less than 45 mm, and in the best examples of less than 40 mm.

In an embodiment of the steel strip it has a yield strength (Rp) is at least 800 MPa, and preferably of at least 850 MPa. In an embodiment of the steel strip it has a yield strength (Rp) of maximum 960 MPa, and preferably of maximum 950 MPa.

The steel strip has an ultimate tensile strength (Rm) of at least 950 MPa. In an embodiment of the steel strip it has an ultimate tensile strength of at least 960 MPa and preferably of at least 980 MPa.

The steel strip has a total elongation (A50) of at least 10%, and preferably of at least 14%.

The required microstructure for this steel strip is achieved by the narrow compositional ranges and by careful control of the manufacturing process, the accelerated cooling of the steel strip on the run-out table (ROT) and the narrow operating window for the coiling temperature (CT) in particular.

The microstructure of the steel strip according to this invention consists of at least 95 vol.% ferrite, and which is precipitation-strengthened with carbide precipitates of titanium and molybdenum, optionally with vanadium and/or niobium; and at most 5 vol.% martensite, balance is inevitable amounts of inclusions, and the sum adding up to 100 vol.%. In an embodiment the microstructure has at least 98 vol.% ferrite, and more preferably at least 99 vol.%. In accordance with the invention it is assumed that the presence of some martensite is favourable to blunt any crack tip and reduces crack propagation. In practice the martensite may contain some traces of retained austenite, but retained austenite is preferably not present. The microstructure of the steel strip has preferably at least 0.2 vol.% martensite, and more preferably at least 0.3 vol.%. In a preferred embodiment the steel strip has at most 3 vol.% martensite, and more preferably at most 2 vol.%.

The texture of the overall microstructure of the steel strip is further characterised by a sufficiently high MisOrientation Distribustion (MOD) index of at least 0.65, preferably of at least 0.80, and more preferably of at least 0.83, and in the best examples of at least 0.90. The MOD index is to be measured at quarter-thickness of the steel strip. The texture of the overall microstructure of the steel strip is further characterized by an area fraction of at most 55%, and more preferably of at most 52%, with a Kernel Average Misorientation (KAM) of 0-1, and an area fraction of at least 45%, and more preferably of at least 48%, with a Kernel Average Misorientation (KAM) of 1-5.

The overall microstructure of the steel strip is further characterised by a (5° to 65°) grain boundary length (ΣGB) of at least 1400 per mm/mm ² or 1/mm, and preferably of at least 1500 1/mm. A higher ΣGB indicates a higher degree of grain refinement resulting in a desired reduced crack sensitivity and suppression or arrest of the crack propagation resulting in a favourable shorter ATCL, for instance when under compression.

Carbon is present in an amount between 0.02 and 0.13 wt.%. To achieve sufficient strength, a suitable minimum C content is 0.04 wt.%, and in a preferred embodiment at least 0.070 wt.%. In a preferred embodiment the C content is at most 0.12 wt.% and is beneficial to suppress the effect of cooling rate dependence on the homogeneity of the final microstructure and to promote high hole-expansion capacity. Furthermore, C is an essential element to achieve precipitation strengthening in combination with carbide-forming micro- alloying elements like titanium, niobium (if added) or vanadium (if added), and to scavenge C to suppress cementite formation in the final microstructure. By optimizing other alloying elements, including Ti, Nb, and/or V, it is possible to obtain an almost uniform ferritic or bainitic-ferritic microstructure with substantially no cementite.

The steel strip has Mn in a range of 1.20 wt.% to 3.50 wt.% to achieve sufficient hardenability and grain refinement. In an embodiment the Mn content is in a range of 1.40 wt.% to 2.40 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.% to obtain sufficient grain refinement improving the fracture toughness and reducing crack susceptibility. In an embodiment the Mn content is maximum 2.20 wt.%, and more preferably maximum 2.0 wt.%. A too high Mn content may lead to segregation during casting which adversely affects the required balance in properties.

Silicon is present in an amount of 0.10 to 1.0 wt.% to improve the strength of the steel by substitutional solid solution strengthening of the iron lattice. Furthermore, Si is beneficial to suppress carbide precipitation (cementite and other carbides). However, when using higher amounts of Si the weldability and coatability of the steel deteriorates, hence the amount of Si is preferably at most 0.95 wt.%, and in a preferred embodiment at most 0.70 wt.% and more preferably at most 0.60 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 0.02 wt.%, the effects of suppression of carbide formation are negligible. Values of aluminium lower than 0.02 wt.% are deemed to be residuals from the deoxidation step during steelmaking, and therefore a minimum value of about 0.02 wt.% is preferred. On the other hand, when Al is above 0.70 wt.% there can be excessive oxide formation during thermomechanical processing (slab reheating, hot rolling, coiling etc.) of the steel. Also, Al increases the ferrite to austenite transformation temperature demanding the need for hot rolling the steel at higher temperatures to finish hot-rolling in the austenitic phase since at lower temperatures intercritical ferrite appears. Higher amounts of 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 exceeds 0.70 wt.% 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 0.02 to 0.70 wt.%, preferably 0.02 to 0.60 wt.% and more preferably in the range of 0.030-0.50 wt.%.

Titanium is another essential alloying element and is present in an amount between 0.04 and 0.25 wt.% as it provides hardenability and 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, at least 0.04 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. When more than 0.25 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 lead to a deterioration of the hole expansion capacity of the steel. Preferably 0.07 to 0.22 wt.% Ti is present. In an embodiment the Ti content does not exceed 0.20 wt.%, preferably it does not exceed 0.15 wt.%, and more preferably it does not exceed 0.13 wt.%.

The steel strip has Mo as a purposive alloying element present in a range of about 0.05 wt.% to 1.0 wt.% to achieve hardenability. In an embodiment the Mo content is at least about 0.10 wt.%, and preferably at least about 0.20 wt.%, and more preferably at least 0.25 wt.%, and most preferably at least 0.30 wt.%. In an embodiment the Mo-content does not exceed about 0.80 wt.%, and preferably does not exceed about 0.55 wt.%, and more preferably does not exceed 0.40 wt.%. The presence of Mo may improve also the weldability of the steel strip.

It is preferred to keep the atomic ratio of Mo with respect to the total amount of micro- alloying elements Nb, Ti, and V, as represented in wt.% by the equation: in the range of 0.75 to 1.25 to increase the thermal stability of carbide precipitates leading to reduced loss in precipitation strengthening from precipitation coarsening due to thermal ageing for example during coil cooling or any subsequent thermal heat treatment of the steel (e.g. during coating or welding).

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 0.010 wt.% N. Amounts higher than these are detrimental for mechanical properties, formability, toughness, 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.% and more preferably only up to 0.0025 wt.%. 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 0.0065 wt.% (65 ppm). A suitable minimum N content is 0.0010 wt.% (10 ppm).

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, and preferably up to 0.0030 wt.% (30 ppm). B is very effective to enhance the hardenability of the steel, which means that a low carbon content and/or lower cooling rates can be used on the run-out table while no or only little pro-eutectic ferrite is formed. B is also an alloying element that is very suitable to increase the yield strength.

Cr can be added to the steel strip in an amount of up to about 0.70 wt.%, and preferably up to 0.25 wt.%, and increases the hardenability of the steel. In a preferred embodiment the steel strip has Cr as a tolerable impurity element. It practice this means it can be present up to 0.10%, preferably up to 0.050%, and more preferably up to 0.03%, as it can impair corrosion performance of the uncoated steel strip substrate.

Copper, when present up to 1.5 wt.%, increases the strength of the steel by both solid solution strengthening as well as precipitation hardening through copper precipitates. In an embodiment the Cu content does not exceed 0.6 wt.%. In an embodiment the Cu is not added as a purposive alloying element and can be present up to 0.10 wt.%, and more preferably up to 0.05 wt.%. Nickel up to about 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 the Ni is not added as a purposive alloying element and can be present up to 0.3 wt.%, preferably up to 0.10 wt.%, and more preferably up to 0.050 wt.%.

Vanadium can be present in the steel up to an amount of about 0.30 wt.%, and preferably up to about 0.20 wt.%, and more preferably up to 0.17 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. In a preferred embodiment V is purposively added in an amount of at least 0.05 wt.%, and more preferably of at least 0.08 wt.%. The combined addition of Ti and V is favoured as Ti offers a catalytic effect for V precipitation, making V precipitation more effective.

Niobium can be present in the steel up to 0.10 wt.%. Nb improves the strength of the steel partly by precipitation hardening but foremost by grain refinement. However, for high amounts of Nb these effects are saturated. Therefore, preferably at most 0.08 wt.% Nb is present, and more preferably at most 0.060 wt.% As Nb is a rather expensive alloying element, in an embodiment Nb is not purposively added and is present as an impurity element and does not exceed 0.03 wt.%, and preferably it does not exceed 0.02 wt.%, and more preferably it does not exceed 0.0050 wt.%. In a most preferred embodiment Nb is an inevitable impurity resulting from the ironmaking and steelmaking process. It has been found that by keeping the Ti at a sufficient high level, the use of Nb can be overcome while still arriving at the aimed balance of formability, mechanical properties and improved fracture or edge-cracking characteristics. Also, Nb 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 and increased sheared-edge quality.

Furthermore, it has been found that cementite formation can be suppressed and the favourable formation of a small fraction of martensite plus retained austenite in the microstructure can be better controlled if the amount of and carbide-forming elements Ti, Nb, V, and Mo represented in wt.% satisfy the equation of: with Ti_sol defined as the amount of free Ti in solution and expressed as: with the amount of Ti and N expressed in wt.%. Preferably, the lower limit of this equation is 0.75, more preferably 0.80. In an embodiment the upper limit is preferably 1.8, and more preferably 1.5, to further suppress the presence of cementite and/or to control the amount of martensite. In a preferred embodiment the formula is in a range of 0.9 to 1.1.

In an embodiment the steel strip has a composition consisting of: 0.02 to 0.13 wt.% C, 1.20 to 2.0 wt.% Mn, 0.10 to 0.60 wt.% Si, 0.01 to 0.70 wt.% Al, 0.04 to 0.25 wt.% Ti, 0.05 to 0.80 wt.% Mo, up to 0.10% Cr, preferably up to 0.050% Cr, up to 0.010 wt.% N, up to 0.02 wt.% P, preferably up to 0.015 wt.% P, up to 0.01 wt.% S, preferably up to 0.0025 wt.% S, up to 0.0050 wt.% B, optionally one or more elements selected from the group consisting of: (up to 0.10% Cu, up to 0.10% Ni, up to 0.30% V, up to 0.10 wt.% Nb, preferably up to 0.03% Nb), and balance Fe and inevitable impurities resulting from the ironmaking and steelmaking process, and with more preferred ranges as herein described and claimed.

The 0.2% offset proof strength or yield strength (Rp), ultimate tensile strength (Rm), uniform elongation (Ag) and tensile elongation (A50) were determined from quasistatic (strain rate 3 x 10' ⁴ s' 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 (Rp or YS). The ratio of yield strength to ultimate tensile strength (Rp/Rm) is expressed as the yield ratio.

The stretch-flangeability of the steel strip or the hole expansion capacity (HEC) was determined by hole expansion tests. Specimens of dimension 90 mm x 90 mm x 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 df was measured when a through-thickness crack formed. The hole expansion ratio λ was calculated using the formula below with d _o = 10 mm:

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

The average total crack length (ATCL) is used to assess the susceptibility of crack formation in situations similar to industrial applications. The ATCL parameter is determined in a common laboratory cylindrical deep drawing tests, using a punch, a draw die, and a blank holder. 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. IB. 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. The blank is a square, measuring 90 x 90 mm. The four corners of the square are 10 mm cut in the direction of the two diagonals as shown in Figure 1A. During the initial stage of the cylindrical deep drawing test, four regions at the edge of the blank are plastically deformed due to high local compressive stresses during drawing. 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. IE). 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. ID) of the deep-drawn cup is measured using a magnifying glass of lOx 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.

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 μm. To obtain a fully deformation free surface, the final polishing step was conducted with colloidal silica (OPS).

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 quarter- 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 μm 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. The EBSD scans were captured using the TexSEM Laboratories (TSL) software: "Orientation Imaging Microscopy (OIM) Data Collection version 7.2". Typically, the following data collection settings were used: Hikari camera at 5 x 5 binning combined with background subtraction (standard mode). 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 100 x 100 μm, with a step size of 0.1 μm, and a scan rate of approximately 100 frames per second. Fe(a) and Fe(y) was used to index the Kikuchi patterns. The Hough settings used during data collections were: Binned pattern size of circa 96; theta set size of 1; rho fraction of circa 90; maximum peak count of 10; minimum peak count of 5; Hough type set to classic; Hough resolution set to low; butterfly convolution mask of 9 x 9; peak symmetry of 0.5; minimum peak magnitude of 10; maximum peak distance of 20.

The EBSD scans were evaluated with TSL OIM Analysis software version "8.0 x64 [12-14- 16]". Typically, the data sets were 90° rotated over the RD axis to get the scans in the proper orientation with respect to the measurement orientation. A standard grain dilation clean-up was performed (Grain Tolerance Angle (GTA) of 5°, a minimum grain size of 5 pixels, criterion used that a grain must contain multiple rows for a single dilation iteration clean-up). Next to this, a pseudo-symmetry clean-up (GTA 5, axis ang 30°@lll) was applied.

The EBSD Image Quality (IQ) maps were used to determine the amount of martensite. Area with a low IQ were identified as MS areas. For the given experimental conditions, typically the low IQ threshold was ~ 0.4 of the peak-maximum position in the IQ histogram. The low IQ threshold was however manually checked for every scan to prevent including grain boundaries from granular bainite or upper bainitic areas in the martensite area fraction.

For calculation of EBSD Kernel Average Misorientation (KAM) maps the fifth nearest neighbour was used with a maximum misorientation of 5° (all points in kernel were used for KAM calculation). The Kernel Average Misorientation is regarded as a signature for the type of ferrite since the Kernel Average Misorientation is a measure for the internal dislocation density. Areas with a relatively low internal dislocation density will predominantly correspond with areas that have a KAM value between 0 and 1°. Areas with a relatively high internal dislocation density will predominantly correspond with areas that have KAM value between 1- 5° .

From the EBSD scans on a 100 x 100 μm area (0.01 mm ² ) on the RD-ND plane at 1/4- thickness, also the sum (ΣGB) of the total grain boundary length of grain boundaries (in mm) with misorientation angles of 5° to 15° (ΣGB5-15) and 15° to 65° (ΣGB15-65) has been measured. The value of ΣGB is expressed in mm' ¹ and is a measure for the density of high-angle grain boundaries. High-angle grain boundaries are effective to arrest crack propagation. Hence, an increased value for ΣGB will be beneficial for increased fracture toughness and reduced crack susceptibility.

In an aspect of the invention it relates to a method of manufacturing a 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°C and 1260°C, preferably for a time of 30 minutes or more, and more preferably of 60 minutes or more, 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 between 820°C and 940°C, preferably between 850°C and 940°C, and most preferably between 850°C and 920°C, and above the Ar3 temperature of the steel. The finish hot rolling temperature (FRT) is above the Ar3 temperature of the steel, 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:

Ar3 = 910°C - 203 x [C] ^1/2 + 44.7 x [Si] - 30 x [Mn] +31.5 x [Mo] ; accelerated cooling the hot rolled steel strip with a run-out table cooling rate between 20 to 250°C/s, and preferably between 40 to 200°C/s, to a temperature on the run-out-table between 560°C and 620°C; coiling the hot-rolled and cooled strip at a temperature between 550°C and 600°C, and preferably between 550°C and 595°C, and more preferably between 550°C and 590°C; allowing the coiled hot-rolled steel strip to further cool to ambient temperature; and pickling the hot-rolled steel strip, 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, an Al-based alloy layer, to provide improved corrosion resistance in service. The metallic coating layer is preferably applied by means heat-to-coat or hot-dip coating,

The method of manufacturing herein described and claimed results in the desired microstructure providing for the aimed improved balance of formability, mechanical properties and fracture characterstics. The invention is also embodied in a steel strip manufactured by the method described herein and claimed having said microstructure and improved balance of formability, mechanical properties and fracture characteristics.

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 mm and 350 mm, and typically of 225 mm to 250 mm, as well as a thin-slab having a cast thickness of between 50 mm and 150 mm in direct strip plant. Schematic examples of a process involving a conventional hot strip mill and of a thin slab casting/direct rolling mill are shown in figure 2A and 2B respectively. For conventional thick-slab casting, reheating of the slab is necessary to reheat the slab from ambient temperatures (usually the think 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 microalloying 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 at FRT>Ar3. Often this requires a (slab) reheating temperature of between 1050°C up to about 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 to 1160 °C. This would also prevent any precipitates from forming when microalloying 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. According to the invention the reheating or the homogenisation time for the thin slab casting route is preferably 30 minutes or more.

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 FRT is above (Ar3 + 30°C), e.g. typically above 850°C, to avoid hot rolling locally below Ar3 at colder edges or the tail of the strip. The FRT should not be too high in the austenite region as a lower FRT will promote more austenite deformation and hence contribute to increased grain refinement and increased ΣGB. A FRT of above 950°C will result in increased edge crack sensitivity. Furthermore, a not too high FRT will also promote textures in the final microstructures that are beneficial for toughness (e.g. {332}<113>) and supresses those that are detrimental (e.g. {001}<110> rotated Cube). For that reason FRT should not exceed about 940°C, preferably it does not exceed about 920°C, and more preferably it does not exceed about 910°C. After hot rolling, the steel strip is accelerated cooled on a run-out table (ROT) to a temperature between 560°C and 620°C. An accelerated cooling rate is desired to suppress recovery and loss of internal stored energy in the austenite in order to promote grain refinement of the final microstructure and increased ΣGB. The cooling rate should be high enough to avoid austenite-to-ferrite phase transformation at elevated temperatures and to preferably promote austenite-to-ferrite phase transformation at relatively low temperatures of about 560 to 630°C on the run-out-table. Increased cooling rate will promote grain refinement, increased ΣGB 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>). There is no critical run-out table cooling rate (ROT-CR) as long as the herein mentioned cooling rate is exceeded through-thickness of the steel strip from a microstructure point of view. However, an unnecessarily high ROT-CR may affect the flatness of the strip after cooling and cause control problems to stop at the correct cooling stop temperature and therefore a suitable maximum ROT-CR is about 250°C/s, preferably about 200°C/s and more preferably about 150°C/s. A practical ROT-CR range is about 20 to 100°C/s, and more preferably about 40 to 100°C/s, as this is achievable through air cooling, laminar cooling or water jet cooling depending on the thickness of the steel strip. For practical reasons the run-out table cooling rate (ROT-CR) is defined as the average cooling rate of the surface of the steel strip.

Next, the hot-rolled steel strip is coiled at a temperature between 550°C and 600°C, and preferably between 550°C and 600°C, and more preferably between about 560°C and 600°C. In an embodiment the hot-rolled strip is coiled at a temperature not exceeding 595°C, and more preferably not exceeding 590°C. The coiling temperature of the steel strip is a key process parameter to arrive at the required microstructure of the steel strip providing for the improved balance in mechanical properties as herein described.

When the coiling temperature is too low 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 not promote any martensite as second-phase constituent in the final microstructure. A too high coiling temperature will also reduce the MOD index, the fraction KAM 0-1 will increase and the fraction KAM 1-5 will be too low. During coil cooling some further precipitation may take place, as well as some further phase transformation. Undesirably, precipitates once formed may coarsen during coil cooling. The alloy composition in combination with the claimed coiling temperature supresses this phenomenon. This coiling temperature will help to promote small grains of ferrite formed during coiling or coil cooling, suppress coarsening of precipitates that strengthen the ferrite matrix, and also promote the formation of a small amount of martensite.

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 ΣGB 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 550°C and 600°C, and preferably between 550°C and 595°C, and more preferably between 550°C and 590°C, in order to allow the austenite-to- ferrite phase transformation to continue at relatively low temperatures to promote fine- grained ferrite, which is precipitation strengthened with carbide precipitates comprising Ti and Mo, and optionally Nb and/or V. The grain refinement and increased ΣGB provides for improved fracture toughness and reduced (edge-) crack susceptibility. Coiling below 550°C will lead to insufficient ferrite formation and loss in precipitation. Furthermore, it may lead to a too high martensite fraction.

Though patent document EP1338665-A1 discloses a coiling temperature in the range of 550 to 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 to 930°C. In accordance with the present invention it has been found that too high coiling temperatures leads amongst others to increased crack susceptibility, in particular edge-crack susceptibility, and also reduced fracture toughness. After the steel strip has cooled to room temperature, the oxides (scale) on the hot- rolled steel strips are removed either by pickling in an acid solution (e.g. HCI) at warm temperatures (80-120°C) or by a combination of pickling and mechanical brushing of the strip surface. This step is 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.

In an embodiment the thickness of the 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 mm.

The hot-rolled steel strip product can be a bare product or uncoated product or it can be provided on one or both of its main surfaces with a thin metallic coating layer, typically up to about 100 g/m ² per side of the steel strip, and preferably up to about 50 g/m ² per side. 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).

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.

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 0.3 to 4.0 wt.% Mg and 0.05% to 6.0 wt.% Al, preferably 0.1 to 5.0 % Al, and optionally at most 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: 0.5 to 3.8 wt.% Al, 0.5 to 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.

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 to 11 wt.% of silicon and at most 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 1 and 4 wt.%, more preferably at least 2 wt.%.

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

In an aspect of the invention it relates to an automotive component, in particular an automotive chassis part, incorporating or made from the hot-rolled high-strength steel strip according to this invention and taking benefit from amongst others improved balance of strength, formability and the improved fracture toughness and reduced (edge-) crack susceptibility. The steel strip can be shaped into an automotive component in a cold forming operation, warm forming and hot forming operation as are known in the art. The automotive component includes 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 efficient and high yields. The high- strength hot-rolled steel product according to this invention can be used also for engineering applications.

BRIEF DESCRIPTION OF THE FIGURES

The invention will now be explained by means of the following, non-limiting figures.

Fig. 1A to Fig IE shows several features of the method to determine to average total crack length (ATCL) of a steel strip product using common cylindrical deep drawing tests for drawing cups.

A schematic drawing of a hot-rolling mill for processing thick cast steel slabs is shown in Figure 2A, and a thin slab casting facility with a direct rolling mill is shown in Figure 2B. The invention will now be illustrated with reference to non-limiting comparative and examples according to the invention.

EXAMPLE

Steel ingots of five inventive (Inv.) chemistries A-E and a comparative (Comp.) steel F of dimensions 320 x 100 x 100 mm were cast by melting charges in a vacuum induction furnace. The chemical compositions of these steels are given in Table 1. In Table 1 also the following two ratios A and B are also listed:

All the ingots were reheated for 1 hour at 1240°C, and rough-rolled to 35 mm thickness. Then, the strips were reheated again to 1220°C for 40 minutes, and hot rolled to their final thickness of about 3.2 to 3.6 mm in 5 rolling passes with FRT's above 870°C and below 910°C which are in the austenitic phase field for all these steels. After the final rolling pass, the hot rolled steels were transferred to the run-out table with a start temperature run-out table (TSTART) in the range of 840 to 880°C and actively cooled from the austenitic phase field with a mixture of water and air to an end temperature in the ferritic phase field at the run-out table (TEND) in the range of 565 to 615°C at a cooling rate in the range of 30 to 70°C/s. Next, the steels were transferred to a furnace to replicate slow coil cooling. This was done with furnace temperatures (CT - coiling temperature) of 540, 580, and 610°C (see Table 2).

Prior to testing of mechanical properties, hole expansion capacity and average total crack length in accordance with the methods herein described, the hot rolled sheets were sand blasted to remove the oxide layer. The microstructures of the hot rolled strips were determined with EBSG according to the methods as herein described. The results of these tests as function of the processing parameters applied are listed in Table 2. Table 1. Chemical composition of the steels, in wt.%. "Inv." is an example according to the invention, and "Comp." is a comparative example.

From the results of Table 2 it can be seen that steel 17 having a composition outside the claimed range and processed according to the invention (CT 580°C) has a far too low Rm. Increasing the coiling temperature (CT 610°C) provides an increased Rm but the edge-crack susceptibility of steel 16 is significantly decreased. This decrease is likely the resultant of a microstructure with too low MOD index and insufficient degree of grain refinement as reflected by the ΣGB. A too low coiling temperature (CT 540°C) for steel 18 results in even lower Rm compared to steel 17.

Steels 1-3 (Alloy A) have all a composition according to the invention and wherein steel

2 has been processed according to the invention (CT 580°C). Steel 2 provides a very good balance of high strength (Rm 1011 MPa), high elongation (A50 15.8%), good hole expansion ratio (HEC 42%) and very good edge-crack resistance (ATCL 28.9 mm). This due to the composition, processing and the microstructure. The microstructure is characterised by a very high degree of grain refinement (ΣGB 1854 1/mm) and high MOD index (1.20). Increasing the CT to 610°C for steel 1 gives still a high Rm, but the edge-crack susceptibility is reduced to unacceptable levels (ATCL 49.3 mm). This is the result of a change of the microstructure as reflected amongst other by too low MOD index (0.52), too high KAM 0-1 and low KAM 1-5, and too little grain refinement as reflected by too low ΣGB. Lowering the CT to 540°C for steel

3 leads to a significant decrease in Rp and Rm.

A similar trend as for steels 1-3 (alloy A) and be seen in steels 4-6 (alloy B). Steels 7-9 (alloy C) have all a composition according to the invention and have a purposive addition of V. Steel 8 has been processed according to the invention with CT 580°C and provides a very good balance of high strength (Rm 1005 MPa), high elongation (A50 14.7%), good hole expansion ratio and very good edge-crack resistance (ATCL of 34.8 mm). This due to the composition, processing and the microstructure. The microstructure is characterised by a very high degree of grain refinement. The microstructure has about 2.6 vol.% of martensite and it is believed that this contributes to the good edge-crack resistance represented by ATCL of 34.8 mm. Increasing the CT to 610°C for steel 7 a slightly increased Rm and Rp but the edge-crack susceptibility of steel 7 is significantly decreased. This decrease is likely the resultant of a change of the microstructure as reflected by the MOD index and KAM values and an insufficient degree of grain refinement as reflected by the ΣGB (ΣGB 1134 1/mm). A too low coiling temperature (CT 540°C) for steel 9 results in a lower Rm compared to steel 8 and a significantly lower Rp and far too low hole expansion ratio.

Alloy D (steels 10-12) has an increased Al-content compared to alloy B (steels 4-6). The comparison of steels 5 and 11 (both processed according to the invention) shows that an increased strength (Rm) will be obtained with the addition of Al. Steel 11 also shows a favourable microstructure as illustrated by the MOD index, KAM values and sufficient degree of grain refinement as illustrated by ΣGB and consequently provides the desired balance of properties.

Alloy E (steels 13-15) has a purposive addition of V compared to alloy D (steels 10-12). Steel 14 provides a further reduced edge-crack susceptibility compared to steel 11. Comparing steel 14 with steel 15 shows that a lower coiling temperature (CT 540°C) results amongst others in a significantly reduced Rp. And comparing steel 13 with steel 14 shows that increasing the CT to 610°C for steel 13 gives still a high Rm and Rp, but the edge-crack susceptibility is reduced to unacceptable levels (ATCL 56.3 mm). This is reflected amongst other by too low MOD index, too high KAM 0-1 and low KAM 1-5, and low ΣGB.

Having now fully described the invention, it will be apparent to one of ordinary skill in the art that many changes and modifications can be made without departing from the spirit or scope of the invention as herein described. Table 2. Process parameters applied and the resultant formability characteristics, mechanical properties, microstructure and texture components of the microstructure. ro o