CN118574944A - Hot rolled high strength steel strip - Google Patents

Abstract

The present invention relates to a hot rolled high strength steel strip having a composition consisting of, in weight %, 0.02-0.13% C, 1.20-3.50% Mn; 0.10-1.00% Si; 0.05-1.0% Mo; 0.02-0.70% Al, 0.04-0.25% Ti, up to 0.010% N, up to 0.02% P, up to 0.01% S, up to 0.005% B, optionally one or more elements selected from the group consisting of: (up to 1.5% Cu, up to 0.70% Cr, up to 0.50% N i, up to 0.30% V, up to 0.10% Nb), and the balance Fe and unavoidable impurities; and wherein the steel strip has a microstructure of at least 95% by volume ferrite and at most 5% by volume martensite, and preferably at least 0.2% by volume martensite; having a misorientation distribution (MOD) index of at least 0.65 at 1/4 thickness; up to 55% area fraction with a core average misorientation (KAM) of 0-1; at least 45% area fraction with a core average misorientation (KAM) of 1-5; and a total grain boundary length (ΣGB[5°-65°]) of at least 1400 mm ^-1 ; and wherein the steel strip has at least the following mechanical properties: an ultimate tensile strength (Rm) of at least 950 MPa, preferably at least 960 MPa, a total elongation (A50) of at least 10%, and a hole expansion ratio (λ) value of at least 40%. The invention also relates to a method for producing such a hot rolled steel strip and to an automotive part comprising the hot rolled steel strip.

Description

Hot rolled high strength steel strip

Technical Field

The present invention relates to a hot rolled high strength steel strip, which is particularly suitable for automotive components. The present invention also relates to a method for producing such a hot rolled high strength steel strip. Furthermore, the present invention relates to an automotive part incorporating the hot rolled high strength steel strip.

Background Art

It is well known that with the increase of the strength of hot rolled (HR) steel strip, the formability decreases. The main application areas of HR steel in transportation and automotive applications are chassis and suspension (C&S). Other areas include frame rails of trucks, bumper beams or battery boxes of electric vehicles. The typical thickness of HR steel used for these applications is less than 4.5mm. Thicker HR steel strips (e.g., up to 12mm) can be used for engineering applications, such as crane booms, or for transportation applications of heavy truck frames. From the perspective of weight reduction, the above applications must adopt higher strength steel so that the specifications of the steel strip can be reduced. Therefore, ultra-high strength steels with an ultimate tensile strength (Rm) of generally more than 950MPa will be used for this purpose. These applications of HR steel require mechanical properties that are difficult to coordinate. In addition to high strength, the steel strip should also have good formability for manufacturing components via, for example, cold forming, because cold forming is an energy-saving manufacturing route compared to hot forming. In addition, good impact and fracture toughness or energy absorption capacity is also required for applications such as bumper beams, battery housings, crane arms or frame rails. In order to assemble the component, good weldability is also required. However, as the tensile strength of the steel increases, the formability parameter decreases. Formability is a general term for steel sheets, which is considered to be a combination of the material's behavior during several mechanical operations such as stretching, bending, drawing and hemming. Depending on the component geometry, any one or a combination of two or more properties of the material are important during sheet metal forming.

In order to reduce the weight of the component, a common approach is to apply high-strength steel and reduce the thickness of the steel strip used to reduce the weight. However, this may result in a loss of stiffness, which is undesirable for some applications in automotive parts. The inherent stiffness loss caused by reducing the thickness of the steel strip used to make automotive parts can be regained by optimizing the component geometry, such as creating deeper flanges or flanges with increased stretch or bending. In order to allow for increased component stiffness via geometry optimization, high-strength steel strips require excellent formability in terms of tensile elongation and hole expansion capability.

Single-phase precipitation-hardened 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 capability or HEC) with both formability modes at high levels may have low fracture toughness values and increased edge crack sensitivity under certain conditions (e.g., under compression), and are prone to increased susceptibility to unstable brittle fracture behavior and delamination during shearing, which impairs shear edge fatigue.

Patent document EP1616970-A1 discloses a method for manufacturing a high-strength hot-rolled steel plate, comprising the following steps: reheating a steel slab at a temperature range of 1150°C to 1300°C, wherein the steel slab is composed of, in terms of weight %, 0.04 to 0.15 mass % C, 1.5 mass % or less Si, 0.5 to 1.6 mass % Mn, 0.04 mass % or less P, 0.005 mass % or less S, 0.04 mass % or less Al, 0.03 to 0.15 mass % Ti, 0.03 to 0.5 mass % Mo and the balance Fe and inevitable impurities; hot rolling the reheated steel slab into a hot rolled steel sheet at a finishing temperature of Ar3 transformation temperature or higher; primary cooling the hot rolled steel sheet at an average cooling rate of 20°C/second or higher in a temperature range of 700 to 850°C; holding the primary cooled steel sheet at a temperature of 680°C or higher for more than 1 second; and secondary cooling the steel sheet at a temperature of 550°C or lower at an average cooling rate of 30°C/second or higher, followed by coiling the steel sheet. Preferably, the hot rolled steel sheet is not only primary cooled to a temperature range of 700 to 850°C, but also primary cooled to a temperature range of (SRT/3+300) to (SRT/8+700)°C, where SRT represents the reheating temperature of the steel slab. In the embodiment of EP1616970-A1, this results in a primary holding temperature of 667 to 860°C in practice. The processing conditions are such that a microstructure consisting of precipitates containing ferrite, a second phase of bainite and/or martensite, and other phases is obtained, wherein the percentage of precipitates containing ferrite is 40 to 95% and the percentage of other phases is 5% or less.

Patent document EP1338665-A1 discloses a method for manufacturing high-strength hot-rolled steel plates, which includes the following steps: producing a steel slab, which is basically composed of 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 Fe, in terms of weight %, and satisfies the equation 0.8≤(C/12)/[(Ti/48)+(Mo/96)]≤1.3; producing a hot-rolled steel plate by hot rolling the steel slab at a temperature of Ar3 transformation point or higher; and coiling the hot-rolled steel plate at a temperature of 550 to 700°C. The processing conditions provide a microstructure consisting essentially of a matrix of a ferrite structure single phase and fine precipitates as composite carbides containing Ti and Mo dispersed in the matrix having a grain size of less than 10 nm, wherein the fine precipitates are dispersed in an amount of 5×10 ⁴ /μm ³ or more per unit volume.

There is a need for hot rolled high strength steel strip having high formability characteristics and with reduced crack sensitivity.

DETAILED DESCRIPTION OF THE INVENTION

As understood herein, with respect to any description of alloy compositions or preferred alloy compositions, all references to percentages are by weight unless otherwise stated.

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 understood by those skilled in the art.

As used herein, the terms "up to" and "up to about" expressly include, but are not limited to, the possibility of zero weight percent of the particular alloying component to which they refer. For example, up to 0.03% Cr may include a steel strip composition having no Cr.

An object of the present invention is to provide a hot rolled high strength steel strip having a high total elongation (A50 or A80) and a high hole expansion capability.

An object of the present invention is to provide a hot rolled high strength steel strip having a high total elongation (A50 or A80) and a high hole expansion capability and having reduced crack sensitivity (in particular reduced edge crack sensitivity).

Another object of the present invention is to provide a method for producing such hot rolled high strength steel strip having an improved balance of total elongation, hole expansion capability and reduced crack sensitivity, in particular reduced edge crack sensitivity.

These and other objects and further advantages are met by the present invention, which provides a hot rolled high strength steel strip having a composition, in weight %, consisting of:

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);

One or more elements selected from the following:

(up to 1.5 wt. % Cu, preferably up to 0.6 wt. % Cu, and more preferably up to 0.10 wt. % Cu,

at most 0.70 wt. % Cr, preferably at most 0.40 wt. % Cr, and more preferably at most 0.25 wt. % Cr,

at most 0.50 wt. % Ni, preferably at most 0.3 wt. % Ni, and more preferably at most 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),

The balance is Fe and unavoidable impurities arising from the iron and steel making processes; and

- wherein the steel strip has a microstructure of at least 95% by volume of ferrite and at most 5% by volume of martensite, and preferably at least 0.2% by volume of martensite;

- having a MisOrientation Distribution (MOD) index of at least 0.65 at 1/4 thickness;

- up to 55% area fraction with a kernel average misorientation (KAM) of 0-1;

- an area fraction of at least 45% with a kernel average misorientation (KAM) of 1 to 5;

- and a total grain boundary length (ΣGB[5°-65°]) 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 at least 960 MPa, and more preferably at least 980 MPa;

A total elongation (A50) of at least 10%, and preferably at least 14%; and

A hole expansion ratio (λ) value of at least 40%.

According to the present invention, it has been found that hot rolled steel strips having these narrow alloy composition ranges in combination with a microstructure provide an improved balance of high strength (Rm), total elongation (A50) and hole expansion. Another important finding of the present invention is that the microstructure of the steel strip results in reduced edge crack sensitivity during forming operations. Reduced crack sensitivity, in particular reduced edge crack sensitivity, can be objectively expressed as an average total crack length (ATCL) according to the measurement method described herein. In a preferred embodiment, the steel strip according to the present invention has an ATCL of less than 45 mm, and in the best example less than 40 mm.

In one embodiment of the steel strip, its yield strength (Rp) is at least 800 MPa, preferably at least 850 MPa. In one embodiment of the steel strip, its yield strength (Rp) is at most 960 MPa, preferably at most 950 MPa.

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

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

The desired microstructure of the steel strip is achieved by a narrow composition range and by careful control of the manufacturing process, accelerated cooling of the steel strip on the run-out table (ROT) and in particular a narrow operating window of the coiling temperature (CT).

The microstructure of the steel strip according to the invention consists of at least 95% by volume of ferrite, which is precipitation strengthened with carbide precipitates of titanium and molybdenum, optionally with vanadium and/or niobium; and at most 5% by volume of martensite, the remainder being inclusions in an unavoidable amount, and the sum totaling 100% by volume. In one embodiment, the microstructure has at least 98% by volume, more preferably at least 99% by volume of ferrite. According to the invention, it is believed that the presence of some martensite is beneficial for blunting any crack tips and reducing crack propagation. In practice, the martensite may contain some traces of retained austenite, but preferably no retained austenite is present. The microstructure of the steel strip preferably has at least 0.2% by volume, and more preferably at least 0.3% by volume of martensite. In a preferred embodiment, the steel strip has at most 3% by volume, and more preferably at most 2% by volume of martensite.

The overall microstructural texture of the steel strip is further characterized by a sufficiently high misorientation distribution (MOD) index of at least 0.65, preferably at least 0.80, more preferably at least 0.83, and in the best case at least 0.90. The MOD index is measured at one quarter of the 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%, more preferably at most 52%, having a core average misorientation (KAM) of 0-1 and an area fraction of at least 45%, more preferably at least 48%, having a core average misorientation (KAM) of 1-5.

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

Carbon exists with the amount of 0.02 to 0.13 % by weight.In order to obtain enough intensity, suitable minimum C content is 0.04 % by weight, and is at least 0.070 % by weight in a preferred embodiment.In a preferred embodiment, C content is at most 0.12 % by weight, and is conducive to suppressing the influence of cooling rate dependency on the uniformity of final microstructure and promoting high hole expansion ability.In addition, C is to realize precipitation strengthening and remove C to suppress the necessary element that cementite in the final microstructure forms with the microalloying element such as titanium, niobium (if adding) or vanadium (if adding) that forms carbide.By optimizing other alloying elements, including Ti, Nb and/or V, it is possible to obtain almost uniform ferrite or bainite-ferrite microstructure without cementite substantially.

The steel strip has 1.20 wt% to 3.50 wt% Mn to achieve sufficient hardenability and grain refinement. In one embodiment, the Mn content is 1.40 wt% to 2.40 wt% to improve the balance of strength, corrosion resistance, fracture toughness and edge crack sensitivity. Preferably, the Mn content is at least 1.50 wt% to obtain sufficient grain refinement, improve fracture toughness and reduce crack sensitivity. In one embodiment, the Mn content is a maximum of 2.20 wt%, more preferably a maximum of 2.0 wt%. Too high a Mn content may lead to segregation during casting, which adversely affects the desired property balance.

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. In addition, Si helps to inhibit carbide precipitation (cementite and other carbides). However, when higher amounts of Si are used, the weldability and coatability of the steel deteriorate, so the amount of Si is preferably at most 0.95 wt %, and in a preferred embodiment at most 0.70 wt %, more preferably at most 0.60 wt %.

In the steel according to the invention, the performance of aluminum is comparable to that of Si. Aluminum slows down the carbide precipitation kinetics and inhibits the formation of cementite. When Al is less than 0.02 wt%, the effect of inhibiting carbide formation is negligible. Values of aluminum below 0.02 wt% are considered to be residues from the deoxidation step during steelmaking, so a minimum value of about 0.02 wt% is preferred. On the other hand, when Al is above 0.70 wt%, there may be excessive oxide formation during the thermomechanical processing of steel (slab reheating, hot rolling, coiling, etc.). In addition, Al increases the transition temperature from ferrite to austenite, requiring hot rolling of the steel at higher temperatures to complete hot rolling in the austenite phase because intercritical ferrite appears at lower temperatures. Higher amounts of oxidation may occur at higher temperatures. These oxide scales are harmful to hot rolling, pickling, coating and overall surface appearance. In addition, when Al exceeds 0.70 wt%, the rolling force during hot rolling increases, which, combined with the presence of Si, makes the steel very brittle and more difficult to hot roll. Therefore, Al is present in the present invention in an amount of 0.02 to 0.70 wt%, preferably 0.02 to 0.60 wt%, and more preferably 0.030-0.50 wt%.

Titanium is another essential alloying element and is present in an amount of 0.04 wt % to 0.25 wt % because it provides hardenability and acts as a carbide forming element, inhibiting 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 carbon sulfides, depending on the specific chemical composition of the steel. Therefore, there is at least 0.04 wt % Ti to combine substantially all N and S in the steel, and there is enough excess Ti to combine with C in the steel. When there is more than 0.25 wt % Ti, coarse Ti nitrides, carbonitrides and carbides may be formed, which are difficult to dissolve during the slab reheating process before hot rolling. In addition, these coarse Ti nitrides, carbonitrides and carbides cause the deterioration of the hole expansion ability of the steel. Preferably, there is 0.07 to 0.22 wt % Ti. In one embodiment, the Ti content is no more than 0.20 wt %, preferably no more than 0.15 wt %, and more preferably no more than 0.13 wt %.

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

The atomic ratio of Mo relative to the total amount of microalloying elements Nb, Ti and V is preferably expressed in weight % by the following equation:

Maintained in the range of 0.75 to 1.25 to improve thermal stability of carbide precipitates, resulting in reduced precipitation strengthening losses from precipitation coarsening caused by, for example, thermal aging during coil cooling or any subsequent heat treatment of the steel, such as during coating or welding.

Nitrogen, sulfur and phosphorus are residual elements present in steel due to steelmaking and refining processes. Their amounts are limited to at most 0.01% by weight of S, at most 0.02% by weight of P and at most 0.010% by weight of N. Amounts above these are harmful to mechanical properties, formability, toughness and weldability. In one embodiment, P is present only at most 0.015% by weight. In one embodiment, S is present only at most 0.005% by weight, more preferably at most 0.0025% by weight. N forms titanium nitride with Ti, which acts as a dispersoid for austenite grain size control during reheating. However, too high N may result in too much coarse TiN particles, which may impair hole expansion capability. Preferably, N content is at most 0.0065% by weight (65ppm). Suitable minimum N content is 0.0010% by weight (10ppm).

Boron is not required to obtain the desired balance of properties of the steel strip, but may be present in an amount of up to 0.005 wt. % (therefore up to 50 ppm) and preferably up to 0.0030 wt. % (30 ppm). B is very effective in increasing the hardenability of the steel, which means that low carbon contents and/or lower cooling rates can be used on the run-out table with no or only minimal formation of proeutectic ferrite. B is also an alloying element very suitable for increasing yield strength.

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

Copper improves the strength of steel by solution strengthening and precipitation hardening by copper precipitates when present at most 1.5 wt %. In one embodiment, the Cu content is no more than 0.6 wt %. In one embodiment, Cu is not added as a target alloying element and may be present at most 0.10 wt %, and more preferably at most 0.05 wt %. Up to about 0.50 wt % and preferably at most 0.30 wt % nickel improves impact toughness and offsets any hot brittleness that may occur during hot working of steel strip due to the presence of copper. In one embodiment, Ni is not added as a target alloying element and may be present at most 0.3 wt %, preferably at most 0.10 wt %, and more preferably at most 0.050 wt %.

Vanadium can be present in steel with about 0.30 % by weight at the most, preferably about 0.20 % by weight at the most and more preferably 0.17 % by weight at the most. But V is a relatively expensive alloying element, which is mainly used to replace the precipitation strengthening effect of Ti, and avoids cementite formation by forming vanadium carbide. In a preferred embodiment, V is purposefully added with at least 0.05 % by weight, more preferably at least 0.08 % by weight. It is favourable for the combination of Ti and V to add, because Ti is for V to separate out and provide catalytic effect, so that V is separated out more effectively.

Niobium can be present in the steel at up to 0.10 wt%. Nb improves the strength of the steel partly by precipitation hardening, but most importantly by grain refinement. However, for large amounts of Nb, these effects are saturated. Therefore, it is preferred that up to 0.08 wt% Nb is present, and more preferably up to 0.060 wt%, because Nb is a relatively expensive alloying element, in one embodiment, Nb is not purposefully 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 the most preferred embodiment, Nb is an unavoidable impurity resulting from the ironmaking and steelmaking processes. It has been found that by keeping Ti at a sufficiently high level, the use of Nb can be overcome while still achieving the target balance of formability, mechanical properties and improved fracture or edge cracking characteristics. In addition, due to centerline segregation, Nb has a high tendency to segregate and form coarse NbC particles, thereby avoiding the use of Nb leads to improved hole expansion ability and improved shear edge quality.

Furthermore, it has been found that cementite formation can be suppressed and the favorable formation of a small portion of martensite plus retained austenite in the microstructure can be better controlled if the amounts of carbide-forming elements Ti, Nb, V and Mo expressed in weight % satisfy the following equation:

where Ti_sol is defined as the amount of free Ti in the solution and is expressed as:

Wherein the amount of Ti and N is expressed in weight %. Preferably, the lower limit of the equation is 0.75, more preferably 0.80. In one embodiment, the upper limit is preferably 1.8, more preferably 1.5, to further suppress the presence of cementite and/or control the amount of martensite. In a preferred embodiment, the formula is in the range of 0.9 to 1.1.

In one 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 wt. % Cr, preferably up to 0.050 wt. % Cr, up to 0.010 wt. % N, up to 0.02 wt. % P, preferably up to 0.0 15 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 following: (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 the balance Fe and inevitable impurities generated by ironmaking and steelmaking processes, and the more preferred ranges are as described and claimed herein.

The 0.2% offset proof strength or yield strength (Rp), ultimate tensile strength (Rm), uniform elongation (Ag) and tensile elongation (A50) are determined by quasi-static (strain rate 3×10 ^-4 s ^-1 ) tensile testing at room temperature with an A50 specimen geometry according to EN 10002-1/150 6892-1 using a tensile test parallel to the rolling direction. The geometry of the tensile specimen comprises a 50 mm gauge length in the rolling direction, a 12.5 mm width and a thickness depending on the final specification. The strength of the steel at a 0.2% offset strain is measured as the yield strength (Rp or YS). The ratio of the yield strength to the ultimate tensile strength (Rp/Rm) is denoted as the yield ratio.

The extended flange properties or hole expansion capacity (HEC) of steel strips are determined by hole expansion tests. Specimens with dimensions of 90 mm x 90 mm x final thickness of the strip are cut from the coiled steel. A hole with a diameter of 10 mm is punched in the middle of the specimen and the hole expansion test is carried out according to ISO/TS16630:2003(E) standard. The hole expansion test of the specimen is carried out with upper deburring. A 60° conical punch is pushed upward from below and the hole diameter d _f is measured when a through crack is formed. The hole expansion ratio λ is calculated using the following formula, where 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 evaluate the sensitivity of crack formation in situations similar to industrial applications. The ATCL parameter is determined in a common laboratory cylindrical deep drawing test using a punch, a die and a blank holder. In this procedure, the punch has a diameter of 50 mm and a punch radius of 7 mm. The die has an inner diameter of 62 mm and a radius of 8 mm. The setup is schematically shown in Figure 1B. The inner diameter is large enough to allow the edge of the formed cup to move freely. The gap (i.e., the distance between the punch wall and the die wall) is 6 mm. The force of the blank holder is set to 50 kN. The blank is square and measures 90×90 mm. As shown in Figure 1A, the four corners of the square are cut 10 mm in the direction of the two diagonals. During the initial stage of the cylindrical deep drawing test, four areas at the edge of the blank are plastically deformed due to the high local compressive stresses during the stamping process. This results in local wrinkling of the edge. At the end of the test and when the force of the blank gripper is released, the four compression areas are subjected to reverse loading due to springback as these areas of the formed cylindrical cup begin to lose contact with the blank gripper. This reverse loading due to springback can lead to nucleation and crack growth in the four compression and wrinkling areas of the punched cup (see, for example, Figure 1E). The cracks can pass through the entire thickness of the steel strip and be visible on both sides of the punched cup (i.e., inside and outside), or only on one of the two sides of the cup. The lengths of all visible cracks on the inside and outside of the four compression edges of the deep punched cup (as shown in Figure 1D) were measured using a 10× magnifying glass equipped with a proportional grid. The sum of the lengths of all observable cracks on the inside and outside of the cup wall was averaged for three of the four punched cups and reported as the average total crack length (ATCL) in mm.

The microstructure of the steel strips was analyzed by electron backscatter diffraction (EBSD), a technique well known in the art, which in turn allows the quantification of the area or volume fractions of the various components. EBSD measurements were performed on cross sections parallel to the rolling direction (RD-ND plane) mounted in a conductive resin and mechanically polished to 1 μm. In order to obtain a completely deformation-free surface, a final polishing step was performed with colloidal silica (OPS).

The scanning electron microscope (SEM) used for EBSD measurements was a Zeiss Ultra 55 machine (FEG-SEM) equipped with a field emission gun and an EDAX PEGASUS XM4HIKARI EBSD system. EBSD scans were collected on the RD-ND plane of the plate at one-quarter thickness. The sample was placed in the SEM at a 70° angle. The accelerating voltage was 15 kV and the high current option was turned on. A 120 μm aperture was used and the usual working distance during scanning was 17 mm. To compensate for the high tilt angle of the sample, dynamic focus correction was used during scanning.

EBSD scans were captured using TexSEM Laboratories (TSL) software: "Orientation Imaging Microscopy (OIM) Data Collection version 7.2". Typically, the following data collection settings were used: a 5×5 pixel binned Hikari camera combined with background subtraction (standard mode). In all cases, the scan area was located at 1/4 of the sample thickness, and care was taken to avoid including non-metallic inclusions in the scan area as much as possible. The EBSD scan size was 100×100 μm in all cases, with a step size of 0.1 μm and a scan rate of approximately 100 frames/second. Fe(α) and Fe(γ) were used to index the Kikuchi patterns. The Hough settings used during data collection were: binning pattern size of approximately 96; theta set size of 1; rho score of approximately 90; maximum peak count of 10; minimum peak count of 5; Hough type set to typical; Hough resolution set to low; butterfly convolution mask of 9 × 9; peak symmetry of 0.5; minimum peak amplitude of 10; and maximum peak distance of 20.

EBSD scans were evaluated using TSL OIM Analysis Software Version 8.0x64 [12-14-16]. Typically, the data set was rotated 90° in the RD axis to obtain a scan in the proper orientation relative to the measured orientation. Standard grain expansion cleaning was performed (grain tolerance angle (GTA) of 5°, minimum grain size of 5 pixels, and the criterion used was that the grains must contain multiple rows for a single expansion iteration to clean). Next, pseudo-symmetric cleaning was applied (GTA 5, axis angle 30°@111).

The amount of martensite is determined using the EBSD image quality (IQ) map. Regions with low IQ are identified as MS regions. For given experimental conditions, the low IQ threshold is typically the peak-maximum position in the IQ histogram of ≈0.4. However, the low IQ threshold is manually checked for each scan to prevent grain boundaries from granular bainite or upper bainite regions from being included in the martensite area fraction.

To calculate the EBSD kernel average misorientation (KAM) map, the fifth nearest neighbor is used with a maximum misorientation of 5° (all points in the kernel are used for the KAM calculation). The kernel average misorientation is considered characteristic of the ferrite type because the kernel average misorientation is a measure of the internal dislocation density. Regions with relatively low internal dislocation density will mainly correspond to regions with KAM values between 0 and 1°. Regions with relatively high internal dislocation density will mainly correspond to regions with KAM values between 1-5°.

The sum of the total grain boundary lengths (in ^mm ) of grain boundaries with misorientation angles of 5° to 15° (ΣGB _5-15 ) and 15° to 65° (ΣGB _15-65 ) has also been measured (ΣGB) from EBSD scans over a 100×100 μm area (0.01 mm 2 ) in the RD-ND plane at ¼ thickness. The value of ΣGB is expressed in mm ⁻¹ and is a measure of the density of high-angle grain boundaries. High-angle grain boundaries are effective in arresting crack propagation. Therefore, increased ΣGB values will benefit from increased fracture toughness and reduced crack sensitivity.

In one aspect of the invention it relates to a method of manufacturing a steel strip as described and claimed herein, the method comprising the following steps in this order:

- casting a slab followed by a step of reheating the solidified slab to a temperature of 1050° C. to 1260° C., preferably for 30 minutes or more, more preferably for 60 minutes or more, and hot rolling said slab, or casting a slab or strip followed by a step of hot rolling said slab or strip;

- hot rolling the steel slab or strip and completing the hot rolling at a finishing temperature of 820°C to 940°C, preferably 850°C to 940°C, most preferably 850°C to 920°C and above the Ar3 temperature of the steel. The finishing temperature (FRT) is above the Ar3 temperature of the steel, where Ar3 is the temperature at which austenite begins to transform into ferrite during cooling. As known in the art, the Ar3 temperature can be calculated according to the following equation:

Ar3＝910℃-203×[C] ^1/2 +44.7×[Si]-30×[Mn]+31.5×[Mo];

- accelerated cooling of the hot rolled steel strip to a temperature on the run-out roller table of 560° C. to 620° C. at a run-out roller table cooling rate of 20 to 250° C./s, preferably 40 to 200° C./s;

- coiling the hot rolled and cooled strip at a temperature of 550°C to 600°C, preferably 550°C to 595°C, and more preferably 550°C to 590°C;

- further cooling the coiled hot rolled steel strip to ambient temperature; and

- pickling the hot rolled steel strip,

- Optionally providing the hot rolled steel strip with a metallic coating, preferably selected from: a Zn layer, a Zn-based alloy layer, an Al-based alloy layer, in order to provide improved corrosion resistance in use. The metallic coating is preferably applied by hot coating or hot dipping.

The manufacturing method described and claimed herein produces a desired microstructure that provides a targeted balance of improved formability, mechanical properties, and fracture characteristics. The invention is also embodied in a steel strip produced by the method described and claimed herein having the microstructure and the improved balance of formability, mechanical properties, and fracture characteristics.

The present invention is not limited by the casting method. Steel can be cast in a direct strip mill as a conventional thick slab with a casting thickness of 150mm to 350mm, and typically 225mm to 250mm, and a thin slab with a casting thickness of 50mm to 150mm. Schematic examples of processes involving conventional hot strip mills and thin slab casting/ingot mills are shown in Figures 2A and 2B, respectively. For conventional thick slab casting, the reheating of the slab is necessary to reheat the slab from ambient temperature (typically, in the slab yard, the thick cast slab has been cooled from the casting temperature to ambient temperature) and to homogenize the composition of the slab, so the reheating temperature should be higher than about 1050°C to dissolve any precipitates when there are microalloying elements, and to make the slab reach a temperature that allows the final hot rolling in the finishing mill to still be carried out under FRT>Ar3. This generally requires a (slab) reheating temperature of 1050°C to about 1260°C. For thin slab casting, the cast slab is subjected to a homogenization treatment in a homogenization furnace immediately after casting the thin slab, wherein the homogenization temperature should be above about 1050° C., and is typically about 1100 to 1160° C. This will also prevent the formation of any precipitates when microalloying elements are present (if any), and also bring the thin slab to a temperature such that the final hot rolling in the finishing mill can still be carried out at FRT>Ar3. According to the present invention, the reheating or homogenization time for the thin slab casting route is preferably 30 minutes or more.

Hot rolling of steel must be carried out in the austenite phase to control the final microstructure. In industrial-scale rolling, FRT should be kept above Ar3 temperature. In a preferred embodiment, FRT is higher than (Ar3+30°C), for example, usually higher than 850°C, to avoid local hot rolling below Ar3 at the cooler edge or tail of the strip. In the austenite region, FRT should not be too high, because lower FRT will promote more austenite deformation, and thus contribute to improved grain refinement and improved ΣGB. FRT higher than 950°C will lead to increased edge crack sensitivity. In addition, not too high FRT will also promote textures that are beneficial to toughness in the final microstructure (e.g., {332}<113>) and suppress those harmful textures (e.g., {001}<110> rotated cube). Therefore, FRT should not exceed about 940°C, preferably, FRT does not exceed about 920°C, and more preferably, FRT does not exceed about 910°C.

After hot rolling, the steel strip is accelerated cooled on the run-out table (ROT) to a temperature of 560°C to 620°C. In order to promote grain refinement and improved ΣGB of the final microstructure, an accelerated cooling rate is required to suppress the recovery and loss of internal stored energy in austenite. The cooling rate should be high enough to avoid the phase transformation of austenite to ferrite at an elevated temperature, and preferably promotes the phase transformation of austenite to ferrite at a relatively low temperature of about 560 to 630°C on the run-out table. The increased cooling rate will promote grain refinement, improved ΣGB, and therefore promote improved fracture toughness and reduced crack sensitivity. The increased cooling rate will also suppress texture randomization and therefore suppress the strength loss of those textures (e.g., {332}<113>) developed by the deformed austenite that promotes toughness. From a microstructural point of view, there is no critical run-out table cooling rate (ROT-CR) as long as the cooling rate mentioned herein exceeds the penetration thickness of the steel strip. However, an unnecessarily high ROT-CR can affect the flatness of the strip after cooling and cause control problems in stopping at the correct cooling stop temperature, so a suitable maximum ROT-CR is about 250°C/sec, preferably about 200°C/sec, and more preferably about 150°C/sec. A practical ROT-CR range is about 20 to 100°C/sec, more preferably about 40 to 100°C/sec, as this can be achieved by air cooling, laminar cooling, or water jet cooling, depending on the thickness of the steel strip. For practical reasons, the run-out roller cooling rate (ROT-CR) is defined as the average cooling rate of the steel strip surface.

Next, the hot rolled steel strip is coiled at a temperature of 550° C. to 600° C., preferably 550° C. to 600° C., more preferably about 560° C. to 600° C. In one embodiment, the hot rolled strip is coiled at a temperature not exceeding 595° C., more preferably not exceeding 590° C. The coiling temperature of the steel strip is a key process parameter to achieve the desired microstructure of the steel strip that provides a balance of improved mechanical properties as described herein.

When the coiling temperature is too low, the kinetics of precipitation are insufficient, so a low strength level will be achieved. When the coiling temperature is too high, the grain refinement is insufficient, resulting in reduced fracture toughness and increased edge crack sensitivity. Too high a coiling temperature will not promote any martensite as a second phase component in the final microstructure. Too high a 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 occur, as well as some further phase transformations. Undesirably, the precipitates once formed may coarsen during coil cooling. The alloy composition combined with the claimed coiling temperature suppresses this phenomenon. This coiling temperature will help promote small grains of ferrite formed during coiling or coil cooling, suppress the coarsening of precipitates that strengthen the ferrite matrix, and also promote the formation of a small amount of martensite.

In the patent document EP1616970-A1, it is disclosed that a holding temperature below 680°C results in insufficient driving force for ferrite transformation, which subsequently results in too low a fraction of ferrite-containing precipitates. For the present invention, the phase transformation from austenite to ferrite is carried out at a temperature below 680°C in order to obtain improved grain refinement and improved ΣGB for improved fracture toughness and reduced crack sensitivity while still having sufficient precipitation kinetics. In addition, the patent document EP1616970-A1 discloses that after holding the steel strip at a temperature of 680°C or higher for more than 1 second, it is necessary to apply secondary cooling to a coiling temperature of 550°C or lower, preferably 450°C or lower, more preferably 350°C or lower at an average cooling rate of 30°C/second or higher, preferably 50°C/second or higher, and to coil to form a secondary phase of bainite and/or martensite and suppress the formation of other phases to 5 volume % or less. For the present invention, the coiling temperature is relatively high, with a value of 550°C to 600°C, preferably 550°C to 595°C, more preferably 550°C to 590°C, in order to allow the phase transformation from austenite to ferrite to continue at a relatively low temperature to promote fine-grained ferrite, which is precipitation-strengthened with carbide precipitates containing Ti and Mo and optionally Nb and/or V. Grain refinement and increased ΣGB provide improved fracture toughness and reduced (edge) crack sensitivity. Coiling below 550°C will result in insufficient ferrite formation and precipitation losses. In addition, it may result in too high a martensite fraction.

Although patent document EP1338665-A1 discloses coiling temperatures of 550° C. to 700° C., steels having a tensile strength of at least 950 MPa and a hole expansion capacity of at least 40% are produced at coiling temperatures exceeding 600° C. The final rolling temperature of all these steels is 880 to 930° C. According to the invention, it has been found that excessively high coiling temperatures lead in particular to increased crack sensitivity, in particular edge crack sensitivity, and also to reduced fracture toughness.

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

In one embodiment, the hot rolled steel strip has a thickness of about 1.5 mm to 8 mm, more preferably about 1.8 mm to 6 mm, and most preferably about 1.8 mm to 4 mm.

The hot rolled steel strip product may be a bare or uncoated product, or it may be provided with a thin metallic coating on one or both major surfaces thereof, typically up to about 100 g/ ^m2 per side of the steel strip, preferably up to about 50 g/ ^m2 per side. The metallic coating is preferably selected from aluminium alloy coatings (e.g. Al-Si alloys or Al-Zn alloys), zinc coatings and zinc alloy coatings (e.g. Zn-Al alloys, Zn-Mg alloys, Zn-Fe alloys, Zn-Al-Mg alloys or Zn-Mg-Al alloys).

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

Other zinc coatings may also be applied. One example includes a zinc alloy coating according to patent document WO2008/102009-A1, which is incorporated herein by reference, in particular a zinc alloy coating consisting of 0.3 to 4.0 wt. % Mg and 0.05 to 6.0 wt. % Al, preferably 0.1 to 5.0 wt. % Al, and optionally up to 0.2 wt. % of one or more additional elements and unavoidable impurities (the remainder is zinc). A preferred Zn bath comprising Mg and Al as main alloying elements has the following composition: 0.5 to 3.8 wt. % Al, 0.5 to 3.0 wt. % Mg, optionally up to 0.2 wt. % of one or more additional elements; the balance is zinc and unavoidable impurities. The additional elements, which are usually added in small amounts of less than 0.2 wt. %, may be selected from 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.%, more preferably at most 0.1 wt.%. For typical applications, these small amounts of additional elements do not modify the properties of the coating or bath to any significant extent. Preferably, when one or more additional elements are present in the coating, each element is present in an amount of at most 0.02 wt.%, preferably each element is present in an amount of at most 0.01 wt.%. The additional elements are usually only added to prevent the formation of dross in a bath with molten zinc alloy for hot dip galvanizing, or to form spangles in the coating layer.

In another embodiment, metallic coating comprises (commercially pure) aluminum layer or aluminum alloy layer.The typical metal bath for hot-dip coating of this type of aluminum layer comprises aluminum alloy alloyed with silicon, for example with about 8 to 11 % by weight of silicon and 4 % by weight of iron at the most, optionally 0.2 % by weight of one or more additional elements (such as calcium), inevitable impurity alloyed aluminum alloy (surplus is aluminum) at the most.There is silicon to prevent the formation of the thick iron-metallic intermetallic compound layer that reduces adhesion and formability.Iron preferably exists with 1 to 4 % by weight, more preferably at least 2 % by weight.

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

In one aspect of the present invention, it relates to an automotive component, in particular an automotive chassis component, which comprises a hot-rolled high-strength steel strip according to the present invention or is made from a hot-rolled high-strength steel strip according to the present invention, and in particular benefits from improved strength, formability and a balance of improved fracture toughness and reduced (edge) crack sensitivity. The steel strip can be formed into automotive components in cold forming operations, warm forming and hot forming operations known in the art. Automotive components include suspension arms, reinforcement members, body-in-white frame members, side members, seat frames, seat rails, bumper beams, battery boxes for electric vehicles, all of which have complex shapes. By using the hot-rolled high-strength steel strip, these components can be manufactured with high quality, cost-effectiveness and high yield. The high-strength hot-rolled steel product according to the present invention can also be used in engineering applications.

BRIEF DESCRIPTION OF THE DRAWINGS

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

1A-1E illustrate several features of a method for determining the average total crack length (ATCL) of a steel strip product using a common cylindrical deep draw test for a punched cup.

A schematic diagram of a hot rolling mill for processing thick cast steel slabs is shown in FIG. 2A , and a thin slab casting facility with an ingot mill is shown in FIG. 2B .

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

Example

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

All ingots were reheated at 1240°C for 1 hour and rough rolled to a thickness of 35 mm. Subsequently, the strip was reheated again to 1220°C for 40 minutes and hot rolled to its final thickness of about 3.2 to 3.6 mm in 5 rolling passes at a FRT above 870°C and below 910°C (which is in the austenite phase region of all these steels). After the final rolling pass, the hot rolled steel was transferred to a run-out table, where the start temperature (T _START ) of the run-out table was 840°C to 880°C, and actively cooled from the austenite phase region to the final temperature (T END ) in the ferrite phase region at the run-out table at 565°C to 615°C with a mixture of water and air at a cooling rate of 30°C/sec to 70°C/ _sec . Next, the steel was transferred to a furnace to repeat the slow coil cooling. This was carried out at furnace temperatures (CT—coiling temperature) of 540, 580°C and 610°C (see Table 2).

Before testing mechanical properties, hole expansion capability and average total crack length according to the methods described herein, the hot rolled plates were sandblasted to remove the oxide layer. The microstructure of the hot rolled strip was determined with EBSG according to the methods described herein. The results of these tests as a function of the processing parameters employed are listed in Table 2.

Table 1. Chemical composition of 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, which has a composition outside the claimed range and is processed according to the invention (CT 580°C), has too low an Rm. Increasing the coiling temperature (CT 610°C) provides an increased Rm, but the edge crack sensitivity of Steel 16 is significantly reduced. This reduction is probably due to a microstructure with a too low MOD index and an insufficient degree of grain refinement as reflected by ΣGB. The too low coiling temperature of Steel 18 (CT 540°C) results in an even lower Rm compared to Steel 17.

Steels 1-3 (alloy A) have all compositions according to the invention, and of these 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 (HEC 42%) and very good edge cracking resistance (ATCL 28.9 mm). This is due to composition, processing and microstructure. The microstructure is characterized by very high grain refinement (ΣGB 18541/mm) and a high MOD index (1.20). For Steel 1, increasing the CT to 610°C still results in a high Rm, but the edge crack sensitivity is reduced to an unacceptable level (ATCL 49.3 mm). This is a result of microstructural changes, especially as reflected 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 of Steel 3 to 540°C resulted in a significant decrease in Rp and Rm.

A similar trend to that of Steels 1-3 (Alloy A) can be seen in Steels 4-6 (Alloy B).

Steels 7-9 (alloy C) have all compositions according to the invention with purposeful additions of V. Steel 8 has been processed according to the invention at CT 580°C and provides a very good balance of high strength (Rm 1005 MPa), high elongation (A50 14.7%), good hole expansion and very good edge cracking resistance (ATCL of 34.8 mm). This is due to composition, processing and microstructure. The microstructure is characterized by very high grain refinement. The microstructure has about 2.6 volume % martensite and it is believed that this contributes to the good edge cracking resistance indicated by the ATCL of 34.8 mm. Increasing the CT of Steel 7 to 610°C slightly increases Rm and Rp, but the edge cracking sensitivity of Steel 7 is significantly reduced. This reduction is probably due to the microstructural changes reflected by the MOD index and KAM value and the insufficient grain refinement reflected by ΣGB (ΣGB 1134 1/mm). The too low coiling temperature (CT 540° C.) of Steel 9 resulted in a lower Rm and a significantly lower Rp as compared to Steel 8, as well as a too low hole expansion ratio.

Alloy D (steels 10-12) has an increased Al content compared to alloy B (steels 4-6). Comparison of steels 5 and 11 (both processed according to the present invention) shows that the addition of Al results in an increased strength (Rm). Steel 11 also shows a favorable microstructure as shown by the MOD index, KAM value and sufficient grain refinement as shown by ΣGB, and thus provides the desired balance of properties.

Compared to alloy D (steels 10-12), alloy E (steels 13-15) has a purposeful addition of V. Compared to steel 11, steel 14 provides a further reduction in edge crack sensitivity. Comparison of steel 14 with steel 15 shows that the lower coiling temperature (CT 540°C) leads in particular to a significantly reduced Rp. And comparison of steel 13 with steel 14 shows that increasing the CT of steel 13 to 610°C still results in high Rm and Rp, but the edge crack sensitivity is reduced to an unacceptable level (ATCL 56.3 mm). This is reflected in particular by the too low MOD index, the too high KAM 0-1 and the low KAM 1-5, and the low ΣGB.

The invention now being fully described, it will be apparent to those skilled in the art that numerous changes and modifications can be made without departing from the spirit or scope of the invention described herein.

Table 2. Process parameters employed and resulting formability characteristics, mechanical properties, microstructure and textural components of the microstructure

Claims (15)

1. A hot rolled high strength steel strip having a composition consisting of, in weight %,: 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; Up to 0.02 wt.% P, Up to 0.01 wt.% S, Up to 0.005 wt.% B, One or more elements selected from the following: Up to 1.5 wt.% Cu, Up to 0.70 wt.% Cr, Up to 0.50 wt.% Ni, Up to 0.30 wt.% V, Up to 0.10 wt.% Nb, and the balance Fe and unavoidable impurities; and wherein the steel strip has a microstructure of at least 95 volume % ferrite and at most 5 volume % martensite; having a misorientation distribution (MOD) index of at least _{0.65 at 1/4} ^thickness ; Up to 55% area fraction, with a kernel average misorientation (KAM) of 0-1; At least 45% area fraction, with a kernel average misorientation (KAM) of 1-5; and a total grain boundary length (ΣGB[5°-65°]) of at least 1400 1/mm; And wherein the steel strip has at least the following mechanical properties: Ultimate tensile strength (Rm) of at least 950 MPa; A total elongation (A50) of at least 10%; and A hole expansion ratio (λ) value of at least 40%. 2. The hot rolled high strength steel strip according to claim 1, wherein the steel strip has at least 98 volume % ferrite, preferably at least 99 volume % ferrite. 3. The hot rolled high strength steel strip according to claim 1 or 2, wherein the overall microstructure of the steel strip has 0.2 to 4 volume %, preferably 0.2 to 3 volume %, more preferably 0.2 to 2 volume % martensite. 4. The hot rolled high strength steel strip according to any one of claims 1 to 3, wherein the microstructure has a misorientation distribution (MOD) index of at least 0.80, more preferably a MOD index of at least 0.83. 5. The hot rolled high strength steel strip according to any one of claims 1 to 4, wherein the microstructure has a total grain boundary length (ΣGB[5°-65°]) of at least 1500 1/mm. 6. The hot rolled high strength steel strip according to any one of claims 1 to 5, wherein the composition has a carbon content of 0.45 to 2.2, preferably 0.55 to 2.1. , and among them Ti_sol is defined as 7. The hot rolled high strength steel strip according to any one of claims 1 to 6, wherein the steel strip has a Mo content of 0.10 to 0.80 wt. %, preferably 0.25 to 0.80 wt. %, more preferably 0.30 to 0.80 wt. %. 8. Hot rolled high strength steel strip according to any one of claims 1 to 7, wherein the steel strip has a Cr content of at most 0.10 wt.-%, preferably at most 0.050 wt.-%, more preferably at most 0.03 wt.-%. 9. The hot rolled high strength steel strip according to any one of claims 1 to 8, wherein the steel strip has a V content of 0.05 to 0.30 wt. %, preferably 0.05 to 0.30 wt. %, more preferably 0.08 to 0.20 wt. %. 10. The hot rolled high strength steel strip according to any one of claims 1 to 9, wherein the steel strip has a Mn content of 1.40 to 2.40 wt. %, preferably 1.40 to 2.20 wt. %. 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 45 mm, preferably less than 40 mm. 12. The hot rolled high strength steel strip according to any one of claims 1 to 11, wherein the steel strip has a metallic coating, and the metallic coating is preferably selected from: a Zn layer, a Zn-based alloy layer, and an Al-based alloy layer. 13. A method for manufacturing a hot rolled high strength steel strip according to any one of claims 1 to 12, the method comprising the following steps: - a step of casting a slab followed by reheating the solidified slab to a temperature of 1050° C. to 1260° C. and hot rolling said slab, or casting a slab or strip followed by direct hot rolling of said slab or strip; - hot rolling the steel slab or strip, and completing the hot rolling at a final rolling temperature of 820°C to 940°C, preferably 850°C to 940°C and higher than the Ar3 temperature of the steel; - accelerated cooling of the hot rolled steel strip to a temperature on the run-out roller table of 560°C to 620°C at a run-out cooling rate of 20 to 250°C/s, preferably 40 to 200°C/s; - coiling the hot rolled and cooled strip at a temperature of 550°C to 600°C, preferably 550°C to 595°C, and more preferably 550°C to 590°C; - further cooling the coiled hot rolled steel strip to ambient temperature; - pickling the hot rolled steel strip, 14. The method according to claim 13, wherein the pickled hot rolled steel strip is provided with a metallic coating, preferably selected from: a Zn layer, a Zn-based alloy layer, an Al-based alloy layer, preferably applied by hot coating or hot dipping. 15. Automotive component comprising a hot rolled high strength steel strip according to any one of claims 1 to 12 or obtained by the method according to any one of claims 13 to 14.