High strength steel sheet with excellent hole expandability and method of producing the same

Abstract

The invention relates to a cold-rolled continuously annealed steel strip or sheet having a tensile strength of 980 MPa or more, a total elongation of 8% or more and a hole expansion ratio of 50% or more, and to a method of producing the same.

Description

[0001] HIGH STRENGTH STEEL SHEET WITH EXCELLENT HOLE EXPANDABILITY AND METHOD OF PRODUCING THE SAME FIELD OF THE INVENTION The invention relates to a high strength cold-rolledsteel strip or sheet with excellent hole expandability and to a method of producing the same. BACKGROUND TO THE INVENTION The demand for advanced high strength steels (AHSS) with higher strengths is increasing in the automotive industry. Automotive parts are often cold formed and the manufacturing process includes shear cutting and flanging. Problems often may occur when the sheared edges are stretched during pressing as they are susceptible to crack formation. The flangeability as characterized by hole expansion capacity or ratio (HEC) tends to decrease with an increase in the strength of a steel. There are two types of formability, viz. drawability and stretch flangeability. The drawability is governed by global ductility, characterized by elongation in tensile tests, the stretch flangeability is governed by local ductility, characterized by the hole expansion ratio in hole expansion tests or by the bending angle in bending tests, respectively. Global ductility and local ductility are incompatible. The total elongation depends mainly on the strain localization in the matrix phase while the local ductility depends on the micro-scale uniformity: the more uniform the microstructure is, the better the stretch flangeability is, although stretch flangeability is also known to decrease with increasing strength. Bendability can also be construed to be similar to stretch flangeability, as a fracture due to a large local deformation, and can be adjusted by the scatter of hardness (i.e. the distribution of hard and soft phases in the microstructure). As steels became increasingly stronger, they simultaneously became increasingly difficult to form into automotive parts. While there have been major improvements recently in the trade-off between ductility and strength, the stretch flangeability of AHSS remains a critical issue. Dual Phase (DP) steels and low alloy transformation-induced plasticity (TRIP) steels are known to show high elongation. However, since these steels are characterized by a large difference in hardness between the main phases, which consists of soft polygonal ferrite and the hard second phase of martensite, these types of steel are inferior in stretch flangeability (local ductility) because of voids that form at the interphase between these phases during deformation processes such as punching. In contrast, Complex Phase (CP) steels yield a lower global ductility as a result of the reduced amount of the soft ferrite and a higher local ductility as a result of the reduced hardness difference among the main phases and together with a more uniform distribution of microstructure. In practice, the application of AHSS steels (DP, CP and TRIP) to car components is still limited by their formability. Therefore, improving formability and manufacturability is an important issue for AHSS applications. To remedy this problem, bainite single phase and bainitic ferrite single phase steel sheets have been developed in order to reduce the hardness difference between the main and second phases, but it cannot be said that the balance of strength and hole expansion ability is adequate, particularly in materials with tensile strength of 980 MPa class and higher. 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 0.015% Mo may include a steel strip composition having no Mo. It is an objective of the present invention to provide a cold-rolled high strength steel. It is also an object of the invention to provide a cold-rolled high strength steel with high strength and an improved formability and HEC value in particular. It is also an object of the invention to provide a cold-rolled high strength steel excellent in hole expandability, aiming at the application for automobiles, such as passenger cars and trucks, etc., for industrial machines, or the like. It is also an object of the invention to provide a process for producing these cold-rolled high strength steels. These objects and further advantages are met or exceeded by the present invention according to claim 1 and a method for producing a cold-rolled continuously annealed steel strip or sheet according to claim 9, and with preferred embodiments set out in the dependent claims and the description. In order to achieve these objects, the present invention proposes, in a first aspect, a cold-rolled continuously annealed steel strip or sheet having a tensile strength (Rm) of 980 MPa or more, a total elongation (A) of 8% or more as measured in accordance with JIS 5 in the JIS- Z2241 specification and a hole expansion ratio (HEC) of 50% or more, preferably of 60% or more, as measured in accordance with the Japan Iron and Steel Federation Standards JFS T 1001, and having a composition comprising, in wt.%: C: 0.150-0.280%; Mn: 1.50-3.00%; Si: 0.02-1.60%; Al: 0.01-0.70%; Al + Si: 0.50-2.00%; P: up to 0.050%; S: up to 0.030%; N: up to 0.010%; the steel composition optionally also comprising one or more elements selected from: B: up to 0.0050%; Ti: up to 0.10%; Nb: up to 0.10%; V: up to 0.10%; Mo: up to 0.20%; Cr: up to 0.60%; Cu: up to 0.20%; Ni: up to 0.20%; Ca: up to 0.0030%; rare earth elements (REM): up to 0.003%; and the balance being Fe and inevitable impurities resulting from the ironmaking and steelmaking process; and having a microstructure consisting of: bainitic ferrite (BF) and / or partitioned martensite (PM), ferrite phases (HF+LF) and second phase comprising fresh martensite (FM) and carbides, and retained austenite (RA), in the following amounts, in vol.%: BF+PM at least 70%; HF+LF up to 20%; HF up to 5%; carbides+FM: up to 3%; and RA 3 to 10%. In the context of this invention the following abbreviation for microstructural components will be used: BF: bainitic ferrite is formed during austempering in the temperature range between Ms and Bs. PM: partitioned martensite is formed during austempering first at a temperature below the Ms point and then heated to a higher temperature for overaging, viz. the temperature range between Ms and Bs. HF: ferrite formed at intercritical temperatures (high temperature ferrite). LF: ferrite formed at cooling below Ac1 after annealing (low temperature ferrite). F: Σ(HF + LF). RA: retained austenite. FM: fresh martensite refers to the martensite formed during cooling after the austempering treatment. In accordance with the invention it has been found that the cold-rolled continuously annealed steel strip having these narrow alloy compositional ranges in combination with the microstructure provides for an improved balance of high strength, total elongation and excellent hole expansion ratio by forming a uniform microstructure with a lower non-relaxed strain gradient as will be explained herein below. The required microstructure for this cold-rolled continuously annealed steel strip is achieved by the narrow compositional ranges and by careful control of the manufacturing process, in particular by obtaining a microstructure of containing bainitic ferrite and partitioned martensite as matrix by performing a quenching treatment to a lower temperature between the Mf+100°C and 50°C (Mf is the martensite transformation finish temperature) before austempering in the annealing process. Furthermore, the amount of the second phases including fresh martensite (FM), carbide and retained austenite (RA) must be controlled to below a certain level. When in the context of this invention use is made of the term “strip”, then this term is deemed to encompass a sheet or any sheet-like material produced from such a strip. The cold-rolled continuously annealed steel strip has an ultimate tensile strength (Rm) of at least 980 MPa. In an embodiment of the steel strip it has an ultimate tensile strength of at least 1040 MPa, and in the best examples preferably of at least 1080 MPa. The cold-rolled continuously annealed steel strip has a hole expansion ratio (HEC) of 50% or more as measured in accordance with the Japan Iron and Steel Federation Standards JFS T 1001. In an embodiment of the steel strip it has a HEC of 60% or more, and in the best examples of 65% or more. In an embodiment the cold-rolled continuously annealed steel strip has a favourable high yield ratio, viz. the ratio Rp / Rm of more than 0.69. In an embodiment the yield ratio is at least 0.80, and in the best examples of at least 0.85. The cold-rolled continuously annealed steel strip has a total elongation (A) of at least 8%, and in the best examples preferably of at least 10%, as measured in accordance with JIS 5 in the JIS-Z2241 specification. An important feature of steel strip according to this invention is its microstructure consisting of, in volume percentages or vol.%, bainitic ferrite (BF) and / or partitioned (PM), ferrite phases (HF+LF) and second phase comprising fresh martensite (FM) and carbides, and retained austenite (RA), in the following amounts, in vol.%: BF+PM at least 70%, preferably >86%; HF+LF up to 20%; HF up to 5%; carbides+FM: up to 3%; and RA 3 to 10%. It is a key feature of the invention that the total volume fraction of bainitic ferrite and partitioned martensite (BF+PM) is 70 vol.% or more, and preferably >86 vol.%, and more preferably >90 vol.%. The BF and PM in the steels according to this invention are different from the bainite and tempered martensite in the conventional steels. In the steels according to this invention, the formation of carbides is suppressed or retarded, due to the presence of 0.50 wt.% or more of the sum of Si and Al in the steel composition during further heat treatment. As a result, the carbon partitions from bainite or martensite to austenite which leads to carbon enriched retained austenite with higher stability. Therefore, in a limited time frame the BF+PM can be considered as carbide-free. The matrix of BF or a mixture of PM+BF ensures a good balance in the strength and ductility of the steels. It is difficult to achieve both a tensile strength of at least 980 MPa or more and a hole expansion ratio of 50% or more when the total amount of BF+PM is less than 70 vol.%. Therefore, the fraction of BF+PM is 70 vol.% or more, and preferably of >86 vol.%, and more preferably >90 vol.%. In an embodiment the total amount of BF and PM is not more than 96 vol.%. The amount of ferrite (F) is at most 20 vol.%. The presence of F can increase the C content of the untransformed austenite and increase the stability of the retained austenite, which increases the total elongation but decreases strength and formability. F can form during austenization process in the intercritical temperature range, referred as high temperature ferrite (HF) and during cooling at high temperatures after austenization, referred as low temperature ferrite (LF). In the present invention, these two kinds of ferrite are distinguished. The LF contains a relatively high carbon content and has a smaller grain size than HF does. Therefore, LF has a higher hardness, which is beneficial to obtain high HEC property. Therefore, a controlled amount of LF is allowable but HF should be avoided. It is difficult to achieve both a TS of 980 MPa or more and a hole expansion ratio of 50% or more when the content of ferrite is greater than 20 vol.%. Therefore, the content of ferrite is 0 to 20 vol.%, in which the total HF is less than 5.0 vol.%, and preferably less than 1.0 vol.%, and more preferably 0 vol.%. The total amount of fresh martensite (FM) and carbides is kept in a range of 0 to 5 vol.%. Carbides may be contained in the form of pearlite and / or separately existed cementite or alloying carbides during cooling to austempering treatment when the cooling rate is not high enough. Carbides may also form during austempering treatment if the overaging temperature is too high and / or the overaging time is too long. The formation of carbides reduces the amount of the retained austenite and the C content in the retained austenite, which deteriorate both the strength and ductility. Coarser carbides can reduce the hole expansion capacity (HEC). Fresh martensite may be produced during final cooling if the bainitic transformation is incomplete during austempering treatment. The hole expansion ratio is remarkably low when the amount of the FM is increased. In order to achieve good formability, it is necessary that the total amount of carbides and FM be 5 vol.% or less, and preferably of 3 vol.% or less. The amount of retained austenite (RA) should be 10 vol.% or less. Retained austenite is useful for increasing elongation by producing an TRIP (strain-induced transformation) effect. The volume fraction of RA needs to be 3 vol.% or more to achieve such an effect. However, when the volume fraction is greater than 10 vol.%, the RA is not stable enough as the C of the RA is too low. It may transform to fresh martensite at early stage of deformation, and thereby deteriorates the stretch flangeability. Therefore, the volume fraction of retained austenite is preferably in a range of 3 to 10 vol.%, and preferably 4 to 8 vol.%. In an embodiment it is recommended that the C concentration in the RA be 0.9 wt.% or more. This C content greatly affects the characteristics of TRIP, and controlling it to 0.9 wt.% or more is particularly effective for improving elongation but not adversely affecting the stretch flangeability. In an embodiment the C concentration is at least 1 wt.%, more preferably at least 1.2 wt.%. The higher the C content of the RA is, the more preferable it is. However, in practical operation, the upper limit that can be adjusted is considered to be about 1.6 wt.%. For the purpose of this invention samples for microstructure observations, tensile tests and hole expansion tests were machined from the heat-treated strips. Dilatometry was done on the cold rolled samples of 10 mm x 5 mm x 1 mm dimensions (length along the rolling direction). Dilatation tests were conducted on a Bahr dilatometer type DIL 805. All measurements were carried out in accordance with SEP 1680. The critical phase transformation points Ac1, Ac3, Ms and Mf were determined from the quenched dilatometry curves. The phase fractions during annealing for different process parameters were determined from dilatation curves simulating the annealing cycles. The microstructure was determined by optical microscopy (OM), scanning electron microscopy (SEM) and electron backscatter diffraction (EBSD) using a commercially available image-processing program. The cross section in the thickness direction parallel to the rolling direction of a steel sheet is polished then etched with a 3% nital solution. The microstructure is observed at a position located at ¼t (position at ¼ of the thickness from the strip surface) by using an optical microscope at a magnification of 1000 times. Under optical microscopy, F is visibly distinguished as tinted grey areas, in which HF is shown as polygonal blocks and LF is shown as acicular shapes along the original austenite grain boundaries. BF and PM as a mixture is observed as light brown areas in which lath ferrite can be vaguely seen. FM is observed as straw-coloured areas, which have a smaller size. However, BF and PM could not be distinguished as the contrast between them is too small. Carbides are shown as black dots decorated at original austenite grain boundaries. The volume fraction of each microstructure is assumed to be equal to the area fraction of each, which is determined using a commercially available image-processing program. The 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. The EBSD scans were captured using the TexSEM Laboratories (TSL) software OIM (Orientation Imaging Microscopy) data collection. The EBSD scans were evaluated with TSL OIM Analysis software. The EBSD scan area was in all cases 100 x 100 µm, with a step size of 0.1 µm, and a scan rate of approximately 80 frames per second. The retained austenite (RA) was determined by X-ray diffraction (XRD) according to DIN EN 13925 on a D8 Discover GADDS (Bruker AXS) with Co-Kα radiation. The XRD measurements were conducted on the subsurface at ¼ thickness of the steel sheet. The steel sheet is mechanically and chemically polished and is then analysed by measuring the integral intensity of each of the (200) plane, (220) plane, and (311) plane of fcc iron and that of the (200) plane, (211) plane, and (220) plane of bcc iron with an X-ray diffractometer using Co-Kα radiation. The amount of RA and the lattice parameter in the RA were determined using Rietveld analysis. The C-content in wt.% in the RA is calculated using the formula (from D. Dyson and B. Holmes, “Effect of alloying additions on the lattice parameter of austenite”, in J. Iron Steel Inst. 208 (1970), pp.469-474): C =(a[Å]-3.572-0.0012 [Mn] + 0.00157 [Si] - 0.0056 [Al]) / 0.033 (1) where “a” is the lattice parameter of the RA in angstrom, and [Mn], [Si] and [Al] are the concentration of these elements in wt.% in the steel. For the tensile tests – JIS5 test pieces (gauge length = 50 mm; width = 25 mm) were machined from the annealed strips such that the tensile direction was parallel to the rolling direction. Room temperature tensile tests were performed in a Schenk TREBEL testing machine following NEN-EN10002-1:2001 standard to determine tensile properties (yield strength Rp in MPa, ultimate tensile strength Rm in MPa, total elongation A in %, and uniform elongation Ag in %). For each condition, three tensile tests were performed, and the average values of mechanical properties are reported. In the hole expansion test (or stretch flangeability evaluation test) test pieces for testing hole expandability (size: 90×90 mm) were sampled from the obtained rolled strip. In accordance with The Japan Iron and Steel Federation Standard, “Method of Hole Expanding Test,” JFS T 1001, a 10 mm diameter punch hole was punched in the centre of the test piece and a 60° conical punch was pushed up and inserted into the hole and pushed up increasing the hole diameter. When a crack penetrated the strip thickness, the hole diameter d (mm) was measured. The hole expansion ratio λ (%) was calculated by, with d0= 10 mm: λ (%) = { (d−d0) / d0} × 100 (2) Another key feature of the invention is related to the composition (in wt.% or weight percent) of the cold-rolled continuously annealed steel strip. C: 0.150% to 0.280%. Carbon (C) is a necessary element for strength and ductility by forming the required microstructure. When the content of C is less than 0.150%, it is difficult to the secure 65 vol.% or more BF+PM in the microstructures to obtain the required strength and improved formability. On the other hand, when the C content is greater than 0.280%, the weldability is deteriorated. Therefore, the C content is 0.150% to 0.280%, and preferably in a range of 0.170% to 0.250%. Si: 0.02% to 1.60%, Al: 0.01% to 0.70%, and Si+Al: 0.50% to 2.00%. Silicon (Si) is an element which is effective in increasing the strength of the steel by solute strengthening and effective in preventing carbon from precipitating in the form of carbides and increasing the stability of residual austenite. Si has a function of improving the ductility, work hardenability and stretch flangeability through the retardation of austenite grain growth during annealing. Preferably a minimum of at least 0.40% is needed to achieve such effects, and preferably at least 0.75%. When the Si content is higher than 1.60%, surface quality of the steel sheet and the weldability of the steel are deteriorated, and is preferably not exceeding 1.50%. Accordingly, the Si content is in a range of 0.02% to 1.60%, preferably in a range of 0.40% to 1.50%, and more preferably in a range of 0.75% to 1.50%. The primary function of aluminium (Al) is to deoxidise the liquid steel before casting. Furthermore, Al has a similar function as Si to prevent the formation of carbides and to stabilize the retained austenite and therefore is effective in improving strength and ductility balance. In order to have such effects at least 0.01% is required. Too high levels of Al increase the ferrite to austenite transformation point temperature to levels that are incompatible with conventional installations and will therefore not provide a microstructure according to the invention. In addition, the risk of cracking during casting and hot rolling increases as the Al content is increased. Therefore, the Al content is 0.01% to 0.70%, preferably 0.020% to 0.50%. If the amount of aluminium is to be limited, for instance for, but not limited to, specific surface quality reasons, then the aluminium content is preferably between 0.020% and 0.090%, more preferably between 0.020% and 0.050%. When both Si and Al are added, the total amount of Si + Al is 0.50% to 2.00%. Mn: 1.50% to 3.00%. Manganese (Mn) is an element which is indispensable for obtaining hardenability of the steel to achieve the desired strength. Mn is also added to balance the elevated phase transformation point as a result of high amounts of Si and Al. In order to produce such effects, the Mn content needs to be 1.50% or more. However, when the Mn content is more than 3.00%, the ductility and formability of the steel are reduced due to the formation of macro-segregation and band microstructure. Therefore, the Mn content is in a range of 1.50% to 3.00%, and preferably in a range of 1.60% to 2.80%, and more preferably in a range of 1.70% to 2.60%. P: up to 0.050%. Phosphorous (P) is an impurity in the steel. However, as P is also effective for increasing the strength of steel and increases the stability of the RA. In order to produce such an effect, it is prefered that the P content be at least 0.002%. On the other hand, in the case where the P content is more than 0.050%, the steel is embrittled due to grain boundary segregation and, therefore, deteriorates stretch flangeability. Therefore, the P content is preferably in a range of 0.002% to 0.050%. If excellent weldability is required, it is preferable that the P content be 0.020% or less, more preferably 0.015% or less, and even more preferably 0.012% or less. S: up to 0.030%. Sulphur (S) is an impurity in the steel, present in the form of an inclusion such as MnS and deteriorates formability. Hence, the content thereof is preferably as low as possible. Therefore, the content of S is 0.030% or less, and preferable 0.010% or less, in view of manufacturing cost. N: up to 0.010%. Nitrogen (N) is inevitably present in the steelmaking process. N in solid solution can increase the strength but decreases the ductility. N deteriorates the toughness and formability of steel when an excess amount of 0.010% is present due to the formation of coarse nitride compounds such as TiN, AlN and BN or blowholes. Accordingly, the N content in steel is to be 0.010% or less, preferably 0.0060% or less. The practical lower limit of nitrogen content is around 0.0005% because decreasing nitrogen content in steel further significantly increases production costs. A suitable and practical minimum N content is between 0.0010% to 0.0050%. One or more of the following elements can be optionally added: B: up to 0.0050%; Ti: up to 0.10%; Nb: up to 0.10%; V: up to 0.10%; Mo: up to 0.20%; Cr: up to 0.60%; Cu: up to 0.20%; Ni: up to 0.20%. These elements are useful as strengthening elements for steel, and are also elements effective for stabilizing RA and ensuring a predetermined amount. Chrome (Cr) increase the hardenability of steel, suppresses the formation of carbides and increases the stability of retained austenite. A proper amount of Cr can improve the strength- ductility balance. Cr can be added to the steel as a optional element to replace Si or Al. However, if Cr is more than 0.60%, it will lead to an increase in strength but also a remarkable decrease in the balance of strength-flange formability. Therefore, the Cr content should be up to 0.60% if added. Preferably the Cr content does not exceed 0.50%. When Cr is purposively added, Mn may be reduced. It is preferred to maintain the sum of Mn+Cr higher than 1.50% but lower than 3.00%, and preferably lower than 2.80%. In an embodiment of the invention there is no purposive addition of Cr and thus present as an impurity element and be present only up to 0.025%. Cu: up to 0.20%, Ni: up to 0.20%, Mo: up to 0.20%. Each of Copper (Cu), Nickel (Ni), and Molybdenum (Mo) can be present as tramp elements originating from the scrap metals used in the steelmaking process. But these are also useful as strengthening elements for the steel. In order to effectively exert such an effect, Cu should be 0.050% or more, Ni should be 0.050% or more, and / or Mo should be 0.050% or more are recommended to be added. When the content of each of these elements is below these lower limits respectively, it is considered to be an impurity within the steel. When present as an impurity element, Cu is preferably up to about 0.020%, Ni is preferably up to about 0.025%, and Mo is preferably up to about 0.020%. Thus in an embodiment the steel has a purposive addition of Mo in a range of 0.050% to 0.20%, and preferably of 0.050% to 0.140%. In an embodiment the steel has a purposive addition of Ni in a range of 0.050% to 0.20%, and preferably of 0.050% to 0.160%. In an embodiment the steel has a purposive addition of Cu in a range of 0.050% to 0.20%, and preferably of about 0.060% to 0.150%. However, when each of these elements is added in amount exceeding the upper limit, Cu >0.20%, Ni >0.20% and Mo >0.20%, these can segregate at the surface during the hot rolling process, which may cause a surface oxidation problem and also deteriorate the ductility of the steel strip. Therefore, it is preferred that the sum of Cu+Ni+Mo is less than 0.50%. One or more of the following elements can be added to the steel: Ti: up to 0.10%, Nb: up to 0.10%, V: up to 0.10%. Each of Titanium (Ti), Niobium (Nb) and Vanadium (V) can form carbonitrides or carbides which can produce precipitation hardening and refine the microstructure of the steel strip, and are useful for increasing the strength. In order to effectively exhibit such an effect, it is recommended to add each of Ti: 0.0050% or more, Nb: 0.0050% or more and V: 0.010% or more, respectively. However, if the content of each of these elements is below the lower limits respectively, it is considered to be an impurity within the steel. However, if any of the elements is added in excess of the upper limit respectively, a large amount of coarse carbides will form during processing leading to a decrease in the achieved balance of the strength-flange formability. It is preferred that Ti does not exceed 0.050%, and preferably Nb does not exceed 0.050%. It is not recommended to add all these elements together in substantial amounts. It is preferred that the sum of (Ti+Nb+V) should be less than 0.10%. B: up to 0.0050%. Boron (B) is not a key element to the invention, but it is a useful element in suppressing formation and growth of ferrite from austenite grain boundaries and therefore increases the hardenability of the steel. When purposively added the B content should be at least 0.0003% to achieve such effect. When the B content exceeds 0.0050%, the formability of the steel is deteriorated. Therefore, preferably the B content is set to be in the range of 0.0003% to 0.0050%, and more preferably in a range of 0.0005% to 0.0020%, and most preferably in a range of 0.0005% to 0.0015% For B to be able to perform this role effectively, it is highly recommended that no free N is present so that the formation of BN is minimized. A suitable amount of Ti or Al should be added to scavenge nitrogen and the maximize the effect of B. Furthermore, Ca: up to 0.0030% (30 ppm) and / or REM: up to 0.0030% (30 ppm) may be present in the steel. Ca and REM are elements that control the form of sulphide in steel and are effective for improving workability. Here, examples of the rare earth element used in the present invention include Sc, Y, and lanthanide. In order to effectively exert the above effects, it is recommended to add 3 ppm or more, respectively. However, even if it is added in excess of 30 ppm, the above effect is saturated, which is economically useless. It is more preferably at most 25 ppm. The same applies to the amount of Sc, Y, and lanthanides to be added of required. The balance being Fe and inevitable impurities resulting from the ironmaking and steelmaking process In an embodiment the cold-rolled continuously annealed steel strip is provided with a hot- dip metallic coating layer on one or both of its main surfaces, preferably on both of its main surfaces, typically up to about 150 g / m2per side of the steel strip, and preferably up to about 80 g / m2per side. The metallic coating layer 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-Al alloy, Zn-Mg alloy, Zn-Fe alloy, Zn-Al-Mg alloy, or Zn-Mg-Al 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. Other zinc coating layers may also be applied. An example comprises a zinc alloy coating according to patent document WO2008 / 102009-A1 and incorporated herein by reference, in particular a zinc alloy coating layer consisting of about 0.3% to 4.0% Mg and about 0.05% to 6.0% Al, preferably about 0.1% to 5.0% Al, and optionally at most 0.2% 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% Al, 0.5% to 3.0% Mg, optionally at most 0.2% 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%, 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 about 0.2%, and more preferably at most about 0.1%. 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 about 0.02%, preferably each is present in an amount up to about 0.01%. 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% of silicon and at most about 4% of iron, optionally at most about 0.2% of one or more additional elements such as calcium, unavoidable impurities, the remainder being aluminium. Silicon is present in order to prevent the formation of a thick iron-metallic intermetallic layer which reduces adherence and formability. Iron is preferably present in amounts between about 1 and 4%, more preferably at least about 2%. In another aspect of the invention it relates to a method of manufacturing a cold-rolled continuously annealed steel strip according to claim 1. In this aspect of the invention the method comprises the steps of, in that order: - preparing a steel melt in a BOS-process and casting a steel slab or ingot having a composition in accordance with this invention and with preferred compositional ranges as herein described and claimed; - (re)heating the steel slab or ingot to a temperature in the range of 1100-1300°C; - hot-rolling the slab or ingot to a hot-rolled strip wherein the finishing rolling temperature is equal to or higher than the Ar3 transformation point; - cooling the hot-rolled strip at an average cooling rate of 20°C / s or more, preferably 30°C / s or more, to a temperature below 650°C and coiling the hot-rolled strip at temperatures in the range from 200oC to 550°C; - pickling or descaling the hot-rolled strip prior to cold rolling; - cold rolling at a total cold rolling reduction of between 30% and 80% to obtain a cold- rolled strip; - a thermal treatment by continuous annealing the cold-rolled steel strip by: * heating the steel strip to a soaking temperature T2 and soaking the steel strip for a time t2 to obtain a fully austenitic microstructure; * cooling the steel strip at a slow cooling rate V3 to a temperature T3 wherein T3 is between Ar3 and (T2-40)°C; * cooling the steel strip at a cooling rate V4 to a temperature T4 between Mf+100oC and 50oC, wherein Mf is the martensite transformation finish temperature; * heating the steel strip from T4 within a time t4 to a temperature T5 between Ms and Ms+200oC for a period t5 between 15 to 150s; * cooling to below 300°C at a cooling rate of at least 1°C / s followed by cooling to ambient temperature at a cooling rate V8. The critical phase transition temperatures are well known to a person skilled in the art and are dependent on the alloy composition and heating rate. In this invention, the critical temperatures Ac3, Ms and Mf are determined using dilatation test by applying a heating rate of 2°C / s and austenitized at 900°C for 65s and then quenched at a cooling rate of 100°C / s to room temperature. Such a heating profile is frequently used in an existing production line. Ac1 is determined as the temperature at which 2% austenite has been formed during heating at 2°C / s. Ac3 is determined as the temperature at which 98% austenite has been formed during heating at 2°C / s. Ms is determined as the temperature at which 5% martensite has been formed during cooling at 100°C / s from a fully austenization temperature of 900°C. Mf is determined as the temperature at which 90% martensite has been formed during cooling at 100°C / s. The error bar for determining these temperatures is ± 3oC. A steel slab is prepared to have the preferred component composition described above based on a melt from the BOS-process. The slab is preferably manufactured by a continuous casting process for the purpose of preventing macro-segregation. These processes are not particularly restricted and may be carried out according to conventional methods. Thin slab casting, strip casting or the like can also be applied which, for the purpose of this invention, will also be considered a slab. In the case of strip casting it is acceptable for the manufacturing method to skip at least a part of the hot-rolling process. The slab is subjected to hot-rolling and then cold-rolling to obtain a cold-rolled steel strip. Preferable manufacturing conditions of a cold rolled steel strip include: (re)heating the slab to a temperature in the range of 1100°C to 1300°C; hot-rolling the slab to a hot-rolled strip wherein the finishing rolling temperature is at temperature equal to or higher than the Ar3 transformation point and preferably between 800oC and 1050°C; cooling at an average cooling rate of 20°C / s or more, and preferably 30°C / s or more, to a temperature below 650°C and coiling the hot-rolled strip at a temperature in the range of from 200oC to 550°C. The coiled hot-rolled strip is descaled by pickling or the like prior to cold-rolling. The resulting steel sheet is preferably cold-rolled at a total cold rolling reduction of between 30% and 80%. If the rolling reduction of cold rolling is less than 30% then the recrystallization during annealing is incomplete so that a non-recrystallized ferrite phase is retained in a microstructure after the continuous annealing process, which may result in a decrease in formability. In addition, in the case where the rolling reduction of cold rolling is excessively high, since there is an increase in rolling loads during cold-rolling and rolling issues such as chattering and fracturing of strips may occur. Therefore, it is preferable that the rolling reduction of cold rolling is about 80% or less. The cold-rolled steel strip thus obtained is subjected to a thermal treatment in the form of continuous annealing. It is noted that if the strip is cut into sheets that the following heat treatment may be performed on the strip or on the cut sheets. The heat treatment includes an austenitizing process, a fast cooling process, an austempering or overaging process and a final cooling process. The heat treatment process is described with reference to the Fig.1. In an embodiment the cold rolled steel strip or sheet is heated through a segment heating at T1 to a pre-determined austenization temperature T2 and held for a time t2, and then cooled at a low cooling rate V3 to T3 and then at a higher cooling rate V4 to a temperature T4 in the range of the Mf+100oC to 50°C. Thereafter, the steel strip is heated up in a time t4 to a temperature T5 in the range of Ms to Ms+200°C for austempering for a time t5. In the meantime, a metallic coating process at a temperature T5 for a time t6 can be applied. Finally, the steel strip is cooled to room temperature by controlled cooling rates V7 and V8. In an embodiment the cold rolled steel strip or sheet is first heated to a temperature T1 in the range of 680oC to 750°C at a heating rate V1. The heating rate is not limiting, but to limit the length of this heating stage the heating rate is preferably at a rate of 10 to 25°C / s, and then to the temperature T2, at a slow heating rate V2 of 0.5 to 10°C / s, preferably at a heating rate V2 of 1 to 5°C / s. In this second heating stage, the phase transformation of ferrite to austenite and recrystallization of ferrite occur concurrently. If V2 is too low, the resulting austenite grains may become too large, which may deteriorate the formability. If V2 is too high, recrystallization of ferrite is slow and results in a banded-like structure with small austenite grains and large austenite grains, and therefore promotes the formation of ferrite during cooling. The steel strip or sheet is annealed at temperature T2 in the austenite single-phase region for a time t2 to obtain a fully austenitic microstructure. T2 should be in a range of Ac3-30°C to Ac3+50°C, preferably Ac3-20°C to Ac3+30°C. It should be noted that Ac3 is usually tested during heating and is higher than the Ae3 temperature. Therefore when isothermally holding a steel at Ac3-30°C for a certain time, it is fully austenitized. The soaking time t2 is in a range from 1 to 300 seconds, preferably 5 to 200 seconds. Annealing at a temperature in the austenite single-phase region is necessary to avoid the formation of polygonal ferrite, which leads to a lower flange formability due to a lower hardness. If T2 is higher than Ac3+50°C or t2 is longer than 300 seconds, austenite grains will grow, which influences the size and distribution of the retained austenite and also slow down the bainitic transformation kinetics later in the partitioning process. An excess amount of fresh martensite may form during final cooling as a result of this incomplete bainitic transformation, which leads to a higher strength but a low ductility and formability. If T2 is lower than Ac3-30°C or t2 is shorter than 1s, reverse transformation to austenite may not proceed sufficiently and / or carbides in the steel strip or sheet may not be dissolved sufficiently. Accordingly, the annealing temperature needs to be equal to or higher than Ac3-30°C, but should not exceed Ac3+50°C. The annealing time t2 is 1 second to 300 seconds, preferably 5 to 200 seconds. After austenization at T2, the steel strip or sheet is first cooled at a slow cooling rate of V3 to a temperature T3. This slow cooling section is important for the current invention. To obtain steel strip or sheet with high formability, ferrite formed in the high temperature range above Ac1 is avoided. Ferrite formed in the low temperature range below the Ac1 point during fast cooling is allowed. During this slow cooling, the C redistribution occurs in the austenite, leading to the formation of a metastable intermediate structure (MIS) including C-lean parts and the C-rich parts, which acts as a central connection among austenite, ferrite and cementite. The MIS is not a stable bulk phase, but acts as precursors of new phases. The formation of the MIS in the austenite phase will facilitate the formation of ferrite in the fast cooling stage and accelerate the bainitic transformation kinetics during austempering. Moreover, the temperature distribution in the steel strip or sheet becomes more uniform, reducing the variation in the microstructure during fast cooling. Thus, the slow cooling has a significant effect of the microstructure although there is no ferrite formed. To make use of this slow cooling effectively, the T3 temperature should be higher than Ar3 to prevent the formation of high temperature ferrite during the slow cooling step. The upper limited of T3 is set to be T2-40°C. T3 is preferably in the range of 680°C to 850°C, preferably 700°C to 820°C. The V3 does not have a significant effect on the microstructure and properties as long as T3 is above Ar3, preferably V3 is in the range of 0.5-15 °C / s, more preferably 1- 5 °C / s. The steel strip is then cooled to a temperature T4 ranging from Mf+100°C to 50°C, preferably from Mf+50oC to 50oC, and more preferably from Mf+30oC to 60oC, preferably at an average cooling rate of V4 of at least 20°C / s. When V4 is less than 20°C / s, the required microstructure is not obtained amongst others because an excess amount of ferrite is produced during cooling. Therefore, the steel strip needs to be cooled preferably at an average rate of 20°C / s or more, preferably of 30 °C / s or more. The upper limit of the cooling rate V4 is not particularly restricted unless variation in temperature occurs in the steel sheet when the cooling is stopped. The upper limit of the average cooling rate is preferably of 200 °C / s or less because the shape of the steel strip is distorted or it is difficult to control the ultimate cooling temperature T4. The temperature T4 is an essential condition for obtaining the microstructure in accordance with the invention. Austenite is largely transformed into martensite by cooling the steel sheet to temperature T4. The formed martensite is transformed into partitioned martensite and untransformed austenite is transformed into retained austenite, bainitic ferrite or fresh martensite in the following reheating for overaging or plating and final cooling. Temperature T4 should be in between Mf+100oC and 50°C to obtain the required microstructures. It has been found that non-relaxed strain gradient in the microstructure affect thet hole expansion capacity in addition to the microstructural inhomogeneity. The non-relaxed strain gradient in the martensite becomes smaller as the quenching stop temperature is increased. If T4 is higher than Mf+100°C, the amount of PM is insufficient, too much RA with low stability and FM can be formed later, which in turn increase the inhomogeneity of the microstructures, and therefore yield a lower HEC value. If T4 is lower than 50°C, the non-relaxed strain gradient is increased and a lower HEC value is also obtained. Although the time t4 is not critical for the microstructural properties, limitation of the typical available production lines requires a short holding time at t4 such that sufficient time is left for the partition step to complete bainitic transformation and to stabilize the retained austenite. The heating time t4 should be controlled within 1 to 10s, preferably within 1 to 5s. The steel strip is then heated to a temperature T5 ranging from Ms to Ms+200°C and overaged for a period t5 ranging from 15 to 150 seconds for partition process. At T5, the martensite transforms to partitioned martensite as C partitioning and the untransformed austenite continues to transform into carbide-free bainitic ferrite. The average C content in the retained austenite is increased as the time t5 is increased, so that retained austenite is made stable. If T5 exceeds Ms+200°C, carbides may precipitate and the desired microstructure of steel cannot be obtained. Preferably, the T5 is below Ms+150°C. If T5 is too low, below the Ms point, the degree of C partitioning is insufficient and the carbon concentration in retained austenite is not high enough to stabilize it in a limited time, which is a known constrain in typical available production lines. Accordingly, the T5 temperature should be above Ms, preferably above Ms+50°C. The partition time t5 at T5 has to be long enough to allow the non-transformed austenite into bainitic ferrite and to allow C enrichment in the retained austenite but short enough to avoid the formation of carbides. When t5 is less than 15s, the partitioning of martensite is insufficient, the desired microstructure may not be obtained, and thus good formability of the steel strip or sheet may not be sufficiently ensured. When t5 is longer than 150 seconds, carbides tend to precipitate in non-transformed austenite, which decreases the carbon content in the retained austenite. Accordingly, t5 is 15 to 150s, preferably 30 s to 100s. A shorter t5 is applied for a higher T5. It is not necessary that the holding temperature T5 be constant as long as the temperatures are within the ranges described. Actually, T5 might gradually increase due to the latent heat produced by bainitic transformation. A hot-dip coating process, such as a Zn or Zn alloy coating process, can be integrated in the process where a temperature T5 for a time t5 can be applied. The steel is then cooled down to 300°C or below at a cooling rate V7 of at least 1°C / s, preferably at least 5°C / s, after which it is further cooled down to ambient temperature at V8, by either forced cooling or uncontrolled natural cooling. A suitable cooling rate V8 for an annealing production line is 5 to 30°C / s. In a further embodiment of the invention, the steel strip or sheet may further be subjected to a coating process as well known to a person skilled in the art, for example hot-dip galvanizing, galvannealing or electrodeposition. The coating may be applied after cooling to ambient temperature or in between the process steps as described above. In a further embodiment of the invention, the steel strip may further be subjected to a coating process as well known to a person skilled in the art, for example hot-dip galvanizing or galvannealing. The coating may be applied after cooling to ambient temperature or in between the process steps. In an alternative embodiment of the invention, the steel strip may be coated after the partitioning process at a temperature T6 and dipping time t6 for a better coating performance. T6 should be in the range of 440°C to 500°C. The time t6 is preferably in the range of 1 to 30s. During this step, the overall microstructure transformation continues and the untransformed austenite continues to transform to bainitic ferrite. Therefore, the total time t5+t6 should be in the range of 15 to 150 s, preferably in the range of 30 to 100 s to limit the precipitation of carbides in non-transformed austenite and obtain the microstructure according to the invention. In a further embodiment of the invention a temper rolling treatment or a tension levelling treatment may be performed with the annealed and optionally zinc coated steel strip in order to fine tune the tensile properties and modify the surface appearance and roughness depending on the specific requirements resulting from the intended use. The temper rolling or tension levelling reduction comes at the expense of a lower elongation potential of the final strip or sheet, so it is preferable to limit the reduction temper rolling or tension levelling reduction is preferably between about 0.2% and 2.0%. It is preferable that temper rolling or tension levelling is performed with a reduction of about 0.5% or less, more preferably of about 0.35% or less, even more preferably of 0.29% or less. Finally, the resulting steel strip or sheet may be coated with resin or oil. In another embodiment of the invention, heat treatment process may be incorporated in a hot forming press cycle to form parts directly. In this embodiment, blanks suitable for producing parts are prepared from the cold rolled strip. A blank is first heated in an electric furnace at T2 for t2. After austenization, the blank is removed from the furnace and placed on the bottom die. In this transfer phase, the blank temperature decreases from temperature T2 slowly to temperature T3 due to heat losses to the environment. The press forming is then applied to form a part. During press forming, the blank temperature decreases at a rate V4 between 60 and 250 °C / s to temperature T4, which can be controlled by the initial die temperature. Finally, the formed part is moved to a furnace with a predetermined temperature T5 and holding for t5 and the cooling to ambient temperature in air. The metallic coating, e.g. Zn coating, can be applied by a separate process. BRIEF DESCRIPTION OF THE FIGURES The invention will now be explained by means of the following, non-limiting figures. Fig.1 is a schematic diagram showing a heat treatment cycle for explaining an annealing process in the method according to the invention. Fig.2 shows the hole expansion capacity as function of temperature T4 for steel A of the examples. The invention will now be illustrated with reference to non-limiting comparative and examples according to the invention. EXAMPLES Steels having compositions as shown in Table 1 were cast into 25 kg ingots of 200 mm x 110 mm x 110 mm in dimensions using vacuum induction. The following process schedule was used to manufacture cold rolled strips of 1 mm thickness: • Reheating of the ingots at 1225°C for 2 hours; • Rough rolling of the ingots from 140 mm to 35 mm; • Reheating of the rough-rolled ingots at 1200°C for 30 min; • Hot-rolling from 35 mm to 4 mm in 6 passes; • Run-out-table cooling: cool from finish rolling temperature (FRT) about 850oC to 900°C to 600°C at a rate of 40 °C / s; • Furnace cooling: strips transferred to a preheated furnace at 600°C and then cooled to room temperature to simulate the coiling process; • Pickling: The hot-rolled strips were then pickled in HCl at 85°C to remove the oxide layers; • Cold rolling: The hot-rolled strips were cold rolled to 1 mm strips; • Heat treating according to the invention. Cold rolled sheets with suitable size were used to simulate the annealing process in a production line by using a continuous annealing simulator (CASIM). The process parameters are given in Table 2 and the results are given in Table 3. Samples for microstructure observations, tensile tests and hole expansion tests were machined from the heat-treated strips. The critical phase transformation points of the steels, Ac1, Ac3, Ms and Mf, determined using dilatometry as herein explained, and are shown in Table 1. The heating rate in the temperature range from 700oC to 900°C was 1.5°C / s. The temperature time profile of the heat treatment process simulating a production condition is schematically shown in Figure 1. The process parameters including temperature T, time, t and heating or cooling rate V are given in Table 2. The microstructures and properties of the cold rolled strips after heat treatments were determined and shown in Table 3. In Table 3, F represents ferrite, HF represent ferrite formed above T3 temperature, LF represent ferrite formed below T3 temperature, RA represents residual austenite, PM represents partitioned martensite, FM represents fresh martensite, FB represents ferritic bainite. In these tables numerical values out of the range of the present invention are underlined. The examples of process 1 to 5 for steel A demonstrate a key feature of the present invention. The microstructure and properties are also a function of T4. In accordance with the invention it has been found that the HEC value is peaked between about 100oC to 150°C, as shown in Figure 2, which is below the Mf point of the subject steel. Process 1 to 4 for steel A are out of the invention because T4 is too high, and the amounts of RA and FM are too high. Process 5 for steel A is out of the invention as T4 is too low. Process 6 for steel A illustrates the effect of temperature T3. If temperature T3 is is too low, ferrite is obtained in excess amount. Table 1. The chemical composition of steels (in wt.%) and the critical phase transformation temperatures (in °C). Alloys C Mn Si Al Cr Cu Ni Nb Mo S P A 0.210 2.000 1.470 0.037 0.010 0.003 0.005 0.001 0.100 0.003 0.004 B 0.213 1.940 1.470 0.020 0.500 0.002 0.004 0.001 0.003 0.005 0.003 C 0.158 2.509 1.019 0.038 0.010 0.003 0.005 0.001 0.003 0.005 0.002 D 0.176 2.013 1.003 0.040 0.010 0.005 0.004 0.001 0.002 0.005 0.003 E 0.177 2.303 1.003 0.037 0.010 0.002 0.030 0.001 0.002 0.004 0.005 F 0.177 2.271 0.993 0.037 0.010 0.107 0.100 0.001 0.002 0.002 0.002 G 0.211 2.289 0.810 0.303 0.010 0.003 0.003 0.001 0.002 0.003 0.004 Table 1. continued Alloys Ti V B N Ca Ac1 Ac3 Ms Mf A 0.001 0.001 0.0005 0.0042 0.0005 730 875 395 180 B 0.002 0.003 0.0003 0.0025 0.0008 728 863 390 195 C 0.001 0.002 0.0007 0.0020 0.0007 725 857 401 190 D 0.004 0.002 0.0005 0.0030 0.0010 728 862 410 198 E 0.002 0.002 0.0004 0.0040 0.0008 726 858 398 185 F 0.001 0.002 0.0005 0.0030 0.0008 724 856 390 200 G 0.001 0.003 0.0005 0.0030 0.0007 729 865 400 185 Table 2. Process parameters used in heat treatments. Process parameters* V2 T2 V3 T3 V4 T4 t4 T5 t5 Process Alloy (°C / s) (°C) (°C / s) (°C) (oC / s) (°C) (s) (°C) (s) 1 A 1.65 880 3.9 720 24.5 480 5 440 48 2 A 1.65 880 3.9 720 24.5 440 5 440 48 3 A 1.65 880 3.9 720 24.5 380 5 440 41.6 4 A 1.65 880 3.9 720 24.5 300 5 440 38.2 5 A 1.65 880 3.9 720 24.5 25 5 440 27 6 A 1.55 870 5.2 660 55.4 250 5 440 48 7 B 1.40 860 3.4 720 65 200 5 400 48 8 C 1.55 870 4.2 700 62 250 5 440 48 9 D 1.55 870 4.2 700 62 250 5 440 48 10 E 1.55 870 4.2 700 62 250 5 440 48 11 F 1.55 870 4.2 700 62 250 5 440 48 12 G 1.44 860 3.9 700 61.5 250 5 440 48 *Rest of the parameters are fixed for all processes. V1 T1 t2 T6 t6 V7 T7 V8 T8 °C / s °C s °C s °C / s °C °C / s °C 23 720 65 460 17 -6.5 300 -11.3 20

[0002]

Claims

CLAIMS 1. A cold-rolled continuously annealed steel strip or sheet having a tensile strength of 980 MPa or more, a total elongation of 8% or more and a hole expansion ratio of 50% or more, and preferably of 60% or more as measured in accordance with the Japan Iron and Steel Federation Standards JFS T 1001, and having a composition comprising, in wt.%: C: 0.150-0.280; Mn: 1.50-3.00; Si: 0.01-1.60; Al: 0.01-0.70; Al + Si: 0.50-2.00; P: up to 0.050; S: up to 0.030; N: up to 0.010; the steel composition optionally also comprising one or more elements selected from: B: up to 0.0050; Ti: up to 0.10; Nb: up to 0.10; V: up to 0.10; Mo: up to 0.20; Cr: up to 0.60; Cu: up to 0.20; Ni: up to 0.20; Ca: up to 0.0030; rare earth elements (REM): up to 0.0030; the balance being Fe and inevitable impurities; and having a microstructure consisting of: bainitic ferrite (BF) and / or partitioned martensite (PM), ferrite phases (high temperature ferrite (HF) and low temperature ferrite (LF)) and second phase comprising fresh martensite (FM) and carbides, and retained austenite (RA), in the following amounts, in vol.%: BF+PM at least 70%; HF+LF up to 20%; HF up to 5%; carbides+FM: up to 3%; RA 3 to 10%.

2. The steel strip according to claim 1, wherein the steel strip has a tensile strength of at least 1040 MPa.

3. The steel strip according to claim 1 or 2, wherein in the microstructure BF+PM is >86 vol.%, and preferably >90 vol.%.

4. The steel strip according to any one of the preceding claims 1 to 3, wherein in the microstructure HF+LF is up to 5 vol.%, and preferably up to 1 vol.%.

5. The steel strip according to any one of the preceding claims 1 to 4, wherein the C concentration in the RA it at least 0.90%, preferably at least 1.0%, and more preferably at least 1.20%.

6. The steel strip according to any one of the preceding claims 1 to 3, wherein the steel composition comprises one or more of the elements in the following amounts: C: 0.150% – 0.280%, and preferably 0.170% to 0.250%; Mn: 1.60% – 2.80%; Si: 0.40% – 1.80%, and preferably 0.40% to 1.50%; Ti: 0.0050% – 0.10%; Nb: 0.0050% – 0.10%; V: 0.010% – 0.10%.

7. The steel strip according to any one of the preceding claims 1 to 6, wherein in the composition one, more or all of Cr, Ni, Cu, Mo, Ti, V, and Nb are present in the following amounts: Cr: up to 0.025%; Ni: up to 0.050%; Cu: up to 0.050%; Mo: up to 0.050%; Ti: up to 0.005%; V: up to 0.010%; Nb: up to 0.005%.

8. The steel strip according to any one of the preceding claims 1 to 7, wherein steel strip is provided with a hot-dip metallic coating layer, preferably a hot-dip galvannealed zinc or zinc alloy coating layer.

9. Method for producing a cold-rolled continuously annealed steel strip or sheet having a composition comprising, in wt.%:C: 0.150-0.280; Mn: 1.50-3.00; Si: 0.02-1.60; Al: 0.01-0.70; Al + Si: 0.50-2.00; P: up to 0.050; S: up to 0.030; N: up to 0.010; the steel composition optionally also comprising one or more elements selected from: B: up to 0.0050; Ti: up to 0.10; Nb: up to 0.10; V: up to 0.10; Mo: up to 0.20; Cr: up to 0.60; Cu: up to 0.20; Ni: up to 0.20; Ca: up to 0.0030; rare earth elements (REM): up to 0.0030; the balance being Fe and inevitable impurities; and having a microstructure consisting of bainitic ferrite (BF) and / or partitioned martensite (PM), ferrite phases (high temperature ferrite (HF) and low temperature ferrite (LF)) and second phase comprising fresh martensite (FM) and carbides, and retained austenite (RA), in the following amounts, in vol.%: BF+PM at least 70%, preferably >86%; HF+LF up to 20%; HF up to 5%; carbides+FM: up to 3%; and RA 3 to 10%; wherein the method comprises the following steps: - preparing a steel melt in a BOS-process and casting a steel slab or ingot; - (re)heating the steel slab or ingot to a temperature in the range of 1100-1300°C; - hot-rolling the slab or ingot to a hot-rolled strip wherein the finishing rolling temperature is equal to or higher than the Ar3 transformation point; - cooling the hot-rolled strip at an average cooling rate of 20°C / s or more, preferably 30°C / s or more, to a temperature below 650 °C and coiling the hot-rolled strip at temperatures in the range from 200 to 550°C; - pickling or descaling the hot-rolled strip prior to cold rolling; - cold rolling at a total cold rolling reduction of between 30% and 80% to obtain a cold-rolled strip; - a thermal treatment by continuous annealing the cold-rolled strip by: * heating the steel strip to a soaking temperature T2 and soaking the steel strip for a time t2 to obtain a fully austenitic microstructure; * cooling the steel strip at a slow cooling rate V3 to a temperature T3 wherein T3 is between Ar3 and (T2-40)°C;* cooling the steel strip at a cooling rate V4 to a temperature T4 between Mf+100oC and 50oC, wherein Mf is the martensite transformation finish temperature; * heating the steel strip from T4 within a time t4 to a temperature T5 between Ms and Ms+200oC for a period t5 between 15 to 150s; * cooling to below 300°C at a cooling rate of at least 1 °C / s followed by cooling to ambient temperature at a cooling rate V8.

10. Method according to claim 9, wherein the strip is at temperature T4 for a period between 1 and 10s, preferably between 1 and 5s.

11. Method according to claim 9 or 10, wherein temperature T4 is between Mf+50oC and 50oC, and preferably between Mf+30oC and 50oC.

12. Method according to any one of the preceding claims 9 to 11, wherein the strip receives a hot-dip metallic coating, preferably a hot-dip zinc alloy coating.

13. Method according to claim 12, wherein the thermal treatment by continuous annealing is combined with a hot-dip metallic coating step after cooling the strip.

14. Method according to any one of the preceding claims 9 to 13, wherein the continuous annealed steel strip is temper rolled and / or tension levelled.

15. Method according to any one of the preceding claims 9 to 14, wherein: t4 is between 1 to 10 seconds, preferably between 1 and 5 seconds, and / or T5 is between Ms and Ms+200oC for a period t5 between 15 to 150 seconds, and / or V8 is preferably between 5 and 30°C / s.

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