High strength cold rolled steel strip with good resistance to retained austenite decomposition for automotive applications

Abstract

The present invention relates to a high strength cold rolled steel strip or plate, comprising a steel composition (wt%) of 0.15-0.25 C, 0.3-0.5 Si, 2.0-3.0 Mn, 0.5-1.0 Al, 0.005-0.5 Cr, and having thermal stability θ>0, where θ=68-500×C+4×Mn+60×Al-22×Si, and the contents of C, Mn, Si, and Al are in wt%), and mechanical stability (k p ) 5 to 35 and the following conditions, namely tensile strength (R m ) ≥ 980MPa, and optionally the following condition: yield strength (R p0.2 )≧400MPa, yield ratio (R p0.2 / R m )≦0.65, Total elongation (A 25 The present invention relates to a high strength cold rolled steel strip or plate having mechanical properties satisfying a retained austenite (RA)≧10% and a microstructure comprising retained austenite (RA)≧8%. The present invention also relates to a method for producing the steel strip or plate and to an automotive structural part comprising the steel plate.

Description

[Technical field]

[0001] The present invention relates to high strength cold rolled steel strip or plate suitable for automotive applications, in particular to cold rolled steel strip or plate having a tensile strength of at least 980 MPa and good resistance to retained austenite decomposition. [Background technology]

[0002] In a wide variety of applications, increased strength levels are a prerequisite for lightweight construction, particularly in the automotive industry, since reduced vehicle mass results in reduced fuel consumption.

[0003] Automotive body parts are often stamped out of steel sheets to form thin complex structural members. However, such parts cannot be manufactured from conventional high strength steels because the formability of the complex structural parts is very poor. For this reason, multi-phase transformation induced plasticity steels (TRIP steels) have attracted a great deal of interest in recent years, especially for use in automotive body structural parts.

[0004] TRIP steels have a multiphase microstructure that contains a metastable retained austenite phase that can produce the TRIP effect. When the steel is deformed, the austenite transforms to martensite, which provides significant work hardening. This hardening effect acts to resist necking in the material and postpone failure in plate forming operations. The microstructure of TRIP steels can significantly change its mechanical properties.

[0005] The problem is that the retained austenite can decompose when the steel is partitioned at 350-450°C after the final anneal. To alleviate this problem, alloying with Si, Al, and P has been suggested to suppress cementite precipitation, thereby stabilizing the austenite.

[0006] However, steel may be subjected to even higher temperatures (>450°C) after partitioning, for example during the manufacture of automobiles, e.g. during hot-dip galvanizing, galvannealed hot-dip galvanizing or when welding parts together.

[0007] There is therefore a need to provide a steel that has better resistance to retained austenite decomposition at high temperatures (>450° C.). Summary of the Invention [Problem to be solved by the invention]

[0008] The present invention relates to a high strength (TRIP) steel strip or plate having a tensile strength of more than 980 MPa. The steel of the present invention is therefore configured to have good resistance to residual austenite decomposition at high temperatures (>450°C). The aim of the present invention is to provide a steel composition which can be processed into structural parts, such as front and centre pillars, vehicle door frame reinforcements, in particular in the automotive industry involving deep drawing operations. Furthermore, it should be possible to produce the steel strip or plate on an industrial scale in a Continuous Annealing Line (CAL) or a Hot Dip Galvanizing Line (HDG) or a Galvannealing Line. [Brief description of the drawings]

[0009] [Figure 1] FIG. 1 illustrates a typical thermal cycle corresponding to a final continuous annealing line (CAL), a hot dip galvanizing line (HDG), or any continuous annealing line (CAL) prior to a hot dip galvannealing line. [Diagram 2] FIG. 2 is a diagram showing a typical thermal cycle of a final continuous annealing line (CAL). [Diagram 3] FIG. 3 is a diagram showing a typical thermal cycle of a hot-dip galvanizing line (HDG). [Figure 4]FIG. 4 is a diagram showing a typical thermal cycle of a galvannealed coating line. DETAILED DESCRIPTION OF THE PREFERRED EMBODIMENTS

[0010] The invention is set out in the claims.

[0011] In a preferred embodiment, the cold rolled steel strip or plate comprises the following alloying elements (in wt%): C 0.15~0.25 Si 0.3~0.5 Mn 2.0~3.0 A1 0.5~1.0 Cr 0.005~0.5 Nb≦0.1 Ti≦0.1 N ≦ 0.05 Mo≦0.5 B ≦ 0.01 V ≦ 0.2 P ≦ 0.05 Ca ≦ 0.05 Cu≦0.1 Ni ≦ 0.2 O ≦ 0.0003 H ≦ 0.0020 The remainder is iron and impurities.

[0012] The importance of the individual elements and their interactions, as well as the limitations of the chemical composition of the alloys as claimed, are briefly described below. All percentages of the chemical composition of the steel are given throughout the specification in weight percent (wt%). The amount of hard phases is given in volume percent (vol%). The upper and lower limits of the individual elements can be freely combined within the limits set in the claims. The arithmetic precision of the numerical values ​​can be increased by one or two orders of magnitude for all values ​​given in this application. Thus, for example, a value given as 0.1% can also be expressed as 0.10 or 0.100%.

[0013] C: 0.15-0.25% C is important to stabilize austenite and to obtain sufficient carbon inside the retained austenite phase. C is also important to obtain the desired strength level. Generally, an increase in tensile strength of the order of 100 MPa can be expected per 0.1% C. If C is less than 0.15%, it is difficult to achieve a tensile strength of 980 MPa. If C exceeds 0.25%, weldability is hindered. The upper limit may thus be 0.25, 0.24, 0.23, 0.22, 0.21 or 0.20%. The lower limit may be 0.15, 0.16 or 0.17%. The preferred range is 0.15-0.25%.

[0014] Silicon: 0.3 to 0.5% Si acts as a solid solution strengthening element and is important to ensure the strength of thin steel sheets. Si inhibits cementite precipitation and is used for austenite stabilization. However, excessively high content leads to the formation of excessively large amounts of silicon oxides on the strip surface. This can lead to cladding on the rolls in CAL or HDG and surface defects on the subsequently produced steel sheets. After cold rolling, these oxides can also cause unwanted problems in galvanizing. Furthermore, excessively high Si content can reduce the RA stability at high temperatures (>450°C) in later stages of the manufacturing process, such as hot-dip galvanizing and galvannealed hot-dip galvanizing operations, or in post-manufacturing operations, such as welding. Furthermore, Si content >0.5% can cause liquid metal embrittlement (LME) during welding. Therefore, the upper limit is 0.5%, which may be limited to 0.49, 0.48, or 0.47%. The lower limit is 0.3%, and may be limited to 0.31, 0.32, 0.33, 0.34 or 0.35. The preferred range is 0.3 to 0.5%.

[0015] Mn: 2.0-3.0% Manganese is a solid solution strengthening element. s By lowering the temperature, it stabilizes austenite and prevents ferrite and pearlite from forming during cooling. c3Mn reduces the temperature and is important for austenite stability, especially at high temperatures (>450°C). If the content is less than 2.0%, it may be difficult to obtain the desired amount of retained austenite, tensile strength of 980 MPa, and the austenitizing temperature may be too high for conventional industrial annealing lines. In addition, at low contents, it may be difficult to avoid the formation of polygonal ferrite. However, if the amount of Mn is higher than 5.0%, segregation problems may occur because Mn accumulates in the liquid phase and causes banding, which may potentially deteriorate the workability. Therefore, the upper limit may be 3.0, 2.9, 2.8, 2.7, 2.6 or 2.5%. The lower limit may be 2.0, 2.1, 2.2 or 2.3%. The preferred range is 2.0-3.0%.

[0016] Al: 0.5-1.0% Al promotes ferrite formation and is also commonly used as a deoxidizer. Al inhibits cementite precipitation and is used for austenite stabilization. Al has been found to be beneficial for RA stability at high temperatures (>450°C). The addition of Al does not adversely affect coatability. The disadvantage of higher amounts of Al is that the M content increases with increasing Al content. s Temperature and A c3 The temperature rises. Thus, the upper limit is 1.0%, and may be limited to 0.9, 0.8, or 0.75%. The lower limit is 0.5%, and may be further limited to 0.6 or 0.7%. The preferred range is 0.5-1.0%.

[0017] Cr: 0.005~0.5% Cr is effective in increasing the strength of steel plates. Cr is an element that forms ferrite and delays the formation of pearlite and bainite. c3 Temperature and M sThe temperature decreases only slightly with increasing Cr content. Cr results in an increase in the amount of stabilized retained austenite. The Cr content is limited to 0.5%. The upper limit may be limited to 0.45, 0.40, 0.35, 0.30 or 0.25%. The lower limit is 0.005% and may be further limited to 0.01, 0.05, 0.1, 0.11, 0.12, 0.13, 0.14 or 0.15%.

[0018] Nb:≦0.1% Nb is generally used in low alloy steels to improve strength and toughness due to its effect on grain size. Nb enhances the strength-elongation balance by refining the matrix microstructure and the retained austenite phase due to the precipitation of NbC. The steel may contain Nb in an amount of ≦0.1%. Intentional addition of Nb is not required according to the invention. The upper limit may therefore be limited to ≦0.03%. The upper limit may be further limited to 0.01 or 0.004%.

[0019] Mo≦0.5% Molybdenum can be added to improve strength. This may further enhance the benefits of NbC precipitation by reducing carbide coarsening kinetics. The steel may contain Mo in an amount of ≦0.5%. The upper limit may be limited to 0.4, 0.3, 0.2, 0.1 or 0.05%. Intentional addition of Mo is not required according to the invention. Therefore, the upper limit may be further limited to 0.03, 0.02 or 0.01%.

[0020] V:≦0.2% The function of V is similar to that of Nb in that it is involved in precipitation hardening and grain refinement. The steel may contain V in an amount of ≦0.2%. The upper limit may be limited to 0.15, 0.10, 0.05, 0.03 or 0.01%. Intentional addition of V is not required according to the invention. The upper limit may therefore be further limited to ≦0.01%.

[0021] Ti:≦0.1% Ti is commonly used in low alloy steels to improve strength and toughness because it affects grain size by forming carbides, nitrides, or carbonitrides. Specifically, Ti is a strong nitride former and can be used to bind nitrogen in the steel. However, its effect tends to saturate above 0.1%. The upper limit may be limited to 0.09, 0.07, 0.05, 0.03, or 0.01%. Intentional addition of Ti is not required according to the present invention. Therefore, the upper limit may be further limited to ≦0.005%.

[0022] Ca≦0.05 Ca may be used to modify non-metallic inclusions. The upper limit is 0.05%, and may be set at 0.04, 0.03, 0.01%. Intentional addition of Ca is not required according to the present invention. Therefore, the upper limit may be limited to ≦0.004%.

[0023] Cu:≦0.1% Cu is an unwanted impurity element that is limited to ≦0.1% by careful selection of the scrap used. The upper limit may be limited to ≦0.06%.

[0024] Ni:≦0.2% Ni is an unwanted impurity element that is limited to ≦0.2% by careful selection of the scrap used. The upper limit may be limited to ≦0.08%.

[0025] B:≦0.01% B increases hardness, but at the same time reduces bendability, and is therefore undesirable in the steel of the present disclosure. B may also make scrap recycling more difficult, and the addition of B may also deteriorate workability. Therefore, the intentional addition of B is undesirable according to the present invention. The upper limit may therefore be limited to ≦0.0006%.

[0026] Other impurity elements may be present in the steel in amounts normally occurring therein.

[0027] It is also preferable to control the nitrogen content so that N is ≦0.05%, preferably ≦0.01%, with the preferred range being 0.001 to 0.008%. Within this range, stable fixation of nitrogen can be achieved.

[0028] Oxygen and hydrogen further include O: ≦0.0003 H: ≦0.0020 can be limited to.

[0029] The θ factor is an index of the thermal stability of the steel, and the composition should satisfy the following conditions: θ>0, Here, θ = 68-500×%C+4×%Mn+60×%A1-22×%Si. A negative θ coefficient is unfavorable for resistance to retained austenite decomposition.

[0030] The lower limit of θ may be 5, 10, 15, 20, or 25.

[0031] The microstructural components are expressed in volume percent (vol%) as follows:

[0032] The steel comprises a matrix of tempered martensite (TM) and / or bainitic ferrite (BF). The total amount is TM+BF≧50% with retained austenite inclusions embedded in the matrix. The upper limit of TM+BF may be 90%. Polygonal ferrite and fresh martensite may be present in the matrix. Retained austenite (RA) is a prerequisite for obtaining the desired TRIP effect. The amount of retained austenite is important for the present invention and should be ≧8%, preferably 10-20%. The upper limit may be 20, 19, 18, 17, 16 or 15%. The retained austenite is preferably mostly acicular. The amount of retained austenite is measured by the saturation magnetic method detailed in Proc. Int. Conf. on TRIP-aided high strength ferrous alloys (2002), Ghent, Belgium, p.61-64.

[0033] Polygonal ferrite (PF) may be in the range of 0-40%. The upper limit may be 30, 20, 10, 5 or 1%. The steel may be free of polygonal ferrite (PF).

[0034] Fresh martensite (FM) may be in the range of 0-10%. The upper limit may be 7, 5, 3 or 1%. The steel may be free of fresh martensite (FM).

[0035] Microstructural ratios can be obtained by cutting a sample from the steel plate and polishing a cross section of the plate perpendicular to the rolling direction. The sample was ground to 1 / 4 of the plate thickness for the measurements. The surface was etched to facilitate identification of the phases. A scanning electron microscope (SEM) may be used at a magnification of 2000x.

[0036] The mechanical properties of the claimed steel are important and at least one of the following requirements should be met: Yield strength (R p0.2 ) ≧400 MPa, preferably 400 to 700 MPa. Tensile strength (R m ) ≧980 MPa, preferably 980 to 1300 MPa. Total elongation (A 25 ) ≧10%, preferably >12% Yield ratio (R p0.2 / R m )≦0.65

[0037] Preferably, all these requirements are met simultaneously.

[0038] Tensile strength (R m ) may be further limited to 1260, 1240, 1220, 1200, 1180, 1160, 1140, 1120 or 1100 MPa.

[0039] Yield strength (R p0.2 The upper limit of the pressure can be further limited to 680, 660, 640, 620, 600, 580, 560, 540 or 520 MPa. The preferred range is 400 to 600 MPa.

[0040] The lower the yield ratio, the easier it is to cold form the material, so the yield ratio is a maximum of 0.65. p0.2 / R m The upper limit of may be further limited to 0.62, 0.60, 0.58, 0.56, 0.54, 0.52 or 0.50. The lower limit may be 0.30, 0.32, 0.34, 0.36, 0.38 or 0.40.

[0041] R m Value, R p0.2 value and total elongation (A 25 ) is derived according to industry standard ISO 6892-1, with samples taken in the longitudinal direction of the strip.

[0042] Mechanical stability (k p ) is a parameter that indicates the mechanical stability of retained austenite (RA). pFactors influencing the value include the chemical composition of the austenite, mainly via carbon enrichment, the grain size (smaller grain size results in a more stable RA), and the morphology of the RA (coarser RA is less stable than lath or acicular RA). For these reasons, the chemical composition of the steel, as well as the heat treatment parameters, are critical.

[0043] Mechanical stability of steel (k p ) should be in the range of 5 to 35, preferably 10 to 35. p The k value indicates the low stability of retained austenite (RA) against mechanical loads. p If the value is too high, RA will already transform during (stress-assisted) elastic loading or at very low plastic strains, thus not increasing the work-hardening behavior of the steel sufficiently to allow high elongation. The present invention aims for optimal stability. k within the suggested range p The k value improves the stability of the RA against mechanical loads and is beneficial for the resistance to residual austenite decomposition. Specifically, the steel should have a k value between 3 and 35 for improved stability of the residual austenite. p value and a positive θ coefficient (θ>0). Mechanical stability (k p ) is determined by intermittent tensile testing. The tensile specimen is deformed to a specific strain, which is between the yielding and before necking of the specimen. The retained austenite content in the undeformed and deformed states is then determined.

[0044] The following relationship provided by Ludwigson and Berger in J. Iron Steel Inst. 1969, vol. 207, pp. 63 applies:

number

[0045] Matsumura et al. suggested in Scr. Metall. 1987, vol. 21, pp. 1301 that p in TRIP-assisted steels can be assumed to be 1. Therefore, k p The value can be derived from a combination of intermittent tensile tests and retained austenite measurements. True strain is the natural logarithm of the ratio of the instantaneous gauge length to the original gauge length in a tensile test. The retained austenite content can be determined by saturation magnetization measurements. The initial retained austenite content (V γ0 ) can be measured on the final heat-treated product. The content of retained austenite after deformation (V γ ) should be taken from the gage length of the deformed tensile specimen.

[0046] The mechanical properties of the steel strip or plate of the present invention can be controlled to a large extent by the alloy composition and the microstructure, which can be controlled by the heat treatment in the CAL, specifically by the isothermal treatment temperature in the partitioning step.

[0047] The steel of the present disclosure can be produced by the following process. a) forming a steel slab having the above composition by converter refining and secondary refining; b) The slab is hot rolled in the austenitic range to produce a hot rolled strip, preferably by reheating the slab to a temperature between 1000°C and 1280°C and rolling the slab completely in the austenitic range. A hot rolled steel strip is obtained at a hot rolling finish temperature of 850°C or higher. c) The hot rolled strip is then coiled at a coiling temperature of 400-580°C. d) The coiled strip is then batch annealed at a temperature of 500-650°C, preferably 550-650°C, for a period of 5-30 hours. e) Optionally, the coiled strip may be subjected to a descaling process such as pickling, either before or after batch annealing. f) The annealed steel strip is then cold rolled with a reduction of at least 50%, preferably about 50-70%. The thickness of the cold rolled strip is preferably 0.9-2.0 mm.

[0048] The cold rolled strip then undergoes a single or double annealing process.

[0049] In the case of a single annealing process, the cold rolled strip is conveyed to a final continuous annealing line (CAL) or a hot dip galvanizing line (HDG). Figure 2 shows the thermal cycle of a final continuous annealing line (CAL), Figure 3 shows the thermal cycle of a hot dip galvanizing line (HDG), and Figure 4 shows the thermal cycle of a hot dip galvannealing line.

[0050] The CAL process includes the following steps: k)A c1 +(A c3 -A c1 ) / 3 and less than 1000°C and annealing the strip with a dew point in the range of -40°C to +10°C for a time greater than 30 seconds; l) quenching the strip by cooling the strip to a quenching temperature QT between 200°C and 400°C. The quenching rate may be between 20 and 60°C / s; m) heating the strip to a partitioning temperature Pt of 250°C to 450°C and maintaining the strip at this temperature for a partitioning time Pt ​​of 10 seconds to 200 seconds, i.e., the partitioning step; n) the strip can be cooled to room temperature. The cooling rate can be between 5 and 60°C / s; and o) Optionally, forming plates from the strips.

[0051] Optionally, step p) subjecting the strip or plate to zinc electroplating or physical vapor deposition (PVD).

[0052] In step k), the annealing temperature TA is preferably below 950°C, more preferably below 900°C.

[0053] In step m), the partitioning temperature PT may be further limited to a range of 350-450°C.

[0054] In step m), the partitioning temperature PT may optionally be equal to the quenching temperature QT, when the latter is in the range of 350-400° C. In step l) and / or n), the strip may be gas quenched. Furthermore, the cooling rate in step n) may be limited to 20-60° C. / s.

[0055] The hot dip galvanizing line (HDG) is treated similarly to the final CAL process, but includes hot dip coating at the end of the partitioning (step m). The strip is immersed in molten zinc (mainly zinc) at about 460°C. Thus, during hot dip coating the temperature becomes above 450°C. The hot dip galvanizing line (HDG) can be the same line as the CAL with the addition of hot dip coating.

[0056] The galvannealed line is identical to the hot dip galvanized line (HDG) with the addition of an annealing step followed by hot dip coating. That is, it is treated similarly to the CAL process, but includes galvannealed coating at the end of the partitioning (step m). Galvannealed coating is a combination of galvanizing at about 480-560 °C and an annealing step to promote a higher concentration of Fe in the ZnFe coating. The annealing step after HDG is exaggerated in Figure 4 to make it more visible. In reality its duration is in the range of seconds.

[0057] In the double annealing process, the cold rolled strip is first conveyed to a continuous annealing line (CAL), which includes the following steps: g) A c1 +(A c3 -Ac1 ) / 1.5 and less than 950°C and annealing the strip at a dew point within the range of -40°C to +10°C for a time greater than 30 seconds; h) Quench to QT < 350℃. The quenching rate may be 20-60℃ / s; i) maintaining the QT for at least 10 seconds; and j) Cooling the strip to room temperature. The cooling rate may be 5-60°C / s.

[0058] In step g) and / or step j), the strip may be gas quenched. Furthermore, the cooling rate in step j) may be limited to 20-60° C. / s.

[0059] The strip then undergoes the same processes as described for the single annealing, i.e. it passes again through a final continuous annealing line (CAL) (or through a second CAL line following the first CAL line), or through a hot-dip galvanizing line (HDG) or a hot-dip galvannealing line. EXAMPLES

[0060] The steels I1-I3 according to the invention and the reference steels R1-R2 were produced by conventional metallurgical techniques by converter refining and secondary refining. The compositions are shown in Table 1. Further elements other than Fe were present only as impurities and below the minimum levels specified herein. Steels I1-I3 and R1 were all within the compositional range of the preferred embodiment, whereas steel R2 had an Al content below the preferred range. The thermal stability (θ coefficient) is positive for steels I1-I3 and R1, but negative for reference steel R2. [Table 1]

[0061] Slabs of steel alloy were produced in a continuous caster. The slabs were reheated and hot rolled to a thickness of 2.8 mm. The hot rolling finish temperature was about 900°C and the coiling temperature was about 500°C. The hot rolled strips were pickled and batch annealed at 620°C for 15 hours to reduce the tensile strength of the hot rolled strips, thereby reducing the cold rolling forces. The strips were then cold rolled in a 5-stand cold rolling mill to a final thickness of 1.4 mm.

[0062] The cold rolled strip was then annealed in a continuous annealing line (CAL). The annealing cycle consisted of heating to the annealing temperature (Table 3) and full austenitization for 150 seconds. The annealed strip was then rapidly cooled to the quenching temperature (Table 3) with a cooling rate of 50°C / s. After quenching, the temperature was increased to the partitioning temperature (Table 3) with a heating rate of 20°C / s and held at the partitioning temperature (Table 3) before quenching to room temperature at 50°C / s. [Table 2]

[0063] It was found that the steels produced according to the invention have excellent mechanical properties as shown in Table 4, in contrast to the inferior mechanical properties of the reference steels R1 and R2. [Table 3]

[0064] All steels have yield strengths above 400 MPa and tensile strengths above 980 MPa. The yield ratios are less than 0.65 for steels I1 to I3, making R1 easier to cold form. The reference steel R2 did not meet the yield ratio requirements. 25 ) was more than 10% for the steel according to the invention and for the reference steel R2, but only 2% for the reference steel R1.

[0065] The mechanical stability (k p) is in the range of 5–35, whereas the mechanical stability (k p ) was 98, outside the expected range.

[0066] The microstructure contained more than 8% retained austenite (RA) for all steels. To test the stability of the retained austenite, the steel samples were heated to 560°C with a heating rate of 20°C / s. As can be seen, the steels of the invention lost at most 20% retained austenite (steel I3) and maintained a retained austenite amount above 8%. Thus, the steels of the invention showed good resistance to retained austenite decomposition.

[0067] Reference steels R1 and R2 lost 60% and 40% of the retained austenite, respectively, both well below the desired minimum of 8%. Thus, the mechanical stability (k p ) performed inferiorly to the steels of the invention (I1-I3) in terms of the stability of retained austenite, and the reference steel R2, which has a negative θ coefficient, also lacked the stability of retained austenite.

[0068] Y.S., T.S., Y.R., T.E., k. p , All RA values ​​were derived according to the methods or specifications disclosed above. [Industrial Applicability]

[0069] The material of the present invention can be widely applied to structural parts in the automotive industry, especially those involving deep drawing, such as front and center pillars, vehicle door frame reinforcements.

Claims

1. A high-strength cold-rolled steel strip or sheet having the following: a) A composition consisting of the following elements (wt%): C 0.15 - 0.25, Si 0.3 - 0.5, Mn 2.0 - 3.0, Al 0.5 - 1.0, Cr 0.005 - 0.5, Nb ≤ 0.1, Ti ≤ 0.1, N ≤ 0.05, Mo ≤ 0.5, B ≤ 0.01, V ≤ 0.2, P ≤ 0.05, Ca ≤ 0.05, Cu ≤ 0.1, Ni ≤ 0.2, O ≤ 0.0003, H ≤ 0.0020, and The balance consisting of iron and impurities; b) Thermal stability θ > 0, Mechanical stability (k ｐ ) 5 to 35 where θ = 68 - 500×C + 4×Mn + 60×Al - 22×Si, The contents of C, Mn, Si, and Al are in wt%; c) Mechanical properties satisfying the following conditions Tensile strength (R ｍ ) ≥ 980 MPa, and Optionally, Yield strength (R ｐ０．２ ) ≥ 400 MPa, Yield ratio (R ｐ０．２ / R ｍ ) ≤ 0.65, and Full extension (A ２５ ) ≥ 10%, preferably > 12% At least one of; d) A microstructure containing the following Retained austenite (RA) ≥ 8%.

2. The high-strength cold-rolled steel strip or sheet according to Claim 1, wherein the microstructure satisfies at least one, preferably all, of the following requirements (vol%): Retained austenite 10 - 20 Fresh martensite 0 - 10 Bainitic ferrite and tempered martensite 50 - 90 Polygonal ferrite 0 - 40.

3. The high-strength cold-rolled steel strip or sheet according to Claim 2, wherein the microstructure contains the following: Retained austenite 10 - 20 Bainitic ferrite and tempered martensite 50 - 90 Fresh martensite ≤ 5 Polygonal ferrite ≤ 5.

4. The high-strength cold-rolled steel strip or sheet according to Claim 1 or 2, wherein the yield ratio is less than 0.

55.

5. The high-strength cold-rolled steel strip or sheet according to Claim 1 or 2, wherein Nb ≤ 0.

01.

6. The high-strength cold-rolled steel strip or sheet according to Claim 1 or 2, wherein Cr ≥ 0.

1.

7. A method for manufacturing the high-strength cold-rolled steel strip or sheet according to Claim 1, comprising the following steps: a) Preparing a steel slab having the composition according to Claim 1, b) Hot-rolling the slab in the austenite region, with the hot-rolling completion temperature being 850°C or higher to obtain a hot-rolled steel strip, c) Coiling the hot-rolled strip at a coiling temperature of 400 - 580°C, d) Batch annealing at a temperature of 500 - 650°C for 5 - 30 hours. e) Optionally, subject the wound strip to a scale removal process such as pickling, before or after the batch annealing. f) Cold roll the annealed steel strip at a reduction ratio of 50% or more. Optionally, the following steps g) to j) g) A ｃ１ +(A ｃ３ -A ｃ１ ), heating the strip to an annealing temperature TA higher than +(A - A) / 1.5 and less than 950 °C and annealing the strip at a dew point within the range of -40 °C to +10 °C for a time exceeding 30 seconds h) Quench the strip to a quenching temperature QT < 350 °C. i) Maintain the quenching temperature QT for at least 10 seconds, and j) Cool the strip to room temperature. k) A ｃ１ +(A ｃ３ -A ｃ１ ), heating the strip to an annealing temperature TA that is higher than (A - A) / 3 and less than 1000 °C, and annealing the strip at a dew point within the range of -40 °C to +10 °C for a time exceeding 30 seconds l) Quench the strip by cooling it to a quenching temperature QT of 200 °C to 400 °C. m) Heat the strip to a partitioning temperature PT of 250 °C to 450 °C and maintain the strip for a partitioning time Pt of 10 seconds to 200 seconds at this temperature, i.e., the partitioning process, and n) Cool the strip to room temperature. o) Optionally, form a plate from the strip.

8. The method according to claim 7, wherein step m) includes a hot dip coating or an alloying hot dip galvanizing process at the end of the partitioning.

9. The method according to claim 7, wherein the quenching temperature QT is in the range of 350 to 400 °C and the partitioning temperature PT in step m) is the same as the quenching temperature QT in step l).

10. The method according to claim 7, wherein the annealing temperature TA in step k) is less than 950 °C, preferably less than 900 °C.

11. An automotive structural part comprising the high-strength cold-rolled steel according to claim 1.

12. The automotive structural part according to claim 11, wherein the structural part is a front pillar, a center pillar, or a vehicle door frame reinforcement of an automobile.

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