A HIGH STRENGTH STEEL STRIP OR SHEET, AND A METHOD FOR PRODUCING THE SAME

Description

CROSS-REFERENCE TO RELATED APPLICATION(S)

This is a National Stage Entry into the United States Patent and Trademark Office from International Patent Application No. PCT/EP2024/057887, filed on Mar. 22, 2024, which relies on and claims priority to Swedish Patent Application No. 2350337-8, filed on Mar. 24, 2023, the entire contents of which both of which are incorporated herein by reference.

FIELD OF THE INVENTION

The present invention relates to high strength steel strip or sheets suitable for applications in automobiles. In particular, the invention relates to cold rolled steel strip or sheets which have been produced by a process comprising double annealing of the cold rolled strip.

BACKGROUND OF THE INVENTION

For a great variety of applications increased strength levels are a pre-requisite for light-weight constructions in the automotive industry, since car body mass reduction results in reduced fuel consumption.

Automotive body parts are often stamped out of sheet steels, forming complex structural members of thin sheet. However, such parts cannot be produced from conventional high strength steels, because of a too low formability for complex structural parts. For this reason, multiphase Transformation Induced Plasticity aided steels (TRIP steels) have gained considerable interest in the last years, in particular for application in auto body structural parts.

TRIP steels possess a multi-phase microstructure, which includes a meta-stable retained austenite phase, which is capable of producing the TRIP effect. When the steel is deformed, the austenite transforms into martensite, which results in remarkable work hardening. This hardening effect acts to resist necking in the material and postpone failure in sheet forming operations. The microstructure of a TRIP steel can greatly alter its mechanical properties.

Quenching & Partitioning (Q&P) is an annealing cycle which consists of a quenching and a partitioning step. In the quenching step, fully austenitized or intercritically annealed steels are quenched to a temperature between the martensite start temperature M_S and the martensite finish temperature M_F to reach a partial martensitic transformation. The quenched steels are then held at a temperature either at or above the initial quenching temperature. Austenite that prevails after quenching is stabilized through carbon partitioning from martensite into the austenite during the partitioning.

TRIP-aided sheet steels with annealed martensite matrix can be produced by double annealing of the cold rolled strip. The martensitic microstructure leads to a predominantly lamellar structure after the second annealing step, and these steels possess a large amount of plate-like retained austenite along annealed martensite lath boundary. Due to the lamellar structure of the austenite, the stability of the retained austenite in these steels is generally high. The high stability of the retained austenite improves the TRIP effect and the global ductility. The homogeneous lamellar structure of the austenite benefits local ductility.

US 2016/0177414 A1 disclose a cold rolled steel that is annealed twice. After the first annealing the steel is cooled to 320-500° C. and held at 320-500° C. for 30 seconds or more. After the second annealing the steel is cooled to 120-320° C. and thereafter held at 320-500° C. for 30 seconds or more. The steel has a microstructure comprising 3-20% ferrite, 5-20% retained austenite, 5-20% martensite, and remained being bainite and/or tempered martensite.

US 2017/0327924 A1 disclose a cold rolled steel that is annealed twice. After the first annealing the steel is cooled to room temperature or to a controlled temperature above room temperature, preferably below the martensite finish temperature, to achieve a predominantly martensitic structure. The second annealing comprises soaking in an intercritical range from 720 to 850° C. and followed by holding at 370 to 430° C. The steel has a microstructure comprising of primarily ferrite (50 up to 80% or higher), 5-25% retained austenite and 0-15% fresh martensite.

US2020/0392598 A1, US 2020/392610 A1, US 2020/00440421 A1, US2019/203316 A1, JP 2004238679 A2, WO 2022/123289 A1, US 2020/354823 A1, US2021/207236 A1 are further examples where multiple annealing has been suggested.

An object of the invention is to further improve the steel properties, preferably properties related to local and/or global ductility, of double annealed steels.

SUMMARY OF THE INVENTION

The present invention is directed to steel strip or sheets having a tensile strength of 780-1350 MPa. The steel strip or sheets are produced in a double annealing process of which the first annealing cycle includes quenching to a temperature between M_S-20 and M_F and partitioning at a temperature above the quenching temperature, and the second annealing cycle includes quenching to a temperature in the range of 150-500° C. and partitioning at a temperature above the quenching temperature or isothermally. The characteristic of the retained austenite leads to an improvement in local and global ductility. Thereby providing improved formability and crashworthiness of safety related components manufactures from the steels of the invention.

BRIEF DESCRIPTION OF THE DRAWING

FIG. 1 shows schematically the annealing cycles of the invention.

DETAILED DESCRIPTION OF EMBODIMENT(S) OF THE INVENTION

In a preferred embodiment the strip or sheet has a composition consisting of the following alloying elements (in wt. %):

	C	0.1-0.2
	Si	0.3-0.9
	Mn	1.4-3.0
	Al	0.03-1.0
	Si + Al	0.5-1.5
	Optionally
	Ti	≤0.1
	Nb	≤0.1
	V	≤0.1
	Cr	≤0.5
	B	≤0.005
	Mo	≤0.2
	balance Fe apart from impurities.

The composition is excluding any coatings applied to the strip or sheet.

The importance of the separate elements and their interaction with each other as well as the limitations of the chemical ingredients of the claimed alloy are briefly explained in the following. All percentages for the chemical composition of the steel are given in weight % (wt. %) throughout the description. Upper and lower limits of the individual elements can be freely combined within the limits set out in the claims. The arithmetic precision of the numerical values can be increased by one or two digits for all values given in the present application. Hence, a value of given as e.g. 0.1% can also be expressed as 0.10 or 0.100%. The amounts of the microstructural constituents are given in volume % (vol. %).

C : 0.1 - 0.2 %

C stabilizes the austenite and is important for obtaining sufficient carbon content within the retained austenite phase. C is also important for obtaining the desired strength level. Generally, an increase of the tensile strength in the order of 100 MPa per 0.1% C can be expected. To achieve sufficient tensile strength C should be at least 0.10%. The upper limit may be 0.20, 0.19 or 0.18%. The lower limit may be 0.10, 0.11, 0.12, 0.13, 0.14, or 0.15%. A preferred range is 0.13-0.18.

Si : 0.3 - 0.9 %

Si acts as a solid solution strengthening element and is important for securing the strength of the thin steel strip. Si suppresses the cementite precipitation and is essential for austenite stabilization. However, if the content is too high, then too much silicon oxides will form on the strip surface, which may lead to cladding on the rolls in the continuous annealing line (CAL) and, as a result there of, to surface defects on subsequently produced steel sheets. The upper limit may be 0.9, 0.8, 0.7, or 0.6%. The lower limit may be 0.3, 0.4, or 0.5%. A preferred range is 0.3-0.8.

Mn : 1.4 - 3. %

Manganese is a solid solution strengthening element, which stabilises the austenite by lowering the M_S temperature and prevents ferrite and pearlite to be formed during cooling. In addition, Mn lowers the A_c3 temperature and is important for the austenite stability. At a content of less than 1.4% it might be difficult to obtain the desired amount of retained austenite and a sufficient tensile strength. Furthermore, the required austenitizing temperature might be too high for conventional industrial annealing lines. In addition, at lower contents it may be difficult to avoid the formation of polygonal ferrite. If the amount of Mn is too high problems with segregation may occur because Mn accumulates in the liquid phase and causes banding, resulting in a potentially deteriorated workability. The upper limit may be 3.0, 2.9, 2.8, 2.7, 2.6, 2.5, or 2.4%. The lower limit may be 1.4, 1.5, 1.6, 1.7, or 1.8%. A preferred range is 1.7-2.4.

Al : 0.03 - 1. %

Additions of Al can increase the carbon content in the retained austenite. Al can also be used as a deoxidizer. Al, like Si, is not soluble in the cementite can therefore delay cementite formation during bainite formation and martensite tempering. An addition of Al can further improve galvanization and reduce the susceptibility to Liquid metal embrittlement. However, the M_S temperature is also increased with increasing Al content. A further drawback of Al is that it results in an increase of the A_c3 temperature. The upper level may be 1.0, 0.9, 0.8, or 0.7%. The lower limit may be set to 0.01, 0.02, 0.03, 0.04, 0.05, 0.06, 0.07, 0.08, 0.09, or 0.1%. A preferred range is 0.03-0.8. If Al is used for deoxidation only then the upper level may then be 0.09, 0.08, 0.07 or 0.06%.

Si + Al : 0.5 - 1.5 %

Si and Al suppress the cementite precipitation during bainite formation. Their combined content is therefore preferably at least 0.5%. A preferred range is 0.6-1.3.

Optional Elements

Ti : ≤ 0 . 1 ⁢ %

Ti is commonly used in low alloyed steels for improving strength and toughness, because of its influence on the grain size by forming carbides, nitrides or carbonitrides. In particular, Ti is a strong nitride former and can be used to bind the nitrogen in the steel. However, the effect tends to be saturated above 0.1%. The upper limit may be restricted to 0.09, 0.07, 0.05, 0.03, 0.01, or 0.005%. A deliberate addition of Ti is not necessary according to the present invention.

Nb : ≤ 0 . 1 ⁢ %

Nb is commonly used in low alloyed steels for improving strength and toughness, because of its influence on the grain size. Nb increases the strength elongation balance by refining the matrix microstructure and the retained austenite phase due to precipitation of NbC. The steel may contain Nb in an amount of ≤0.1%. The upper limit may be restricted to 0.09, 0.07, 0.05, 0.03, 0.01, or 0.005%. A deliberate addition of Nb is not necessary according to the present invention.

V : ≤ 0 . 1 ⁢ %

The function of V is similar to that of Nb in that it contributes to precipitation hardening and grain refinement. The steel may contain V in an amount of ≤0.1%. The upper limit may be restricted to 0.09, 0.07, 0.05, 0.03, or 0.01%. A deliberate addition of V is not necessary according to the present invention.

Cr : ≤ 0.5 %

Cr is effective in increasing the strength of the steel sheet. Cr is an element that forms ferrite and retards the formation of pearlite and bainite. The A_c3 temperature and the M_S temperature are only slightly lowered with increasing Cr content. Cr results in an increased amount of stabilized retained austenite. When above 0.5% it may impair surface finish of the steel, and therefore the amount of Cr is limited to 0.5%. The upper limit may be 0.50 or 0.40, 0.30, 0.20, 0.10 or 0.05%. The lower limit may be 0.01, 0.03, 0.05, 0.07, 0.10, 0.15, 0.20 or 0.25%. A deliberate addition of Cr is not necessary according to the present invention.

B : ≤ 0.005 %

B suppresses the formation of ferrite and improves the weldability of the steel sheet. In order to have a noticeable effect at least 0.001% should be added. However, excessive amounts of B deteriorate the workability. B increases hardness but may come at a cost of reduced bendability and can make scrap recycling more difficult. A deliberate addition of B is not necessary according to the present invention.

Mo ≤ 0.2 %

Molybdenum is a powerful hardenability agent. It may further enhance the benefits of NbC precipitates by reducing the carbide coarsening kinetics. The steel may therefore contain Mo in an amount up to 0.2%. Mo delays the decomposition of austenite and stabilizes the retained austenite. Amounts of more than 0.2% results in high costs. The upper limit may be restricted to 0.2, 0.1, 0.05, 0.01%. The lowest amount may be set to 0.001, 0.005, 0.01, 0.02, 0.03, 0.04 or 0.05%. A deliberate addition of Mo is not necessary according to the present invention.

Impurities

The following impurities may optionally be limited as disclosed below.

Ca : ≤ 0.05 %

Ca could be used for the modification of the non-metallic inclusions. The upper limit is 0.05% and may be set to 0.04, 0.03, 0.01, or 0.005%. A deliberate addition of Ca is not necessary according to the present invention.

Cu : ≤ 0.1 %

Cu is an undesired impurity element that is restricted to ≤0.1% by careful selection of the scrap used. The upper limit is 0.1% and may be further restricted to 0.05%.

Ni : ≤ 0.2 %

Ni is also an undesired impurity element that is restricted to ≤0.2% by careful selection of the scrap used. The upper limit is 0.2% and may be further restricted to 0.1 or 0.05%.

Other impurity elements may be comprised in the steel in normal occurring amounts. However, it is preferred to limit the amounts of P, S, As, Zr, Sn to the following optional maximum contents:

P : ≤ 0.05 , ≤ 0.04 , ≤ 0.03 or ≤ 0.02 % S : ≤ 0.05 , ≤ 0.03 , ≤ 0.01 , ≤ 0.005 or ≤ 0.001 % As : 0.02 , or ≤ 0.01 % Zr : ≤ 0.01 , or ≤ 0.006 % Sn : ≤ 0.03 , or ≤ 0.015 %

Oxygen and Hydrogen can Further be Limited to

O : ≤ 0.001 , or ≤ 0.0003 % H : ≤ 0.005 , or ≤ 0.002 %

It is also preferred to control the nitrogen content to the range:

• • N: ≤0.015%, preferably 0.001-0.008%

Mechanical Properties

The steel should fulfil the following condition on the tensile strength:

• • TS, Tensile Strength (R_m) 780-970 MPa, preferably 800-900 MPa

And optionally one or more of the following mechanical properties:

YS, Yield Strength (Rp0.2)	400-800 MPa, preferably 500-700 MPa
HER, Hole Expansion Ratio (λ)	≥40%, preferably 50-90%
TE, Total Elongation (A80)	≥15%, preferably 20-30%
TS*TE*HER/100	≥11 000 MPa %, preferably 12 000-20 000 MPa %

Preferably, all these requirements are fulfilled at the same time.

TS, YS, TE are examples of properties related global ductility. HER is a property related to local ductility.

The upper limit of TS, Tensile Strength (R_m) can be restricted to 970, 960, 950, 940, 930, 910, 900, 890, 880, 870 or 860 MPa.

The R_m, R_p0.2 values as well as the total and/or ultimate elongation are derived in accordance with the Industrial Standard DIN EN ISO 6892-1, wherein the samples with a gauge length of 80 mm are taken in the longitudinal direction of the strip.

The hole expansion ratio (λ) is determined by the hole expansion test according to ISO/WD 16630:2009 (E). In this test a conical punch having an apex of 60° is forced into a 10 mm diameter punched hole made in a steel sheet having the size of 100×100 mm². The test is stopped as soon as the first crack is determined, and the hole diameter is measured in two directions orthogonal to each other. The arithmetic mean value is used for the calculation.

The hole expanding ratio (λ) in % is calculated as follows:

λ = ( Dh - Do ) / Do × 100

wherein Do is the diameter of the hole at the beginning (10 mm) and Dh is the diameter of the hole after the test.

The strip or sheet thickness of the final product may be 0.1-4 mm, preferably 0.2-3 mm. The strip or sheet width may be 500-2000 mm, preferably 700-1750 mm in non-slit condition.

Microstructure

The microstructural constituents are in the following expressed in volume % (vol. %).

	double tempered martensite	20-70
	tempered martensite	0-70
	bainite	0-70
	fresh martensite + carbides	0-10
	retained austenite (RA)	5-25
	bainite + tempered martensite	10-70
	other phases	0-10

Each of the phases or combinations of phases in the list above may balance the microstructure.

The double tempered martensite is tempered martensite from the first quenching and partitioning (step i and j) that is tempered once more in the second partitioning step (step n). The lower limit of double tempered martensite may be 20, 25, 30, 35, 40 or 45 vol. %. The upper limit may be 70, 65, 60, or 55 vol %. A preferred range is 40-60 vol %.

Tempered martensite is martensite from the second quenching step that is tempered in the second partitioning step. The lower limit of tempered martensite may be 0, 10, or 15 vol. %. The upper limit may be 70, 60, 50, 40, or 30 vol %. A preferred range is 10-40 vol. %.

Fresh martensite can be formed upon final cooling after the second quenching and partitioning (step m and n). Small amount of carbides may precipitate in the matrix. The upper limit of fresh martensite+carbides may be 10 or 5 vol. %. A preferred range is 0-5 vol. %.

The steel may contain bainite. The upper limit of bainite may be 70, 60, 50, 40, 30, or 20 vol %. The lower limit may be 0, 5, or 10 vol %. A preferred range is 0-30 vol %.

The amount of tempered martensite+bainite is preferably 10-70 vol. %, more preferably 20-40 vol %.

Retained austenite is a prerequisite for obtaining the desired TRIP effect. The amount of retained austenite should be in the range of 5-25 vol %, preferably 10-20 vol %. The retained austenite has a predominantly homogenous lamellar structure which gives an optimal mechanical stability improving the TRIP effect, improved global ductility, and improved local ductility.

The steel may optionally contain up to 10 vol. % of other phases than the above mentioned, for example polygonal ferrite. Preferably, less than 5 vol. % of other phases, most preferably the steel does not contain any other phases.

The microstructure, including the amount of each phase, can be identified in scanning electron microscope (SEM) using 2000 times magnification. Preferably by cutting out a sample from a steel plate and polishing a cross section of a plate parallel to the rolling direction. The microstructure should be taken from 14 of the thickness. The surface can be etched to make the phases easier to identify.

Electron Backscatter Diffraction, EBSD, can be used to perform quantitative microstructural analyses in the Scanning Electron Microscope. For instance, double tempered martensite can be identified by a local misorientation 0° up to and including 0.5° and tempered martensite by a local misorientation of above 0.5° up to and including 1.2°.

The metallurgical reason for this difference stems from a lower lattice distortion as carbon further diffuses away from the body-centered-tetragonal (bct) lattice of martensite during tempering to form carbides or to partition into RA. Furthermore, the annihilation of dislocations during tempering, resulting in a lower dislocation density, that decreases the local misorientation of the double tempered martensite.

However, the amount of retained austenite is preferably determined by means of the saturation magnetization method described in detail in Proc. Int. Conf. on TRIP-aided high strength ferrous alloys (2002), Ghent, Belgium, p. 61-64.

The improved characteristic of the retained austenite can be determined by one or more of the parameters described below.

Mechanical stability (k_p) is a parameter that describes the mechanical stability of the retained austenite (RA). Factors influencing the k_p value include the chemical composition of the austenite mainly via carbon enrichment, the grain size—smaller grain size leads to more stable RA, the morphology of the RA—globular RA is less stable than lath or needle shaped RA. For these reasons, the chemical composition of the steel as well as the heat treatment parameters are decisive.

The RA should have a mechanical stability (k_p) in the range of 10-35, preferably 15-25. A k_p-value above 35 indicates lower stability of retained austenite (RA) against mechanical loading, If the k_p-value is too high RA already transforms during elastic loading (stress-assisted) or at very low plastic strains and therefore does not sufficiently increase the work hardening behaviour of a steel and enables high elongations. The invention aims for an optimal stability of RA. A k_p-value in the suggested range improves the stability of RA against mechanical loading and is beneficial for withstandability to RA decomposition. The upper limit may be 35, 30, 25, or 20. The lower limit may be 10, 12, 14, 15 or 16.

The mechanical stability (k_p) can be determined using interrupted tensile testing. Tensile samples are deformed to a certain strain that lies between yielding and before necking of the specimen. Subsequently the retained austenite content in the undeformed and deformed state is determined.

Following relation given by Ludwigson and Berger in J. Iron Steel Inst. 1969, vol. 207, pp. 63 is applied:

V γ0 - V γ V γ = k p × ε p

• • V_γ0 . . . initial retained austenite content
  • V_γ . . . retained austenite content after deformation
  • ε . . . true strain
  • p . . . constant related to autocatalytic effect
  • k_p . . . indication for retained austenite stability

Matsumura et al. suggested in Scr. Metall. 1987, vol. 21, pp. 1301 that in TRIP aided steels p can be assumed to be 1. Therefore, the k_p-value can be derived from the combined interrupted tensile testing 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 measurement. The initial retained austenite content (V_γ0) can be measured in the final heat treated sample. The sample for the retained austenite after deformation (V_γ) should be taken out of the gauge length of the deformed tensile specimen. Samples with a gauge length of 80 mm are taken in the longitudinal direction of the strip.

The carbon content, x_cγ, in the retained austenite may optionally fulfil (in wt. %):

x c ⁢ γ 0.7 - 1.

In this range of carbon content in the retained austenite the optimal combination of global and local formability can be reached. A preferred range is 0.8-0.95 wt %. The upper limit may further be restricted to 0.90.

The carbon content in the retained austenite can be calculated from the lattice parameter α_γ obtained from the γ reflections in X-ray diffraction (XRD) measurements using the following equation described in N. H. van Dijk, A. M. Butt, L. Zhao, J. Sietsma, S. E. Offerman, J. P. Wright, and S. van der Zwaag, Thermal stability of retained austenite in TRIP steels studied by synchrotron X-ray diffraction during cooling, Acta Materialia 53 (2005) 5439-5447:

a γ = 3.556 + 0.0453 x c ⁢ γ + 0.00095 x Mn + 0.0056 x Al

where x_cγ, x_Mn, and x_Al are the content of carbon, manganese, an aluminium in austenite (in wt. %). The contents for x_Mn and x_Al are assumed to be equal to the nominal content of the alloy.

The amount of retained austenite (RA vol. %), the mechanical stability of the austenite (k_p), and the carbon content (x_cγ) of the retained austenite should fulfil the following relation:

RA*k_p/x_cγ≥200, preferably 220-400

The mean equivalent circle diameter (ECD) in μm and/or the mean aspect ratio (AR) of the retained austenite (RA) may optionally fulfil the following relations:

	AR	≥2.0, preferably 2.03-2.30
	ECD	≤0.57, preferably 0.45-0.55

Equivalent circle diameter (ECD) in μm and aspect ratio (AR) of the retained austenite (RA) can be determined by means of Electron backscatter diffraction (EBSD). ECD=2√(A/π), where A is the measured area of the grain. Aspect ratio=m/n, where m is the average length of the major axis of the grain and n the average length of the minor axis of the grain. The measurement methods are described in Li, M., Wilkinson, D. and Patchigolla, K. (2005) Comparison of Particle Size Distributions Measured Using Different Techniques. Particulate Science and Technology, 23, 265-284.

Further Definitions

Temperatures are given in degrees Celsius throughout the description.

Ae1 and Ae3 represent the equilibrium transformation temperatures. Austenite is completely stable above Ae3 and partially unstable between Ae3 and Ae1. Ae1 and Ae3 are calculated by means of ThermoCalc 2022 TCFE 12.

M_S temperature was calculated using the M_S formula found in S. Kaar, K. Steineder, R. Schneider, D. Krizan and C. Sommitsch: “New M_S-formula for exact microstructural prediction of modern 3rd generation AHSS chemistries”, Scr. Mater., Vol. 200, 2021, 113923.

Ms = 692 - 502 * ( C + 0.86 N ) ^ 0.5 - 37 * Mn - 14 * Si + 20 * Al - 11 * Cr

The M_S formula uses the content in weight % of each element.

M_F formula is derived from Koistinen-Marburger equation found in (Koistinen, D. and Marburger, R: “A general equation prescribing the extent of the austenite-martensite transformation in pure iron-carbon alloys and plain carbon steels”, Acta Metall, 7, 1959, pp. 59-60):

M F = Ms + [ ln ⁡ ( 1 - f M ) ] / 0.011 ,

where the fraction of martensite f_M is 0.95.

Production of a Cold Rolled Strip.

A cold rolled steel strip may be produced by the following steps:

• • a) making steel slabs of the conventional metallurgy by converter melting and secondary metallurgy with the composition suggested above.
  • b) The slabs are hot rolled in austenitic range to a hot rolled strip. Preferably by reheating the slab to a temperature between 1000° C. and 1280° C. Preferably rolling the slab completely in the austenitic range wherein the hot rolling finishing temperature is greater than or equal to 850° C. to obtain the hot rolled steel strip.
  • c) Thereafter, the hot rolled strip may be coiled at a coiling temperature in the range of 400-700° C.
  • d) Optionally annealing at a temperature in the range of 500-950° C. Preferably batch annealing at 500-700° C., for a duration of 5-30 h. The strip may alternatively be continuously annealed at temperature in the range of 650-950° C. for 10-200 s.
  • e) Optionally subjecting the coiled strip before and/or after the batch annealing to a scale removal process, such as pickling.
  • f) Thereafter cold rolling the annealed steel strip at a reduction rate between 20-90%, preferably around 50-70% reduction.

The steps a)-f) of producing the cold rolled strip described above is an example on how the cold rolled strip can be produced. The invention may be applied to cold rolled strips that are produced by other known steps.

According to the invention the cold rolled strip is subjected to double annealing process disclosed in step a) to p) below. FIG. 1 show the heat cycle of the double annealing process.

The first annealing cycle described in step g) to k) is a Quench and Partitioning (Q&P) process, in which the steel is annealed and thereafter quenched to a temperature between M_S-20° C., and M_F followed by partitioning at a temperature above the quenching temperature, and finally cooled to room temperature. The second annealing cycle is described in step 1) to p).

The First Annealing Cycle:

• • g) Providing a cold rolled steel strip to a continuous annealing line, the steel strip having a composition as suggested above.
  • h) heating the strip at a rate (HR1) of to 1-20° C./s to a first annealing temperature (Tan1) between 800° C. and 1000° C., and soaking for 10-300 s (tan1), preferably the first annealing temperature (Tan1) is above Ae3.
  • i) cooling the strip at a rate (CR1) of 10-100° C./s to a first quenching temperature (TQ1) between M_F and (M_S-20) ° C. The lower limit may further be restricted to the highest of M_F and value chosen from 150, 160, 170, and 180° C. The upper limit may further be restricted to M_S-30, M_S-40, M_S-50, M_S-60, M_S-70, M_S-80, M_S-90, or M_S-100° C.
  • j) heating the cooled strip at a rate (HR2) of 1-100° C./s to a first partitioning temperature (Toa1) in the range of the first quenching temperature (TQ1)+10° C. to 500° C. and partitioning the strip for 20-1000 s (toa1). Preferably the first partitioning temperature (Toa1) is above M_S. The upper limit may be restricted to 500, 480, 450, 430, or 410° C. The lower limit may be restricted to 260, 280, 300, 320, 340, 360, or 380° C.
  • k) cooling the strip to a temperature below 50° C., preferably to room temperature, at a rate (CR2) of 1-50° C./s.

After the first annealing cycle the matrix of the steel comprises of tempered martensite with a comparably large amount of carbides and RA. The larger amounts of carbides and RA prior to the second annealing ensures more nucleation sites for the austenite formation. This results in the second annealing cycle in the formation of a higher amount of RA with an optimal stability against strain induced martensitic transformation (SIMT), having a fine microstructure with a lamellar morphology that improves global and local formability. This combination of global and local formability is better than all known double annealed concepts.

The second annealing cycle following the first annealing cycle (step g to k):

• • l) heating the strip at a rate (HR3) of 1-20° C./s to a second annealing temperature (Tan2) between 700° C. and 900° C. and soaking for 10-300 s (tan2). Preferably, the second annealing temperature (Tan2) is in the intercritical range, i.e. between Ae1 and Ae3. However, the second annealing temperature (Tan2) may be allowed to be above Ae3.
  • m) cooling the strip at a rate (CR3) of 10-100° C./s to a second quenching temperature (TQ2) in the range of 150-500° C. The second quenching temperature (TQ2) may be below M_S, preferably below (M_S-20) ° C., and must be above M_F.
    • The upper limit may be restricted to 500, 450, 400, 350, 300, or 250° C. The lower limit may be restricted to 150, 170, 200, 220, 240, 260, 280 or 300° C.
  • n) partitioning the strip for 20-1000 s (toa2) at a second partitioning temperature (Toa2) in the range of the second quenching temperature (TQ2) to 500° C. The upper limit may be restricted to 500, 480, 450, or 430° C. The lower limit may be restricted to 150, 200, 250, 300, or 350° C. The heating rate (HR4) to the second partitioning temperature (Toa1) may be in the range of 1-100° C./s. In case of isothermal tempering, there is obviously no heating rate from the second quenching temperature (TQ2) to the second partitioning temperature (Toa1), i.e. in this case the heating rate (HR4) is 0° C./s.
  • o) cooling the strip to a temperature below 50° C., preferably to room temperature, at a rate of 1-50° C./s (CR4).
  • p) optionally making sheets from the strips.

Optional Coating

The steel sheet or strip may optionally be coated and comprise a zinc or a zinc-alloy coating. The coating can e,g. be applied by:

• • Electrogalvanizing (EG) including electroplating.
  • Physical Vapor Deposition (PVD).
  • Hot Dip Galvanizing (HDG) in a Hot Dip Galvanizing Line, in which the strip is immersed in a molten zinc bath at the end of the final partitioning (step n). A Hot Dip Galvanizing Line can be the same line as a Continuous Annealing Line with added hot dip coating.
  • Galvannealing (GA) in a Galvannealing line, which is the similar as the Hot Dip Galvanizing Line with the addition of an annealing step following the hot dip coating. I.e. processed the same way as the Continuous Annealing Line, but including galvannealing at the end of the final partitioning (step n). Galvannealing is a combination of galvanizing and annealing around 480-560° C. in order to facilitate a higher degree of Fe in the ZnFe coating.

A zinc alloy coating may comprise in weight %

	Mg	0.1-10
	Al	0.1-10

Optionally one or more of:

• • Bi, Pb, Sn, Sb, Si, Ti, Ca, Mn, La, Ce, Cr, Ni and Zr in a total amount of 0.01-1.0 Balance Zn and impurities.

A galvannealed coating may contain 5-20 wt. % of diffused Fe.

Other coating composition known in the art can be applied.

EXAMPLES

Two alloys L1 and L2 were produced by conventional metallurgy by converter melting and secondary metallurgy. The compositions of the alloys (elements are in [wt %]) are shown in table 1, further elements were present only as impurities, and below the lowest levels specified in the present description. Determined values for M_S, M_F, Ae1 and Ae3 (in ° C.) are also shown in Table 1.

Unless otherwise specified, parameter values throughout the examples are determined by the methods given previously in the description.

TABLE 1
Alloy	C	Si	Mn	Al	Other	MS	MF	Ae1	Ae3
L1	0.142	0.57	1.87	0.05	0.37 Cr	420	148	681	820
					0.02 Nb
L2	0.166	0.38	2.35	0.60	0.23 Cr	404	132	668	855
					0.02 Nb

The alloys L1 and L2 were continuously cast and cut into slabs. The slabs were reheated and hot rolled in austenitic range to a thickness of about 2.8 mm. The hot rolling finishing temperature was about 900° C. The hot rolled steel strips where thereafter coiled at a coiling temperature of 630° C. The coiled hot rolled strips were pickled and batch annealed at about 624° C. for 10 hours in order to reduce the tensile strength of the hot rolled strip and thereby reducing the cold rolling forces. The strips were thereafter cold rolled in a five stand cold rolling mill to a final thickness of about 1.4 mm.

All steels were subjected to two final annealing cycles in a continuous annealing line, where Table 2a show the process values of the first annealing cycle and Table 2b the process values of the second annealing cycle. The annealing cycle is schematically shown in FIG. 1. The reference samples are denoted by Rn, where n=0, 1, 2, 3, and the inventive samples are denoted Sn, where n=1 . . . 5.

TABLE 2a
		HR1,	Tan1,	tan1	CR1	TQ1	HR2	Toa1,	toa1
Alloy	Sample	° C./s	° C.	s	° C./s	° C.	° C./s	° C.	s

L1	R0	3	820	120	50	225	40	380	450
L1	R1	4	900	105	80	25	—	—	—
L1	R2	4	900	105	80	25	—	—	—
L2	R3	3	950	150	60	25	—	—	—
L1	S1	3	850	130	28	200	19	350	450
L1	S2	3	850	120	30	200	22	350	400
L1	S3	3	850	130	28	200	20	350	450
L1	S4	3	850	130	28	200	20	350	450
L2	S5	3	950	150	28	200	18	400	500

TABLE 2b
		HR3	Tan2	tan2	CR3	TQ2	HR4	Toa2	toa2
Alloy	Sample	° C./s	° C.	s	° C./s	° C.	° C./s	° C.	s
L1	R0	—	—	—	—	—	—	—	—
L1	R1	18	810	45	80	240	30	420	70
L1	R2	18	810	45	70	280	20	420	70
L2	R3	8	850	64	17	200	26	430	100
L1	S1	3	820	130	18	330	9	400	450
L1	S2	3	790	105	25	325	13	400	360
L1	S3	18	810	45	65	320	Isoth.	320	70
L1	S4	18	790	45	65	310	Isoth.	310	70
L2	S5	8	845	60	20	210	30	420	100

All samples, except for sample R0, were subjected to double annealing. Sample R0 was single annealed. After the first annealing, the reference samples R1-R3 were quenched to room temperature without any partitioning prior to the second annealing. The inventive samples S1-S5 were quenched to temperatures below M_S-20° C. but above M_F after the first annealing. Thereafter the inventive samples S1-S5 were partitioned at temperatures above the quenching temperatures prior to the second annealing. The reference sample R0 was quenched to a temperature below M_S-20° C. but above M_F after the first annealing and thereafter partitioned at a temperature above the quenching temperature.

After the second annealing, the samples were either quenched followed by partitioning at higher temperatures than the quenching temperature, or they were quenched and isothermally held after the quenching.

The mechanical properties and details of the retained austenite are displayed in Table 3. The inventive steels (S1-S5) had similar tensile strength in comparison to the reference steels (R1-R3) of same composition. However, the total elongation and the hole expansion ratio of the inventive steels were larger than the reference steels of the same composition. Consequently, the product from multiplying the tensile strength, the total elongation, and the hole expansion ratio, was considerably larger for the inventive samples compared to reference samples. Criteria for the steel, e.g. TS*TE*HER/100 and RA*k_p/x_cγ, were determined based on the findings. From the result it is evident that none of the reference samples met RA*k_p/x_cγ≥200 or TS*TE*HER/100≥11 000 MPa %.

TABLE 3
		YS	TS	UE	TE	HER	TS*TE*HER/	RA	xcγ		RA*kp/xcγ
Alloy	Sample	MPa	MPa	%	%	%	100 MPa %	vol-%	wt-%	kp-	vol-%/wt-%

L1	R0	558	875	15.4	18.9	45	7442	10	0.96	14	146
L1	R1	590	842	16.6	19.3	50	8120	9	0.95	14	133
L1	R2	584	858	14.3	17.0	53	7731	8	0.93	15	129
L2	R3	642	839	16.1	19.9	61	10179	11	0.96	16	183
L1	S1	571	835	15.5	22.8	86	16373	13	0.87	17	254
L1	S2	588	842	16.0	22.1	74	13770	14	0.83	19	320
L1	S3	575	825	16.8	23.1	64	12197	12	0.90	18	240
L1	S4	576	839	16.1	22.6	64	12135	12	0.91	18	237
L2	S5	631	842	16.7	22.9	76	14654	14	0.88	19	302

The microstructure after the second annealing cycle of inventive sample S3 and S5 were examined. The microstructure comprised a matrix having a lamellar morphology and including double tempered martensite, tempered martensite, bainite and a high amount of RA. The amount of RA was determined to 12 respectively 14 vol %. The results are presented in Table 4.

TABLE 4
		Double				Fresh
		tempered	Tempered		Retained	martensite +	Other
Alloy	Sample	martensite	martensite	Bainite	austenite	carbides	phases

L1	S3	52	21	15	12	0	0
L2	S5	49	16	21	14	0	0

The mean aspect ratio (AR) and the mean equivalent circle diameter (ECD) of the retained austenite (RA) were determined for the reference sample R2 and compared to the inventive sample S5.

The results are shown in Table 5.

TABLE 5
			ECD	RA	xcγ		RA*kp/xcγ
Alloy	Sample	AR	μm	vol-%	wt-%	kp-	vol-%/wt-%

L1	R2	1.97	0.59	8	0.93	15	129
L2	S5	2.08	0.53	14	0.88	19	302

The following conclusions can be drawn from the results shown in Table 5.

• • The amount of RA is higher for the inventive sample.
  • The carbon content x_cγ of the RA is lower for the inventive sample.
  • The ECD of the RA is smaller for the inventive sample.
  • The AR of the RA is higher for the inventive sample.
  • Only the inventive sample fulfils the criterion RA*k_p/x_cγ≥200.

Despite a smaller ECD and higher AR, the lower C content of the RA results in a lower but optimal stability of the RA against strain induced martensitic transformation (SIMT) of the inventive samples compared to the reference once. This leads to a better exploitation of the TRIP effect and thus better global formability. Moreover, the lower C content in RA guarantees a lower hardness difference between the matrix and RA, which also improves local formability in the case of the inventive samples.