EP4689201A1 - A high strength steel strip or sheet, and a method for producing the same - Google Patents

A high strength steel strip or sheet, and a method for producing the same

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Publication number
EP4689201A1
EP4689201A1 EP24715481.8A EP24715481A EP4689201A1 EP 4689201 A1 EP4689201 A1 EP 4689201A1 EP 24715481 A EP24715481 A EP 24715481A EP 4689201 A1 EP4689201 A1 EP 4689201A1
Authority
EP
European Patent Office
Prior art keywords
strip
retained austenite
temperature
mpa
steel
Prior art date
Legal status (The legal status is an assumption and is not a legal conclusion. Google has not performed a legal analysis and makes no representation as to the accuracy of the status listed.)
Pending
Application number
EP24715481.8A
Other languages
German (de)
French (fr)
Inventor
Daniel Krizan
Johannes REHRL
Christoph KICKINGER
Current Assignee (The listed assignees may be inaccurate. Google has not performed a legal analysis and makes no representation or warranty as to the accuracy of the list.)
Voestalpine Stahl GmbH
Original Assignee
Voestalpine Stahl GmbH
Priority date (The priority date is an assumption and is not a legal conclusion. Google has not performed a legal analysis and makes no representation as to the accuracy of the date listed.)
Filing date
Publication date
Application filed by Voestalpine Stahl GmbH filed Critical Voestalpine Stahl GmbH
Publication of EP4689201A1 publication Critical patent/EP4689201A1/en
Pending legal-status Critical Current

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Classifications

    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/46—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
    • C21D1/18—Hardening; Quenching with or without subsequent tempering
    • C21D1/19—Hardening; Quenching with or without subsequent tempering by interrupted quenching
    • C21D1/22—Martempering
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/005—Heat treatment of ferrous alloys containing Mn
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/008—Heat treatment of ferrous alloys containing Si
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0221—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • C21D8/0236—Cold rolling
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0247—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0247—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
    • C21D8/0273—Final recrystallisation annealing
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/04—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for drawing, e.g. for deep-drawing
    • C21D8/0421—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for drawing, e.g. for deep-drawing characterised by the working steps
    • C21D8/0436—Cold rolling
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/04—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for drawing, e.g. for deep-drawing
    • C21D8/0447—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for drawing, e.g. for deep-drawing characterised by the heat treatment
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/12—Ferrous alloys, e.g. steel alloys containing tungsten, tantalum, molybdenum, vanadium, or niobium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/14—Ferrous alloys, e.g. steel alloys containing titanium or zirconium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/26—Ferrous alloys, e.g. steel alloys containing chromium with niobium or tantalum
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/38—Ferrous alloys, e.g. steel alloys containing chromium with more than 1.5% by weight of manganese
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/001—Austenite
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/002—Bainite
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/008—Martensite

Definitions

  • the present invention relates to high strength steel strip or sheets suitable for applications in automobiles.
  • 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.
  • Automotive body parts are often stamped out of sheet steels, forming complex structural members of thin sheet.
  • such parts cannot be produced from conventional high strength steels, because of a too low formability for complex structural parts.
  • multiphase Transformation Induced Plasticity aided steels TRIP 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.
  • 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 is an annealing cycle which consists of a quenching and a partitioning step.
  • the quenching step fully austenitized or intercritically annealed steels are quenched to a temperature between the martensite start temperature Ms and the martensite finish temperature MF 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 Al 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 Al 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.
  • 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.
  • 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 Ms-20 and MF 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.
  • Fig. 1 shows schematically the annealing cycles of the invention.
  • the strip or sheet has a composition consisting of the following alloying elements (in wt. %):
  • composition is excluding any coatings applied to the strip or sheet.
  • 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.
  • 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 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.
  • 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.
  • Mn lowers the A C 3 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.
  • 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.
  • 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 C 3 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 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.
  • 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.
  • Ti is a strong nitride former and can be used to bind the nitrogen in the steel.
  • 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 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 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 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 C 3 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 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.
  • 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.
  • 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 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 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 %.
  • 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:
  • Oxygen and hydrogen can further be limited to
  • N ⁇ 0.015 %, preferably 0.001 - 0.008 %
  • the steel should fulfil the following condition on the tensile strength: TS, Tensile Strength (R m ) 780 - 980 MPa, preferably 800 - 900 MPa
  • TS*TE*HER/100 11 000 MPa%, preferably 12 000 - 20 000 MPa%
  • 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 , Rpo.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 (X) 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 x 100 mm 2 . 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 (X) in % is calculated as follows:
  • X (Dh - Do)/Do x 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.
  • microstructural constituents are in the following expressed in volume % (vol. %).
  • 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 of the thickness. The surface can be etched to make the phases easier to identify.
  • SEM scanning electron microscope
  • Electron Backscatter Diffraction 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 (bet) 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.
  • 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 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 kp-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.
  • 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 7 o) can be measured in the final heat treated sample.
  • the sample for the retained austenite after deformation (V y ) 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 C y, in the retained austenite may optionally fulfil (in wt.%): Xcy 0.7 -1.0
  • the carbon content in the retained austenite can be calculated from the lattice parameter a, obtained from the y 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.
  • the amount of retained austenite (RA vol.% ), the mechanical stability of the austenite (k p ), and the carbon content (x C y) of the retained austenite should fulfil the following relation:
  • the mean equivalent circle diameter (ECD) in pm 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
  • ECD Equivalent circle diameter
  • AR aspect ratio
  • RA retained austenite
  • ECD 2"V(A/TI)
  • A 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.
  • Ael and Ae3 represent the equilibrium transformation temperatures. Austenite is completely stable above Ae3 and partially unstable between Ae3 and Ael. Ael and Ae3 are calculated by means of ThermoCalc 2022 TCFE 12.
  • Ms temperature was calculated using the Ms formula found in S. Kaar, K. Steineder, R. Schneider, D. Krizan and C. Sommitsch: carefulNew Ms-formula for exact microstructural prediction of modern 3rd generation AHSS chemistries”, Scr. Mater., Vol. 200, 2021, 113923.
  • Ms 692 - 502 * (C + 0.86N) A 0.5 - 37 * Mn - 14 * Si + 20 * Al - 11 * Cr
  • the Ms formula uses the content in weight % of each element.
  • MF 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 M s + [ln(l- f M )]/0.011, where the fraction of martensite fw is 0.95.
  • 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.
  • annealing at a temperature in the range of 500 - 950 °C.
  • 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.
  • 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 Ms - 20 °C, and MF 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 upper limit may further be restricted to Ms -30, Ms -40, Ms -50, Ms -60, Ms -70, Ms -80, Ms -90, or Ms -100 °C. j) heating the cooled strip at a rate (HR2) of 1-100 °C/s to a first partitioning temperature (Toal) in the range of the first quenching temperature (TQ1) + 10 °C to 500 °C and partitioning the strip for 20-1000 s (toal).
  • the first partitioning temperature (Toal) is above Ms.
  • 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.
  • 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.
  • SIMT strain induced martensitic transformation
  • 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).
  • the second annealing temperature (Tan2) is in the intercritical range, i.e. between Ael and Ae3.
  • the second annealing temperature (Tan2) may be allowed to be above Ae3.
  • the second quenching temperature (TQ2) may be below Ms, preferably below (Ms-20) °C, and must be above MF.
  • 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 (Toal) 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 (Toal), 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.
  • 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.
  • PVD Physical Vapor Deposition
  • Hot Dip Galvanizing 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 %:
  • a galvannealed coating may contain 5-20 wt.% of diffused Fe.
  • Other coating composition known in the art can be applied.
  • the alloys LI 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.
  • sample RO 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.
  • inventive samples S1-S5 were quenched to temperatures below Ms-20 °C but above MF 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 RO was quenched to a temperature below Ms- 20 °C but above MF after the first annealing and thereafter partitioned at a temperature above the quenching temperature.
  • 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 C y, were determined based on the findings. From the result it is evident that none of the reference samples met RA*kp/ x > 200 or TS*TE*HER/100> 11 000 MPa%. Table 3 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.
  • 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 amount of RA is higher for the inventive sample.
  • the carbon content 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.
  • 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.
  • 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.

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Abstract

The invention relates to a high strength steel strip or sheet comprising in wt. %: 0.1 – 0.2 C, 0.3 – 0.9 Si, 1.4 – 3.0 Mn, 0.03 –1.0 Al, 0.5 – 1.5 Si + Al. The strip or sheet having a tensile strength (Rm) of 780 – 980 MPa, a microstructure including 5-25 vol% retained austenite. The retained austenite fulfil: a mechanical stability (kp) of 10 – 35, a carbon content xcγ of the retained austenite of 0.7 –1.0, and RA * kp / xcγ ≥ 200. The invention also relates to a method including double annealing cycle for producing the strip or sheet.

Description

A HIGH STRENGTH STEEL STRIP OR SHEET, AND A METHOD FOR PRODUCING THE
SAME
TECHNICAL FIELD
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 ART
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 Ms and the martensite finish temperature MF 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 Al 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 Al 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 Al, US 2020/392610 Al, US 2020/00440421 Al, US2019/203316 Al, JP 2004238679 A2, WO 2022/123289 Al, US 2020/354823 Al, US2021/207236 Al 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.
DISCLOSURE 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 Ms-20 and MF 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 DECRIPTION OF THE DRAWING
Fig. 1 shows schematically the annealing cycles of the invention. DETAILED DESCRIPTION
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: 0.10 - 0.20 %
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 - 3.0 %
Manganese is a solid solution strengthening element, which stabilises the austenite by lowering the Ms temperature and prevents ferrite and pearlite to be formed during cooling. In addition, Mn lowers the AC3 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 - 1.0 %
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 Ms temperature is also increased with increasing Al content. A further drawback of Al is that it results in an increase of the AC3 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 AC3 temperature and the Ms 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.
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: < 0.020, or < 0.010%
Zr: < 0.010, or < 0.006%
Sn: < 0.030, or < 0.015%
Oxygen and hydrogen can further be limited to
O: < 0.001, or < 0.0003 %
H: < 0.0050, or < 0.0020 %
It is also preferred to control the nitrogen content to the range:
N: < 0.015 %, preferably 0.001 - 0.008 %
Mechanical
The steel should fulfil the following condition on the tensile strength: TS, Tensile Strength (Rm) 780 - 980 MPa, preferably 800 - 900 MPa
And optionally one or more of the following mechanical properties:
YS, Yield Strength (Rpo.2) 400 - 800 MPa, preferably 500 - 700 MPa
HER, Hole Expansion
Ratio (X) > 40 %, preferably 50 - 90 %
TE, Total
Elongation (Ago) > 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 (Rm) can be restricted to 970, 960, 950, 940, 930, 910, 900, 890, 880, 870 or 860 MPa.
The Rm, Rpo.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 (X) 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 x 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 (X) in % is calculated as follows:
X = (Dh - Do)/Do x 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 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 (bet) 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 (kp) is a parameter that describes the mechanical stability of the retained austenite (RA). Factors influencing the kp 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 (kp) in the range of 10 - 35, preferably 15 - 25. A kp- value above 35 indicates lower stability of retained austenite (RA) against mechanical loading, If the kp-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 kp-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 (kp) 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:
V7o ■ • .initial retained austenite content
V7 . . .retained austenite content after deformation
8 . . . true strain p . . . constant related to autocatalytic effect kp . . . 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 kp-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 (V7o) can be measured in the final heat treated sample. The sample for the retained austenite after deformation (Vy) 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, xCy, in the retained austenite may optionally fulfil (in wt.%): Xcy 0.7 -1.0
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 a, obtained from the y 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: a7= 3.556+0.0453 xcy + 0.00095 XMn+ 0.0056 XAI where xCy, XMn, and XAI are the content of carbon, manganese, and aluminium in austenite (in wt.%). The contents for x\iri and XAI 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 (kp), and the carbon content (xCy) of the retained austenite should fulfil the following relation:
RA* kp / xc > 200, preferably 220 - 400
The mean equivalent circle diameter (ECD) in pm 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 pm and aspect ratio (AR) of the retained austenite (RA) can be determined by means of Electron backscatter diffraction (EBSD). ECD =2"V(A/TI), 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.
Ael and Ae3 represent the equilibrium transformation temperatures. Austenite is completely stable above Ae3 and partially unstable between Ae3 and Ael. Ael and Ae3 are calculated by means of ThermoCalc 2022 TCFE 12.
Ms temperature was calculated using the Ms formula found in S. Kaar, K. Steineder, R. Schneider, D. Krizan and C. Sommitsch: „New Ms-formula for exact microstructural prediction of modern 3rd generation AHSS chemistries”, Scr. Mater., Vol. 200, 2021, 113923.
Ms = 692 - 502 * (C + 0.86N)A0.5 - 37 * Mn - 14 * Si + 20 * Al - 11 * Cr The Ms formula uses the content in weight % of each element.
MF 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):
MF = Ms + [ln(l- fM)]/0.011, where the fraction of martensite fw 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 Ms - 20 °C, and MF 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 (Tanl) between 800 °C and 1000 °C, and soaking for 10 - 300 s (tanl), preferably the first annealing temperature (Tanl) is above Ae3. i) cooling the strip at a rate (CR1) of 10 - 100 °C/s to a first quenching temperature (TQ1) between MF and (Ms-20) °C. The lower limit may further be restricted to the highest of MF and value chosen from 150, 160, 170, and 180 °C. The upper limit may further be restricted to Ms -30, Ms -40, Ms -50, Ms -60, Ms -70, Ms -80, Ms -90, or Ms -100 °C. j) heating the cooled strip at a rate (HR2) of 1-100 °C/s to a first partitioning temperature (Toal) in the range of the first quenching temperature (TQ1) + 10 °C to 500 °C and partitioning the strip for 20-1000 s (toal). Preferably the first partitioning temperature (Toal) is above Ms. 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 Ael 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 Ms, preferably below (Ms-20) °C, and must be above MF. 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 (Toal) 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 (Toal), 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 LI 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 Ms, MF, Ael 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
The alloys LI 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
Table 2b
All samples, except for sample RO, were subjected to double annealing. Sample RO 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 Ms-20 °C but above MF 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 RO was quenched to a temperature below Ms- 20 °C but above MF 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*kp/ xCy, were determined based on the findings. From the result it is evident that none of the reference samples met RA*kp/ x > 200 or TS*TE*HER/100> 11 000 MPa%. Table 3 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
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
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 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 * kp / > 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.

Claims

1. A high strength steel strip or sheet having: a) a composition consisting of the following 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; b) a tensile strength (Rm) 780 - 980 MPa wherein Rm is derived according to the European norm EN 10002 Part 1, wherein the samples are taken in the longitudinal direction of the strip; c) a microstructure comprising of in vol. %: retained austenite 5 - 25 wherein the amount of retained austenite is measured 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, and the other phases are determined by use of scanning electron microscope (SEM) at 2000 times magnification; d) wherein the retained austenite fulfil: a mechanical stability (kp) 10 - 35, kp as defined by Ludwigson and Berger in J. Iron Steel Inst. 1969, vol. 207, 10 pp. 63, where p=l: V7o ■ • .initial retained austenite content
V7 . . .retained austenite content after deformation
8 . . . true strain; a carbon content of the retained austenite (in wt.%):
Xc 0.7 -1.0, wherein xCy is determined by the formula: a7= 3.556+0.0453 xcy + 0.00095 +in + 0.0056 XAI where %Mn, and XAI are the content of manganese, and aluminium in austenite (in wt.%) and wherein the contents for XMn and XAI are assumed to be equal the nominal content, and the lattice parameter a, is obtained from the y 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;; and RA * kp / Xcy >200.
2. The steel strip or sheet according to claim 1 , wherein the microstructure fulfils at least one of the following requirements in vol. %, preferably all the requirements: 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.
3. The steel strip or sheet according to claim 1 or 2, wherein the microstructure fulfils at least one of the following requirements in vol. %, preferably all the requirements: double tempered martensite 40 - 60 tempered martensite 10 - 40 bainite 0 - 30 fresh martensite + carbides 0 -5 retained austenite (RA) 10 - 20 other phases 0 - 5
4. The steel strip or sheet according to any one of the preceding claims, wherein one or more of the following mechanical properties are fulfilled:
YS, Yield Strength (Rpo.2) 400 - 800 MPa, preferably 500 - 700 MPa
HER, Hole Expansion
Ratio (X) > 40 %, preferably 50 - 90 %
TE, Total
Elongation (Ago) > 15 %, preferably 20 - 30 %
TS*TE*HER/100 > 11 000 MPa%, preferably 12 000 - 20 000 MPa%
5. The steel strip or sheet according to any one of the preceding claims, wherein one or more of the following mechanical properties are fulfilled:
TS, Tensile Strength (Rm) 800 - 900 MPa YS, Yield Strength (Rpo.2) 500 - 700 MPa HER, Hole Expansion Ratio (X) 50 - 90 %
TE, Total Elongation (Ago) 20 - 30 % TS*TE*HER/100 12 000 - 20 000 MPa%
6. The steel strip or sheet according to any one of the preceding claims, wherein the retained austenite fulfils one or more of the following:
AR > 2.0
ECD < 0.57 mechanical stability (kp) 15 - 25 carbon content in retained austenite (x ) 0.8 - 0.95
RA* kp / Xcy 220 - 400.
7. The high strength steel strip or sheet according to any one of the preceding claims having a composition consisting of the following alloying elements (in wt. %):
C 0.13 - 0.18
Si 0.3 - 0.8 Mn 1.7 - 2.4
Al 0.03 -0.8
Si+ Al 0.6 - 1.3
Optionally
Ti < 0.1
Nb < 0.1
V < 0.1
Cr < 0.5
B < 0.005
Mo < 0.2 balance Fe apart from impurities.
8. A method for producing a high strength steel strip or sheet according to any one of claims 1 -
7, comprising: g) Providing a cold rolled steel having a composition as defined in item a) of the claims: h) heating the strip at a rate (HR1) of to 1 - 20 °C/s to a first annealing temperature (Tanl) between 800 °C and 1000 °C, and soaking for 10 - 300 s (tanl), i) cooling the strip at a rate (CR1) of 10 - 100 °C/s to a first quenching temperature (TQ1) between MF and (Ms-20) °C, j) heating the cooled strip at a rate (HR2) of 1-100 °C/s to a first partitioning temperature (Toal) in the range of the first quenching temperature (TQ1) + 10 °C to 500 °C, and partitioning the strip for 20-1000 s (toal), k) cooling the strip to a temperature below 50 °C at a rate (CR2) of 1 - 50 °C/s, l) heating the strip at a rate (HR3) of to 1 - 20 °C/s to a second annealing temperature (Tan2) between 700 °C and 900 °C and soaking for 10 - 300 s (tan2), 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, 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, o) cooling the strip to a temperature below 50 °C, p) optionally making sheets from the strips.
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