EP4136265A1 - Procédé de fabrication d'une bande d'acier à structure multiphasée et bande d'acier associée - Google Patents

Procédé de fabrication d'une bande d'acier à structure multiphasée et bande d'acier associée

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Publication number
EP4136265A1
EP4136265A1 EP21718871.3A EP21718871A EP4136265A1 EP 4136265 A1 EP4136265 A1 EP 4136265A1 EP 21718871 A EP21718871 A EP 21718871A EP 4136265 A1 EP4136265 A1 EP 4136265A1
Authority
EP
European Patent Office
Prior art keywords
steel strip
steel
annealing
weight
temperature
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
EP21718871.3A
Other languages
German (de)
English (en)
Inventor
Konstantin MOLODOV
Jan ROIK
Ingwer Denks
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.)
Salzgitter Flachstahl GmbH
Original Assignee
Salzgitter Flachstahl 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 Salzgitter Flachstahl GmbH filed Critical Salzgitter Flachstahl GmbH
Publication of EP4136265A1 publication Critical patent/EP4136265A1/fr
Pending legal-status Critical Current

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    • 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
    • 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
    • C21D9/52—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for wires; for strips ; for rods of unlimited length
    • 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
    • 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
    • C21D1/25—Hardening, combined with annealing between 300 degrees Celsius and 600 degrees Celsius, i.e. heat refining ("Vergüten")
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    • 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
    • C21D1/26—Methods of 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
    • C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
    • C21D1/26—Methods of annealing
    • C21D1/28—Normalising
    • 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 by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0205—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips of ferrous alloys
    • 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 by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • C21D8/0226—Hot 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 by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0221—Modifying the physical properties 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 by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0247—Modifying the physical properties 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
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    • 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 by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/04—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for deep-drawing
    • C21D8/0447—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for deep-drawing characterised by the heat treatment
    • C21D8/0473—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
    • 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
    • C21D9/48—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals deep-drawing sheets
    • 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
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/20—Ferrous alloys, e.g. steel alloys containing chromium with copper
    • 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
    • C22C38/22—Ferrous alloys, e.g. steel alloys containing chromium with molybdenum or tungsten
    • 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
    • C22C38/24—Ferrous alloys, e.g. steel alloys containing chromium with vanadium
    • 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
    • C22C38/28—Ferrous alloys, e.g. steel alloys containing chromium with 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
    • C22C38/38—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
    • C21D2211/005—Ferrite
    • 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
    • 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
    • C21D6/001—Heat treatment of ferrous alloys containing Ni
    • 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
    • C21D6/002—Heat treatment of ferrous alloys containing Cr
    • 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
    • C22C38/32—Ferrous alloys, e.g. steel alloys containing chromium with boron

Definitions

  • the invention relates to a method for producing a steel strip with a multiphase structure and a steel strip with a multiphase structure.
  • a steel strip is understood to be a hot-rolled or cold-rolled and annealed steel strip.
  • Usual thicknesses of a hot-rolled steel strip, also referred to as hot strip, are between 2 mm and 8 mm.
  • Cold-rolled, annealed steel strips are referred to as cold strip or thin sheet and usually have thicknesses in the range from 0.5 mm to 2.5 mm.
  • the steel suppliers take the above-mentioned task into account by providing high-strength steels.
  • high-strength steels with a smaller sheet thickness, the weight of the vehicle components can be reduced with the same and possibly even improved component behavior.
  • these newly developed steels must meet the high material requirements in terms of yield strength, tensile strength and elongation at break as well as bake hardening, as well as the high component requirements for toughness, insensitivity to edge cracks, improved bending angle and bending radius, energy absorption and defined hardening via work hardening Have effect.
  • a low yield strength ratio (R e / R m ) of, for example, less than 0.6 combined with a very high tensile strength, strong work hardening and good cold formability is typical for a dual-phase steel and is primarily used for formability in stretching and deep-drawing processes.
  • Dual-phase steels consist of a ferritic basic structure in which a martensitic second phase is embedded. It has been found that in the case of low-carbon, micro-alloyed steels, small proportions of other phases such as bainite and retained austenite have an advantageous effect, for example, on the hole expansion behavior, the bending behavior and the hydrogen-induced brittle fracture behavior. Of the Bainite can be present in different forms, such as upper and lower bainite.
  • a higher yield strength ratio R e / R m is characterized, among other things, by a high one
  • the multiphase structure is characterized by a predominantly ferritic-bainitic basic matrix, with proportions of martensite, tempered martensite, retained austenite and / or pearlite also being present. Delayed recrystallization or precipitation of micro-alloying elements results in strong grain refinement (i.e. fine-grain structure) and thus high strength.
  • these complex or multi-phase steels have higher yield strengths, a higher yield strength or yield strength ratio, less work hardening and a higher hole expansion capacity.
  • Such steels are therefore excellently suited for the production of components with complex geometry, in particular in the case of components subject to crash loads which require a high energy absorption capacity.
  • Multiphase steels are known, for example, from the laid-open specifications DE 102012 002 079 A1 and DE 102015 111 177 A1. With those revealed there
  • the cold-rolled steel strips are usually annealed in the continuous annealing process in a recrystallizing process to form easily deformable sheet for economic reasons.
  • the process parameters such as throughput speed, annealing temperatures and Cooling speed, can be adjusted according to the required mechanical-technological properties with the necessary structure.
  • Grains of different sizes can transform into different phase components when they cool down from the furnace temperature and cause further inhomogeneity.
  • the cold strip is heated in the continuous annealing furnace to a temperature at which the required structure formation (for example dual or complex phase structure) occurs during cooling.
  • the annealing treatment usually takes place in a continuous hot-dip galvanizing plant, in which the heat treatment or annealing and the subsequent galvanizing take place in a continuous process.
  • the required structure is only set during the annealing treatment in the continuous furnace in order to achieve the required mechanical properties.
  • a disadvantage of these multi-phase or complex-phase steels has been found to be that, after the austenitizing annealing of the hot or cold strip in a continuous furnace, a high yield strength ratio can be achieved however, at the expense of a lower elongation at break Aso compared to dual-phase steels. If a high elongation at break Aso is required, a high elongation limit ratio can no longer be set reliably in the process. The reason for this is that during the large-scale continuous annealing process, depending on the alloy concept, the conversion of the austenite back to bainite does not take place completely, since the retained austenite in the holding area is enriched with carbon at temperatures of 200 ° C to 500 ° C and is thereby stabilized.
  • the invention is therefore based on the object of a method for producing a steel strip with a multiphase structure and a steel strip with a
  • the method is intended to compensate for the drop in the yield strength and thus achieve a combination of high yield strength or high yield strength ratio and high elongation at break.
  • a corresponding cold-rolled or hot-rolled steel strip should also be specified.
  • the method for producing a steel strip with a multi-phase structure with the following steps:
  • - First annealing in particular a continuous annealing of the steel strip, in particular the cold-rolled steel strip, at a temperature between 750 ° C up to and including 950 ° C for a total of 10 s to 1200 s, in particular from 50 s to 650 s, and then first cooling the steel strip to a temperature between 200 ° C up to and including 500 ° C with an average cooling rate of 2 K / s to 150 K / s, in particular from 5 K / s to 100 K / s,
  • the yield strength can vary depending on the process parameters variably set on the finish annealing and final cooling, and a high ratio of Rp o, 2-yield strength of the finish annealed steel strip to the tensile strength R m of the finish annealed steel strip are achieved.
  • the steel strip according to the invention has good weldability and has a low tendency towards liquid metal and hydrogen embrittlement.
  • the two process steps “final annealing and final cooling” can be directly related in the course of the process according to the invention for producing a steel strip in terms of time and location or, depending on the circumstances, shifted by hours or days or take place at a different location.
  • the example steels Div, Dv are identical in their alloy composition and are only given a different index for later description.
  • An essential difference between the example steels according to the invention and the reference steels is a lower carbon content, which improves weldability and minimizes the susceptibility to liquid metal and hydrogen embrittlement.
  • the reference steels Ai and B N are not according to the invention because the C content is too high. This results in poorer weldability. In addition, the tensile strength is too low (less than 920 MPa).
  • the reference steels Ai and B N also do not respond as effectively to the treatment according to the invention.
  • the effect of the elements in the steel strip according to the invention with a multi-phase structure is described in more detail below.
  • the multi-phase steels are typically chemically structured in such a way that alloying elements are combined with and without micro-alloying elements.
  • Accompanying elements are unavoidable and are taken into account in the analysis concept with regard to their effect if necessary.
  • Hydrogen (H) is the only element that can diffuse through the iron lattice without creating lattice tension. This means that the hydrogen in the iron lattice is relatively mobile and can be absorbed relatively easily during manufacture. Hydrogen can only be absorbed into the iron lattice in atomic (ionic) form. Hydrogen has a very embrittling effect and diffuses preferentially to energetically favorable locations (defects, grain boundaries, etc.). Defects act as hydrogen traps and can reduce the dwell time of the Significantly increase the amount of hydrogen in the material. Recombination to molecular hydrogen can cause cold cracks. This behavior occurs with hydrogen embrittlement or with hydrogen-induced stress corrosion cracking. Hydrogen is also often cited as the reason for a delayed crack, the so-called delayed fracture, which occurs without external stresses. Therefore, the hydrogen content in the steel should be as low as possible.
  • Oxygen (O) In the molten state, the steel has a relatively high absorption capacity for gases, but at room temperature oxygen is only soluble in very small amounts. Similar to hydrogen, oxygen can only diffuse into the material in atomic form. Because of the highly embrittling effect and the negative effects on aging resistance, attempts are made to reduce the oxygen content as much as possible during production. To reduce the oxygen, there are procedural approaches such as vacuum treatment on the one hand and analytical approaches on the other. By adding certain alloy elements, the oxygen can be converted into less dangerous conditions. A binding of the oxygen via manganese, silicon and / or aluminum is usually common. However, the resulting oxides can cause negative properties as defects in the material. In the case of a fine precipitation, especially of aluminum oxides, on the other hand, a
  • Nitrogen (N) is also an accompanying element in steel production. Steels with free nitrogen tend to have a strong aging effect. The nitrogen diffuses at dislocations even at low temperatures and blocks them. It thus causes an increase in strength combined with a rapid loss of toughness.
  • the nitrogen can be set in the form of nitrides by adding aluminum or titanium to the alloy. For the reasons mentioned above, the optional nitrogen content is limited to ⁇ 0.016% by weight or to amounts that are unavoidable in steel production.
  • sulfur (S) is bound as a trace element in iron ore. It is undesirable in steel (with the exception of free-cutting steels), as it tends to segregate strongly and has a strong embrittling effect. Attempts are therefore made to use the smallest possible amounts of To achieve sulfur in the melt (e.g. by a deep vacuum treatment).
  • the sulfur present is converted into the relatively harmless compound manganese sulfide (MnS) by adding manganese.
  • MnS manganese sulfide
  • the manganese sulfides are often rolled out in lines during the rolling process and act as nucleation sites for the transformation. In the case of diffusion-controlled transformation in particular, this leads to a distinctly lined structure and, if the lined one is very pronounced, can lead to impaired mechanical properties (e.g. pronounced martensite ropes instead of distributed martensite islands, anisotropic material behavior, reduced elongation at break).
  • the sulfur content is limited to ⁇ 0.005% by weight or to amounts that are unavoidable in steel production.
  • Phosphorus (P) is a trace element from iron ore and is dissolved in the iron lattice as a substitution atom. Phosphorus increases hardness through solid solution strengthening and improves hardenability. As a rule, however, an attempt is made to use the
  • the phosphorus content is limited to ⁇ 0.020% or to amounts that are unavoidable in steel production.
  • Alloy elements are usually added to steel in order to specifically influence certain properties.
  • An alloy element can influence different properties in different steels. The relationships are varied and complex. In the following, the effect of the alloying elements will be discussed in more detail.
  • Carbon (C) is considered to be the most important alloying element in steel. It is only through its targeted introduction of up to 2.06% that iron becomes steel. The carbon content is often drastically reduced during steel production.
  • the multiphase steel according to the invention in particular for continuous hot-dip processing, its proportion is 0.085% by weight to 0.149% by weight, preferably up to 0.115% by weight. Due to its comparatively small atomic radius, carbon is dissolved interstitially in the iron lattice.
  • the solubility in a-iron is a maximum of 0.02% and in g-iron a maximum of 2.06%.
  • carbon significantly increases the hardenability of steel. Due to the different solubility, pronounced diffusion processes are necessary during the phase transition, which can lead to very different kinetic conditions.
  • carbon increases the thermodynamic stability of the austenite, which is shown in the phase diagram in an expansion of the austenite area to lower temperatures. As the forced-release carbon content in the martensite rises, the lattice distortions increase and the associated strength of the non-diffusive phase increases. Carbon is also required to form carbides. A representative that occurs in almost every steel is cementite (Fe3C).
  • the minimum C content is therefore set at 0.085% by weight and the maximum C content at 0.149% by weight, preferably 0.115% by weight.
  • Aluminum (AI) is usually added to steel in order to bind the oxygen and nitrogen dissolved in the iron.
  • the oxygen and nitrogen are converted into aluminum oxides and aluminum nitrides.
  • Aluminum nitride is not precipitated when titanium is present in sufficient quantities. Titanium nitrides have a lower enthalpy of formation and are formed at higher temperatures. In a dissolved state, aluminum, like silicon, shifts the formation of ferrite to shorter times and thus enables sufficient ferrite to be formed. It suppresses also the formation of carbide and thus leads to a delayed transformation of the austenite.
  • Al is also used as an alloying element in retained austenitic steels in order to replace part of the silicon with aluminum.
  • the reason for this approach is that Al is somewhat less critical for the galvanizing reaction than Si.
  • the Al content is therefore limited to 0.005% by weight up to a maximum of 0.1% by weight.
  • Silicon (Si) binds oxygen during casting and thus reduces segregation and impurities in the steel.
  • silicon increases the strength and the yield strength ratio of the ferrite with only a slight decrease in elongation at break. Another important effect is that silicon shifts the formation of ferrite to shorter times, thus allowing sufficient ferrite to form before the quenching.
  • the formation of ferrite enriches the austenite with carbon and stabilizes it. At higher contents, silicon stabilizes in the lower temperature range, especially in the area of
  • Bainite formation by preventing carbide formation significantly increases the austenite.
  • silicon content is high, strongly adhering scale can form, which can impair further processing.
  • silicon can diffuse to the surface during annealing and, alone or together with manganese, form film-like oxides. These oxides worsen the galvanization by impairing the galvanizing reaction (iron solution and inhibition layer formation) when the steel strip is immersed in the zinc melt. This manifests itself in poor zinc adhesion and non-galvanized areas.
  • a suitable furnace operation with an adapted moisture content in the annealing gas and / or a low Si / Mn ratio and / or the use of moderate amounts of silicon can, however, ensure good galvanization of the steel strip and good zinc adhesion.
  • the minimum Si content is set to 0.200% by weight and the maximum Si content is set to 0.750% by weight.
  • Manganese (Mn) is added to almost all steels for desulfurization in order to convert the harmful sulfur into manganese sulfides. Manganese also increases the strength of the ferrite through solid solution strengthening and shifts the conversion to lower temperatures.
  • One of the main reasons for adding manganese to the alloy is the significant improvement in hardenability. Because of the diffusion obstruction becomes the pearlite and bainite transformation postponed for longer times and the martensite start temperature lowered. Like silicon, manganese tends to form oxides on the steel surface during the annealing treatment. Depending on the annealing parameters and the content of other alloy elements (in particular Si and Al), manganese oxides (e.g. MnO) and / or Mn mixed oxides (e.g.
  • Mn2Si04 manganese with a low Si / Mn or Al / Mn ratio is to be regarded as less critical, since globular oxides rather than oxide films are formed. Nevertheless, high manganese contents can negatively affect the appearance of the zinc layer and the zinc adhesion.
  • the Mn content is therefore set to 1.6% by weight to 2.9% by weight, preferably up to 2.6% by weight.
  • Chromium (Cr) The addition of chromium mainly improves hardenability. In the dissolved state, chromium shifts the pearlite and bainite transformation for longer times and at the same time lowers the martensite start temperature. Another important effect is that chromium increases the tempering resistance considerably, so that there is almost no loss of strength in the zinc bath. Chromium is also a carbide former. If chromium is in carbide form, the austenitizing temperature must be high enough before hardening to dissolve the chromium carbides. Otherwise, the increased number of germs can lead to a deterioration in hardenability. Chromium also tends to form oxides on the steel surface during the annealing treatment, which can deteriorate the galvanizing quality. The optional Cr content is therefore set to values from 0.05 to 0.500% by weight.
  • Molybdenum Molybdenum is added in a similar way to chromium to improve hardenability. The pearlite and bainite transformation is pushed to longer times and the martensite start temperature is lowered. Molybdenum also increases the tempering resistance considerably, so that no loss of strength is to be expected in the zinc bath and, through solid solution strengthening, increases the strength of the ferrite.
  • the Mo content is added depending on the dimensions, the system configuration and the structure setting. For cost reasons, the optional Mo content is set at 0.05 to 0.5% by weight.
  • Copper The addition of copper can increase the tensile strength and hardenability. In conjunction with nickel, chromium and phosphorus, copper can be a Form a protective oxide layer on the surface, which can significantly reduce the corrosion rate. In connection with oxygen, copper can form harmful oxides at the grain boundaries, which can have negative effects, especially for hot forming processes. The optional copper content is therefore limited to 0.01 to 0.3% by weight.
  • the optional nickel content is therefore limited to 0.01 to 0.050% by weight.
  • Micro-alloy elements are usually only added in very small amounts ( ⁇ 0.1%). In contrast to the alloying elements, they mainly work through the formation of precipitates, but can also influence the properties in a dissolved state. Despite the small additions, micro-alloy elements have a strong influence on the manufacturing conditions as well as the processing and final properties. As a rule, soluble carbide and nitride formers are used as micro-alloying elements in the iron lattice. The formation of carbonitrides is also possible due to the complete solubility of nitrides and carbides in one another. The tendency to form oxides and sulphides is usually most pronounced with the micro-alloying elements, but is usually prevented in a targeted manner due to other alloying elements.
  • micro-alloying elements are aluminum, vanadium, titanium, niobium and boron. These elements can be dissolved in the iron lattice and form carbides and nitrides with carbon and nitrogen.
  • Titanium (Ti) forms very stable nitrides (TiN) and sulfides (TiS2) even at high temperatures. Depending on the nitrogen content, some of these only dissolve in the melt. If the precipitates created in this way are not removed with the slag, they form coarse particles in the material due to the high temperature at which they are formed, which are usually not beneficial for the mechanical properties are. The binding of free nitrogen and oxygen has a positive effect on toughness. Titanium protects other dissolved micro-alloy elements such as niobium from being set by nitrogen. These can then develop their effect optimally. Nitrides, which are only formed at lower temperatures due to the drop in oxygen and nitrogen content, can also effectively hinder austenite grain growth.
  • Titanium that has not set forms titanium carbides at temperatures above 1150 ° C and can thus cause grain refinement (inhibition of austenite grain growth, grain refinement through delayed recrystallization and / or increase in the number of nuclei with a / Y conversion) and precipitation hardening.
  • the optional Ti content therefore has values from 0.005 to 0.060% by weight.
  • Niobium (Nb) causes a strong grain refinement, since it is the most effective of all micro-alloy elements to delay recrystallization and also to inhibit austenite grain growth.
  • the strength-increasing effect is qualitatively higher than that of titanium, evident from the increased grain refinement effect and the larger amount of strength-increasing particles (binding of titanium to TiN at high temperatures).
  • Niobium carbides are formed at temperatures below 1200 ° C. When nitrogen is bonded with titanium, niobium can increase its strength-increasing effect through the formation of small carbides that are effective in terms of their effect in the lower temperature range (smaller carbide sizes).
  • niobium Another effect of niobium is the retardation of the a / g conversion and the lowering of the martensite start temperature in the dissolved state. On the one hand, this is done through the solute drag effect and, on the other hand, through the grain refinement. This causes an increase in the strength of the structure and thus also a higher resistance to the increase in volume during martensite formation.
  • the addition of niobium to the alloy is limited until its solubility limit is reached. This limits the amount of excretions, but if it is exceeded, it mainly causes an early excretion with very coarse particles. Precipitation hardening can thus be particularly effective in steels with a low carbon content (greater supersaturation possible) and in hot forming processes (deformation-induced precipitation).
  • Vanadium (V) The carbide and nitride formation of vanadium is only just beginning Temperatures around 1000 ° C or even after the a / g conversion, i.e. much later than with titanium and niobium. Vanadium has hardly any grain-refining effect due to the small number of precipitates present in the austenite. Austenite grain growth is also not inhibited by the late precipitation of the vanadium carbides. Thus, the strength-increasing effect is based almost entirely on the
  • V content is therefore limited to values from 0.001 to 0.060% by weight.
  • Boron (B) forms nitrides or carbides with nitrogen as well as with carbon; as a rule, however, this is not the aim.
  • B Boron
  • An increase in hardness on the surface is not achieved (with the exception of boriding with the formation of FeB and Fe2B in the edge zone of a workpiece).
  • attempts are usually made to bind the nitrogen with more affine elements. Titanium in particular can ensure that all of the nitrogen is bound.
  • boron in very small amounts leads to a significant improvement in hardenability.
  • boron The working mechanism of boron can be described in such a way that boron atoms, if the temperature is controlled appropriately, attach themselves to the grain boundaries and there, by lowering the grain boundary energy, make the formation of viable ferrite nuclei significantly more difficult.
  • it When controlling the temperature, it must be ensured that boron is predominantly distributed atomically in the grain boundary and is not present in the form of precipitates due to excessively high temperatures.
  • the effectiveness of boron is reduced with increasing grain size and increasing carbon content (> 0.8%).
  • An amount above 60 ppm also causes a decrease in hardenability, since boron carbides act as nuclei on the grain boundaries.
  • boron diffuses extremely well and has a very high affinity for oxygen, which can lead to a reduction in the boron content in areas near the surface (up to 0.5 mm).
  • annealing at temperatures over 1000 ° C is not recommended. This is also recommended because boron can lead to a strong formation of coarse grains at annealing temperatures above 1000 ° C.
  • Boron is an extremely critical element for the process of continuous hot-dip galvanizing with zinc, as it can be used alone or together with manganese during the Annealing treatment can form film-like oxides on the steel surface. These oxides passivate the strip surface and prevent the galvanizing reaction (iron dissolution and inhibition layer formation).
  • Whether film-like oxides are formed depends both on the amount of free boron and manganese and on the annealing parameters used (e.g. moisture content in the annealing gas, annealing temperature, annealing time). Higher manganese contents and long annealing times tend to lead to more globular and less critical oxides. An increased moisture content in the annealing gas also makes it possible to reduce the amount of boron-containing oxides on the steel surface. For the reasons mentioned above, the B content is limited to values from 0.0001 to 0.0060% by weight.
  • the restoration according to the invention of the R p o, 2-proof strength of the steel strip through the annealing and the final cooling is carried out by utilizing one or more of the following conditions:
  • the Hollomon-Jaffe parameter contains the natural logarithm ln (x).
  • the maximum temperature TH used is the highest temperature that is reached on the surface of the steel strip during the final annealing.
  • This maximum temperature TH is decisive for the Hp value according to the invention and the effect of increasing the elongation limit or the occurring metal-physical processes. Therefore, lower temperatures are neglected during the heating phase of the final annealing.
  • the total duration is defined as the duration of the final annealing. The final cooling is therefore not taken into account in the total duration. In the event that the final annealing takes place in a furnace, the total duration begins with an entry into the furnace and ends with an exit from the furnace. In a known manner, the final annealing can alternatively also take place inductively or conductively.
  • the Hollomon-Jaffe parameter Hp as a process parameter in addition to the temperature and the total duration, represents a further process condition for the final annealing.
  • Hp restricts the possible combinations of maximum temperature TH and total duration J such that 12 ⁇ 10 3 >Hp> 7.5 x 10 3 , preferably 10.5 x 10 3 >
  • the steel strip is finally annealed in such a way that the finally annealed and finally cooled steel strip has a value of the tensile strength R m of the steel strip after the final cooling which has increased compared to a value of the tensile strength R m of the steel strip before the final annealing and / or the finally annealed and finally cooled Steel strip has a value of the tensile strength R m of the steel strip after final cooling, which is retained compared to a value of the tensile strength R m of the steel strip before the final annealing in the sense of not less than before the final annealing.
  • the finally annealed and finally cooled steel strip advantageously has a
  • the steel strip is finally annealed at a maximum temperature of above 200 ° C. and / or at a maximum temperature of up to 400 ° C. and / or for a total duration of 10 s to 500 s.
  • an intermediate annealing in particular a continuous annealing, at a temperature between 200 ° C up to and including 500 ° C for a total of 10 s is subjected to 430 s.
  • the steel strip is cooled to a subcooling temperature below 50 ° C. and optionally down to room temperature.
  • the steel strip is cooled down after the first annealing and before the first cooling to an intermediate temperature greater than 600 ° C.
  • the steel strip is cooled at an average cooling rate of 0.1 K / s to 30 K / s over a period of 5 s to 300 s.
  • the steel strip is finally annealed in several stages (for example in several successive furnaces). If the final annealing is carried out in stages, TH, T and the Hp value must be calculated as follows:
  • the total duration of the n-stage annealing is calculated as:
  • the hot or cold rolled steel strip is produced from the aforementioned steel but with a C content of 0.085 to 0.115% by weight and / or from the aforementioned steel but with an Mn content of 1.6 to 2 , 6% by weight is produced.
  • the steel strip before the final annealing with a rolling force F [N]> (0.5 x ⁇ ), where ⁇ is the width of the steel strip in mm, with a maximum degree of rolling of 1.5%.
  • a steel strip with a multiphase structure consisting of the following elements in% by weight: C: from 0.085 to 0.149; AI: from 0.005 to 0.1; Si: from 0.2 to 0.75; Mn: from 1.6 to 2.9;
  • P ⁇ 0.02; S: ⁇ 0.005; and optionally from one or more of the following elements in% by weight: Cr: 0.05 to 0.5; Mo: 0.05 to 0.5; Ti: 0.005 to 0.060; Nb: 0.005 to 0.060; V: 0.001 to 0.060; B: 0.0001 to 0.0060; N: 0.0001 to 0.016; Ni: 0.01 to 0.5; Cu: 0.01 to 0.3; Remainder iron including usual steel-accompanying elements, characterized in that the steel band of greater than 5600 MPa%, in particular of greater than 7200 MPa%, having a product of R p o, 2- yield strength and elongation at rupture A80.
  • This steel strip is advantageously manufactured using the manufacturing method described above.
  • the steel is alloyed with Cr and Mo, where Mn + Cr + 4 ⁇ Mo> 2.5% by weight and 0.1% by weight ⁇ Mo ⁇ 0.5% by weight apply.
  • the steel strip particularly preferably has a minimum tensile strength of 920 MPa, in particular 980 MPa and / or a bake hardening value BH2 of> 25 MPa and / or a residual austenite content of less than 10%, in particular less than 5%.
  • the steel band m, a ratio of R p o, 2- yield strength of the finish annealed steel strip and endabgekühlten to the tensile strength R of the finish annealed steel strip and endabgekühlten comprising greater than 0.68 up to and including 0.97.
  • the structure of the finally annealed and finally cooled steel strip advantageously has the following composition: Ferrite: less than 60%; Bainite + martensite: 30% to 98%; Retained austenite: less than 10%, in particular less than 5%.
  • the percentages given for the structural components relate to parts of the area that are usually also taken over as proportions by volume.
  • At least 1% fresh martensite is present in the structure of the steel strip before the final annealing.
  • the presence of fresh martensite makes the present invention particularly effective, since the fresh martensite ensures a reduction in the yield strength, which is compensated for by the heat treatment according to the invention.
  • the more fresh martensite is present the more this beneficial effect of the heat treatment increases.
  • the structure of the finally annealed and finally cooled steel strip is advantageously characterized in that the structure has a KGs characteristic value of less than 0.4, in particular less than 0.3.
  • room temperature is understood to mean a temperature between 10 ° C to 40 ° C, preferably 15 ° C to 25 ° C.
  • the method for producing a high-strength steel strip according to the invention with a multiphase structure takes place from a cold or hot rolled steel strip of different thickness via a continuous annealing system or optionally via a hot-dip galvanizing system.
  • a first annealing process the cold or hot rolled strip is continuously annealed at a temperature between 750 ° C and 950 ° C for a total of 10 s to 1200 s in order to set the desired degree of austenitization.
  • a phase portion of recovered and / or recrystallized ferrite is retained.
  • the tendency to recovery and / or recrystallization can be controlled by the optional elements such as Mo, Ni, Ti and V, with higher contents of these elements leading to delayed recrystallization kinetics.
  • intermediate annealing follows in this temperature range between 200 ° C to 500 ° C inclusive for a total of 10 s up to 430 s with the aim of converting austenite into bainite. Hot-dip finishing can be carried out as an option.
  • Mn, Mo, Cr, Ni, Nb and B in particular can be added.
  • the austenite is not completely transformed, since the retained austenite is enriched with carbon and thus stabilized.
  • the remaining austenite can only be converted to martensite by cooling to a subcooling temperature of less than 100 ° C., preferably less than 50 ° C. with an average cooling rate of 1 K / s to 50 K / s.
  • glissile dislocations are generated in the surrounding structure, which is expressed from a technological point of view in a lowering of the R p o, 2-yield strength.
  • a heat treatment is necessary after cooling below 100.degree. C., preferably below 50.degree.
  • the tetragonism of the martensitic tetragonal body-centered phase is broken down, as carbon diffuses into the surrounding structural areas and glissile (slidable) dislocations become sessile (immobile) dislocations as a result of Cotrell clouds.
  • the final structure of the multiphase steel according to the invention is composed of ⁇ 60% ferrite, 30 to 98% bainite and martensite (fresh or tempered before the final annealing and tempered after the final annealing), with at least 1% fresh martensite present before the final annealing, as well as a low content of retained austenite less than 10%, preferably less than 5%.
  • the individual annealing treatments can be designed in several stages or additional annealing treatments can be provided in relation to the overall process.
  • Tables 2a and 2b the relevant process parameters of the continuous annealing are listed for an exemplary selection of temperature cycles la to VII of the continuous annealing, which are used for the production of the steel strip according to the invention. The following process parameters are listed in Tables 2a and 2b:
  • T IA maximum annealing temperature in the intercritical area (first annealing)
  • ti A duration of annealing (first annealing)
  • THD temperature hot-dip coating (intermediate annealing)
  • the final annealing is described as the last step of the continuous annealing by the Hollomon-Jaffe parameter Hp described above.
  • Laboratory tests and large-scale tests with the temperature cycles specified in Tables 2a and 2b were carried out and the steel strip produced was then characterized with regard to mechanical-technological parameters.
  • Laboratory tests each relate to the last step of the final annealing after To has been reached and were simulated on a previously large-scale steel strip in a continuous annealing on a laboratory scale in order to determine the dependence of the final properties on the Hp value.
  • table 3 divided into tables 3a and 3b - mechanical characteristics in the longitudinal direction (rolling direction) of the reference steels Ai and B and inventive example steels Cm, Div, Dv, Evi, Fvn and Gvm before and after the final annealing as well as the relative change of the R p o, 2-proof stress specified value Hp by final annealing at an appropriate.
  • the following mechanical parameters are listed in Tables 3a and 3b:
  • AR m change in tensile strength due to final annealing
  • D Rpo , 2 / R P O , 2 ° Relative increase in the yield strength through final annealing
  • the reference steels and steels of the invention possess example before the final annealing a comparable R p o, 2-proof stress (R p o, 2 °).
  • R p o, 2 ° a comparable R p o, 2-proof stress
  • RR P o, 2 f / R m f according to the invention can be achieved as steel of 0.93 (see, for example, temperature cycle purple).
  • the example steel according to the invention retains a high elongation at break of>
  • the tensile strength of the steels according to the invention also increases as a result of the final annealing, so that a final tensile strength R m f of> 920 MPa is achieved, which is significantly higher than the tensile strength R m f of the reference steels Ai and BN.
  • the example steels Cm, Div and Dv treated according to the temperature cycles Ulf, IVe, Vg, Vh have not been assessed as inventive, since the Hp value is less than or equal to 7.5 and the increase in the yield point is less than 5% .
  • 1 is a achieved by the inventive annealing relative increase the R p o
  • 2-proof strength of the steel sheet is represented as a function of the Jaffe Hollomon parameter Hp by diagram.
  • a ratio DR p o , 2 / R p o , 2 ° of the change in the R p o , 2 yield point (DR p o , 2) of the steel strip due to the final annealing to the R p is shown in an x / y diagram o
  • 2-proof stress of the steel strip before final annealing R P o , 2 °
  • the Hollomon-Jaffe parameter Hp TH (ln ([) + 20) [10 3 ] with TH in K (maximum final annealing temperature after cooling to the subcooling temperature To) and in h with values from 6 to 11 [10 3 ] on the x-axis.
  • the Hollomon-Jaffe parameter Hp can be used to characterize the conditions of the final glow over the final glow duration and the maximum final glow temperature TH.
  • five curves are drawn for the reference steel Ai with the temperature cycle group la-f, for the reference steel Bn with the temperature cycle group lla-e, for the example steel according to the invention with the temperature cycle group II la-f and the invention Example steels Div and Dv with temperature cycle groups IVa-e and Va-h.
  • the curves were determined using measurement data from the experiments (see Table 3) by an adapted Johnson-Mehl-Avrami-Kolmogorow equation (see, for example, A. Kolmogoroff; Izv. Akad. Nauk SSSR Ser. Mat.
  • a significantly higher increase in the yield strength compared to the reference steels Ai and Bn can be seen in the example steels Cm, Div and Dv according to the invention.
  • the increase is of the R p o, 2-yield strength for steel Cm, processed through temperature cycle IIIa-f, for steel Div, processed through temperature cycle IVa-e, and for steel Dv, processed over temperature cycle Va-h already over 20%, while the reference steels Ai and B N are ⁇ 10%.
  • the example steels according to the invention therefore show a significant increase in the yield strength even at lower Hp values, which is due to their composition, in particular the increased Si content, whereby cementite precipitations are avoided and the carbon necessary for increasing the yield strength remains dissolved.
  • the C content of reference steels A and B is significantly higher, the increase in the yield strength is significantly lower compared with the steels according to the invention.
  • Table 4 lists the structural components for the steels Ai-Gvm.
  • the structural components were determined in the longitudinal section perpendicular to the roll surface on the basis of measurements by means of electron backscatter diffraction with the aid of the Kikuchi strip contrast and light-optical recordings.
  • the particle diameters were determined from the measurements by means of electron backscatter diffraction, where a grain is defined by the fact that it has a grain boundary with a disorientation angle of> 15 ° (so-called large-angle grain boundary - GWKG, see G. Gottstein, Physical Basics of Material Science, Springer-Verlag Berlin Heidelberg, 2007).
  • the structure of the steels according to the invention Cm-Gvm is composed of ⁇ 60% ferrite, 30 to 98% bainite and martensite (fresh or tempered martensite before the final annealing and tempered martensite after the final annealing), with at least 1% fresh martensite before the final annealing , as well as a content of retained austenite ⁇ 10%, in particular ⁇ 5%.
  • the structure of the steels according to the invention Cm to Gvm has a KGs characteristic value ⁇ 0.4, preferably ⁇ 0.3.
  • fresh martensite Due to its formation mechanism, fresh martensite has a high dislocation density and high hardness. In the case of electron backscatter diffraction, such areas appear darker in the Kikuchi band contrast than other structural components, since the diffraction condition is violated by a disturbed crystal lattice. From this, the proportion of fresh martensite can be determined quantitatively. Alternatively, the formation of fresh martensite can be determined with the help of dilatometry based on the change in volume when a sample is cooled.
  • the KGs value does not change during the final annealing.
  • the percentages given for the structural components relate to parts of the area that are usually also taken over as proportions by volume.
  • martensite is defined as tempered if the fresh martensite was subsequently annealed again at least at a minimum temperature of 100 ° C after its formation.
  • the minimum temperature of 100 ° C. corresponds to the minimum temperature of the final annealing according to the invention.
  • the fresh martensite before the final annealing is then understood as tempered martensite after the final annealing.
  • Fresh martensite is therefore a conversion product of austenite, which is created when it is cooled and is not tempered.
  • a temperature cycle according to the invention requires an Hp value of> 7.5.
  • a KGs value of ⁇ 0.3 does not necessarily mean that any temperature cycle is successful, but it is an advantageous criterion for the final annealing to be successful from Hp> 7.5.
  • the KGs value does not change during the final annealing. Due to the restriction of the form factor to F ⁇ 3, very elongated, irregular structural components from the rolling process are irrelevant when considering the grain sizes.
  • the KGs parameter thus correlates with coarse structural components that are newly formed during cooling after the first annealing.
  • the structural components newly formed after the first annealing are decisive for the yield strength, since the formation of fresh martensite in these areas reduces the yield strength.
  • short diffusion paths are necessary, which preferably requires the smallest possible grain size and thus a low KGs parameter ⁇ 0.4, advantageously ⁇ 0.3.
  • FIG. 1 An exemplary comparison of the microstructure of the reference steel Bn (left microstructure) with a KGs characteristic value of 0.58 and example steel Div (right microstructure) with a KGs characteristic value of 0.1 is shown in FIG.
  • Grains with an equivalent diameter d> 5 qm and form factor F ⁇ 3 are marked in gray in Figure 2, the remaining fine structure is shown in white.
  • the lowest possible proportion of grains shown in gray is advantageous for the invention, which is reflected by the KGs characteristic value.

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Abstract

L'invention vise à mettre au point un procédé de fabrication d'une bande d'acier présentant une structure multiphasée et permettant de produire des géométries complexes offrant une capacité élevée d'absorption d'énergie et une résistance élevée à la fissuration des bords, compensant en particulier la chute de la limite d'élasticité, et permettant par conséquent d'obtenir une combinaison d'une limite d'élasticité élevée ou d'un rapport d'élasticité élevé et d'un allongement à la rupture élevé, le procédé étant proposé comme comprenant les étapes suivantes : production d'une bande d'acier laminée à chaud ou à froid à partir d'un acier constitué des éléments suivants en % en poids : C : de 0,085 à 0,149 ; Al : de 0,005 à 0,1 ; Si : de 0,2 à 0,75 ; Mn : de 1,6 à 2,9 ; N : < 0,02, S : ≤ 0,005 et éventuellement d'un ou de plusieurs des éléments suivants en % en poids : Cr : 0,05 à 0,5 ; Mo : 0,05 à 0,5 ; Ti : 0,005 à 0,060 ; Nb : 0,005 à 0,060 ; V : 0,001 à 0,060 ; B : 0,0001 à 0,0060 ; N : 0,0001 à 0,016 ; Ni : 0,01 à 0,5 ; Cu : 0,01 à 0,3 ; le reste étant du fer, y compris des éléments d'accompagnement classiques en sidérurgie ; premier recuit, en particulier un recuit continu, de la bande d'acier, en particulier la bande d'acier laminée à froid, à une température comprise entre 750 °C et 950 °C, inclus, pendant toute la durée de 10 s à 1200 s, en particulier de 50 s à 650 s, puis premier refroidissement de la bande d'acier à une température comprise entre 200 °C et 500 °C, inclus, à une vitesse moyenne de refroidissement de 2 K/s à 50 K/s, en particulier de 5 K/s à 100 K/s ; refroidissement supplémentaire de la bande d'acier jusqu'à une température de sur-refroidissement inférieure à 100 °C à une vitesse moyenne de refroidissement de 1 K/s à 50 K/s ; recuit final, en particulier recuit continu, de la bande d'acier avec un paramètre Hollomon-Jaffe Hp = TH * (ln (Ʈ) +20) supérieur à 7,5 x 103, la température maximale TH, en K, allant de 100 °C à 470 °C, inclus, et la durée totale Ʈ, en h, allant de 2 s à 1000 s, inclus ; et refroidissement final de la bande d'acier jusqu'à la température ambiante à une vitesse moyenne de refroidissement de 1 K/s à 160 K/s, en particulier de 1 K/s à 30 K/s. L'invention concerne également une bande d'acier à structure multiphasée produite par le présent procédé.
EP21718871.3A 2020-04-15 2021-04-14 Procédé de fabrication d'une bande d'acier à structure multiphasée et bande d'acier associée Pending EP4136265A1 (fr)

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PCT/EP2021/059672 WO2021209503A1 (fr) 2020-04-15 2021-04-14 Procédé de fabrication d'une bande d'acier à structure multiphasée et bande d'acier associée

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CN115698347A (zh) 2023-02-03
WO2021209503A1 (fr) 2021-10-21
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