Classifications

• • 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
  • 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
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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
  • 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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  • 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
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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
  • C21D6/00—Heat treatment of ferrous alloys
  • C21D6/005—Heat treatment of ferrous alloys containing Mn
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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
  • C21D7/00—Modifying the physical properties of iron or steel by deformation
  • C21D7/02—Modifying the physical properties of iron or steel by deformation by cold working
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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
  • C21D7/00—Modifying the physical properties of iron or steel by deformation
  • C21D7/02—Modifying the physical properties of iron or steel by deformation by cold working
  • C21D7/10—Modifying the physical properties of iron or steel by deformation by cold working of the whole cross-section, e.g. of concrete reinforcing bars
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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
  • C21D7/00—Modifying the physical properties of iron or steel by deformation
  • C21D7/13—Modifying the physical properties of iron or steel by deformation by hot working
• • 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
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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
  • 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
  • C21D8/0263—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment following hot rolling
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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
  • 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
  • C21D8/0421—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 working steps
  • C21D8/0426—Hot rolling
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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
  • 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
  • C21D8/0421—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 working steps
  • C21D8/0436—Cold rolling
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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
  • 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
  • C21D8/0463—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 following 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
  • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
  • C21D9/08—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for tubular bodies or pipes
• • 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
  • C22C38/06—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
  • C22C38/26—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
  • C22C38/32—Ferrous alloys, e.g. steel alloys containing chromium with boron
• • 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
  • C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
  • C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
  • C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
  • C23C2/04—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor characterised by the coating material
  • C23C2/12—Aluminium or alloys based thereon
• • C—CHEMISTRY; METALLURGY
  • C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
  • C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
  • C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
  • C23C2/26—After-treatment
  • C23C2/28—Thermal after-treatment, e.g. treatment in oil bath
• • C—CHEMISTRY; METALLURGY
  • C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
  • C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
  • C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
  • C23C2/34—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor characterised by the shape of the material to be treated
  • C23C2/36—Elongated material
  • C23C2/40—Plates; 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
  • C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
  • C21D1/62—Quenching devices
  • C21D1/673—Quenching devices for die 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
  • C21D2211/00—Microstructure comprising significant phases
  • C21D2211/004—Dispersions; Precipitations
• • 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
  • C25—ELECTROLYTIC OR ELECTROPHORETIC PROCESSES; APPARATUS THEREFOR
  • C25D—PROCESSES FOR THE ELECTROLYTIC OR ELECTROPHORETIC PRODUCTION OF COATINGS; ELECTROFORMING; APPARATUS THEREFOR
  • C25D5/00—Electroplating characterised by the process; Pretreatment or after-treatment of workpieces
  • C25D5/34—Pretreatment of metallic surfaces to be electroplated
  • C25D5/36—Pretreatment of metallic surfaces to be electroplated of iron or steel
• • C—CHEMISTRY; METALLURGY
  • C25—ELECTROLYTIC OR ELECTROPHORETIC PROCESSES; APPARATUS THEREFOR
  • C25D—PROCESSES FOR THE ELECTROLYTIC OR ELECTROPHORETIC PRODUCTION OF COATINGS; ELECTROFORMING; APPARATUS THEREFOR
  • C25D7/00—Electroplating characterised by the article coated
  • C25D7/06—Wires; Strips; Foils
  • C25D7/0614—Strips or foils

Definitions

• “Flat steel products” are defined here as rolled products whose length and width are each significantly greater than their thickness. These include, in particular, steel strips, steel sheets, and blanks obtained from them, such as blanks and the like. Flat steel products of the type in question are used for cold forming followed by quenching and tempering to adjust the mechanical properties of the resulting component, or for hot forming to adjust the mechanical properties of the resulting component.
• Hot forming is also referred to as “die hardening” or “press hardening.” Strictly speaking, press hardening refers to the hardening of a workpiece or component in a cooled tool, while hot forming also includes the preliminary shaping in a heated state. However, these three terms are often used synonymously.
• Quenching and tempering refers to a treatment known per se from the state of the art, which initially involves heating to a temperature at which the steel of the processed flat steel product (component) has a completely austenitic structure. This heating serves to bring the respective component to the appropriate temperature. This heating is performed as a separate work step on the component previously cold-formed from the flat steel product. After heating, the component is cooled at an accelerated rate so that the steel of the flat steel product from which the component is formed forms a hardened structure, resulting in the component achieving significantly increased strength. After quenching, the component can be tempered to reduce the internal stresses that can arise in the component's structure during the quenching process. Generally, the highest possible temperatures are aimed for during quenching and tempering to shorten cycle times and thus reduce costs.
• the invention is based on the object of creating a weight-reduced component which, in the tempered and/or hot-formed state, has an optimal combination of strength and toughness and is, as such, suitable for applications which place the highest demands on mechanical properties or resistance to abrasive wear.
• a method according to the invention which achieves the above-mentioned object comprises at least the work steps specified in claim 9. It goes without saying that a person skilled in the art, when carrying out the method according to the invention, will supplement the work steps not explicitly mentioned here, which he knows from his practical experience are regularly used when carrying out such methods.
• a component according to the invention is accordingly formed from a flat steel product which consists of a steel which, in mass%, consists of C: 0.1 - 0.6%, in particular 0.10 - 0.60% C, Mn: 0.1 - 2.0%, in particular 0.1 - 2.0% Mn, Al: 0.05 - 0.20%, in particular 0.050 - 0.20% Al, Nb: 0.01 - 0.06%, in particular 0.010 - 0.060% Nb, B: 0.0005 - 0.005%, Cr: 0.05 - 0.8%, Si: up to 0.8%, Mo: up to 1.5%, Cu: up to 0.5%, Ni: up to 1.5%, V up to 0.2%, REM up to 0.05%, Ti: up to 0.02%, in particular 0.020% Ti, Ca: up to 0.005%, the remainder being iron and unavoidable impurities, the impurities including contents of up to 0.03% P, up to 0.03% S, up to 0.01% N, less than 0.05% Sn, less than
• a flat steel product according to the invention has a structure in which, in a homogeneous distribution across the strip thickness, at most 150 area ppm of particles of high hardness are present, the mean circle-equivalent particle size of which is 0.2 - 10 ⁇ m and which consist of Al compounds on an oxide basis, of AIN, TiN or of conglomerates formed on the basis of these particles.
• Carbon "C” is contained in the steel of the flat steel product from which a component according to the invention is formed, as a mandatory element in contents of 0.1 - 0.6 mass%, in particular 0.10 - 0.60 mass%.
• the presence of C controls the level of hardening potential.
• both the martensite content and the hardness of the martensite obtained in the microstructure of a component according to the invention increase after austenitization and accelerated cooling, with a single-phase martensite structure representing the target microstructure of the finished processed component.
• the increase in hardness is synonymous with an increase in strength in the tensile test.
• a C content of at least 0.1 mass%, in particular at least 0.10 mass% is required.
• the beneficial effects of the presence of C can be achieved particularly reliably in the flat steel product according to the invention with a C content of at least 0.12 mass%, in particular at least 0.15 mass%.
• the hardness or strength after accelerated quenching would be so high that the toughness of the flat steel product from which a component according to the invention is formed, as reflected in the elongation at break or area at break, would be significantly reduced.
• the tendency to crack formation would increase and weldability would deteriorate.
• Negative effects of the presence of C can be particularly reliably prevented by limiting the C content to a maximum of 0.55 mass%, in particular a maximum of 0.50 mass%. Therefore, the optimal C content is 0.12–0.55 mass%, in particular 0.15–0.50 mass%.
• Si can optionally be present in the steel of the flat steel product from which a component according to the invention is formed, in amounts of up to 0.8 mass%.
• Si inhibits cementite and pearlite transformation and thereby increases the martensite hardenability of the flat steel product.
• Si reduces the cooling rate critical for the desired martensite formation and thus increases the hardening of a flat steel product produced according to the invention.
• Si also exhibits opposite segregation behavior than Mn and thus overall improves the segregation resistance of the steel from which the flat steel product from which a component according to the invention is formed. Minimizing segregation across the cross-section is particularly important when a component according to the invention is pipe or similar.
• Si contents of at least 0.1 mass%, in particular at least 0.15 mass% can be provided.
• excessively high Si contents could impair the wetting behavior of the flat steel product from which a component according to the invention is formed, particularly if alloyed steel flat products according to the invention are to be hot-dip coated.
• Si tends to form external oxides during the annealing of the flat steel product carried out in this case.
• the Si content of a flat steel product used for a component according to the invention is at most 0.8 mass%. Negative effects of the presence of Si can be particularly reliably avoided if the Si content is limited to a maximum of 0.5 mass%.
• the optimal Si content is therefore 0.1 - 0.8 mass%, in particular 0.15 - 0.5 mass%.
• Manganese "Mn” is present in the steel of the flat steel product from which a component according to the invention is formed, in contents of 0.1 - 2 mass%, in particular 0.10 - 2.0 mass%. Mn increases the hardenability of the steel by lowering the A3 transformation temperature (i.e. the Ac3 and/or Ar3 temperature) from ferrite to austenite. As a result, during the heat treatment of the flat steel product from which a component according to the invention is formed, the furnace temperature for complete transformation to austenite can be reduced during heating. In particular, the formation of the diffusion-controlled transformation phases ferrite, pearlite, and bainite is shifted to longer times. Therefore, in this respect, manganese is a similarly effective alloying element to carbon.
• A3 transformation temperature i.e. the Ac3 and/or Ar3 temperature
• manganese Compared to carbon, manganese has the advantage of achieving greater deformability in the hardened state, which is expressed, for example, in higher notched impact toughness.
• the reduction in the critical cooling rate with increasing manganese content is also associated with an increase in hardening capacity. Fluctuations in the cooling conditions or different contact conditions during the cooling of components manufactured from steel material alloyed according to the invention can be better compensated for, and the property scatter is limited.
• excessive Mn contents increase C segregation behavior, which can lead to inhomogeneous hardening behavior across the cross-section of the respective product and the formation of hardening cracks.
• Mn contents increase the risk of external Mn oxides or Mn-based mixed oxides forming on the surface of the product produced from the steel material alloyed according to the invention. As with excessive Si contents, this would trigger the risk of a deterioration in the wetting behavior of a flat steel product made from the steel material alloyed according to the invention during hot-dip coating.
• excessive Mn contents on the hot or cold strip surface would also lead to undesirable discoloration or so-called "manganese haze" due to the formation of manganese oxides.
• the Mn content of a flat steel product intended for forming a component according to the invention must be limited to a maximum of 2% by mass.
• the Mn content is at least 0.4 mass%, in particular at least 0.40 mass%.
• the Mn content is 0.4 - 1.5 mass%, preferably 0.40 - 1.50 mass%, in particular 0.6 - 1.3 mass% or 0.6 - 1.2 mass%, preferably 0.60 - 1.30 mass% or 0.60 - 1.20 mass%.
• Phosphorus is one of the unavoidable steel components present during production. P segregates particularly at the grain boundaries and reduces grain boundary strength. Higher P contents would therefore contribute to a weakening of the microstructure, which in turn would lead to a deterioration in the toughness of the material.
• the P content of a flat steel product intended for forming a component according to the invention is therefore limited to a maximum of 0.03 mass%, with the P content being set as low as possible. Therefore, the P content of the flat steel product is preferably a maximum of 0.025 mass%, in particular a maximum of 0.02 mass%.
• Sulfur "S” is also an accompanying element, the presence of which is fundamentally undesirable in the steel flat product intended for forming a component according to the invention. Due to the Mn contents provided for in the invention, non-metallic MnS precipitates would form at higher S contents. After rolling the steel flat product, these precipitates would be present in an elongated form due to their low hardness and would negatively influence the fracture behavior. During deformation, initial microscopic material separations could form, expand, and develop on elongated MnS due to crack initiation and crack propagation. grow together until they impair the material's macroscopic behavior in the form of reduced impact toughness and increasing material anisotropy. To eliminate the negative effects of the presence of S in the steel alloyed according to the invention, the S content is limited to a maximum of 0.03 mass%, with low S contents of less than 0.006 mass%, in particular less than 0.003 mass%, being particularly favorable.
• Aluminum is present in the steel of the flat steel product from which a component according to the invention is formed, in contents of 0.05–0.2 mass%, in particular 0.050–0.20 mass%.
• Al traditionally serves as a deoxidizer, for which purpose it is typically alloyed in practice in contents of 0.02–0.05 mass%.
• increased Al contents of 0.050–0.20 mass%, in particular 0.050–0.20 mass%, in combination with optionally low Ti contents of up to 0.02 mass%, in particular up to 0.020 mass% are provided in the steel alloyed according to the invention.
• the Al content can be limited to at least 0.06 mass%, in particular at least 0.060 mass%, or at least 0.07 mass%, in particular at least 0.070 mass%.
• contents of more than 0.2 mass%, in particular more than 0.20 mass% Al there would be a risk of external Al oxides forming on the surface of a product made from a steel material alloyed according to the invention, which would impair the wetting behavior during hot-dip coating.
• the minimum Nb content of a flat steel product intended for forming a component according to the invention is 0.010 mass%, with Nb contents of at least 0.015 mass% or at least 0.020 mass% having proven particularly advantageous.
• UG Al_Nrest is equal to 0.070 mass%, in particular equal to 0.075 mass%, preferably equal to 0.080 mass%, in particular equal to 0.081 mass%
• OG Al_Nrest is equal to 0.150 mass%, in particular equal to 0.135 mass%, preferably equal to 0.125 mass%, particularly preferably equal to 0.121 mass%. Accordingly, according to a particularly advantageous embodiment of the invention: 0 , 081 Masse ⁇ % ⁇ % Al ⁇ 27 / 14 * % Nrest ⁇ 0,121 Masse ⁇ %
• Mo Molybdenum
• Mo can optionally be present in the steel of the steel product according to the invention in contents of up to 1.5 mass%.
• Mo suppresses the formation of ferrite and pearlite during cooling and enables increased martensite or bainite formation even at lower cooling rates, thereby increasing hardenability.
• the hardenability-enhancing effect of Mo is significantly greater than that of Cr.
• Mo can effectively increase strength in large thicknesses and cross-sections where, due to dimensional or structural reasons, only relatively low cooling rates are possible.
• Mo also reduces temper embrittlement of heat-treatable steels.
• Mo is also a strong carbide former and can thus also contribute to increasing strength through precipitation formation.
• Vanadium “V” can also optionally be present in the steel of the flat steel product from which a component according to the invention is formed to effect precipitation strengthening. Suitable V contents for this purpose are up to 0.2 mass%, whereby the effect of V can be utilized by optional contents of at least 0.03 mass%.
• Endogenous or exogenous inclusions that arise during steel production generally lead to a reduction in the degree of purity, which can lead to premature failure of components. This can be an increasing problem, especially for high-strength components. This is especially true when such components are exposed to cyclic or dynamic loads.
• Exogenous inclusions are generally isolated cases and originate, for example, from ladle slag or refractory material, but play no role here and are therefore not considered here.
• one of the objectives of the inventive adjustment of the steel alloy of a component according to the invention was to reduce the proportion of coarse and hard TiN-, AlN-, and oxidic Al-based particles as well as conglomerates of these compounds for toughness reasons, while still reliably binding the nitrogen present in each case in order to achieve complete transformation to martensite, even with larger strip thicknesses and component cross-sections, via the strong transformation-retarding effect of interstitially dissolved B, even at relatively low cooling rates of at least 30 °C/s to a maximum of 120 °C/s.
• Al and the optionally present Ti form hard precipitates, which, in components formed from flat steel products alloyed according to the invention, could be the source of cracks and their propagation due to the notch effect and the stress fields surrounding the particles.
• the angular and cube-shaped TiN particles in particular prove to be harmful here simply due to their shape and size.
• the invention has matched the contents of the alloying elements and the conditions during the production of flat steel products intended for forming components according to the invention in such a way that in the microstructure of a flat steel product according to the invention and of a component produced therefrom, at most up to 150 area ppm of hard TiN particles and Al-based oxide particles as well as AIN with an average, circle-equivalent particle size of 0.2 - 10 ⁇ m are present, homogeneously distributed over the strip thickness.
• hard particles here essentially includes particles of AlN, Al2O3 , and Al2O3 - based spinels, as well as TiN particles and conglomerates formed on the basis of these particles.
• Such particles each exhibit a high Mohs hardness of approximately 9. Due to their high hardness, they are hardly deformable during rolling or forming processes and lead to local stress fields in their environment, which can contribute to premature material failure.
• Conglomerates are particularly referred to here as particle composites in which additional particles form through heterogeneous nucleation on already existing particles, e.g., Al2O3 with MnS, where the basis represents one of the previously mentioned hard particle types.
• the alloying concept according to the invention has achieved that the total number of hard TiN-based precipitates and their mixed forms falling within this particle size range in a component formed from a flat steel product alloyed according to the invention is reduced to less than 30% of the particles in the 0.2 - 10 ⁇ m size class present in the structure of a component.
• the absolute number of precipitates falling within the relevant particle size range is reduced compared to conventional flat steel products, for example, made from a steel with higher Ti contents, whereby the average distance between the 0.2 - 10 ⁇ m sized precipitates in the component formed from a flat steel product alloyed according to the invention is significantly increased.
• the proportion of hard TiN particles and their mixed forms accounts for more than 45 - over 80% of the volume fraction of the particles present in the size class 0.2 - 10 ⁇ m. Due to this high proportion, a reduction in the Ti mass fraction makes sense, which correspondingly leads to a reduction in the proportion of hard TiN particles in the inventive concept. Due, among other things, to the optional, but in any case limited, addition of titanium alloying, coarse particles such as TiN occur much less frequently in a flat steel product intended for forming a component according to the invention than is the case with conventional concepts in which higher Ti contents are provided. By reducing the proportion of coarse precipitates, an improvement in toughness is achieved, which prevents the formation and propagation of cracks.
• the moderate increase in the Al mass content does not, in turn, lead to a significant increase in the proportion of similarly hard, oxide-based Al-based precipitates, as well as AlN and their conglomerates.
• the risk of premature material failure is reduced in the flat steel products intended for forming a component according to the invention.
• the optimization of toughness achieved by the invention is noticeable in an improvement in the area of reduction at fracture of the component according to the invention in the hot-formed, tempered state, which is of particular interest to the component manufacturer.
• a microstructure consisting entirely of martensite in the technical sense is created through full austenitization followed by quenching and optional tempering. This naturally includes, according to expert understanding, the possibility that up to 5% of other components may be present in the microstructure of a component according to the invention, which, however, are ineffective with regard to the properties of a component according to the invention determined by the martensite content.
• Nb and Al in dissolved and precipitated form reduce austenite grain growth during the production and heat treatment of the flat steel product made from the alloyed steel according to the invention and the component manufactured therefrom, and after transformation, the martensite package size is reduced.
• further relevant precipitates in the flat steel product such as NbN, NbC, and AlN, which, as monolithic particles without nucleation on previously formed precipitates, generally only reach a maximum size of approximately 100 nm. In this way, more homogeneous precipitate fractions with narrower particle size ranges are achieved. These prove particularly effective with regard to controlling austenite grain size.
• the steel used to produce the flat steel product from which a component according to the invention is formed has an austenite grain size during austenitization that is up to half an ASTM grain size finer than that of conventional steel concepts belonging to the category of the steel according to the invention.
• the grain sizes of a flat steel product alloyed and processed according to the invention are within a narrower range, i.e., with a reduced standard deviation.
• Recrystallization and grain growth occur, i.e., they predominate in the microstructure shortly before the start of quenching.
• the finer this austenite grain size on average the finer the resulting martensite package size and the more beneficial it is for the toughness of the martensite and thus of the material or component.
• the grain size of the austenite fluctuates only slightly, thus resulting in only minimal local fluctuations in the martensite hardness after transformation. This also helps prevent springback effects on a press-hardened or tempered component due to locally inhomogeneous microstructures.
• the endogenous inclusions can be influenced in terms of type, size, and distribution. In addition to solidification, this influence extends particularly to the hot rolling production stage, as explained below.
• a melt is produced which, according to the above explanations, is composed of the alloy of the steel of a flat steel product intended for forming a component according to the invention. Alloying of this melt, the information given above regarding advantageous designs of the steel of a flat steel product intended for forming a component according to the invention naturally applies equally to the melt produced and processed in the course of the method according to the invention.
• Nucleation is understood to mean the formation of precipitations on previously formed precipitates based on heterogeneous nucleation, as opposed to homogeneous precipitation without extraneous nucleation sites.
• the alloy of the steel used in the flat steel products processed according to the invention and the components formed therefrom reduces, on average, nucleation effects on previously formed precipitates.
• the nucleation of TiC, NbN, NbC, AlN on TiN would reduce the probability of formation These precipitates decrease at lower formation temperatures, thus impairing their effectiveness in achieving the microstructure refinement sought by the invention.
• austenite grain size refinement can be achieved even at the intended, comparatively low Nb contents.
• the preheating temperatures used according to the invention are 1100–1350 °C, and preferably 1150–1280 °C. Below 1100 °C, coarsening and nucleation effects of the particles during preheating must be expected. Temperatures above 1350 °C should be avoided to limit coarsening of the austenite grain, reduce material loss due to scaling, and, from an economic perspective, reduce energy costs.
• the holding times for preheating the precursors are required for the complete dissolution of the precipitates present in the precursors to be preheated after casting.
• the total holding time for slabs envisaged according to the invention is 150–400 minutes, with the total holding time including the time required for heating to the respective target preheating temperature and for thorough heating of the precursors. With total holding times of less than 150 minutes, there is a risk that the relevant microalloy precipitates will not dissolve completely. However, holding times of more than 400 minutes should also be avoided to limit austenite grain coarsening. Thin slabs are preheated in an equalizing furnace for significantly shorter times of 10–90 minutes.
• Strips produced by strip casting are generally not preheated, but are hot-rolled directly in one or more hot-roll stands to final hot-rolled thicknesses of 1 - 4 mm.
• Hot rolling may include roughing, in which the slabs are typically rolled in a so-called “roughing stand” in a reversing manner to an intermediate thickness of approximately 35 to 60 mm.
• the roughed slab then enters a multi-stand finishing hot rolling mill, where it is continuously hot-rolled step by step into a hot strip.
• pre-rolling is unnecessary. It can be fed directly into the finishing hot rolling mill after preheating (if necessary).
• hot rolling in step d) is terminated at hot rolling end temperatures that are at least 50 °C higher than the Ar3 temperature of the steel, but no more than 150 °C above this temperature. Hot rolling is thus terminated at a hot rolling end temperature at which the resulting hot strip still has a fully austenitic microstructure.
• Such a rolling strategy is referred to as "normalizing rolling.”
• the hot rolling end temperature is selected such that the tendency of Nb and Al to form deformation-induced precipitates is reduced and a larger proportion of precipitation potential is available for inhibiting grain growth during austenitization in the subsequent quenching and tempering process or during hot forming.
• Suitable hot rolling end temperatures for the alloying concept according to the invention are typically above 830 °C.
• the final rolling temperature is at least 60 °C and at most 130 °C higher than the Ar3 temperature, whereby hot rolling final temperatures that are at most 110 °C above the Ar3 temperature have proven to be particularly practical in order to To limit austenite grain growth.
• Normalizing rolling is preferred here because it involves comparatively low hot rolling forces and avoids the precipitation of deformation-induced, relatively coarse precipitates.
• the precipitation potential for reducing the original austenite grain size can be maximized in subsequent austenitizing stages of quenching and tempering and hot forming. This has a positive effect on toughness.
• step e) In order to preserve the precipitation potential of Al and Nb in the hot strip obtained after hot rolling for subsequent process steps, it is necessary in step e) to cool the hot strip after hot rolling in the temperature range of 800 °C to 650 °C at a cooling rate of more than 20 °C/s to the coiling temperature.
• the actual coiling temperature achieved is determined by the cooling in the cooling section. According to the invention, it is significantly below the A1 temperature of the steel from which the flat steel product according to the invention is produced in order to avoid relatively coarse pearlite precipitation in the hot strip.
• the temperature "A1" is the temperature at which austenite decomposes into pearlite from high temperatures.
• A1 is 723 °C, whereby this transformation occurs at carbon contents > 0.02 mass%, which is the case with the steel concepts according to the invention.
• the A1 temperature is based on empirical formulas that reflect the influence of the alloying elements on A1 (see for example Hougardy, HP “Material Science Steel Volume 1: Fundamentals", Verlag Stahleisen GmbH, Düsseldorf, 1984, p. 229 ), at 722 - 727 °C and thus in a narrow range.
• coiling temperatures of ⁇ 720 °C are used in particular.
• the solution state and the precipitation form of the carbon are influenced in such a way that a finely distributed C precipitation is achieved for subsequent quenching and tempering or hot forming treatments in order to ensure the C dissolution for the To accelerate the hardening process.
• the rapid cooling of the resulting hot strip in the temperature range of 800–650°C according to the invention thus suppresses the precipitation of Nb and Al. This can be ensured in particular by a cooling rate of at least 20°C/s. It should be noted that during cooling after hot rolling, reheating of up to 30°C can occur due to the phase transformation. In practice, the strip can be sprayed with water for the inventively controlled cooling downstream of the hot rolling mill in which the hot rolling takes place. Cooling sections known from the prior art, which combine laminar and spray cooling devices, are particularly suitable for this purpose. These should be capable of achieving cooling rates of preferably more than 20°C/s, in particular at least 50°C/s, and a maximum of 200°C/s, especially in the temperature range of 800–650°C.
• the coiling temperature to which the hot strip is cooled after hot rolling and at which the hot strip is coiled into a coil in step f) is 450 - 720 °C.
• the upper limit of 720 °C is advantageous in order to be able to set a sufficiently low tensile strength for subsequent cold forming at C contents ⁇ 0.4 %.
• the coiling temperature is particularly preferably lower than 650 °C in order to further suppress the precipitation formation of Nb and Al and to achieve the most finely distributed C dissolution state possible.
• An upper coiling temperature of 650 °C has proven to be particularly advantageous because coarse-structured pearlite formation can then be largely avoided.
• the flat steel product obtained according to the invention as a hot-rolled strip after coiling typically has a tensile strength of less than 700 MPa. Only through the subsequent quenching and tempering treatment or through the processing completed during hot forming are the largely fully martensitic microstructure according to the invention and, consequently, the optimized mechanical properties of a component according to the invention achieved.
• step g the hot-rolled strip can be subjected to pickling for further processing to remove any adhering scale.
• This processing step is advantageous if the hot-rolled strip is formed in a cold-forming tool and contamination or damage to the tool can be avoided by abrasion of the scale. There are no special requirements for pickling. It can be carried out using any method known for this purpose.
• the hot strip can be subjected to batch annealing in step h) to reduce the strength of the steel for subsequent cold forming.
• the core temperatures of the coiled steel flat product during batch annealing are 500 - 720 °C.
• a core temperature of at least 500 °C is required to ensure sufficient strength reduction.
• Annealing temperatures of However, temperatures exceeding 720 °C would ensure that the formation of new pearlite by exceeding the A1 temperature in all areas of the coil can be reliably avoided during batch annealing.
• a batch annealing time of at least 5 h at core temperature level is required in order to significantly reduce the strength level, i.e. to ⁇ 700 MPa tensile strength.
• HNX annealing results in longer total annealing times of up to 50 hours, as heat transfer is slower than with a pure H2 atmosphere.
• the core temperature of the bell annealing should be below 720 °C, in particular around 680 °C, but in any case below the A1 temperature of the steel from which the flat steel product is made. This restriction prevents the formation of new pearlite during the annealing process. Instead, cementite particles (carbide particles) are partially formed from the hot-rolled strip microstructure present at the beginning of annealing, particularly through carbon diffusion and carbon redistribution. At the same time, coagulation can lead to coarsening of the microstructure.
• Two alternative approaches are available to form components according to the invention from hot-rolled or cold-rolled strip produced according to the invention, which possess an optimized combination of high strength and toughness.
• a blank cut from the respective hot-rolled or cold-rolled strip is heated and press-hardened according to steps I.1 - I.3 of the inventive method.
• the blank is first cold-formed and then tempered according to steps m.1 - m.3 of the inventive method.
• the austenitizing temperature maintained in steps I.1 and m.2 is accordingly within a range from (Ac3 - 100 °C) to 950 °C, in particular (Ac3 - 75 °C) to 950 °C, or, particularly advantageously, from (Ac3 - 100 °C) to 950 °C, with austenitizing temperatures of Ac3 - 950 °C being particularly practical.
• a total heating time of typically 1 second to 20 minutes is required for the through-heating of the blank or component, although in practice, total times of at least 10 seconds, especially at least 1 minute, are suitable to reliably achieve through-heating.
• the total heating time includes the time required to heat to the austenitizing temperature.
• the respective flat steel product is placed after austenitizing within a transfer time of 1 - 20 seconds into a hot forming device known for this purpose from the prior art, in which it is then press-hardened into a component in an equally known manner, the average cooling rate to room temperature being 30 - 120 °C/s.
• the component heated to the austenitizing temperature is quenched to room temperature after austenitizing, also at an average cooling rate of 30–120 °C/s.
• the component can be immersed in a suitable quenching medium in a conventional manner or exposed to the quenching medium using equally conventional devices, such as nozzles or jet devices. If a continuous heating device of the type described above, particularly an inductive one, is used to heat the component, the section of the blank heated to the austenitizing temperature can also be cooled in a continuous flow using a suitable quenching device upon exiting the respective heating device.
• the microstructure of the flat steel products in question and the components according to the invention made from them was investigated as follows: The proportions of hard oxide and nitride particles in the microstructure of a flat steel product are given in area ppm, unless otherwise stated. The exact procedure for determining them is described below. According to ASTM E2142 from 2008, the area fraction of inclusions can be equated with the volume fraction. Likewise, the phase fractions of the microstructure given in this text refer to the evaluated ground surface and are therefore given in area %.
• the inclusions were examined on longitudinal sections across the strip thickness using a scanning electron microscope (SEM) from Zeiss (model GeminiSEM 500) equipped with the EDX system "Oxford Xmax" from the manufacturer "Oxford Instr.” for energy-dispersive element analysis.
• SEM scanning electron microscope
• the data was evaluated using the software "Aztec 3.3 SPI, Feature Analysis” from “Oxford Instr.” Inclusions with a size of approximately 0.2 ⁇ m or larger were recorded.
• the element contents of the precipitates were determined using calibration samples.
• the inclusions were classified using the Stoichiometry of the known precipitates, classified into oxides, sulfides, and TiN. Quantification and standardization of the measured elements, excluding Fe, C, and Ag, were performed.
• the recorded elements were converted to oxides (excluding S, P, Cl, F) and standardized to 100%. Additionally, the subsystem ⁇ Al2O3-SiO2-CaO> was calculated and standardized to 100%.
• Computer - aided classification tables of the analyzed inclusions were then created from the raw data obtained. Inclusions that could not be clearly classified were listed separately. These inclusions were individually examined. Particle size was idealized as a circle-equivalent diameter, regardless of particle shape.
• the homogeneity of the microstructure of the former austenite and the distribution of its constituents was determined using electron backscatter diffraction (EBSD) in the fully martensitic state after tempering or press hardening on longitudinal sections across the strip thickness.
• the samples were polished with the polishing compound "OP-S Suspension" from the manufacturer "Struers.”
• a measuring field measuring 140 ⁇ m x 140 ⁇ m was positioned at different layers across the strip thickness and scanned with a step size of 0.15 ⁇ m.
• Several layers across the strip thickness were also examined (1/6, 1/3, 1/2) to obtain information about the homogeneity of the microstructure.
• the mechanical parameters of flat steel products or components made from them mentioned in this text are the tensile test values (tensile strength, yield strength, modulus of elasticity, uniform elongation and elongation at break) determined according to DIN EN ISO 6892-1.
• tf denotes the thickness of the thinnest points in the necking strain range of the fracture cross-section, determined from four measurements across the sample width.
• the "absolute strain in the thickness direction” or “necking at fracture” was measured on tensile specimens after quenching and tempering using an optical system (microscope).
• the thickness tf was determined in the fracture cross-section at four points across the width (1 mm to the right of the left edge, center, minimum, 1 mm to the left of the right edge).
• Three parallel tensile specimens were tested in each case to obtain a representative statement for the respective condition examined. A total of six fracture cross-sections were measured. The average value for one sample was calculated from the six measured values.
• the tensile specimens were oriented longitudinally to the rolling direction.
• Steel grades 1–6 were each melted and cast into slabs. The slabs were then heated to a preheating temperature and then hot-rolled into a hot strip. The hot strip obtained during hot rolling was cooled to a coiling temperature at which it was coiled into a coil. The coil was then cooled to room temperature.
• the flat steel product produced in this way as unpickled hot-rolled strip from steel 2 was made available for further processing without further treatment.
• the hot strips produced from steels 1 and 3 - 6 were subjected to a pickling treatment in order to remove any scale adhering to them.
• the hot-rolled strips produced from steels 1 and 4 were then cold-rolled into cold-rolled strips without intermediate annealing.
• the resulting cold-rolled strips each underwent continuous annealing, were coated with an AlSi layer by hot-dip coating, and finally skin-pass rolled.
• the flat steel products, available as cold-rolled strips with an AlSi coating, were prepared for further processing into components.
• the hot-rolled strips produced from steels 3, 3a and 6 were subjected to batch annealing and were made available in this condition as flat steel products for further processing.
• the flat steel product produced from steel 5 as hot-rolled strip was prepared for further processing after pickling.
• Table 1 lists the chemical compositions of steels 1 - 6.
• the contents of the elements P, S, and N, which are present during production but are attributable to impurities, are given here because they are of particular importance for the quality of the steels produced according to the invention and, particularly in the case of steels 1 - 3a according to the invention, it must be ensured that the contents of these elements comply with the requirements of the invention.
• the final thickness D of the steel flat products produced from steels 1 - 6 is given in Table 2. This means that for the steel strips produced from steels 1 and 4, the thickness D in the finished cold-rolled and AlSi-coated state and for the hot-rolled steel strips produced from steels 2, 3, 3a, 5 and 6, the thickness after coiling (from the hot-rolled strip produced from steel 2) or hot-rolled strip produced after descaling (from steels 3, 3a, 5 and 6).
• Table 2 shows for steels 1 - 6 the result of the equation %Ti-48/14*%N, the ratio %Ti/%N, the content %Nrest of the nitrogen not bound by Ti, the result of the equation %Al-27/14*%Nrest, the ratio %Al/%N and the result of the equation %Al//%N*14/27, where %Ti denotes the Ti content, %N the N content and %Al the Al content of the respective steel.
• Each of the steels was alloyed with B, with B contents of at least 0.001 mass%.
• the inventive steels 1 - 3a each have Ti contents that are insufficient, or at best only just sufficient, to bind the N content present in the respective steel.
• the %Ti/%N ratio is significantly below this value.
• the %Ti/%N ratio is still below the stoichiometric ratio of 3.43.
• the %Ti/%N ratio for the inventive steels was less than 4.
• all comparative steels 4 - 6 had a %Ti/%N ratio > 5.
• the Al content of the inventive steels 1-3a was increased in order to achieve AlN precipitation through the higher Al contents, i.e., through a higher precipitation pressure, and to avoid BN formation.
• the melts composed according to the invention were cast into slabs in a conventional continuous caster. After being heated to a preheating temperature "VWT" over a holding time "LIZ,” the slabs were first pre-rolled in a conventional hot strip mill to form a pre-strip with a thickness "VBD.” The pre-strips leaving the roughing mill at a temperature "VBT” were then finished hot-rolled in a continuous, conventional hot-rolling process to form hot strip with a hot strip thickness "WBD.” The finished hot-rolled hot strips leaving the hot rolling mill were cooled to a coiler temperature (HT) of less than 650 °C, with a cooling rate (ABK) of at least 50 °C/s being set in the temperature range of 800 - 650 °C.
• HT coiler temperature
• ABSK cooling rate
• the hot strips produced from the inventive steel 1 and the comparative steel 4 were rolled in cold rolling mills to their final thickness "D".
• the degree of cold rolling achieved by cold rolling is not a decisive factor. It is determined solely by the given hot strip thickness and the required cold strip thickness, so that cold rolling can be carried out according to the usual procedure in the prior art. Cold rolling causes the strip to undergo plastic deformation, which results in strong hardening and a reduction in further This results in a loss of formability. Therefore, after cold rolling, a recrystallization annealing is also carried out in a conventional manner, which softens the strip in question and makes it suitable for forming into a component again.
• hot-dip coating is to be applied, as in the example of the cold-rolled strip produced from steel 1, the annealing can be integrated in the usual way into the hot-dip coating process, which is usually carried out in a continuous process. Alternatively, a batch annealing can also be used. Likewise, an electrolytic coating can be carried out instead of hot-dip coating.
• samples of the flat steel products in question were subjected to a simulation of a conventional quenching and tempering or hot forming process.
• the samples were each heated to an austenitizing temperature "T_aust" that was approximately 60 °C higher than the Ac3 temperature of the respective steels 1-6.
• the austenitizing time required for heating and soaking at the austenitizing temperature T_aust was 7-10 minutes, including the heating time in a salt bath furnace.
• the samples were quenched in oil at an average cooling rate of 70-120 °C/s to room temperature.
• the samples were tempered at 170 - 200 °C for a period of 20 minutes. This tempering corresponds to both a tempering
• the heat treatment typically completed after the coating process, as well as the conditions prevailing during cathodic dip coating in a typical automotive painting process, are also common in industrial practice. Temperatures of 150–700 °C combined with holding times of 5–60 minutes are also common for annealing a component.
• the mechanical tensile test parameters "Young's modulus”, “yield strength Rp0.2”, “tensile strength Rm”, “uniform elongation Ag”, and "elongation at break A” were determined on the specimens produced from steels 1 - 6 in the manner described above in accordance with DIN EN ISO 6892-1.
• the elongation at break A refers to specimen shape 2 with cross-sections 20 mm wide and an initial gauge length of 80 mm.
• an initial gauge length of 50 - 65 mm proportional tensile specimens
• the KG grade is calculated by multiplying the mean former austenite grain size by the standard deviation of the diameter of the former austenite grain size.
• the smaller the KG grade the more favorable the effects on toughness and local elongation. Toughness is known to improve with decreasing grain size.
• a smaller grain size dispersion ensures increased homogeneity of the deformation behavior and thus a delayed onset of instability due to fracture reduction, since there are fewer local differences.

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