EP4367276A1 - Bande d'acier à deux phases revêtue à haute résistance et son procédé de production - Google Patents

Bande d'acier à deux phases revêtue à haute résistance et son procédé de production

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Publication number
EP4367276A1
EP4367276A1 EP22747320.4A EP22747320A EP4367276A1 EP 4367276 A1 EP4367276 A1 EP 4367276A1 EP 22747320 A EP22747320 A EP 22747320A EP 4367276 A1 EP4367276 A1 EP 4367276A1
Authority
EP
European Patent Office
Prior art keywords
strip
steel strip
hot
dual phase
optionally
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
EP22747320.4A
Other languages
German (de)
English (en)
Inventor
Joost Willem Hendrik VAN KREVEL
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.)
Tata Steel Ijmuiden BV
Original Assignee
Tata Steel Ijmuiden BV
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 Tata Steel Ijmuiden BV filed Critical Tata Steel Ijmuiden BV
Publication of EP4367276A1 publication Critical patent/EP4367276A1/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
    • 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
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/58—Ferrous alloys, e.g. steel alloys containing chromium with nickel 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
    • 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
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    • 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
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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
    • 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
    • C—CHEMISTRY; METALLURGY
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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
    • C21D8/0278—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips involving a particular surface treatment
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties 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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    • 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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    • 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
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    • 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
    • C21D8/0478—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 involving a particular surface treatment
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    • 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
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    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/001—Ferrous alloys, e.g. steel alloys containing N
    • 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
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/44—Ferrous alloys, e.g. steel alloys containing chromium with nickel 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
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    • C22C38/46—Ferrous alloys, e.g. steel alloys containing chromium with nickel with vanadium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
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    • C22C38/00—Ferrous alloys, e.g. steel alloys
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    • C22C38/48—Ferrous alloys, e.g. steel alloys containing chromium with nickel 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
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    • 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
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    • C23C2/022—Pretreatment of the material to be coated, e.g. for coating on selected surface areas by heating
    • C23C2/0224—Two or more thermal pretreatments
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    • 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
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    • C23C2/06—Zinc or cadmium or alloys based thereon
    • C—CHEMISTRY; METALLURGY
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    • 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
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    • C23C2/12—Aluminium or alloys based thereon
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    • 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
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    • C21D2211/00—Microstructure comprising significant phases
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    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/008—Martensite
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    • C22C38/38—Ferrous alloys, e.g. steel alloys containing chromium with more than 1.5% by weight of manganese

Definitions

  • the invention relates to a high strength coated dual phase steel strip having improved formability, such as used in the automotive industry.
  • the invention also relates to a method for producing such steel strip.
  • (advanced) high strength steel sheets are increasingly used in car components to reduce weight and fuel consumption.
  • a series of (advanced) high strength steels such as HSLA, Dual phase (DP), Ferritic-bainitic (FB) including stretch-flangeable (SF), Complex phase (CP), Transformation-induced plasticity (TRIP), Hot-formed, Twinning-induced plasticity (TWIP) has been developed to meet the growing requirements.
  • AHSS sheet steels cannot be applied easily to a wide variety of car components because their formability is relatively poor.
  • the product is preferrably supplied as a coil to an end customer who cuts blanks and shapes these to end parts.
  • Dual Phase steels offer a good combination of strength and stampability (or pressability) as a result of their microstructure, in which main constituents of a (mixture of) hard martensitic and/or bainitic phase is dispersed in a ductile ferritic matrix. Small amounts of retained austenite (1-3%) and cementite may be present. These steels have high strain hardenability. This gives them not only good strain redistribution capability and thus stampability, but also finished part mechanical properties, including yield strength, that are far superior to those of the flat metal. The yield strength of Dual Phase steels is further increased by the paint baking (also called Bake Hardening, BH) process.
  • paint baking also called Bake Hardening, BH
  • Y minimum yield
  • T minimum tensile
  • Table l Mechanical properties according to VDA 239-100.
  • BH2 (MPa): Increase in yield strength between a reference condition after a 2% plastic pre-strain and the condition obtained after a heat treatment (170°C-20 minutes).
  • nio-20 (-): Strain hardening exponent determined between 10 and 20 % plastic strain or Ag if Ag ⁇ 20 %.
  • n4-6 (-): Strain hardening exponent determined between 4 and 6 % plastic strain.
  • Rpo.2 Yield strength at 0.2 % plastic extension (MPa). This is obtained after temper rolling.
  • Al 0.200 - 0.800 N: at most 0.015 and optionally one or more of the elements selected from:
  • V at most 0.200
  • Ca at most 0.0050 the remainder being iron and unavoidable impurities, wherein the steel has a tensile strength Rm of 580 - 720 MPa and a yield strength Rp of 310 - 430 MPa.
  • the steel strip is optionally hot-dip galvanised, electrogalvanized or electroplated.
  • Dual-phase steels have of a multi-phase microstructure. Martensite islands are embedded in a ferritic matrix, the proportion of which is used to adjust the strength of the respective grade.
  • the microstructure consists mainly of ferrite and martensite and small amounts of bainite and retained austenite.
  • the microstructure and the desired mechanical properties are adjusted by means of the chemical composition and the annealing process optionally combined with hot-dip galvanizing.
  • Dual-phase steels show a very low yield strength to tensile strength ratio (Rp/Rm) combined with high tensile strength and strong work hardening capacity. These dual phase steels are therefore well suited for forming operations with high stretching requirements.
  • Cementite is preferably absent from the microstructure.
  • hot-dip galvanising includes hot-dip aluminizing (hot-dip coating with commercially pure aluminium or an aluminium alloy).
  • the annealed strip may be provided with a metallic coating by electrodeposition.
  • electroplating and electrodeposition are synonyms and are electrolytic processes by which a metallic coating is deposited on another metal.
  • Electrogalvanizing is a species of electrodeposition generally associated with electrodeposition of a zinc or zinc alloy coating on another metal.
  • the use of the main alloying elements C-Mn-Si-AI-B, with the optional addition of optional elements, is a cost effective solution for achieving the grade CR330Y590T-DH as prescribed in the European norm VDA 239-100.
  • Some requirements for this grade according to VDA 239-100 are shown in table 1.
  • the steel according to the invention complies with VDA 239-100.
  • the enhanced formability gives it added value in the case of forming critical parts during cold forming operations such as drawing, while retaining the welding performance equivalent to the conventional dual phase grades. Thanks to their mechanical properties, these grades are suitable for automotive safety parts for crash resistance. By limiting the amount of substitutional elements (such as Mn) good Hole Expansion Coefficient (HEC) values could be obtained.
  • substitutional elements such as Mn
  • HEC Hole Expansion Coefficient
  • This limitation of Mn also has a beneficial effect on the amount of the post-uniform elongation.
  • Strengthening of the steel is further achieved by the addition of boron and silicon.
  • a disadvantage of using silicon is the risk of cracks in spot welds like for example through liquid metal embrittlement (LME) or though thermal shrink pressure and hence the addition of aluminium is necessary to suppress cracking during spot welding. This is quite surprising as aluminium in general is believed to contribute to LME, but in this carefully balanced chemistry the beneficial effect of the aluminium as a ferrite former on the phase transformation prevails.
  • the inventors have found that by a careful selection of the amounts of the main constituting elements of the steel, being carbon, manganese, silicon, aluminium, chromium and boron, a high strength hot dip galvanised steel strip can be produced that has the required formability, homogeneity, processability, strength and elongation, while at the same time providing a sufficient weldability, coatability and surface quality.
  • Carbon is an essential cost-effective alloying element in the steel grade to achieve the strength levels in a conventional continuous annealing/galvanising line. Carbon has to be present in an amount that ensures the required strength and elongation. To ensure hardenability and the resulting formation of martensite at the cooling rates available in a conventional annealing/galvanising line was found to require at least 0.090 wt.%. A maximum level of 0.140 wt.% has been found to be essential to secure sufficient weldability and elongation.
  • the C is at least 0.094 wt.%, more preferably 0.098 wt.%, more preferably 0.104 wt.% and even more preferably at least 0.110 wt.%.
  • the carbon content is preferably at most 0.130 wt.%, and more preferably at most 0.125 wt.%.
  • C is from 0.090 wt.% to 0.130 wt.%.
  • Mn 1.200 - 1.900 wt.%.
  • Manganese improves the hardenability of the steel so that it promotes the formation of bainite and martensite. Manganese strengthens the ferritic matrix by solid solution strengthening and lowers the transformation temperature range of the steel, thus lowering the required annealing temperature.
  • Mn is known to adversely affect the HEC-values and through coil ductility variation as well as the rolling forces during hot- and cold-rolling and therefore Mn-values should be below 1.900 wt.% and preferably below 1.800 wt.%.
  • a suitable maximum value for Mn is 1.750 wt.%.
  • Mn values below 1.700 wt.% or even below 1.600 wt.% further improve HEC-values and ductility. These maximum levels are also given in view of the stronger segregation during casting and centre line segregation in the strip at higher values.
  • the product should further be processable in continuous annealing lines with both RTF heating and faster flame heating.
  • the minimum Mn value is 1.200 wt.%, preferably at least 1.300 wt.%, and more preferably at least 1.400 wt.% and most preferably at least 1450 wt.%.
  • Mn-level is too low then the required strength level is not achieved while DP properties with low Rp/Rm ratio are not achieved at lower values. This lower limit is possible in view of the addition of other elements, such as boron. These values are reached more easily at higher values of Mn, but the costs and elongation values deteriorate, so that a proper balance must be chosen.
  • Mn is from 1.400 wt.% to 1.750 wt.%.
  • Si 0.200 - 0.800 wt.%. Silicon suppresses cementite formation in the end product and stabilizes retained austenite formation thereby improving the ductility of the product. Silicon effectively reinforces the steel matrix by solid solution to achieve the required strength level and increases the work-hardening rate of the ferrite component, which improves local formability (HEC) and a high total elongation.
  • the maximum Si- content in the steel according to the invention is 0.800 wt.%. It was however observed that Si adversely affects the spot-weldability and may be prone to liquid metal embrittlement of the coated strip and therefore the maximum is preferably limited to 0.750 or even 0.700 wt.% or most preferably up to 0.650 wt.%.
  • an excess silicon may lead to a too high degree of internal and grain boundary oxidation during coil cooling. This may be circumvented by using a very low coiling temperature and batch annealing but it is outside the scope of a cost-effective solution in terms of process steps.
  • To achieve a minimum strengthening effect at least 0.200 wt.% is required, and preferably at least 0.250 wt.%.
  • a minimum Si-level of at least 0.400 wt.% and even more preferably at least 0.450 wt.% may be preferred.
  • Si is from 0.500 wt.% to 0.700 wt.%.
  • Al 0.200 - 0.800 wt.%. Aluminium is primarily added to liquid steel for the purpose of de-oxidation in relatively small amounts. The excess Al also reacts with nitrogen and forms AIN precipitates which hence prevents boron from reacting with nitrogen. As a result B segregates on the grain boundaries and hence acts as a hardenability element. To achieve at least the binding of free oxygen and nitrogen, a lower level of Al of at least 0.060 wt.% is proposed.
  • Al slows bainite transformation, thus determining the optimum bainite formation within the time constraints imposed by the annealing section of a continuous annealing/galvanising line.
  • the Al is given as the total aluminium content. Aluminium also retards the formation of carbides thus keeping carbon in solution, which in turn causes partitioning of the carbon to austenite during overaging and promotes the stability of austenite. Unlike Silicon, Al hardly increases the partitioning time for austenite carbon enrichment in the continuous annealing line overage section. This improves the elongation values and the hardenability and also allows a minimum or no overage time in the continuous annealing line while still achieving the required. A maximum level of 0.800 wt.% is imposed for castability, since high aluminium contents lead to poisoning of the casting mould slag and consequently an increase in mould slag viscosity, leading to incorrect heat transfer and lubrication during casting.
  • aluminium segregates resulting in carbon, silicon and manganese enriched segregation bands with intermediary lower Mn high aluminium rich segregation bands. Both effects increase the elongation and the hardenability value (n4_6 (which is the strain-hardening exponent in the strain range of 4 to 6%) and n-values (which is the strain-hardening coefficient)) of the end product while keeping the yields stress low.
  • n4_6 which is the strain-hardening exponent in the strain range of 4 to 6%
  • n-values which is the strain-hardening coefficient
  • aluminium affects intercritical annealing and the phase distribution during intercritical annealing.
  • This can be understood by aluminium being a ferrite former.
  • Al is preferably above 0.200 wt.% and more preferably a minimum Al content of 0.25 wt.%.
  • the aluminium content is at most 0.750 wt.%, more preferably at most 0.650 wt.%, and even more preferably at most 0.400 wt.%.
  • Al is from 0.200 wt.% to 0.400 wt.%.
  • B 0.0010 - 0.0050 wt.% (5-50 ppm by weight).
  • Boron is added as a hardenability element to achieve the required strength level.
  • the addition of boron allows a saving on the use of more expensive alloying elements. Compared to more expensive solid solution alloying elements, boron affects the ability to recrystallisation much less. This allows for addition static or dynamic recrystallisation in the hot rolling line, which reduces mean flow stress build up. It also increases the cold rolling recrystallisation capability during the heating section of the continuous annealing thus allowing sufficient grain growth and achieving the DH600 properties both in a direct firing and slow tube heating annealing line.
  • the boron amount is 0.0010 - 0.0050 wt.% (10-50 ppm).
  • a suitable lower boundary level is 0.0012 wt.% or even 0.0015 wt.%. Above a certain boron level zinc coating is adversely affected while hardenability reduces.
  • a suitable upper boundary level is 0.0040 wt.%, preferably 0.0035 wt.%, more preferably 0.0030 or even 0.0025 wt.%.
  • B is from 0.0015 to 0.0025 wt.%.
  • N at most 0.015 wt.% (150 ppm). Nitrogen is present during casting. Free nitrogen in the steel should be avoided as it can lead to severe ageing and affects the capability of boron to function as a hardenability element. However, in the chemical composition of the steel according to the invention there is a surplus of aluminium and boron and consequently all nitrogen is bound to Al or B as AIN or BN. The N-content is limited to 0.015 wt.%, typical for BOS-steelmaking and continuous casting plants. However, it is preferably to reduce the nitrogen to a value of at most 0.012 wt.%, preferably to at most 0.010 wt.%.
  • a more preferred maximum level is 0.008 wt.% or even 0.006 wt.%. It is not economic to remove all nitrogen from the steel and therefore in a preferable embodiment N is from 0.0001 wt.% to 0.006 wt.% (10 to 60 ppm).
  • P at most 0.050.
  • Phosphorus interferes with the formation of carbides, and therefore some phosphorus in the steel can be advantageous.
  • phosphorus positively affects picklability and strengthens the steel.
  • a maximum of 0.050 wt.% is allowable.
  • phosphorus can make steel brittle upon welding, so the amount of phosphorus is preferably limited to 0.040 wt.% to maintain sufficient hot ductility and to avoid failure by peel during tension-shear tests carried out on spot-welded assemblies. More preferably the maximum level is of P is 0.030 wt.%, even more preferably 0.020 mass %. In a preferable embodiment P is at most 0.015 wt.%.
  • S at most 0.050 wt.%.
  • Sulphur (S) is maximised at a value of 0.050 wt.% and preferably maximised at a value of 0.020 wt.% or even at a value of 0.010 wt.%. Even though sulphur may somewhat improve picklability, sulphur is preferably avoided completely but is inevitably present as a result of the steelmaking process and the raw materials used therein. Sulphur precipitates in the form of manganese and/or calcium sulphides that greatly reduce the formability.
  • the amount of sulphur is 0.0001 - 0.005 wt.%, more preferably at most 0.003 wt.% or more preferably at most 0.002 wt.%.
  • a suitable minimum amount is 0.0002 wt.%.
  • S is from 0.0001 wt.% to 0.002 wt.%.
  • one or more elements selected from Ti, V, Cr, Mo, Nb, Ni, Cu, Ca, may be added to the steel composition or be present as impurities.
  • the addition of Calcium (Ca) modifies the morphology of manganese sulphide inclusions.
  • Elongated inclusions also called stringers, may act as planes of weakness along which lamellar tearing and delamination fracture can occur.
  • stringers may act as planes of weakness along which lamellar tearing and delamination fracture can occur.
  • the avoidance of stringers is beneficial for forming processes of steel sheets which entail the expansion of holes or the stretching of flanges and promotes isotropic forming behaviour.
  • Calcium treatment also prevents the formation of hard, angular, abrasive alumina inclusions in aluminium deoxidised steel types, forming instead calcium aluminate inclusions which are softer and globular at rolling temperatures, thereby improving the material's processing characteristics.
  • REM Rare Earth Metals
  • Chromium (Cr) is added to increase hardenability. It promotes formation of ferrite. A maximum level of 0.500 wt.% is imposed to ensure that not too much martensite forms at the cost of retained austenite. It is also possible to add no Cr, in which case it is absent completely or present at most at a residual element or impurity level. Cr is quite expensive. From that perspective the amount of Cr is preferably at most 0.300 wt.%, more preferably at most 0.250 wt.% and even more preferably at most 0.200 wt.%. From a mechanical property point of view, if Cr is added as an alloying element, the preferable range is between 0.020 and 0.400 wt.%.
  • the desired mechanical properties and costs is obtained when Cr is between 0.020 and 0.150 wt.%, and preferably between 0.020 and 0.120 wt.%. If added as an alloying element, a suitable minimum amount of Cr is 0.030 wt.%, or even 0.040 wt.%. Values below 0.020 are deemed to be at residual element level. When an element is at residual element level, the elements are deemed not to affect the properties or processability of the steel in a significant way, and in these cases the costs of removing these elements further exceed the expected benefits of a further reduction, even if it were technically feasible to reduce the level further.
  • Titanium (Ti) is mainly added to strengthen the steel, and the allowable level is at most 0.060 wt.%. It is also possible to add no Ti, in which case it is absent completely or present at most at a residual element or impurity level. If added, then preferably the amount of Ti is between 0.010 and 0.060 wt.%, more preferably between 0.02 and 0.050 wt.%. Values below 0.010 wt.% are deemed to be at residual element level.
  • Vanadium may be added up to 0.200 wt.%. If added a preferred range is 0.005 - 0.200 wt.% When the vanadium content is less than 0.005% by weight vanadium precipitation strengthening is insufficient. When the vanadium content is greater than 0.200 wt.% the precipitation occurs at an early stage during or after the hot-rolling in the form of fine precipitates, which reduces the dimensional window or coarsen during continuous annealing which reduces precipitate strengthening.
  • the vanadium amount is 0.005 - 0.200 wt.%, preferably 0.010 - 0.200 wt.%, more preferably 0.030 - 0.200 wt.%, most preferably 0.040 - 0.150 wt.%. It is also possible to add no V, in which case it is absent completely or present at most at a residual element or impurity level. Values below 0.010 wt.% are deemed to be at residual element level.
  • elements such as molybdenum, which retards the bainitic transformation and promotes solid-solution hardening, may be added in amounts not exceeding 0.40 wt.%.
  • the amount of Mo is preferably 0.005 - 0.20 wt.%, more preferably 0.005-0.10 wt.%, so as to limit the cost of the steel and keep the dimensional window as large as possible.
  • Mo can be added to improve to strength and improves the quality of a zinc-based coating. Mo also helps to strengthen the steel through precipitate formation. However, it is also possible to add no Mo, in which case it is absent completely or present at most at a residual element or impurity level. Values below 0.015 wt.% are deemed to be at residual element level.
  • Niobium (Nb) may be added in an amount of preferably 0.001 - 0.060 wt.%, more preferably 0.001 - 0.050 wt.%, and most preferably 0.001 - 0.030 wt.%.
  • the addition of Nb increases the strength by complementary carbonitride precipitation but adds to hot-rolling forces which reduce the dimensional window. If added, then preferably the amount of Nb is between 0.010 and 0.060 wt.%, more at most 0.050, or at most 0.030 wt.%. Values below 0.010 wt.% are deemed to be at residual element level.
  • a dual phase steel strip according to any one of the preceding claims, wherein one or more or all of Cr, Mo, V, Nb and Ti is present only as an impurity this means that, if applicable, the levels of one or more or all of Cr is below 0.020 wt.%, Ti is below 0.010 wt.%, V is below 0.010 wt.%, Mo is below 0.015 wt.% and Nb is below 0.010 wt.% respectively.
  • Nickel and copper are not added to the steel according to the invention and these elements are either absent completely or present at most at a residual element or impurity level. Values below 0.060 wt. % are considered to be at residual element level.
  • the dual phase steel strip according the invention has a tensile strength Rm of 590 to 720 MPa and/or a yield strength Rp of 320 to 430 MPa.
  • the steel according to the invention has a yield strength Rp of at least 330 MPa.
  • the steel according to the invention has a tensile strength Rm of at most 700 MPa and a yield strength Rp of at most 430 MPa.
  • a method for producing a coated and optionally temper-rolled dual phase steel strip with a tensile strength Rm of 580 - 720 MPa and a yield strength Rp of 310 - 430 MPa comprising the following steps:
  • V at most 0.200
  • Ca at most 0.0050 the remainder being iron and unavoidable impurities, to a hot-rolled strip having a thickness of 2.0 - 4.5 mm, wherein finish-rolling is performed while the strip has an austenitic microstructure;
  • the strip is a), provided with a metallic coating by hot-dip coating between step iv. and v., optionally followed by galvannealing the hot-dip coated strip, or b). with a metallic coating by electrodeposition after step v.; • Optionally temper rolling or tension levelling of the coated steel strip with a reduction of at most 0.70%;
  • the coated and optionally temper-rolled steel strip has a tensile strength Rm of 580 - 720 MPa and a yield strength Rp of 310 - 430 MPa; ⁇ Coiling the coated steel strip or cutting the coated steel strip into sheets or blanks;
  • the steel melt is preferably produced in a BOS-process (Basic Oxygen
  • unavoidable impurities and residual elements mean the same in that the level of unavoidable impurities and the level of residual elements are determined by the technical or economic incapability to lower the level of an element below these levels.
  • the steel melt is continuously cast into slabs on a conventional caster. After casting the strip cools in the continuous caster. During cooling of the continuous cast slab goes to a specific temperature range while being bent in which the ductility of the steel is reduced which is known as the temperature ductility trough. In this temperature range precipitates may form.
  • One of the precipitate forming element is Al, which forms AIN at grain boundaries. This should partly be suppressed as it may lead to sliver formation later in the process.
  • B AIN is partially suppressed as BN forms preferentially.
  • the cast slab may be a thick slab (150-350 mm thick) or a thin slab (50 - 150 mm). Inhomogeneous excessive grain growth can be suppressed by the formation of BN and/or AIN (mixed) precipitates. Part of the BN precipitates are dissolved while AIN precipitates form.
  • Hot-rolling Subsequently the slab is hot-rolled.
  • the thick slabs are hot-rolled in a conventional hot-strip-mill (HSM).
  • the thin slabs are usually directly hot-rolled after casting in a thin slab caster and (direct) rolling facility (TSCR)
  • TSCR direct rolling facility
  • TSCR direct rolling facility
  • the hot-rolling reduction is performed above the non-recrystallisation temperature (TNR) and that full recrystallisation occurs in the first stand of the finishing mill.
  • TNR non-recrystallisation temperature
  • at least 40% and more preferably at least 50% of the hot-rolling reduction (transfer bar thickness to hot-rolled strip thickness) in the tandem finishing mill is performed above TNR.
  • a higher reduction above TNR leads to a finer austenite structure due to repeated static or dynamic recrystallisation and thereby to a finer final structure of the hot-rolled coil with a more evenly distributed fraction of retained austenite in the ferrite matrix.
  • the equation used is (compositions in wt.%):
  • TNR is preferably lower than 1050 °C and the finish rolling temperature (FRT) is preferably at least 880 °C.
  • the steel must still be in its austenitic state (i.e. above Ac3) during the last hot-rolling reduction step to promote the r-value and hence local ductility by optimizing the final ferrite crystallographic grain orientation.
  • the coiling temperature specified in this invention relates to the aimed coiling temperature of the body of the strip. Due to the effects of cooling of the coiled hot-rolled strip the head and tail of the strip cools faster on the coil than the body of the strip. The head and tail form the inner and outer wraps of the coil. To compensate for the faster cooling of the head and tail of the hot-rolled strip it may be preferable to aim for a higher coiling temperature of the head and tail by employing a so called U-type cooling pattern.
  • the body coiling temperature after hot-rolling is preferably below 610 °C to avoid surface grain boundary oxidation and preferably below 600°C, or even below 590 or 580 °C.
  • the lower the coiling temperature the lower the risk of surface grain boundary oxidation.
  • pearlite/cementite starts to form in band-like morphology and this is also preferably avoided.
  • the mid-coil strength shows a considerable drop and the alloy composition may have the risk of not achieving the right strength at mid-coil position.
  • too low a coiling temperature leads to hard edges and a reduction of elongation at the coil head and tail edges.
  • the body coiling temperature is 500 °C or higher and preferably 520 °C or higher. Any adverse tail and head effects may be compensated by a suitably chosen higher head- and or tail coiling temperature.
  • the strip is pickled and subsequently cold-rolled. Due to the alloying elements in solid solution the rolling forces are considerable.
  • the cold-rolling reduction is therefore at most 80 %, preferably at most 75 % and more preferably at most 70 % and even more preferably at most 65%. Higher reductions lead to high rolling forces, and thus to a higher risk of shape defects and a smaller dimensional window. A smaller dimensional window means a lower width to thickness ratio while the mechanical properties of the steel over the width of the strip remain suitable.
  • the cold rolling reduction also influences the final microstructure after annealing. If the cold rolling reduction exceeds a certain (composition dependent) threshold the microstructure is deformed to such an extent that the microstructure loses its potential to form retained austenite during the annealing treatment. This small amount of retained austenite is crucial for obtaining the desired formability and hence the cold rolling reduction is critical. In addition, high cold rolling reductions result in smaller ferrite grains which may negatively affect n4_6 values.
  • the continuous annealing step is a last and crucial step in obtaining the desired final properties.
  • the cold rolled strip is heated at an average heating rate of between 3-30 °C/s in a Non-Oxidising Furnace (NOF), a direct fired furnace (DFF), an induction furnace, a radiation tube furnace (RTF) or heating by hot gas.
  • NOF Non-Oxidising Furnace
  • DFF direct fired furnace
  • RTF radiation tube furnace
  • heating it is important that overlap between recrystallisation of the cold-rolled microstructure and the austenite formation is low. At least a substantial part and preferably all of the deformed microstructure has recrystallised before the transformation to austenite starts.
  • the un recrystallised ferrite grains recover prior to recrystallisation, and carbon in these recovered grains tends to diffuse to the low-angle grain boundaries in the recovered grains.
  • the recrystallised grains do not "cling" to its carbon, and the carbon diffuses to the austenite that has formed and due to the lower carbon content of the ferrite, the recrystallised grains tend to transform before the unrecrystallised grains.
  • the annealing cycle is an intercritical annealing cycle, this means that the last remaining ferrite is the unrecrystallised ferrite with the higher carbon content.
  • This carbon does not contribute to the formation of martensite during cooling after annealing, and the carbon does also not help in stabilising retained austenite, which is deemed essential for the formability.
  • the unrecrystallised ferrite moreover adversely affects the formability.
  • the strip surface is pre-oxidised in the section of the furnace to allow the formation of an oxide layer on the surface with sufficient adhesion which is suitable for galvanizing.
  • the composition of the steels according to the invention suffer from poor wettability of the molten metal during hot-dip galvanising. By oxidising the surface with e.g. steam, HNX, water vapour or by a controlled dew point.
  • the iron oxide film also contains some manganese oxide.
  • the other oxides from the alloying elements such as Si and Al, which are not soluble in iron/manganese oxides, are rejected at the steel/oxide interface.
  • the oxidised alloying elements do not totally cover the surface but are located in nodules, which enables better wettability and coating adhesion.
  • the maximum annealing temperature (T_top (T2)) is chosen between (Acl + 50 °C) to (Ac3 - 30 °C).
  • the phase transformation temperatures during reheating (Acl and Ac3) and during cooling (e.g. bainite and martensite transformation temperatures) can be simply determined by means of simulating the thermal profile as schematically depicted in Figures 5 to 8 in a dilatometer.
  • a preferred temperature range for T_top (T2) is between 790 °C and 880 °C.
  • the strip may be held at T_top (T2) for a soaking period of up to 300 seconds. Preferably the soaking period is at most 90 s.
  • T_top may be a peak temperature wherein the temperature is held above 790 °C for a soaking period of up to 300 seconds and wherein the temperature is slowly raised to T_top and slowly decreased from T_top (T2). After the soaking period the strip is slowly cooled to T_slow (T3) which is between 9 °C to 100 °C lower than the top temperature.
  • T_slow T3 (cooling from T2 to T3 is the slow cooling section) is higher than 650 °C so that enough austenite remains which can be transformed to bainite, martensite or further enriched in retained austenite.
  • the cooling rate after the slow cooling section (CR2) is between 5 and 70 °C/s and preferably at least 30 °C/sec so that austenite is prevented from forming carbides.
  • a direct cooling from T_slow (T3) to the molten metal bath may be used (see Figure 8), followed by hot-dip galvanizing, a gas jet wiper and cooling to ambient temperature.
  • the hot-dip galvanised strip may optionally be annealed prior to cooling to ambient temperature (galvannealing).
  • the galvanizing may be followed by an annealing treatment, known as galvannealing.
  • Overageing may have the advantage that additional bainite is formed and that the existing microstructure is tempered. Carbon partitioning may take place for retained austenite formation. Also hydrogen dissolved in the strip may be released, at least partly and strip tensions may also be partly relieved.
  • the strip is cooled from T-slow (T3) to and held at an overageing temperature (T_oa) at a cooling rate (CR2) of at least 15 °C/s between 370 and 470 °C for an overageing period of up to 50 seconds, subsequently optionally heated or cooled to the molten metal bath temperature, hot-dip galvanised by dipping the strip in the molten metal bath (such as depicted in Figure 5, 6 and 7). After dipping the galvanised strip is cooled to ambient temperature or optionally annealed and cooled to ambient temperature.
  • T3 T-slow
  • T_oa overageing temperature
  • CR2 cooling rate
  • the molten metal bath may contain molten zinc, a molten zinc-alloy, molten aluminium or a molten aluminium alloy.
  • the zinc alloy may comprise 0.3-4.0 wt.% Mg and 0.3-6.0 wt.% Al; optionally at most 0.2 wt.% of one or more additional elements, unavoidable impurities; the remainder being zinc.
  • the alloying element contents in the zinc-alloy coating layer shall be 1.0 - 2.0 % Mg and 1.0 -3.0 % Al, optionally at most 0.2% of one or more additional elements, unavoidable impurities and the remainder being zinc.
  • the zinc alloy coating comprises at most 1.6% Mg and between 1.6 and 2.5% Al, optionally at most 0.2% of one or more additional elements, unavoidable impurities and the remainder being zinc.
  • the steel strip, sheet or blank is provided with a (commercially pure) aluminium layer or an aluminium alloy layer.
  • a typical metal bath for a hot dip coating such an aluminium layer comprises of aluminium alloyed with silicon e.g. aluminium alloyed with 8 to 11 wt.% of silicon and at most 4 wt.% of iron, optionally at most 0.2 wt.% of one or more additional elements such as calcium, unavoidable impurities, the remainder being aluminium.
  • Silicon is present in order to prevent the formation of a thick iron-intermetallic layer which reduces adherence and formability.
  • Iron is preferably present in amounts between 1 and 4 wt.%, more preferably at least 2 wt.%.
  • the annealed strip can also be electrogalvanized with a zinc or zinc alloy coating instead of hot-dip galvanising.
  • the annealed strip may be electroplated with nickel, cobalt or chrome and combinations thereof.
  • the coated steel strip may optionally be subjected to temper rolling or tension levelling with a reduction of the coated steel strips thickness of at most 0.70%. Preferably the reduction is at most 0.45%. Preferably the reduction is at least 0.05%. The maximum of the reduction is governed by the need to secure a sufficiently high value of n4_6. A minimum reduction may be used to improve the surface texture of the coated strip and/or to suppress the yield point elongation (YPE) and/or to increase the yield strength (Rp) of the material.
  • YPE yield point elongation
  • Rp yield strength
  • Table 2 the composition of the comparative (C) and inventive (I) steels is shown.
  • Table 3 to 5 provide the process settings for the experiments. It is noted that table 4 relates to industrial scale trials and table 5 to laboratory scale trials.
  • Table 2 Comparative (C) and Inventive (I) examples (in m%, B and N in ppm).
  • Table 3a Hot rolling and cold rolling process setting of the lab cast alloys.
  • Table 3b Hot rolling and cold rolling process setting of the line-trialed alloys
  • Table 6 and 7 give the results of the industrial and laboratory scale trials. It should be noted that the results of the inventive examples all meet the requirements of the VDA-239 as presented in table 1.
  • the values for Rp and Rm meet the requirements of VDA-239 of Rp between 330 and 430 MPa and Rm between 590 and 700 MPa
  • HEC -2— - - x 100% d 0
  • Stretch flangeability is evaluated on the basis of the maximum HEC-value and is deemed satisfactory when HEC > 35%
  • Figure 1 shows a schematic diagram of a continuous annealing and hot-dip coating line equipped with a pot to hold a molten metal (in this case zinc).
  • a molten metal in this case zinc.
  • the location indicated with the asterisk is the location where the optional galvannealing section is to be found.
  • the hot-dip coating is optional, so the zinc pot can be bypassed if so desired.
  • FIG. 2 shows a schematic diagram of a galvannealing section located between the zinc pot and the turning roll at the top.
  • the galvannealing section consists of a heating section, a holding section and a cooling section.
  • Figure 3 shows a comparison of the results of a tensile test where the sample is taken under 0° (i.e. longitudinal), under 90° (transversal) and under 45° to the rolling direction of the strip.
  • the Rm and Rp are relatively independent of the direction, but the total strain, uniform strain are not.
  • Figure 4 shows a microstructure of alloy 4 (L) at 1 ⁇ 4t.
  • the microstructure consists of over 90% of ferrite (F) > 90% (light grey matrix), 4-6% retained austenite (off- white), some martensite (dark grey/black), some bainite (grey) and the remainder (if any) is cementite.
  • Figure 5 to 8 show various (non-limitative) options for annealing the steels according to the invention:
  • Figure 9 shows a schematic representation of the U-type coiling temperature profile of a hot-rolled strip prior to coiling.
  • Table 6 A80 tensile properties and HEC values - Results from line trials.

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  • Chemical & Material Sciences (AREA)
  • Engineering & Computer Science (AREA)
  • Materials Engineering (AREA)
  • Mechanical Engineering (AREA)
  • Metallurgy (AREA)
  • Organic Chemistry (AREA)
  • Physics & Mathematics (AREA)
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  • Crystallography & Structural Chemistry (AREA)
  • Chemical Kinetics & Catalysis (AREA)
  • Oil, Petroleum & Natural Gas (AREA)
  • Heat Treatment Of Sheet Steel (AREA)

Abstract

L'invention concerne une bande d'acier à deux phases revêtue et facultativement laminée à froid avec une résistance à la traction Rm de 580 à 720 MPa et une limite d'élasticité Rp de 310 à 430 MPa, ayant une aptitude améliorée au formage. L'invention concerne également un procédé de fabrication d'une telle bande d'acier et son utilisation.
EP22747320.4A 2021-07-07 2022-07-07 Bande d'acier à deux phases revêtue à haute résistance et son procédé de production Pending EP4367276A1 (fr)

Applications Claiming Priority (2)

Application Number Priority Date Filing Date Title
EP21184350 2021-07-07
PCT/EP2022/069016 WO2023281035A1 (fr) 2021-07-07 2022-07-07 Bande d'acier à deux phases revêtue à haute résistance et son procédé de production

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EP4367276A1 true EP4367276A1 (fr) 2024-05-15

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EP (1) EP4367276A1 (fr)
KR (1) KR20240032894A (fr)
CN (1) CN117677717A (fr)
BR (1) BR112023027416A2 (fr)
WO (1) WO2023281035A1 (fr)

Family Cites Families (4)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US20140205858A1 (en) * 2011-09-13 2014-07-24 Tata Steel Ijmuiden B.V. High strength hot dip galvanised steel strip
KR20210096595A (ko) * 2018-11-29 2021-08-05 타타 스틸 네덜란드 테크날러지 베.뷔. 우수한 딥 드로잉성을 가진 고강도 강 스트립을 제조하는 방법과 그에 따라 제조된 고강도 강
JP6897874B2 (ja) * 2019-01-09 2021-07-07 Ｊｆｅスチール株式会社 高強度冷延鋼板及びその製造方法
CN112301293B (zh) * 2019-07-26 2022-06-28 宝山钢铁股份有限公司 一种冷轧热镀锌钢及其制造方法

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CN117677717A (zh) 2024-03-08
WO2023281035A1 (fr) 2023-01-12
BR112023027416A2 (pt) 2024-03-12

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