US20140034196A1 - Hoherfester mehrphasenstahl mit ausgezeichneten umformeigenschaften high strength multi-phase steel having excellent forming properties - Google Patents

Hoherfester mehrphasenstahl mit ausgezeichneten umformeigenschaften high strength multi-phase steel having excellent forming properties Download PDF

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US20140034196A1
US20140034196A1 US13/981,870 US201113981870A US2014034196A1 US 20140034196 A1 US20140034196 A1 US 20140034196A1 US 201113981870 A US201113981870 A US 201113981870A US 2014034196 A1 US2014034196 A1 US 2014034196A1
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steel
strip
content
cooling
steel strip
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Andreas Wedemeier
Thomas Schulz
Michael Pohl
Phillip Wüllner
Jörg Heinecke
Christian Schlegel
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Salzgitter Flachstahl GmbH
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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
    • C22C38/001—Ferrous alloys, e.g. steel alloys containing N
    • 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
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
    • C21D1/18—Hardening; Quenching with or without subsequent tempering
    • C21D1/185—Hardening; Quenching with or without subsequent tempering from an intercritical temperature
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0221—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • C21D8/0226—Hot rolling
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0221—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • C21D8/0236—Cold rolling
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0247—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C1/00—Making non-ferrous alloys
    • C22C1/02—Making non-ferrous alloys by melting
    • 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
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/22—Ferrous alloys, e.g. steel alloys containing chromium with molybdenum or tungsten
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/26—Ferrous alloys, e.g. steel alloys containing chromium with niobium or tantalum
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/28—Ferrous alloys, e.g. steel alloys containing chromium with titanium or zirconium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/32—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
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/002—Bainite
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/005—Ferrite

Definitions

  • the invention relates to a high-strength multiphase steel with dual bainite or complex phase microstructure and excellent forming characteristics in particular for the lightweight vehicle construction and a method for producing a steel flat product according to the preamble of patent claim 1 .
  • the invention further relates to a method for producing such a steel according to claim 14 .
  • the hotly contested automobile market forces manufacturers to constantly seek solutions for lowering the fleet consumption while at the same time maintaining a greatest possible comfort and occupant protection.
  • weight reduction of all vehicle components plays an important role, but on the other hand also an optimal behavior of the individual components in case of high static and dynamic stress during operation and also in the case of a crash.
  • the distributors of raw material provide high strength and ultra-high-strength steels to thereby allow reducing the sheet thickness of the vehicles at improved forming and component behavior during manufacture and operation. These steels therefore have to meet relatively high standards regarding their strengths and ductility, energy absorption and during their processing such as for example during coaching, hot and cold forming, welding and/or surface finishing (for example metallic finishing, organically coating).
  • Newly developed steels thus besides satisfying the required weight reduction have to satisfy the high material demands regarding stretch limit, tensile strength, and elongation at break at good formability as well as the components requirements for high tenacity edge break resistance, energy absorption and strength via the Work-Hardening-Effect and the Bake-Hardening-Effect.
  • dual-phase steels are increasingly used in vehicle construction, which consist of a ferritic base structure in which a martensitic second phase and possibly a further phase with bainite and residual austenite is incorporated.
  • the processing properties of the dual phase steels which determine the steel types such as a very low stretch limit ratio at very high tensile strength, a strong strain hardening and a good cold formability are sufficiently known.
  • multiphase steels such as complex phase steels, ferritic-bainitic steels, bainitic steels and martensitic steels which are characterized by different microstructure compositions.
  • Complex phase steels of the hot or cold rolled type are steels which contain small proportions of martensite, residual austenite and/or perlite in a ferritic/bainitic base structure, wherein an extreme grain refinement as a result a delayed recrystallization or of precipitation of micro alloying elements.
  • Ferritic-bainitic steels in hot-rolled configuration are steels, which contain bainite or hardened bainite in a matrix of ferrite and/or hardened ferrite.
  • the hardening of the matrix is caused by a high dislocation density, by grain refinement and by the precipitation of micro-alloying elements.
  • Bainitic steels in hot-rolled or cold-rolled configuration are steels, which are characterized by a very high stretch limit and tensile strength at sufficiently high stretching for cold forming processes.
  • the chemical composition results in a good weldability.
  • the microstructure typically consists of bainite. In some cases, smaller proportions of other phases such as martensite and ferrite can be contained.
  • Mastensitic steels in hot-rolled configuration are steels, which as a result of thermo-mechanical rolling contain small proportions of ferrite and/or bainite in a basic structure of martensite. These steel types are characterized by a very high stretch limit and tensile strength at sufficiently high stretching for cold forming processes. Within the group of multiphase steels, the martensitic steels have the highest tensile strengths.
  • TRB®'s With the currently known alloys and available continuous annealing systems, the production of TRB®'s is not possible without limitations for strongly varying sheet thicknesses with multiphase microstructure, such as heat treatment before cold-rolling. Due to the temperature gradient occurring in the conventional process windows in regions of different sheet thickness, no uniform multiphase microstructure can be established in cold-rolled and hot rolled steel strips.
  • cold-rolled steel strips are usually re-crystallizingly annealed in the continuous annealing method to well formable fine steel sheet.
  • process parameters such as throughput speed, annealing temperatures and cooling speed are adjusted according to the required mechanical—technological properties with the microstructure required therefore.
  • the hot or cold strip is heated in the continuous annealing furnace to such a temperature, that the required microstructure is formed during the cooling.
  • the hot or cold strip is heated in the continuous annealing furnace to such a temperature, that the required microstructure is formed during the cooling.
  • the same is true for adjusting a steel with complex-phase microstructure, martensitic, ferritic-bainitic and purely bainitic microstructure.
  • the annealing usually occurs in a continuous annealing furnace located upstream of the galvanizing bath.
  • the required microstructure is only adjusted in the annealing in the continuous furnace depending on the alloying concept in order to realize the required mechanical properties.
  • a narrow process window means here that depending on the cross-section of the strip to be annealed, the process parameters have to be adjusted in order to achieve the required microstructure and the mechanical—technological properties through homogenous temperature distribution in the strip and during cooling. Increased process windows also enable the required strip properties at different cross-sections of the strips to be annealed while using same process parameters.
  • the deciding process parameter is thus the adjustment of the speed during continuous annealing because the phase transition occurs in dependence on temperature and time.
  • a method for producing a steel strip with different thickness over the strip length is described for example in DE 100-378-67A1.
  • the regions with small sheet thickness have either excessive strengths due to great martensite proportions as a result of the transformation processes during cooling, or the regions with greater sheet thickness attain too low strengths as a result of insufficient martensite proportions.
  • Homogenous mechanical—technological properties over the strip length or strip width cannot practically be achieved with the known alloy concepts in the continuous annealing.
  • the goal to achieve the final mechanical—technological properties in a narrow region over strip width and strip length through the controlled adjustment of the volume proportions of the microstructure phases has the highest priority and is therefore only possible through an increased process window.
  • the known alloy concepts for multiphase steels are characterized by a too narrow process window and are therefore not suited for solving the present problem in particular in the case of flexibly rolled strips. With the known alloy concepts only steels of one strength class with defined cross-sectional regions can be produced so that for different strength classes and or cross-sectional regions different alloy concepts are required.
  • the invention is therefore based on the object, to propose a different alloy concept for a high-strength multiphase steel with a broad spectrum of different microstructure compositions with which the process window for the continuous annealing of hot or cold strips can be increased so that besides strips with different cross-sections also steel strips having varying thickness over strip length and strip width can be produced with mechanical—technological properties that are as homogenous as possible. Further, an alloy concept is proposed with which also allows meeting the demand for different strength classes. In addition, a method for producing a strip that is produced from this steel is proposed.
  • this object is solved by a steel with the following contents in weight %:
  • Remainder iron including common steel tramp elements not mentioned above.
  • the steel according to the invention has the advantage of a significantly increased process window compared to the known steels. This results in an increased process reliability in the continuous annealing of cold and hot strip with multiphase microstructure. Thus, more homogenous mechanical—technological properties can be ensured in the strip for continuously annealed hot or cold strips even in the case of different cross-sections and otherwise same process parameters.
  • stress-optimized components can be advantageously produced from these materials by forming.
  • the produced material can be produced as cold strip as well as hot strip via a galvanizing line or a pure continuous annealing system in the dressed and undressed state (intermediate annealing).
  • the steel strips produced with the alloy composition according to the invention are characterized by a significantly broader process window with regard to temperature and throughput speed when manufacturing a multiphase or bainitic steel in the inter-critical annealing between A C1 and A C3 or in an austenizing annealing over A C3 with final controlled cooling compared to the known alloy concepts.
  • Annealing temperatures of 700 to 950° C. and cooling speeds of 15 to 100° C./s to a temperature of 420 to 470° C. with holding at an intermediate temperature of 200 to 250° C. and also with prior optional reheating have proven advantageous, with which the required multiphase microstructures can be evenly adjusted over the strip length. This is particularly advantageous when annealing flexibly rolled strips or in subsequent annealing of strips with different cross sections, which thus allows achieving very uniform material properties.
  • the prerequisite for achieving a broad process window is the combination according to the invention of the micro-alloy elements titanium, niobium and boron with optional addition of molybdenum.
  • Fine titanium precipitations act in the same way as niobium-carbides and together enhance this effect. Titanium binds nitrogen, which therefore is no longer available for the formation of boron-nitride, through which the boron alloy can act. In this case, the addition of boron, which is present as free boron, causes an increase of the hardenability.
  • Boron is one of the elements that besides a high hardening is also characterized by a high hardness penetration.
  • the microstructure becomes more isotropic because differences in the cooling rates, which are caused by the process conduction or the geometry of the strip, have a smaller influence, which also leads to a greater process window.
  • Free boron is able to produce a relatively homogenous microstructure (same microstructure proportions) viewed over the sheet thickness. The same is true for the less pronounced influence of temperature gradients, which occur over the length of the strip or with regard to its width.
  • the proportions of ferrite are more or less reduced by increased proportions of bainite depending on the process parameters.
  • the combination of the three micro-alloy elements, enables the previously described material diversity.
  • Carbide and nitride formation only starts above temperatures of about 1000 C or after the a/ ⁇ transformation, i.e., significantly later than in the case of titanium and niobium.
  • vanadium has hardly a grain refining effect due to the low number of precipitations in the austenite.
  • growth of austenite grains is not inhibited by the late precipitation of the vanadium carbides.
  • the strength increasing effect is almost solely based on the precipitation hardening.
  • vanadium is the high solubility in the austenite and the large volume proportion of fine precipitations caused by the low precipitation temperature.
  • Tramp elements are elements, which are already present in the iron ore or are introduced into the steel due to manufacture. Due to their predominantly negative effect however they are usually not desired. It is sought to remove them to a tolerable content or to convert them into less deleterious forms.
  • Hydrogen(H) is the only element that can diffuse through the iron lattice without causing lattice stresses. As a result, hydrogen is relatively mobile in the iron lattice and can be absorbed relatively easily during manufacture. Hydrogen can only be taken up in the iron lattice in atomic (ionic) form.
  • Hydrogen has a strong embrittling effect and preferably diffuses to energetically favorable sites (gaps, grain boundaries etc). Gaps act as hydrogen traps and can significantly increase the retention time of the hydrogen in the material.
  • Recombination to molecular hydrogen can lead to cold cracks. This behavior occurs in the hydrogen embrittlement or hydrogen-induced tension crack corrosion. Hydrogen is also often identified as the cause in the so-called delayed fracture, which occurs without outer tensions.
  • Oxygen(O) in the molten state, the steel has a relatively high absorption capacity for gases, at room temperature however, oxygen can only diffuse into the material in atomic form. Due to the strong embrittling effect and the negative effect on ageing resistance, oxygen is sought to be reduced as much as possible during manufacture.
  • the oxygen content in the steel should be as low as possible.
  • Nitrogen(N) is also a tramp element in steel production. Steels with free nitrogen are prone to a strong aging effect. Nitrogen already diffuses at low temperature at dislocations and blocks the dislocations. With this, it causes a strength increase associated with a rapid loss of tenacity. Binding of nitrogen in form of nitrides is possible by adding aluminum or titanium to the alloy.
  • sulfur(S) is bound in the iron ore as trace element. It is undesired in the steel (exception automate steels) because it has a strong tendency for segregation and acts strongly embrittling. It is therefore sought to achieve an amount of sulfur in the melt that is as small as possible (for example by ultra vacuum treatment). Further, the present sulfur is converted into the relatively harmless component manganese sulfide (MnS) by adding manganese.
  • MnS manganese sulfide
  • the manganese sulfides are rolled out row-like during the rolling process and act as nuclei for the transformation. Especially in the case of diffusion-controlled transformation this leads to a microstructure that is configured band-like and can lead to poorer mechanical properties in the case of strongly pronounced microstructure banding (for example pronounced martensite bands instead of distributed martensite islands, no isotropic material properties, decreased elongation at break).
  • Phosphorus (P) is a trace element from the iron ore and is dissolved in the iron lattice as substitution atom. Through solid solution hardening, phosphorus increases strength and improves hardenability.
  • phosphorus is sometimes also used in dual-phase steels as strength carrier.
  • Alloy elements are usually added to the steel in order to influence certain properties in a targeted manner.
  • An alloy element can influence different properties in different steels. Generally, the effect strongly depends on the amount and the solution state of the material.
  • chromium in dissolved form, even small amounts of chromium can already significantly increase the hardenability of steels.
  • chromium carbides it can cause a direct increase of strength as a result of particle hardening.
  • the hardenability is reduced.
  • Carbon(C) is regarded as the most important alloy element in the steel. Its presence causes the iron to become steel in the first place. In spite of this fact, the carbon content is drastically lowered during steel production. In dual-phase steels for a continuous hot-dip coating, its proportion according to DIN EN 10346 depending on grade is maximally 0.23%, a minimal value is not given.
  • Carbon is also required to form carbides.
  • An example is cementite (Fe3C), which is present in almost every steel.
  • significantly harder special carbides can form together with other metals such as for example chromium, titanium, niobium, vanadium.
  • chromium, titanium, niobium, vanadium Not only the type but also the distribution and size of the precipitations is of significant importance for the resulting strength increase.
  • the minimal C-content is 0.060% and the maximal C-content is 0.115%.
  • Silicon (Si) binds oxygen during casting and thus decreases segregations and contaminations in the steel.
  • silicone increases the strength and the stretch limit ratio of the ferrite through hybrid Crystal hardening at only minimally decreasing elongation at break. They further important effect is that silicone shifts the formation of ferrite to shorter times and thus enables the formation of sufficient ferrite prior to the quenching.
  • the austenite is enriched with carbon and stabilized.
  • silicon stabilizes the austenite in the lower temperature range particularly in the range of the bainite formation by preventing carbide formation (no depletion of carbon).
  • silicon can diffuse to the surface during the annealing where it can lead to silicon oxides.
  • silicone oxides can interfere with the formation of a closed adhesive layer between steel and zinc (inhibition layer). This results in a poor zinc adhesion and un-galvanized sites.
  • the minimal Si content is 0.100% and the maximal Si-content is 0.500%.
  • Manganese(Mn) is added to almost all steels for de-sulfurization in order to convert the deleterious sulfur into manganese sulfide.
  • manganese increases strength of the ferrite by solid solution strengthening and shifts the a ⁇ / ⁇ transformation to low temperatures.
  • a main reason for adding manganese into the alloy in dual-phase steels is the significant improvement of the hardness penetration. Due to the impeding of diffusion, the perlite and bainite transformation is shifted to longer times and the martensite start temperature is lowered.
  • manganese can lead to manganese oxides at the surface at high concentrations, which can negatively influence the adhesion properties of zinc and the surface aesthetics.
  • the manganese content is therefore set to 1.300 to 2.500%.
  • Chromium(Cr) the addition of Chromium in dual-phase steels mainly improves the hardness penetration. In the dissolved state, chromium shifts the perlite and bainite transformation to longer times and lowers at the same time the martensite stat temperature.
  • a further important effect is that chromium significantly increases the tempering resistance so that almost no strength losses occur in the zinc bath.
  • chromium is a carbide former.
  • the austenization temperature has to be selected to be sufficiently high to dissolve the chromium carbides. Otherwise the increased number of nuclei can lead to a deterioration of the hardness penetration.
  • the Cr-content is therefore set to values of 0.280 to 0.480%.
  • Molybdenum also significantly increases the tempering resistance so that in the zinc bath no strength losses are to be expected and causes through solid solution strengthening a strength increase of the ferrite.
  • the Mo-content is optionally added to the alloy depending on the dimensions the system configuration and the microstructure adjustment, wherein then the minimal addition should be 0,050% to achieve an effect. For reasons of costs, the Mo content is set to maximally 0.150%.
  • Copper (Cu) the addition of copper can increase the tensile strength and the hardness penetration. Together with nickel, chromium and phosphorous copper can form a protective oxide layer at eh surface, which significantly reduces the corrosion rate.
  • copper can from deleterious oxides at the grain boundaries, which can have negative effects, especially for hot forming processes.
  • the content of copper is therefore limited to unavoidable tramp elements.
  • Ni nickel
  • Sn tin
  • Micro-alloy elements are normally only added in very low amounts ( ⁇ 0.1%). In contrast to the alloy elements, they act mainly through formation of precipitations but can also influence the properties in dissolved state. In spite of the small amounts added, the micro-alloy elements strongly influence the production conditions and the processing and end properties.
  • Typical micro-alloy elements are aluminum, vanadium, titanium, niobium and boron. These elements can be dissolved in the iron lattice and together with carbon and nitrogen form carbides or nitrides due to a reduction in the free enthalpy.
  • Aluminum (Al) is normally added to the steel to bind the oxygen that is dissolved in the iron. Oxygen and nitrogen are thus converted to aluminum oxides and aluminum nitrides. These precipitations can cause a grain refinement via an increase of the nuclei and thus increase the tenacity properties and strength values.
  • Titanium nitrides do not precipitate when titanium is present in sufficient amounts. Titanium nitrides have a lower binding enthalpy and therefore form at higher temperatures.
  • the Al-content is therefore limited to 0.020 to maximally 0.060%.
  • Titanium (Ti) already forms very stable nitrides (TiN) and sulfides (TiS2) at high temperatures. These dissolve in dependence on the nitrogen content partly only in the melt. When the thus generated precipitations are not removed with the slag, they form relatively large particles due to the high formation temperature and are normally not beneficial for the mechanical properties.
  • a positive effect on the tenacity is caused by the binding of free nitrogen and oxygen.
  • titanium protects other micro-alloy elements such as niobium from being bound by oxygen. These can then optimally exert their effect.
  • Nitrides which due to the lowering of the oxygen and nitrogen content are generated only at lower temperatures, can cause an effective impediment of the austenite growth.
  • titanium forms titanium carbides at temperatures above 1150° C. and can thus cause a grain refinement (inhibition of the austenite growth, grain refinement through delayed recrystallization and/or increase of the number of nuclei at a/ ⁇ transformation) and a precipitation hardening.
  • Ti-content has therefore values of more than 0.005 and less than 0,050%.
  • Niobium (Nb) causes a strong grain refinement because it is most effective among all micro-alloy elements in delaying recrystallization and in addition inhibits the austenite growth.
  • the strength increasing effect is however higher than that of titanium due to the increased grain refining effect and the larger amount of strength increasing particles (bonding of the titanium to TiN at high temperatures).
  • Niobium carbides form above about 1200° C. In Connection with titanium, which as mentioned binds nitrogen, niobium can increase its strength-increasing effect through carbide formation in the lower temperature range (smaller carbide sizes).
  • a further effect of niobium is the delay of the a/ ⁇ -transformation and the lowering of the martensite start temperature in the dissolved state. On one hand, this occurs through the solute drag effect and on the other hand through the grain refinement. This causes a strength increase of the microstructure and thus also a higher resistance against the expansion in the martensite formation.
  • niobium is limited by the relatively low solubility limit. The latter limits the amount of precipitations however causes an early precipitation formation with relatively coarse particles.
  • the precipitation hardening can thus become effective mostly in steels with low C-content (greater over saturation possible) and in hot forming processes (deformation induced precipitation).
  • Vanadium(V) the carbide and nitride formation of vanadium only starts above temperatures of about 1000° C. or still after the a/ ⁇ transformation, i.e., significantly later than in the case of titanium and niobium. Vanadium has thus hardly a grain refining effect due to the low number of precipitations that are present in the austenite. Also the austenite growth is not inhibited by the late precipitation of the vanadium carbides.
  • vanadium is not required in the present alloy concept, the content of vanadium is limited to unavoidable steel tramp amounts.
  • nitrogen is more affine to beryllium, aluminum, cerium, titanium and zirconium.
  • titanium can ensure binding of the entire nitrogen.
  • Aluminum is less capable for this.
  • boron atoms preferably accumulate at the grain boundaries where they inhibit the diffusion and the grain boundary sliding by lowering the grain boundary energy.
  • the reduction of the precipitation formation at the grain boundaries results in a decrease of the nuclei.
  • Boron has a very high affinity for oxygen, which can lead to a decrease of the boron content in regions close to the surface (up to 0.5 mm). In this context, it is advised against an annealing of over 1000° C. This is to be recommended because boron can lead to a strong coarse grain formation at annealing temperatures above 1000° C.
  • the B-content is limited to values of 0.0005 to 0.0060%.
  • the values are however below 0.0050 or 0.0040%.
  • the annealing temperatures for the steel according to the invention lie between 700 and 950° C., with this a partial austenitic (two phase region) or a full austenitic structure (austenitic region) is achieved depending on the microstructure to be achieved (complex-phase microstructure).
  • the hot-dip coated material can be manufactured as hot strip, as cold temper rolled hot strip or cold strip in the dressed (cold after rolled) or stretch bent straightened state (undressed).
  • Steel strips in the present case as hot strip, cold temper rolled hot strip or cold strip made of the alloy composition according to the invention are further characterized by a high resistance against edge proximate fracture formation during further processing.
  • the material can also be used transverse, longitudinal and diagonally to the rolling direction.
  • the hot strip is according to the invention produced with finishing rolling temperatures in the austenitic range above A C3 and coiling temperature above the recrystallization temperature.
  • FIG. 1 schematically the process chain for the production of the steels according to the invention
  • FIG. 2 results of the hole-widening test
  • FIG. 3 examples for analytical differences of the steel according to the invention relative to the state of the art
  • FIG. 5 schematically the time temperature course of the process steps hot rolling and continuous annealing
  • FIG. 6 ZTU-diagram for a steel according to the invention
  • FIG. 7 mechanical characteristics in case of variation of the rolling degrees
  • FIG. 8 overview over the strength classes that can be adjusted with the alloy concept according to the invention.
  • FIG. 9 temperature time curve (schematic).
  • FIG. 1 shows schematically the process chain for the production of the steels according to the invention. Shown are the different process routes that relate to the invention. Up to position 5 (pickling) the process route is the same for all steels, after that the corresponding processing occurs according to the desired results.
  • the pickled hot strip can be galvanized or cold-rolled and galvanized. Or it can be soft annealed, cold-rolled and galvanized.
  • FIG. 2 shows results of the hole-widening test (relative values in comparison). Shown are the results of the hole widening tests for the steel according to the invention compared to the standard grades. All materials have a sheet thickness of 2.00 mm. On the left hand partial picture, the results for the test ISO TS 16630 are shown, on the right hand side the results for the KWI-test (Kaiser Wilhelm Institute). It can be seen that the steels according to the invention achieve the best widening values in punched holes independent of the type of processing.
  • Process 1 corresponds here to an annealing for example on a fire galvanization with combined directly fired furnace and radiation tube furnace.
  • Process 2 corresponds for example to a process control in a continuous annealing system.
  • a reheating of the steel by means of induction furnace can be optionally achieved prior to the galvanization bath.
  • FIG. 3 shows the relevant alloy elements of the steel according to the invention compared to steels of the same grade, which correspond to the state of the art.
  • the main difference is the carbon content, which lies in the ultra peritectic range.
  • FIG. 4 shows the mechanical characteristics of the steel according to the invention compared to those of the state of the art. All characteristics correspond to the normative standard.
  • FIG. 5 shows schematically the time temperature course of the process steps hot rolling and continuous annealing of strips having the alloy composition according to the invention. Shown is the time and temperature dependent transformation for the hot rolling process and also for a heat treatment after the cold rolling. Of particular interest is the shift of the ferrite phase to later times. This enables the potential for complex phase steels and bainitic steels.
  • FIG. 6 shows a ZTU-Diagram for a steel according to the invention.
  • the determined ZTU diagram with corresponding chemical composition and the A C1 and A C3 temperature is shown.
  • FIG. 7 shows the mechanical characteristics with same parameters of continuously annealed strips at variation of the rolling degrees of different material thickness. Shown are the characteristics tensile strength, stretch limit and elongation at break in dependence on selected rolling degrees. Only the tensile strength slightly increases. All values lie within the standard for an HCT780XD an show that also in the case of different sheet thicknesses practically same mechanical properties are present after the continuous annealing.
  • FIG. 8 shows schematically the temperature time courses during the annealing treatment and cooling with 3 different variants.
  • Variant 1 shows the annealing and cooling of produced cold or hot rolled steel strip in a continuous annealing system.
  • the strip is heated to a temperature of 700 to 950° C.
  • the annealed steel strip is subsequently cooled from the annealing temperature with a cooling speed between 15 and 100° C./s to an intermediate temperature of 200 to 250° C.
  • the steel strip is cooled on air with a cooling speed of 2 and 30° C. until reaching room temperature or the cooling speed between 15 and 100° C./s is maintained until reaching room temperature, i.e., the intermediate temperature corresponds to the room temperature.
  • Variant 2 ( FIG. 9 b ) shows the process according to variant 1 , however, the cooling of the steel strip is shortly interrupted during passage through the hot dip container for the purpose of a hot-dip coating, in order to subsequently continue the cooling with a cooling speed between 15 and 100° C./s until reaching an intermediate temperature of 200 to 250° C. Subsequently, the steel strip is cooled on air with a cooling speed of 2 and 30° C./s until reaching room temperature.
  • Variant 3 ( FIG. 9 c ) also shows the process according variant 1 in a hot-dip coating, however, the cooling of the steel strip is interrupted by a short break (1 to 20 s) at an intermediate temperature of 200 to 250° C. and reheated to the temperature that is required for the hot-dip coating. Subsequently, the steel strip is cooled to an intermediate temperature of 200 to 250° C. With a cooling speed of 2 and 30° C./s the final cooling of the steel strip is carried out on air until reaching room temperature.

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US10519525B2 (en) 2012-03-20 2019-12-31 Salzgitter Flachstahl Gmbh High strength multi-phase steel, and method for producing a strip from said steel
US11512364B2 (en) 2015-07-17 2022-11-29 Salzgitter Flachstahl Gmbh Method for producing a hot strip of a bainitic multi-phase steel having a Zn—Mg—Al coating, and a corresponding hot strip
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EP3816319B1 (de) 2019-10-29 2022-09-14 Salzgitter Flachstahl GmbH Verfahren zur herstellung eines hochfesten stahlbandes mit verbesserter haftung zinkbasierter schmelztauchüberzüge
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CN107208232A (zh) * 2014-11-18 2017-09-26 萨尔茨吉特法特尔有限公司 具有优异加工性能的空气硬化的超高强度多相钢以及由其制备带材的方法
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CN107208232B (zh) * 2014-11-18 2019-02-26 萨尔茨吉特法特尔有限公司 具有优异加工性能的空气硬化的超高强度多相钢以及由其制备带材的方法
US10626478B2 (en) 2014-11-18 2020-04-21 Salzgitter Flachstahl Gmbh Ultra high-strength air-hardening multiphase steel having excellent processing properties, and method for manufacturing a strip of said steel
US10640855B2 (en) * 2014-11-18 2020-05-05 Salzgitter Flachstahl Gmbh High-strength air-hardening multiphase steel having excellent processing properties, and method for manufacturing a strip of said steel
US20190127821A1 (en) * 2015-07-10 2019-05-02 Salzgitter Fachstahl GmbH Ultra high strength multi-phase steel and method for producing a cold-rolled steel strip therefrom
US11512364B2 (en) 2015-07-17 2022-11-29 Salzgitter Flachstahl Gmbh Method for producing a hot strip of a bainitic multi-phase steel having a Zn—Mg—Al coating, and a corresponding hot strip
CN110283972A (zh) * 2019-07-02 2019-09-27 天津市琨泰机械制造有限公司 一种扁铁(多条)热镀锌及钢带热镀锌综合连续生产线
CN110527795A (zh) * 2019-07-02 2019-12-03 天津市琨泰机械制造有限公司 一种扁铁(多条)热镀锌连续生产线
CN119980074A (zh) * 2025-04-09 2025-05-13 本钢板材股份有限公司 一种车轮轮辋用含稀土复相热轧钢及其制备方法

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