US20190127821A1 - Ultra high strength multi-phase steel and method for producing a cold-rolled steel strip therefrom - Google Patents

Ultra high strength multi-phase steel and method for producing a cold-rolled steel strip therefrom Download PDF

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US20190127821A1
US20190127821A1 US15/743,229 US201615743229A US2019127821A1 US 20190127821 A1 US20190127821 A1 US 20190127821A1 US 201615743229 A US201615743229 A US 201615743229A US 2019127821 A1 US2019127821 A1 US 2019127821A1
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steel
content
strip
approximately
cold
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Thomas Schulz
Sascha Kluge
Christian Meyer
Peter Matthies
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Salzgitter Fachstahl GmbH
Salzgitter Flachstahl GmbH
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Salzgitter Fachstahl GmbH
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Assigned to SALZGITTER FLACHSTAHL GMBH reassignment SALZGITTER FLACHSTAHL GMBH ASSIGNMENT OF ASSIGNORS INTEREST (SEE DOCUMENT FOR DETAILS). Assignors: KLUGE, Sascha, Matthies, Peter, MEYER, CHRISTIAN, SCHULZ, THOMAS
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Definitions

  • the invention relates to an ultra high strength multi-phase steel with a dual-phase microstructure or complex-phase microstructure and with small proportions of residual austenite having improved properties for production and processing, in particular for lightweight vehicle construction according to the preamble of claim 1 .
  • the invention also relates to a method for producing cold-rolled steel strips from such a steel according to claim 18 and steel strips produced therewith according to claim 34 .
  • the invention relates to steels having a tensile strength in the region of at least 980 MPa for producing components which have improved deformability, such as e.g. increased hole expansion (Xpand®) and increased bending angles and improved welding properties.
  • improved deformability such as e.g. increased hole expansion (Xpand®) and increased bending angles and improved welding properties.
  • the fiercely competitive car market means that producers are constantly forced to find solutions for reducing fleet fuel consumption whilst maintaining the highest possible level of comfort and passenger protection.
  • the weight saving of all of the vehicle components plays a decisive role as does, on the other hand, the most favourable possible behaviour of the individual components in the event of high static and dynamic loading during operation and also in the event of a crash.
  • the suppliers of semi-finished products attempt to take this requirement into account in that, by providing ultra high strength steels having a smaller sheet thickness, the weight of the vehicle components can be reduced whilst the behaviour of components remains the same or is possibly even improved.
  • multi-phase steels such as complex-phase steels, ferritic-bainitic steels, bainitic steels and martensitic steels which have different microstructure compositions.
  • Complex-phase steels in hot-rolled or cold-rolled form are steels which contain small proportions of martensite, residual austenite and/or perlite in a ferritic/bainitic basic microstructure, wherein an extreme grain refinement is produced by means of delayed recrystallisation or by precipitation of microalloy elements.
  • TRB®s having a multi-phase microstructure are not possible without limitations, such as e.g. for heat treatment prior to cold rolling.
  • a homogeneous, multi-phase microstructure cannot be set in cold-rolled and hot-rolled steel strips by reason of a temperature gradient which occurs in the established process windows.
  • the cold-rolled steel strips are typically annealed in the continuous annealing method in a recrystallising manner to produce thin sheet which can be deformed in an effective manner.
  • the process parameters such as throughput speed, annealing temperatures and cooling rate, are set corresponding to the required mechanical-technological properties with the microstructure required for this purpose.
  • the degree of thinning by rolling in cold rolling describes the percentage difference in the hot strip starting thickness with respect to the cold strip end thickness based on the hot strip starting thickness.
  • the degree of thinning by rolling is related to the strip thickness and therefore is to be understood in terms of a thickness reduction.
  • the degrees of thinning by rolling in cold rolling are relatively constant, they are up to approximately 40% in the case of thicker strips of more than 2 mm and are up to approximately 60% in strips of up to 1 mm in thickness.
  • a degree of cold rolling of 50% is required in the case of continuous annealing treatment in order to ensure normal recrystallisation.
  • deviations to higher or lower amounts lead to fluctuating technological characteristic values, as described in the case of the TRB®'s.
  • the critical threshold for recrystallisation cannot be overcome and the microstructure cannot be refined.
  • the different grain sizes in the cold strip also give rise to different grain sizes in the final microstructure, which results in fluctuations in the characteristic values.
  • grains of different sizes can convert into different phase components and ensure further inhomogeneity.
  • the cold strip is heated in the continuous annealing furnace to a temperature at which, during cooling, the required microstructure formation (e.g. dual-phase or complex-phase microstructure) is achieved.
  • the required microstructure formation e.g. dual-phase or complex-phase microstructure
  • the annealing treatment is typically performed in a continuous hot-galvanising installation, in which the heat treatment or annealing and the downstream galvanising take place in a continuous process.
  • This relates not only to flexibly rolled strips having different sheet thicknesses over the strip length but also primarily to strips having a different thickness and/or different width which have to be annealed in succession.
  • a homogeneous temperature distribution can be achieved only with difficulty specifically in the case of different thicknesses in the transition region from one strip to the other.
  • alloy compositions with excessively small process windows this can result in e.g. the thinner strip being moved too slowly through the furnace and as a result the productivity is reduced, or in the thicker strip being moved too quickly through the furnace and the required annealing temperature for the desired microstructure is not achieved. As a result, there is an increased amount of scrap.
  • the corresponding process parameter for material having a relatively constant degree of thinning by rolling during cold rolling is the setting of speed during continuous annealing because the phase conversion proceeds in dependence upon temperature and time. Therefore, the more insensitive the steel in relation to the uniformity of the mechanical properties during changes in the temperature and time profiles during continuous annealing, the greater the process window.
  • a method for producing a steel strip having a different thickness over the strip length is described e.g. in DE 100 37 867 A1.
  • the narrow process window during the continuous annealing of strips of different thickness means that it is only possible with difficulty to achieve uniform mechanical properties over the entire length of the strip.
  • Complex-phase steels also have an even narrower process window than dual-phase steels.
  • the setting of relatively homogeneous mechanical-technological properties of different cold strips having variable degrees of thinning by rolling during cold rolling cannot be achieved in practical terms with the known alloy concepts during continuous annealing.
  • the degree of thinning by rolling during cold rolling which is required for recrystallisation annealing results in a very clear restriction in the flexibility of the material production within the entire process chain.
  • the final cold strip thickness establishes the thickness of the hot strip and therefore the hot strip manufacturing parameters.
  • the excessively small process window means that the regions having a smaller sheet thickness owing to the conversion processes during cooling have either excessively high levels of strength as a result of excessively large martensite proportions or the regions having a greater sheet thickness achieve excessively low levels of strength as a result of excessively small martensite proportions.
  • Homogeneous mechanical-technological properties over the strip length or width cannot be achieved in practical terms with the known alloy concepts during continuous annealing.
  • the known alloy concepts for multi-phase steels are characterised by an excessively narrow process window and therefore are unsuitable in particular for cold strip manufacture with variable degrees of rolling during cold rolling, and for flexibly rolled strips.
  • Laid-open document DE 10 2012 002 079 A1 discloses an ultra high strength multi-phase steel having minimum tensile strengths of 950 MPa which, even though it already has a very wide process window for the continuous annealing of hot or cold strips, it has been shown that even with this steel it is not possible to achieve variable degrees of thinning by rolling with a single hot strip thickness (master hot strip thickness) whilst producing uniform material properties.
  • the material processing properties deteriorate to an increasing extent, e.g. during welding, deforming and hot-dip finishing.
  • a low yield strength ratio (Re/Rm) is typical of a dual-phase steel and is used primarily for deformability in stretching and deep-drawing procedures.
  • a higher yield strength ratio (Re/Rm), as is typical of complex-phase steels, is also characterised by the resistance to edge cracks. This can be attributed to the smaller differences in the strengths of the individual microstructure components, which has a favourable effect upon homogeneous deformation in the region of the cut edge.
  • the analytical landscape for achieving multi-phase steels having minimum tensile strengths of 980 MPa is very diverse and comprises very large alloy ranges in the strength-increasing elements of carbon, manganese, phosphorous, aluminium and chromium and/or molybdenum, as well as in the addition of microalloys individually or in combinations, and in the material-characterising properties.
  • the dimension spectrum is broad and is in the thickness range of 0.50 to 3.00 mm, wherein the range between 0.80 to 2.10 mm is relevant in terms of quantity.
  • the object of the invention is to provide a novel alloy concept for an ultra high strength multi-phase steel, a method for producing a steel strip from this ultra high strength multi-phase steel and to provide a steel strip which is produced according to this method and with which the process window for the continuous annealing of cold strips can be extended such that from a specified hot strip thickness (master hot strip thickness) it is possible to manufacture different cold strip thicknesses or from different hot strip thicknesses it is possible to manufacture a cold strip thickness (master cold strip thickness), i.e. variable degrees of thinning by rolling can be used instead of relatively constant degrees of thinning by rolling during the cold rolling procedure. In this case, the most uniform possible material properties are to be achieved independently of the set degree of thinning by rolling.
  • the process window for the annealing, in particular continuous annealing, of steel strips which are rolled to an end thickness is to be extended such that, in addition to strips having different cross-sections, it is also possible to produce steel strips having a thickness which varies over the strip length and optionally strip width, with the most homogeneous mechanical-technological properties possible.
  • this object is achieved by an ultra high strength multi-phase steel having a minimum tensile strength of 980 MPa with the following contents in wt. %:
  • the mechanical-technological properties are reliably achieved in a narrow range for cold strips having variable degrees of thinning by rolling during cold rolling.
  • the degree of thinning by rolling is related to the strip thickness and therefore is to be understood in terms of a thickness reduction. What is significant for this is a selected, closely maintained alloy composition with the emphasis on a restricted and very limited boron content which has proven to be very effective for achieving uniform material properties with different degrees of cold rolling.
  • the mechanical-technological properties which can be produced are achieved in a narrow range over the strip width and strip length by setting the volume proportions of the microstructure phases in a controlled manner.
  • the previous production philosophy that the final cold strip thickness (end thickness) determines the necessary hot strip thickness can be disregarded to the extent that only one selected master hot strip thickness is required for different cold strip thicknesses.
  • the steel according to the invention offers the advantage of a considerably increased process window in comparison with the known steels. This results in an increased level of process reliability in the continuous annealing of cold strip having a multi-phase microstructure. Therefore, for continuously annealed cold strips it is possible to ensure more homogeneous mechanical-technological properties in strips having variable degrees of thinning by rolling during cold rolling, and in the strip or in the transition region of two strips even with different cross-sections and otherwise identical process parameters.
  • the inventive steel to produce a steel strip, in which the multi-phase steel is used to produce a hot strip, from the hot strip the steel strip is cold-rolled with the end thickness to be achieved and subsequently the steel strip is annealed, in particular continuously annealed.
  • the properties of the multi-phase steel render it possible that, proceeding from a selected master hot strip having a specific thickness or selected hot strips having different thicknesses in a wide range of degrees of thinning by rolling of 10% to 70% steels strips are cold-rolled with the end thickness to be achieved.
  • the chemical composition of the multi-phase steel is selected in dependence upon the end thickness of the steel strip to be obtained. Therefore, within selectable thickness graduations of the cold strip to be obtained, it is possible to produce, from a master hot strip having a thickness, corresponding cold strips having one or more end thicknesses or else to produce from different hot strip thicknesses a master cold strip having a consistent thickness.
  • the steel strip is cold-rolled to an end thickness of 0.50 mm to 3.00 mm and the chemical composition of the multi-phase steel is selected beforehand in dependence upon the end thickness to be achieved, as follows:
  • end thickness 0.50 mm to 1.00 mm inclusive sum of Mn+Si+Cr ⁇ 2.500 to ⁇ 3.100 wt. %
  • end thickness above 1.00 mm to 2.00 mm inclusive sum of Mn+Si+Cr ⁇ 2.650 to ⁇ 3.150 wt. %
  • end thickness above 2.00 mm to 3.00 mm inclusive sum of Mn+Si+Cr ⁇ 2.700 to ⁇ 3.250 wt. %. Therefore, the end thickness of the steel strip to be achieved is associated with the alloy composition of the hot strip produced from the multi-phase steel.
  • CEV(IIW) % C+% Mn/6+(% Cu+% Ni)/15+(% Cr+% Mo+% V)/5.
  • end thickness 0.50 mm to 1.00 mm inclusive Mn content ⁇ 1.900 to ⁇ 2.200 wt. %
  • end thickness over 1.00 mm to 2.00 mm inclusive Mn content ⁇ 2.050 to ⁇ 2.250 wt. %
  • end thickness over 2.00 mm to 3.00 mm inclusive Mn content ⁇ 2.100 to ⁇ 2.350 wt. %.
  • the material produced can be produced as a cold strip via a hot-dip galvanising line or a pure continuous annealing installation, in the temper-rolled and non-temper-rolled state and even in the heat-treated state (ageing) and in the stretch-formed and non-stretch-formed state (stretch-bend-straighten).
  • the alloy composition according to the invention can be used to produce steel strips by means of inter-critical annealing between Ac1 and Ac3 or with austenitising annealing over Ac3 with concluding controlled cooling, which results in a dual-phase or multi-phase microstructure.
  • Annealing temperatures of about 700 to 950° C. have proven to be advantageous. According to the invention, in dependence upon the entire process (only continuous annealing or with additional hot-dip finishing) there are different approaches for heat treatment.
  • the steel strip which is cold-rolled to an end thickness is cooled, starting from the annealing temperature, at a cooling rate of approximately 15 to 100° C./s to an intermediate temperature of approximately 160 to 250° C.
  • cooling can be performed in advance at a cooling rate of approximately 15 to 100° C./s to a previous intermediate temperature of 300 to 500° C. Cooling to room temperature is finally effected at a cooling rate of approximately 2 to 30° C./s (see method 1 , FIG. 7 a ).
  • cooling can be performed at a cooling rate between approximately 15 and 100° C./s from the intermediate temperature of 300 to 500° C. to room temperature.
  • Cooling is stopped as described above prior to entry into the melting bath and is continued only after exit from the bath until the intermediate temperature of approximately 200 to 250° C. is achieved. Depending upon the melting bath temperature, a holding temperature of approximately 400 to 470° C. is provided in the melting bath. Cooling to room temperature is the performed at a cooling rate of approximately 2 to 30° C./s (see method 2 , FIG. 7 b ).
  • the second variant for temperature control in the hot-dip finishing procedure includes maintaining the temperature for approximately 1 to 20 s at the intermediate temperature of approximately 200 to 350° C. and subsequently reheating to the temperature of approximately 400 to 470° C. required for the hot-dip finishing procedure. After the finishing procedure, the strip is then cooled to approximately 200 to 250° C. Cooling to room temperature is then effected at a cooling rate of approximately 2 to 30° C./s (see method 3 , FIG. 7 c ).
  • the carbon equivalent can be reduced, whereby the welding suitability is improved and excessive hard spots during welding are avoided. Furthermore, in the case of resistance spot welding the electrode service life can be considerably increased.
  • Associated elements are elements which are already present in the iron ore or get into the steel as a result of the production process. They are generally undesired by reason of their predominantly negative influences. The attempt is made to remove them to a tolerable level or to convert them into less damaging forms.
  • Hydrogen (H) can diffuse as a single element through the iron lattice, without producing lattice tensions.
  • the hydrogen in the iron lattice is relatively mobile and can be relatively easily absorbed during the processing of the steel.
  • Hydrogen can be absorbed into the iron lattice only in atomic (ionic) form.
  • Hydrogen exerts a significant embrittling effect and diffuses preferably to locations which are favourable in terms of energy (flaws, grain boundaries etc.). Flaws thus function as hydrogen traps and can considerably increase the dwell time of the hydrogen in the material.
  • Cold cracks can be produced by means of a recombination to molecular hydrogen. This behaviour occurs in the event of hydrogen embrittlement or in the event of hydrogen-induced tension crack corrosion. Even in the case of delayed cracking, so-called delayed fracture, which occurs without external tensions, hydrogen is often stated to be the reason instigating this. Therefore, the hydrogen content in the steel should be kept as small as possible.
  • a more uniform microstructure which, in the case of the steel according to the invention, is achieved inter alia by its widened process window, also reduces the susceptibility to hydrogen embrittlement.
  • Oxygen (O) In the molten state, the steel has a relatively large absorbency for gases. However, at room temperature oxygen is soluble only in very small quantities. In a similar manner to hydrogen, oxygen can diffuse only in atomic form into the material. Owing to the highly embrittling effect and the negative effects upon the ageing resistance, every attempt is made during production to reduce the oxygen content.
  • Phosphorous is a trace element from the iron ore and is dissolved in the iron lattice as a substitution atom. Phosphorous increases the hardness and improves the hardenability by means of mixed crystal solidification. However, attempts are generally made to lower the phosphorous content as much as possible because inter alia its low solubility in the solidifying medium means that it exhibits a strong tendency towards segregation and greatly reduces the level of toughness. The attachment of phosphorous to the grain boundaries causes grain boundary fractures. Moreover, phosphorous increases the transition temperature from tough to brittle behaviour up to 300° C. During hot rolling, near-surface phosphorous oxides at the grain boundaries can result in the formation of fractures.
  • phosphorous is used in small quantities ( ⁇ 0.1 wt. %) as a microalloy element, e.g. in higher-strength IF-steels (interstitial free), bake-hardening steels or even in some alloy concepts for dual-phase steels.
  • the steel according to the invention differs from known analysis concepts, which use phosphorous as a mixed crystal forming agent, inter alia in that phosphorous is not alloyed thereto but instead is set as low as possible.
  • the phosphorous content in the steel according to the invention is limited to quantities which are unavoidable in the production of steel.
  • P should be ⁇ 0.020 wt. %.
  • S Sulphur
  • Sulphur is bound as a trace element in the iron ore. Sulphur is not desirable in steel (the exception being machining steels) because it exhibits a strong tendency towards segregation and has a greatly embrittling effect. Therefore, attempts are made to achieve, where possible, a very low content of sulphur in the melt, e.g. by means of vacuum treatment. Furthermore, the sulphur present is converted by the addition of manganese into the relatively innocuous compound manganese sulphide (MnS). The manganese sulphides are often rolled out in lines during the rolling process and function as nucleation sites for the conversion.
  • the sulphur content in the steel according to the invention is limited to ⁇ 0.0020 wt. % or advantageously to ⁇ 0.0015 wt. %, optimally to ⁇ 0.0010 wt. %.
  • Alloy elements are generally added to the steel in order to influence specific properties in a targeted manner.
  • An alloy element can thereby influence different properties in different steels. The effect generally depends greatly upon the quantity and solution state in the material. Accordingly, the relationships can be very varied and complex.
  • Carbon (C) is considered to be the most important alloy element in steel. Its targeted introduction at an amount of up to 2.06 wt. % turns iron first into steel. The carbon proportion is often drastically reduced during the production of steel. In the case of dual-phase steels for continuous hot-dip finishing, its proportion according to EN 10346 or VDA 239-100 is at most 0.230 wt. %, a minimum value is not specified.
  • Carbon is interstitially dissolved in the iron lattice owing to its comparatively small atomic radius.
  • the solubility is at most 0.02% in the ⁇ -iron and is at most 2.06% in the ⁇ -iron.
  • Carbon, in dissolved form increases the hardenability of steel considerably and is therefore essential for forming a sufficient quantity of martensite.
  • excessively high carbon contents increase the hardness difference between ferrite and martensite and restrict weldability.
  • the steel according to the invention contains carbon contents of ⁇ 0.115 wt. %.
  • Carbon also forms carbides.
  • a microstructure phase which occurs almost in every steel is cementite (Fe 3 C).
  • substantially harder special carbides can be formed with other metals such as e.g. chromium, titanium, niobium but also vanadium. Therefore, it is not only the type but also the distribution and extent of the precipitation which is of crucial significance for the resulting increase in strength. Therefore, in order to ensure, on the one hand, sufficient strength and, on the other hand, effective weldability, improved hole expansion, an improved bending angle and a sufficient resistance to hydrogen-induced crack formation (delayed fracture free), the minimum C content is fixed to 0.075 wt. % and the maximum C content is fixed to 0.115 wt. %, contents having a cross-section-dependent differentiation are advantageous, such as:
  • end thickness 0.50 mm to 1.00 mm inclusive (C ⁇ 0.100 wt. %) end thickness over 1.00 mm to 2.00 mm inclusive (C ⁇ 0.105 wt. %) end thickness over 2.00 mm to 3.00 mm inclusive (C ⁇ 0.115 wt. %)
  • silicon (Si) binds oxygen and is therefore used for calming purposes in the course of deoxidation of the steel. It is important for the subsequent steel properties that the segregation coefficient is considerably less than e.g. that of manganese (0.16 in comparison with 0.87). Segregations generally lead to a lined arrangement of the microstructure components which impair the deforming properties, such as hole expansion and bending capability.
  • the addition of silicon produces a strong mixed crystal solidification.
  • an addition of 0.1% silicon produces an increase in tensile strength by approximately 10 MPa, wherein in the case of an addition of up to 2.2% silicon the elongation is only slightly impaired. This has been examined for different sheet thicknesses and annealing temperatures. The increase from 0.2% to 0.5% silicon produced an increase in strength of approximately 20 MPa in yield strength and approximately 70 MPa in tensile strength. The elongation at fracture decreases by about 2%.
  • a further important effect is that silicon shifts the formation of ferrite towards shorter times and temperatures and therefore permits the production of sufficient ferrite prior to quench hardening.
  • this provides a basis for improved cold-rollability.
  • the accelerated formation of ferrite causes the austenite to be enriched with carbon and thus stabilised. Since silicon hinders the formation of carbide, the austenite is additionally stabilised. Therefore, during the accelerated cooling the formation of bainite can be suppressed in favour of martensite.
  • silicon in the span according to the invention has resulted in further surprising effects which are described hereinafter.
  • the above-described delay in the formation of carbide could also be brought about e.g. by aluminium.
  • aluminium forms stable nitrides so that sufficient nitrogen is not available for the formation of carbonitrides with microalloy elements.
  • Alloying with silicon obviates this problem because silicon does not form either carbides or nitrides. Therefore, silicon indirectly has a positive effect upon the precipitate formation by microalloys which, in turn, have a positive effect upon the strength of the material. Since the increase in the conversion temperatures caused by silicon tends to promote grain coarsening, a microalloy with niobium, titanium and boron is particularly expedient, as is the targeted setting of the nitrogen content in the steel according to the invention.
  • the atmospheric conditions during the annealing treatment in a continuous hot-dip coating installation produces a reduction of iron oxide which can form on the surface e.g. during cold rolling or as a result of storage at room temperature.
  • oxygen-affine alloy components such as e.g. silicon, manganese, chromium, boron
  • the gas atmosphere is oxidising, as a result of which a segregation and selective oxidation of these elements can take place.
  • the selective oxidation can take place both externally, i.e. on the substrate surface, and internally within the metallic matrix.
  • silicon in particular, diffuses to the surface during annealing and forms, on its own or together with manganese, oxides on the steel surface. These oxides can inhibit the contact between the substrate and the melt and can prevent or significantly impair the wetting reaction. As a result, non-galvanised locations, so-called “bare spots” or even large-surface regions without any coating can occur. Furthermore, the adhesion of the zinc or zinc alloy layer on the steel substrate can be reduced by reason of an impaired wetting reaction, resulting in insufficient inhibitor layer formation.
  • the internal oxidation of the alloy elements can be influenced in a targeted manner by setting the oxygen partial pressure of the furnace atmosphere (N 2 —H 2 -protective gas atmosphere).
  • the set oxygen partial pressure must satisfy the following equation, wherein the furnace temperature is between 700 and 950° C.
  • Si, Mn, Cr, B define the corresponding alloy proportions in the steel in wt. % and pO2 defines the oxygen partial pressure in mbar.
  • the furnace region consists of a combination of a direct fired furnace (DFF) or non-oxidising furnace (NOF) and a downstream radiant tube furnace (see method 2 in FIG. 7 b )
  • DFF direct fired furnace
  • NOF non-oxidising furnace
  • the selective oxidation of the alloy elements can likewise be influenced by the gas atmospheres of the furnace regions.
  • the oxygen partial pressure and therefore the oxidation potential for iron and the alloy elements can be set by the combustion reaction in the NOF. This is to be set in such a way that the oxidation of the alloy elements takes place internally below the steel surface and possibly a thin iron oxide layer forms on the steel surface after passage through the NOF region. This is achieved e.g. by reducing the CO value below 4 vol. %.
  • the possibly formed iron oxide layer is reduced and similarly the alloy elements are further oxidised internally.
  • the set oxygen partial pressure in this furnace region must satisfy the following equation, wherein the furnace temperature is between 700 and 950° C.
  • Si, Mn, Cr, B define the corresponding alloy proportions in the steel in wt. % and pO 2 defines the oxygen partial pressure in mbar.
  • the dew point of the gas atmosphere N 2 —H 2 -protective gas atmosphere
  • the oxygen partial pressure are to be set in such a way as to avoid oxidation of the strip prior to dipping in the melting bath.
  • Dew points in the range of ⁇ 30 to ⁇ 40° C. have proven to be advantageous.
  • the method route selected is that involving continuous annealing with subsequent electrolytic galvanising (see method 1 in FIG. 6 a ), it is not necessary to implement any particular measures in order to ensure the galvanising capability. It is known that the galvanisation of more highly alloyed steels can be achieved substantially more easily by means of electro-deposition than by means of continuous hot-dipping methods. In the case of electrolytic galvanising, pure zinc is deposited directly onto the strip surface. In order not to hinder the electron flow between the steel strip and the zinc ions and therefore galvanisation, it is necessary to ensure that no surface-covering oxide layer is present on the strip surface. This condition is generally ensured by a standard reducing atmosphere during annealing and by pre-cleaning prior to electrolysis.
  • the minimum silicon content is fixed to 0.400 wt. % and the maximum silicon content is fixed to 0.500 wt. %.
  • Manganese (Mn) is added to almost all steels for the purpose of desulphurisation in order to convert noxious sulphur into manganese sulphides. Moreover, by means of mixed crystallisation manganese increases the strength of the ferrite and shifts the ⁇ -/ ⁇ -conversion towards lower temperatures.
  • a main reason for alloying manganese in multi-phase steels is the considerable improvement in the potential hardness increase.
  • the perlite and bainite conversion is shifted towards longer times and the martensite starting temperature is decreased.
  • Manganese like silicon, tends to form oxides on the steel surface during the annealing treatment.
  • manganese oxides e.g. MnO
  • Mn mixed oxides e.g. Mn 2 SiO 4
  • manganese is to be considered to be less critical in a small Si/Mn or Al/Mn ratio because globular oxides are more likely to form instead of oxide films.
  • high manganese contents can negatively influence the appearance of the zinc layer and the zinc adhesion.
  • the above-stated measures for setting the furnace regions during continuous hot-dip coating serve to reduce the formation of Mn oxides or Mn mixed oxides on the steel surface after annealing.
  • the manganese content is fixed to 1.900 to 2.350 wt. %.
  • the manganese content is preferably in a range between ⁇ 1.900 to ⁇ 2.200 wt. %, in the case of end thicknesses of 1.00 to 2.00 mm inclusive, said manganese content is in a range between ⁇ 2.050 to ⁇ 2.250 wt. % and in the case of end thicknesses of 2.00 mm to 3.00 mm inclusive, it is in a range between 2.100 wt. % to ⁇ 2.350 wt. %.
  • a further special feature of the invention is that the variation in the manganese content can be compensated for by simultaneously changing the silicon content.
  • the increase in strength (in this case the yield strength, YS) by reason of manganese and silicon is generally described in an effective manner by the Pickering equation:
  • the coefficients of manganese and silicon are approximately equal for the yield strength and for the tensile strength, thus providing the option of substituting manganese by silicon.
  • chromium (Cr) in dissolved form and even in small quantities can considerably increase the hardenability of steel.
  • Cr chromium in the form of chromium carbides effects particle solidification.
  • the associated increase in the number of nucleation sites with a simultaneously lowered content of carbon leads to a reduction in hardenability.
  • chromium In dual-phase steels, the addition of chromium mainly improves the potential hardness increase. Chromium in the dissolved state shifts the perlite and bainite conversion towards longer times and at the same time lowers the martensite starting temperature.
  • a further important effect is that chromium considerably increases the tempering resistance so that in the hot-dip bath there is almost no loss of strength.
  • chromium is a carbide forming agent. Should chromium-iron-mixed carbides be present, the austenitising temperature must be selected, prior to hardening, to be high enough to dissolve chromium carbides. Otherwise, the increased number of nuclei can cause a deterioration in the potential hardness increase.
  • Chromium likewise tends to form oxides on the steel surface during the annealing treatment, as a result of which the hot-dip quality can be impaired.
  • the above-stated measures for setting the furnace regions during continuous hot-dip coating serve to reduce the formation of Cr oxides or Cr mixed oxides on the steel surface after annealing.
  • the chromium content is fixed to contents of 0.250 to 0.400 wt. %.
  • Molybdenum (Mo) The addition of molybdenum results, in a similar manner to the addition of chromium and manganese, in the improvement of hardenability. The perlite and bainite conversion is shifted towards longer times and the martensite starting temperature is decreased. At the same time, molybdenum is a strong carbide forming agent which allows the production of finely distributed mixed carbides, inter alia also with titanium. Moreover, molybdenum considerably increases the tempering resistance so that no losses in strength are to be expected in the hot-dip bath. Molybdenum also acts by mixed crystal hardening but is less effective than manganese and silicon.
  • the content of molybdenum is set between more than 0.200 to 0.300 wt. %.
  • the Mo content is advantageously set to a range between more than 0.200 and 0.250 wt. %.
  • the alloy concept according to the invention has proven to be advantageous for the alloy concept according to the invention to have a total content of Mo+Cr of ⁇ 0.650 wt. %.
  • Copper (Cu): The addition of copper can increase the tensile strength and the potential hardness increase. In conjunction with nickel, chromium and phosphorous, copper can form a protective oxide layer on the surface which can considerably reduce the corrosion rate.
  • copper can form, at the grain boundaries, noxious oxides which can produce negative effects particularly for hot-deformation processes. Therefore, the content of copper is fixed to ⁇ 0.050 wt. % and is thus limited to quantities which are unavoidable in steel production.
  • Vanadium (V) Since, in the case of the present alloy concept, an addition of vanadium is not necessary, the content of vanadium is limited to unavoidable quantities associated with steel.
  • Aluminium is generally alloyed to the steel in order to bind the oxygen and nitrogen dissolved in the iron. Oxygen and nitrogen are thus convertered into aluminium oxides and aluminium nitrides. These precipitations can effect grain refinement by increasing the nucleation sites and can thus increase the toughness properties and strength values.
  • Titanium nitrides have a lower enthalpy of formation and are formed at higher temperatures.
  • aluminium like silicon shifts the formation of ferrite towards shorter times and thus permits the formation of sufficient ferrite in the dual-phase steel. It also suppresses the formation of carbide and thus results in a delayed conversion of the austenite. For this reason, aluminium is also used as an alloy element in residual austenite steels (TRIP steels) in order to substitute a part of the silicon.
  • TRIP steels residual austenite steels
  • the aluminium content is limited to 0.005 to a maximum of 0.060 wt. % and is added for the purpose of calming the steel.
  • Niobium acts differently in steel. During hot rolling on the production line, it delays the recrystallisation by virtue of the formation of very finely distributed precipitations, whereby the nucleation site density is increased and a finer grain is produced after the conversion. The proportion of dissolved niobium also acts to inhibit recrystallisation.
  • the precipitations act to increase strength in the final product. They can be carbides or carbonitrides. They are often mixed carbides, into which titanium is also incorporated. This effect begins from 0.005 wt. % and is most pronounced from 0.010 wt. % niobium. Moreover, the precipitations prevent the grain growth during the (partial) austenitisation in the hot-dip galvanising. Above 0.060 wt. % niobium, no additional effect is to be expected. Contents of 0.025 to 0.045 wt. % have proven to be advantageous.
  • Boron is an extremely effective alloying agent for achieving variable degrees of thinning by rolling.
  • Tests have surprisingly shown that the range for the addition of boron, which is very narrow according to the invention, has a pronounced effect with regard to the uniformity of the mechanical properties of the produced cold strips with a variable degree of thinning by rolling in the subsequent processing.
  • This pronounced effect initially results in the possibility of setting, instead of with a relatively constant degree of thinning by rolling during cold rolling, defined characteristic value ranges after the process steps ( FIG. 7 a , 7 b or 7 c ) also with the material with variable degrees of thinning by rolling on the basis of a master hot strip thickness or on the basis of a master cold strip thickness.
  • boron is an effective element for increasing hardenability which becomes effective even in very small quantities.
  • the martensite starting temperature remains unaffected thereby.
  • boron In order to become effective, boron must be present in a solid solution. Since it has a high affinity to nitrogen, the nitrogen must initially be removed, preferably by the stoichiometrically required quantity of titanium. By reason of its low solubility in iron, the dissolved boron preferably becomes attached to the austenite grain boundaries. At this location, it partially forms Fe—B carbides which are coherent and decrease the grain boundary energy. Both effects act in such a way as to delay the formation of ferrite and perlite and thus increase the hardenability of the steel.
  • the boron content for the alloy concept according to the invention is fixed to values of more than 0.0005 to 0.0010 wt. %, advantageously to values ⁇ 0.0009 or optimally to >0.0006 wt. % to ⁇ 0.0009 wt. %.
  • Nitrogen (N) can be both an alloy element and an associated element from steel production. Excessively high contents of nitrogen produce an increase in strength associated with a rapid loss of toughness as well as ageing effects.
  • N Nitrogen
  • titanium and niobium fine grain hardening can be achieved via titanium nitrides and niobium (carbo)nitrides. Moreover, coarse grain formation is suppressed during reheating prior to hot rolling.
  • the N content is therefore fixed to values of ⁇ 0.0020 to ⁇ 0.0120 wt. %.
  • niobium and titanium For the total contents of niobium and titanium, contents of ⁇ 0.100 wt. % have proven to be advantageous and owing to the basic interchangeability of niobium and titanium up to a minimum niobium content of 0.0010 wt. % and for cost reasons contents of 0.090 wt. % have proven to be particularly advantageous.
  • total contents of ⁇ 0.102 wt. % have proven to be advantageous and total contents of ⁇ 0.092 wt. % have proven to be particularly advantageous. Higher contents no longer have the effect of providing improvement in terms of the invention.
  • Calcium (Ca) An addition of calcium in the form of calcium-silicon mixed compounds causes deoxidation and desulphurisation of the molten phase in steel production. For instance, reaction products are converted into slag and the steel is cleaned. According to the invention, the increased level of purity results in improved properties in the end product.
  • a Ca content of ⁇ 0.0010 to ⁇ 0.0060 wt. % is set. Contents of ⁇ 0.0030 wt. % have proven to be advantageous.
  • the annealing temperatures for the dual-phase microstructure to be obtained are between approximately 700 and 950° C. for the steel according to the invention so that, depending upon the temperature range, a partially austenitic (two-phase region) or a fully austenitic microstructure (austenite region) is achieved.
  • the continuously annealed and occasionally hot-dip-finished material can be manufactured in the temper-rolled (cold-post-rolled) or non-temper-rolled state and/or in the stretch-bend-straightened or non-stretch-bend-straightened state and also in the heat-treated state (ageing).
  • the steel strips consisting of the alloy composition according to the invention are characterised during further processing by high edge crack insensitivity and a high bending angle.
  • the very small characteristic value differences of the steel strip longitudinally and transversely with respect to its rolling direction are advantageous for the subsequent material use.
  • the cutting of blanks from a strip can be performed independently of the rolling direction (e.g. transversely, longitudinally and diagonally or at an angle with respect to the rolling direction) and therefore the waste is minimised.
  • the hot strip is produced according to the invention with end rolling temperatures in the austenitic region above Ar3 and at reeling temperatures above the bainite starting temperature.
  • a blank is cut from a steel strip according to the invention, said blank is then heated to a temperature above Ac3, the heated blank is formed into a component and then hardened in a deforming tool or in air.
  • the steel according to the invention has the property that the hardening occurs even during cooling in stationary air so that separate cooling of the deforming tool can be omitted.
  • the microstructure of the steel is converted to the austenitic range by heating, preferably to temperatures over 950° C. under a protective gas atmosphere. During subsequent cooling in air or protective gas, a martensitic microstructure is formed for a high-strength component.
  • Subsequent tempering facilitates the reduction of intrinsic tensions in the hardened component.
  • the hardness of the component is reduced such that the required toughness values are achieved.
  • FIG. 1 (schematically) the process chain for the production of a strip consisting of the steel according to the invention
  • FIG. 2 the chemical composition (examples 1 to 5) of the steel according to the invention
  • FIG. 3 (schematically) the time-temperature graph of the process steps of hot rolling and cold rolling and continuous annealing (optional hot-dip finishing) by way of example for the steel according to the invention
  • FIG. 4 a the ratio of hot strip thickness to cold strip thickness (end thickness) by way of example for the prior art
  • FIG. 4 b the degree of thinning by rolling in relation to cold strip thickness (end thickness) by way of example for the prior art
  • FIG. 5 a the ratio of hot strip thickness to cold strip thickness (end thickness) by way of example for the steel according to the invention in the case of a master hot strip thickness
  • FIG. 5 b the degree of thinning by rolling in relation to cold strip thickness (end thickness) by way of example for the steel according to the invention in the case of a master hot strip thickness
  • FIG. 6 a the ratio of hot strip thickness to cold strip thickness (end thickness) by way of example for the steel according to the invention in the case of a master cold strip thickness (a specific end thickness of the cold-rolled steel strip to be achieved),
  • FIG. 6 b the degree of thinning by rolling in relation to cold strip thickness (end thickness) by way of example for the steel according to the invention in the case of a master cold strip thickness (a specific end thickness of the cold-rolled steel strip to be achieved),
  • FIG. 7 a method 1 , temperature-time curves (annealing variants shown schematically),
  • FIG. 7 b method 2 temperature-time curves (annealing variants shown schematically),
  • FIG. 7 c method 3 temperature-time curves (annealing variants shown schematically),
  • FIG. 8 a material characteristic values by way of example for the steel according to the invention in the case of a master hot strip thickness of 2.30 mm with a variable degree of thinning by rolling after cold rolling, transversely and longitudinally to the rolling direction (in table form),
  • FIG. 8 b material characteristic values by way of example for the steel according to the invention in the case of a master hot strip thickness of 2.30 mm with a variable degree of thinning by rolling after cold rolling, transversely to the rolling direction (graphically),
  • FIG. 8 c material characteristic values by way of example for the steel according to the invention in the case of a master hot strip thickness of 2.30 mm with a variable degree of thinning by rolling after cold rolling, longitudinally to the rolling direction (graphically) and
  • FIG. 9 a schematic diagram of the sub-peritectic grade chain (basic alloy is HCT500XD, increasing the addition quantities and addition substances brings about the change to other grades or corresponding sheet thickness ranges).
  • FIG. 1 schematically illustrates the process chain for the production of a strip from the steel according to the invention.
  • the different process routes relating to the invention are shown. Up until hot rolling (end rolling temperature) the process route is the same for all steels according to the invention, thereafter, deviating process routes are followed depending on the desired results.
  • the acid-cleaned hot strip can be galvanised or can be cold-rolled and galvanised with different degrees of thinning by rolling.
  • Soft annealed hot strip or soft annealed cold strip can also be cold-rolled and galvanised.
  • Material can also optionally be processed without a hot-dip finishing procedure, i.e. only within the scope of continuous annealing with and without subsequent electrolytic galvanising. From the optionally coated material a complex component can now be produced. Subsequently thereto, the hardening process takes place in which cooling in air is effected according to the invention. Optionally, a tempering step can conclude the thermal treatment of the component.
  • FIG. 2 illustrates, in examples 1 to 5, the alloy compositions of the tested steels, in dependence upon the cold strip thickness (end thickness) to be produced. From a master hot strip thickness of 2.30 mm cold strips of different thicknesses have been produced. In dependence upon the strip thickness to be produced, example 1 shows the alloy composition for a cold strip thickness of 2.00 mm, example 2 for a cold strip with a thickness of 1.80 mm, example 3 for a cold strip thickness of 1.50 mm, example 4 for a cold strip with a thickness of 1.20 mm and example 5 for a cold strip with a thickness 1.00 mm.
  • FIG. 3 schematically illustrates the time-temperature curve of the process steps of hot rolling and continuous annealing of strips made from the alloy composition according to the invention.
  • FIG. 4 illustrates the ratio of hot strip thickness to cold strip thickness (end thickness) ( FIG. 4 a ) and degree of thinning by rolling in relation to cold strip thickness (end thickness) ( FIG. 4 b ) as an example of the prior art.
  • FIG. 5 illustrates the ratio of hot strip thickness to cold strip thickness (end thickness) ( FIG. 5 a ) and degree of thinning by rolling in relation to cold strip thickness (end thickness) ( FIG. 5 b ) as an example of the steel according to the invention in the case of a master hot strip thickness.
  • FIG. 6 illustrates the ratio of hot strip thickness to cold strip thickness ( FIG. 6 a ) and degree of thinning by rolling in relation to cold strip thickness (end thickness) ( FIG. 6 b ) as an example of the steel according to the invention in the case of a master cold strip thickness (of a specific end thickness of the cold-rolled steel strip to be achieved).
  • FIG. 7 schematically illustrates three variations of the temperature-time curves according to the invention in the case of the annealing treatment and cooling and of austenitisation conditions which differ in each case.
  • Method 1 presents the annealing and cooling of the produced steel strip cold-rolled to the end thickness in a continuous annealing installation.
  • the strip is heated to a temperature in the range of about 700 to 950° C. (Ac1 to Ac3).
  • the annealed steel strip is then cooled from the annealing temperature at a cooling rate between approximately 15 and 100° C./s to an intermediate temperature (ZT) of approximately 200 to 250° C.
  • ZT intermediate temperature
  • This schematic illustration does not show a second intermediate temperature (approximately 300 to 500° C.).
  • the steel strip is then cooled in air at a cooling rate between approximately 2 and 30° C./s until room temperature (RT) is reached or cooling at a cooling rate between approximately 15 and 100° C./s is maintained until room temperature is reached.
  • RT room temperature
  • Method 2 ( FIG. 7 b ) shows the process according to method 1 but the cooling of the steel strip is briefly interrupted during passage through the hot-dipping vessel for the purposes of a hot-dip finishing procedure, in order then to continue the cooling at a cooling rate between approximately 15 and 100° C./s to an intermediate temperature of approximately 200 to 250° C.
  • the steel strip is then cooled in air at a cooling rate between approximately 2 and 30° C./s until reaching room temperature.
  • Method 2 corresponds to annealing, e.g. hot-dip galvanising with combined direct fired furnace and radiant tube furnace, as depicted in FIG. 7 b.
  • Method 3 ( FIG. 7 c ) likewise shows the process according to method 1 in a hot-dip finishing procedure but cooling of the steel strip is interrupted by a short pause (approximately 1 to 20 s) at an intermediate temperature in the range of approximately 200 to 400° C. and heating is effected to the temperature (ST) required for the hot-dip finishing procedure (approximately 400 to 470° C.).
  • ST temperature required for the hot-dip finishing procedure
  • the steel strip is then cooled to an intermediate temperature of approximately 200 to 250° C.
  • the subsequent cooling in air of the steel strip takes place at a cooling rate of approximately 2 and 30° C./s until room temperature is reached.
  • Method 3 corresponds e.g. to a process being carried out in a continuous annealing installation, as depicted in FIG. 7 c .
  • reheating of the steel is optionally achieved directly prior to the zinc bath.
  • FIG. 8 shows in table form ( FIG. 8 a ) and graphically the material characteristic values in an exemplified manner for the steel according to the invention in the case of a mast hot strip thickness of 2.30 mm with a variable degree of thinning by rolling after cold rolling, transversely ( FIG. 8 b ) and longitudinally ( FIG. 8 c ) to the rolling direction as an example of processing according to method 3 ( FIG. 7 c ).
  • the degrees of thinning by rolling thus vary from 13% to 35% for cold rolling in a single cold rolling step and 55% (30%+25%) or 68% (35+33%) when cold rolling is effected twice. It is impressively shown that both for very low degrees of deformation and also for the very different degrees of deformation relatively uniform values provided with a conventional fluctuation range are achieved for the tensile strength and yield strength, transversely and longitudinally to the rolling direction.
  • the material was previously hot-rolled at an end rolling desired temperature of 910° C. and reeled at a desired reeling temperature of 650° C. with a master hot strip thickness of 2.30 mm and cold-rolled after the acid-cleaning without additional heat treatment (such as e.g. batch-type annealing) to 2.00 mm in one pass (degree of thinning by rolling 13%).
  • the yield strength ratio Re/Rm in the transverse direction was 65%.
  • the yield strength ratio Re/Rm in the longitudinal direction was 63%.
  • additional heat treatment such as e.g. batch-type annealing
  • the yield strength ratio Re/Rm in the transverse direction was 66%.
  • the yield strength ratio Re/Rm in the longitudinal direction was 64%.
  • additional heat treatment such as e.g. batch-type annealing
  • the yield strength ratio Re/Rm in the transverse direction was 68%.
  • the yield strength ratio Re/Rm in the longitudinal direction was 67%.
  • the material was previously hot-rolled at an end rolling desired temperature of 910° C. and reeled at a desired reeling temperature of 650° C. with a master hot strip thickness of 2.30 mm and cold-rolled after the acid-cleaning without additional heat treatment (such as e.g. batch-type annealing) to 1.20 mm in two passes (degree of thinning by rolling 30% and 25%).
  • the yield strength ratio Re/Rm in the transverse direction was 60%.
  • the yield strength ratio Re/Rm in the longitudinal direction was 60%.
  • additional heat treatment such as e.g. batch-type annealing
  • the yield strength ratio Re/Rm in the transverse direction was 67%.
  • the yield strength ratio Re/Rm in the longitudinal direction was 68%.
  • FIG. 9 shows the schematic diagram of a sub-peritectic grade chain with the basic analysis of a HCT500XD which is based on the fact that by increasing the addition quantities and addition substances a change to other grades with respect to their minimum tensile strength (classes) or corresponding sheet thickness ranges is possible.
US15/743,229 2015-07-10 2016-07-07 Ultra high strength multi-phase steel and method for producing a cold-rolled steel strip therefrom Abandoned US20190127821A1 (en)

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Cited By (5)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
CN111008477A (zh) * 2019-12-06 2020-04-14 马鞍山钢铁股份有限公司 一种基于冷轧镀锌带钢力学性能调整工艺参数的方法
CN111145842A (zh) * 2020-02-27 2020-05-12 西安交通大学 一种面向高速变形过程的多相金属材料本构行为构建方法
CN111270166A (zh) * 2020-03-30 2020-06-12 武汉钢铁有限公司 一种屈服强度大于830MPa级非调质薄规格超高钢及其生产方法
CN111519016A (zh) * 2020-04-30 2020-08-11 江苏省沙钢钢铁研究院有限公司 热镀锌板的制备方法、以及热镀锌板
US10870900B2 (en) * 2017-06-07 2020-12-22 A. Finkl & Sons Co. High toughness martensitic stainless steel and reciprocating pump manufactured therewith

Families Citing this family (10)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
DE102017123236A1 (de) 2017-10-06 2019-04-11 Salzgitter Flachstahl Gmbh Höchstfester Mehrphasenstahl und Verfahren zur Herstellung eines Stahlbandes aus diesem Mehrphasenstahl
MX2020007610A (es) * 2018-01-17 2020-09-03 Jfe Steel Corp Lamina de acero electrogalvano-recocida de alta resistencia y metodo para la fabricacion de la misma.
CN110616303A (zh) * 2018-06-19 2019-12-27 宝山钢铁股份有限公司 一种980MPa级以上冷轧或镀锌双相钢板的制造方法
CN108914000B (zh) * 2018-07-24 2020-10-02 唐山钢铁集团有限责任公司 一种抗拉强度780MPa级冷轧双相钢及其生产方法
DE102018122901A1 (de) * 2018-09-18 2020-03-19 Voestalpine Stahl Gmbh Verfahren zur Herstellung ultrahochfester Stahlbleche und Stahlblech hierfür
DE102020110319A1 (de) 2020-04-15 2021-10-21 Salzgitter Flachstahl Gmbh Verfahren zur Herstellung eines Stahlbandes mit einem Mehrphasengefüge und Stahlband hinzu
DE102021108448A1 (de) 2021-04-01 2022-10-06 Salzgitter Flachstahl Gmbh Stahlband aus einem hochfesten Mehrphasenstahl und Verfahren zur Herstellung eines derartigen Stahlbandes
CN113308649B (zh) * 2021-05-14 2022-10-18 唐山钢铁集团有限责任公司 低屈强比1000MPa级冷轧双相带钢及其生产方法
CN113621887B (zh) * 2021-07-23 2022-06-10 山东钢铁集团日照有限公司 一种低成本打包带用冷轧钢带原料的生产方法
DE102022125128A1 (de) 2022-09-29 2024-04-04 Salzgitter Flachstahl Gmbh Verfahren zur Herstellung eines Stahlbandes aus einem hochfesten Mehrphasenstahl und entsprechendes Stahlband

Citations (4)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US20040238080A1 (en) * 2001-08-29 2004-12-02 Sven Vandeputte Ultra high strength steel composition, the process of production of an ultra high strength steel product and the product obtained
US20070029015A1 (en) * 2003-09-30 2007-02-08 Naoki Yoshinaga High-yield-ratio high-strength thin steel sheet and high-yield-ratio high-strength hot-dip galvanized thin steel sheet excelling in weldability and ductility as well as high-yield ratio high-strength alloyed hot-dip galvanized thin steel sheet and process for producing the same
US20140034196A1 (en) * 2011-01-26 2014-02-06 Andreas Wedemeier Hoherfester mehrphasenstahl mit ausgezeichneten umformeigenschaften high strength multi-phase steel having excellent forming properties
US20190316222A1 (en) * 2014-11-18 2019-10-17 Salzgitter Flachstahl Gmbh Ultra high-strength air-hardening multiphase steel having excellent processing properties, and method for manufacturing a strip of said steel

Family Cites Families (18)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
FR2727431B1 (fr) * 1994-11-30 1996-12-27 Creusot Loire Procede d'elaboration d'un acier au titane et acier obtenu
DE19610675C1 (de) 1996-03-19 1997-02-13 Thyssen Stahl Ag Mehrphasenstahl und Verfahren zu seiner Herstellung
DE10037867A1 (de) 1999-08-06 2001-06-07 Muhr & Bender Kg Verfahren zum flexiblen Walzen eines Metallbandes
EP1832667A1 (de) * 2006-03-07 2007-09-12 ARCELOR France Herstellungsverfahren von Stahlblechen mit hoher Festigkeit, Duktilität sowie Zähigkeit und so hergestellte Bleche.
KR100851189B1 (ko) * 2006-11-02 2008-08-08 주식회사 포스코 저온인성이 우수한 초고강도 라인파이프용 강판 및 그제조방법
DE102006053819A1 (de) * 2006-11-14 2008-05-15 Thyssenkrupp Steel Ag Verfahren zum Herstellen eines Bauteil durch Warmpresshärten und hochfestes Bauteil mit verbesserter Bruchdehnung
EP2031081B1 (de) 2007-08-15 2011-07-13 ThyssenKrupp Steel Europe AG Dualphasenstahl, Flachprodukt aus einem solchen Dualphasenstahl und Verfahren zur Herstellung eines Flachprodukts
ES2387040T3 (es) 2007-08-15 2012-09-12 Thyssenkrupp Steel Europe Ag Acero de doble fase, producto plano de un acero de doble fase de este tipo y procedimiento para la fabricación de un producto plano
JP4924730B2 (ja) * 2009-04-28 2012-04-25 Ｊｆｅスチール株式会社 加工性、溶接性および疲労特性に優れる高強度溶融亜鉛めっき鋼板およびその製造方法
PL2524972T3 (pl) * 2010-01-13 2017-06-30 Nippon Steel & Sumitomo Metal Corporation Sposób wytwarzania blachy stalowej grubej o dużej wytrzymałości i doskonałej odkształcalności
JP5333298B2 (ja) * 2010-03-09 2013-11-06 Ｊｆｅスチール株式会社 高強度鋼板の製造方法
EP2524970A1 (de) * 2011-05-18 2012-11-21 ThyssenKrupp Steel Europe AG Hochfestes Stahlflachprodukt und Verfahren zu dessen Herstellung
EP2729590B1 (de) * 2011-07-10 2015-10-28 Tata Steel IJmuiden BV Heissgewalzter hochfester bandstahl mit verbesserter haz-weichmacherbeständigkeit und herstellungsverfahren für diesen stahl
DE102012002079B4 (de) 2012-01-30 2015-05-13 Salzgitter Flachstahl Gmbh Verfahren zur Herstellung eines kalt- oder warmgewalzten Stahlbandes aus einem höchstfesten Mehrphasenstahl
CN104245971B (zh) * 2012-03-30 2017-09-12 奥钢联钢铁有限责任公司 高强度冷轧钢板和生产该钢板的方法
JP2013227624A (ja) * 2012-04-25 2013-11-07 Jfe Steel Corp 加工性に優れる高強度冷延鋼板の製造方法
DE102012013113A1 (de) * 2012-06-22 2013-12-24 Salzgitter Flachstahl Gmbh Hochfester Mehrphasenstahl und Verfahren zur Herstellung eines Bandes aus diesem Stahl mit einer Mindestzugfestigkleit von 580MPa
DE102013013067A1 (de) * 2013-07-30 2015-02-05 Salzgitter Flachstahl Gmbh Siliziumhaltiger, mikrolegierter hochfester Mehrphasenstahl mit einer Mindestzugfestigkeit von 750 MPa und verbesserten Eigenschaften und Verfahren zur Herstellung eines Bandes aus diesem Stahl

Patent Citations (4)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US20040238080A1 (en) * 2001-08-29 2004-12-02 Sven Vandeputte Ultra high strength steel composition, the process of production of an ultra high strength steel product and the product obtained
US20070029015A1 (en) * 2003-09-30 2007-02-08 Naoki Yoshinaga High-yield-ratio high-strength thin steel sheet and high-yield-ratio high-strength hot-dip galvanized thin steel sheet excelling in weldability and ductility as well as high-yield ratio high-strength alloyed hot-dip galvanized thin steel sheet and process for producing the same
US20140034196A1 (en) * 2011-01-26 2014-02-06 Andreas Wedemeier Hoherfester mehrphasenstahl mit ausgezeichneten umformeigenschaften high strength multi-phase steel having excellent forming properties
US20190316222A1 (en) * 2014-11-18 2019-10-17 Salzgitter Flachstahl Gmbh Ultra high-strength air-hardening multiphase steel having excellent processing properties, and method for manufacturing a strip of said steel

Cited By (5)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US10870900B2 (en) * 2017-06-07 2020-12-22 A. Finkl & Sons Co. High toughness martensitic stainless steel and reciprocating pump manufactured therewith
CN111008477A (zh) * 2019-12-06 2020-04-14 马鞍山钢铁股份有限公司 一种基于冷轧镀锌带钢力学性能调整工艺参数的方法
CN111145842A (zh) * 2020-02-27 2020-05-12 西安交通大学 一种面向高速变形过程的多相金属材料本构行为构建方法
CN111270166A (zh) * 2020-03-30 2020-06-12 武汉钢铁有限公司 一种屈服强度大于830MPa级非调质薄规格超高钢及其生产方法
CN111519016A (zh) * 2020-04-30 2020-08-11 江苏省沙钢钢铁研究院有限公司 热镀锌板的制备方法、以及热镀锌板

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