ULTRA HIGH STRENGTH MULTI-PHASE STEEL AND METHOD FOR PRODUCING A COLD-ROLLED STEEL STRIP THEREFROM
The invention relates to an ultrahigh strength multiphase steel with a dual-phase microstructure or a complex-phase microstructure and with a low content of residual austenite in particular for lightweight vehicle construction, to a method for producing cold-rolled steel strips from such a steel, and to steel strips produced using said method. The aim of the invention is a novel alloy concept with which the processing window for the continuous annealing of cold strips can be extended. According to the invention, this is achieved by a multiphase steel with a minimum tensile strength of 980 MPa containing (contents in wt. %): C≥0.075 to ≤0.115, Si≥0.400 to ≤0.500, Mn≥1.900 to ≤2.350, Cr≥0.250 to ≤0.400, Al≥0.005 to ≤0.060, N≥0.0020 to ≤0.0120, S≤0.0020, Nb≥0.005 to ≤0.060, Ti≥0.005 to ≤0.060, B>0.0005 to ≤0.0010, Mo>0.200 to ≤0.300, Ca>0.0010 to ≤0.0060, Cu≤0.050, Ni≤0.050, and residual iron, including customary steel-accompanying smelting-related impurities, wherein the total content of Mn+Si+Cr is ≥2.500 to ≤3.250 with a view to a processing window which is as wide as possible during the annealing process, in particular during the continuous annealing process, of cold strips of said steel.
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.
In particular, 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.
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. On the one hand, 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.
These newly developed steels must satisfy not only the required weight reduction but also the high material requirements in relation to elasticity limit, tensile strength and elongation at fracture and bake hardening index and also the component requirements according to edge crack insensitivity, energy absorption and defined solidification relating to the work hardening effect and the bake hardening effect. Furthermore, it is necessary to ensure good processability. This relates both to the processes performed by the car producer, e.g. deforming, welding or lacquering and also the manufacturing processes performed by the suppliers of semi-finished products, such as e.g. surface finishing by means of metallic or organic coating.
The combination of properties required of the steel material ultimately represents a component-specific compromise of individual properties. Therefore, in vehicle construction dual-phase steels are being increasingly used, said steels consisting of a ferritic basic microstructure, into which a martensitic second phase and possibly a further phase containing bainite and residual austenite are incorporated. The characteristic processing properties of the dual-phase steels, such as a very low yield strength ratio with very high tensile strength at the same time, strong cold solidification and good cold deformability are sufficiently known.
Also used to an increasing extent in vehicle construction are 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.
These steels are currently being used in structural components, chassis components and crash-relevant components, as well as flexibly cold-rolled strips. This Tailor Rolled Blank lightweight construction technology (TRB®) permits a significant weight reduction as a result of the load-adapted selection of the sheet thickness over the component length.
However, with currently known alloys and available continuous annealing installations for greatly varying sheet thicknesses the production of TRB®s having a multi-phase microstructure is not possible without limitations, such as e.g. for heat treatment prior to cold rolling. In regions of different sheet thickness, 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.
If thin sheets are to be produced, economic reasons dictate that 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. In dependence upon the alloy composition and the strip cross-section, 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.
Typically, 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. In order to achieve the technological characteristic values required by the standards, on average a degree of cold rolling of 50% is required in the case of continuous annealing treatment in order to ensure normal recrystallisation. In the case of classic steels, deviations to higher or lower amounts lead to fluctuating technological characteristic values, as described in the case of the TRB®'s.
In order to achieve a fine-grain microstructure after the continuous annealing procedure, it is known that a minimum degree of cold rolling is set in dependence upon the recrystallisation temperature, in order to set a corresponding dislocation density for the recrystallisation annealing.
If the degree of thinning by rolling is too low (even in local regions), the critical threshold for recrystallisation cannot be overcome and the microstructure cannot be refined. After recrystallisation, 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. During cooling from the furnace temperature, grains of different sizes can convert into different phase components and ensure further inhomogeneity.
In order to achieve the respectively required microstructure, 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.
If, by reason of high corrosion protection requirements, the surface of the cold strip is to be hot-dip galvanised, 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.
In the case of continuous annealing of hot-rolled or cold-rolled steel strips using alloy concepts, which are known e.g. from the documents EP 2 028 282 A1 and EP 2 031 081 A1, for ultra high strength dual-phase steels having minimum tensile strengths of approximately 980 MPa, the problem exists that only a small process window is provided for the annealing parameters. Therefore, even in the case of minimal cross-sectional changes (thickness, width) adaptations of the process parameters are required in order to achieve uniform mechanical properties.
In the case of extended process windows, with the process parameters being the same the required strip properties are possible even in the case of larger cross-sectional changes of the strips which are to be annealed.
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. In the case of alloy compositions with excessively small process windows, during continuous annealing 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.
Therefore, 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.
The problem of an excessively narrow process window becomes particularly serious in the annealing treatment of cold strips which have excessively low or excessively high degrees of thinning by rolling during cold rolling, as well as in the annealing treatment of strips having sheet thicknesses which vary over the strip length for the production of load-optimised components consisting of hot or cold strip.
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.
When the known alloy concepts are applied for the group of multi-phase steels, 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.
In the case of flexibly rolled cold strips consisting of multi-phase steels of known compositions, 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.
Therefore, with the known alloy concepts it is currently possible to produce only steels of one strength class with defined hot strip and cold strip cross-sectional regions (sheet thickness and strip width) so that changed alloy concepts are required for different strength classes and/or cross-sectional regions.
It is also prior art that an increase in strength is achieved by increasing the amount of carbon and/or silicon and/or manganese and an increase in strength is achieved by the structural settings and the mixed crystal solidification (mixed crystal hardening).
However, by increasing the amount of the aforementioned elements, the material processing properties deteriorate to an increasing extent, e.g. during welding, deforming and hot-dip finishing.
However, steel production is witnessing a trend for reducing the carbon and/or manganese content in order to achieve improved cold processing and better usage properties.
One example is the hole expansion test for describing and quantifying the edge crack behaviour. With correspondingly optimised grades, the steel user expects higher values than with the standard material. However, welding suitability, characterised by the carbon equivalent, also comes further into focus.
The automotive industry is increasingly demanding grades of steel which have requirements in terms of yield strength (Re)/elasticity limit (Rp0.2) which differ considerably depending upon the application. This leads to steel developments with a comparatively large yield strength interval with a typical tensile strength interval.
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.
Therefore, 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.
Moreover, 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.
According to the teaching of the invention, 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 remainder is iron, including typical steel-associated, smelting-related impurities, wherein the total content of Mn+Si+Cr is ≥2.500 to ≤3.250 wt. % with regard to a processing window which is as wide as possible during the annealing, in particular the continuous annealing, of cold strips of this steel.
By means of the alloy concept according to the invention, 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.
Furthermore, 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.
Furthermore, 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. However, it is also advantageously possible to produce a cold strip thickness, which is to be achieved, in a similar manner from different hot strip thicknesses. This considerably increases the flexibility in manufacturing and also reduces the production costs.
Moreover, 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.
According to the invention, it is possible to use 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.
In this case, according to the invention 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.
In order to achieve uniform mechanical properties, it has proven to be advantageous that 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.
It has also proven to be advantageous that the chemical composition of the multi-phase steel is selected in dependence upon the end thickness to be achieved, as follows:
end thickness 0.50 mm to 1.00 mm inclusive: C content ≤0.100 wt. % and carbon equivalent CEV (IIW)≤0.62%,
end thickness above 1.00 mm to 2.00 mm inclusive: C content ≤0.105 wt. % and the carbon equivalent CEV (IIW)≤0.64%,
end thickness above 2.00 mm to 3.00 mm inclusive: C content ≤0.115 wt. % and the carbon equivalent CEV (IIW)≤0.66%.
In this case, the carbon equivalent (CEV) (IIW) is calculated according to the formula
CEV(IIW)=% C+% Mn/6+(% Cu+% Ni)/15+(% Cr+% Mo+% V)/5.
It has also been shown to be advantageous that the chemical composition of the multi-phase steel is selected in dependence upon the end thickness to be achieved, as follows:
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. %.
This applies to the continuous annealing of successive strips having different strip cross-sections and also to strips having a varying sheet thickness over the strip length or strip width. For example, it is thus possible to process cold strips having variable degrees of thinning by rolling during cold rolling.
If, according to the invention, higher strength and ultra high strength cold strips, which are produced in the continuous annealing method, are produced from multi-phase steel having varying sheet thicknesses, it is advantageously possible to produce loading-optimised components from this material by deformation technology.
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).
At the same time, it is possible by specifically varying the process parameters to set the microstructure proportions such that steels in different strength classes, e.g. with yield strengths between 550 MPa and 950 MPa and tensile strengths between 980 MPa and 1140 MPa, can be produced.
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.
In the case of a continuous annealing installation without subsequent hot-dip finishing, 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. Optionally, 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,
In the case of a heat treatment as part of a hot-dip finishing procedure, there are two temperature control options. 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,
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,
In the case of the known dual-phase steels, it is not only carbon but also manganese, chromium and silicon which are responsible for the conversion of austenite to martensite. Only the inventive combination of the elements, which are alloyed thereto at the limits indicated, of carbon, silicon, manganese, nitrogen, molybdenum and chromium as well as niobium, titanium and above all boron ensures, on the one hand, the required mechanical properties, such as minimum tensile strengths of 980 MPa with at the same time a considerably widened process window during the continuous annealing procedure.
It is also characteristic of the material that by means of the addition of manganese with increasing percentages by weight the ferrite region is displaced for longer periods of time and at lower temperatures during cooling. The proportions of ferrite are reduced to a greater or lesser extent by increased proportions of bainite depending upon process parameters.
By setting a low carbon content of ≤0.115 wt. %, 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.
The effect of the elements in the alloy according to the invention will be described in greater detail hereinafter. Associated elements are unavoidable and, if necessary, are taken into consideration in the analysis concept in terms of their effect.
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. As a result, 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.
On the one hand, procedural approaches such as vacuum treatment and, on the other hand, analytical approaches are provided in order to reduce the oxygen. By adding specific alloy elements, the oxygen can be converted into less dangerous states. For instance, it is generally conventional to remove the oxygen in the course of deoxidation of the steel with manganese, silicon and/or aluminium. However, the resulting oxides can produce negative properties as flaws in the material.
Therefore, for the reasons stated above the oxygen content in the steel should be kept as small as possible.
Phosphorous (P) 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.
However, in some steels owing to the low costs and significant increase in strength 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.
For the reasons stated above, the phosphorous content in the steel according to the invention is limited to quantities which are unavoidable in the production of steel. Preferably, P should be ≤0.020 wt. %.
Sulphur (S), like phosphorous, 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. Primarily in the case of a diffusion-controlled conversion this produces a microstructure of pronounced lines and, in the case of a highly pronounced line formation, can result in impaired mechanical properties, such as e.g. pronounced martensite lines instead of distributed martensite islands, anisotropic material behaviour, reduced elongation at fracture.
For the reasons stated above, 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.
The effect of the alloy elements will be discussed in greater detail hereinafter.
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 α-iron and is at most 2.06% in the γ-iron. Carbon, in dissolved form, increases the hardenability of steel considerably and is therefore essential for forming a sufficient quantity of martensite. However, excessively high carbon contents increase the hardness difference between ferrite and martensite and restrict weldability.
In order to meet the requirements of e.g. high hole expansion and bending angles as well as improved weldability, the steel according to the invention contains carbon contents of ≤0.115 wt. %.
The different solubility of the carbon in the phases makes pronounced diffusion procedures necessary during the phase conversion, which procedures can result in very different kinetic conditions. Moreover, carbon increases the thermodynamic stability of the austenite, which is demonstrated in the phase diagram in an extension of the austenite region at lower temperatures. As the forcibly dissolved carbon content in the martensite increases, the lattice distortions and, associated therewith, the strength of the phase produced without diffusion increase.
Carbon also forms carbides. A microstructure phase which occurs almost in every steel is cementite (Fe3C). However, 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. %)
During casting, 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.
In a manner characteristic of the material, the addition of silicon produces a strong mixed crystal solidification. Roughly estimated, 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%. The latter situation is attributed inter alia to the fact that silicon reduces the solubility of carbon in the ferrite and increases the activity of carbon in the ferrite, thus preventing the formation of carbides which, as brittle phases, reduce the ductility which, in turn, improves deformability. The low strength-increasing effect of silicon within the span of the steel according to the invention provides the basis for a wide process window.
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. During hot rolling, this provides a basis for improved cold-rollability. During hot-dip finishing, 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.
The addition of 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. However, 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.
During hot rolling, it is known that more highly silicon-alloyed steels will result in the formation of strongly adhering red scale and in an increased risk rolled-in scale, which can have an influence upon the subsequent acid-cleaning result and acid-cleaning productivity. This effect could not be established in the steel according to the invention at 0.400 to 0.500% silicon, when the acid-cleaning is advantageously performed with hydrochloric acid instead of with sulphuric acid.
In relation to the galvanising capability of silicon-containing steels, it is stated inter alia in DE 196 10 675 C1 that steels containing up to 0.800 wt. % silicon or up to 2.000 wt. % silicon cannot be hot-galvanised by reason of the very poor wettability of the steel surface with the liquid zinc.
In addition to the recrystallisation of the roll-hard cold strip, 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. However, for 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.
It is known that 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.
Contrary to this common general knowledge in the art, it has surprisingly been found in tests that effective hot-dip finishing of the steel strip and effective adhesion of the coating can be achieved solely by means of suitable furnace operation during recrystallisation annealing and during passage through the hot-dip bath.
To this end, it is initially necessary to ensure that the strip surface is free of scale residues, acid-cleaning oil or rolling oil or other contaminant particles by performing a chemical-mechanical or thermal-hydromechanical pre-cleaning procedure. Furthermore, in order to prevent silicon oxides from getting onto the strip surface it is necessary to resort to methods which promote the internal oxidation of the alloy elements below the material surface. In this case, different measures are applied depending upon the configuration of the installation.
In the case of one configuration of the installation, in which the annealing process step is performed exclusively in a radiant tube furnace (RTF): (see method 3 in
−12>Log pO2≥−5*Si−0.25−3*Mn−0.5−0.1*Cr−0.5−7*(−ln B)0.5
In this case, Si, Mn, Cr, B define the corresponding alloy proportions in the steel in wt. % and pO2 defines the oxygen partial pressure in mbar.
In the case of a configuration of the installation, in which 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
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. %.
Under the N2—H2-protective gas atmosphere in the downstream radiant tube furnace, 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.
−18>Log pO2≥−5*Si−0.3−2.2*Mn−0.45−0.1*Cr−0.4−12.5*(−ln B)0.25
In this case, Si, Mn, Cr, B define the corresponding alloy proportions in the steel in wt. % and pO2 defines the oxygen partial pressure in mbar.
In the transition region between the furnace→zinc pot (nozzle), the dew point of the gas atmosphere (N2—H2-protective gas atmosphere) and thus 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 above-described measures in the furnace region of the continuous hot-dip coating installation prevent the formation of oxides on the surface and provide uniform, effective wettability of the strip surface with the liquid melt.
If, instead of hot-dip finishing procedure (in this case e.g. hot-dip galvanising), the method route selected is that involving continuous annealing with subsequent electrolytic galvanising (see method 1 in
In order to ensure the widest possible process window during annealing and sufficient galvanising capability, 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 α-/γ-conversion towards lower temperatures.
A main reason for alloying manganese in multi-phase steels, such as e.g. in the case of dual-phase steels, is the considerable improvement in the potential hardness increase. By reason of the inhibition of diffusion, the perlite and bainite conversion is shifted towards longer times and the martensite starting temperature is decreased.
However, at the same time the addition of manganese serves to increase the hardness ratio between martensite and ferrite. Moreover, the line formation of the microstructure is enhanced. A high hardness difference between the phases and the formation of martensite lines result in a lower hole expansion capacity, which is the equivalent of an increased edge crack sensitivity.
Manganese, like silicon, tends to form oxides on the steel surface during the annealing treatment. In dependence upon the annealing parameters and the contents of other alloy elements (in particular silicon and aluminium) manganese oxides (e.g. MnO) and/or Mn mixed oxides (e.g. Mn2SiO4) can occur. However, 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. Nevertheless, 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.
For the reasons stated, the manganese content is fixed to 1.900 to 2.350 wt. %.
In order to achieve the required minimum strengths, it is advantageous to maintain a strip thickness-dependent differentiation of the manganese content.
In the case of an end thickness of 0.50 mm to 1.00 mm inclusive, 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:
YS (MPa)=53.9+32.34[wt. % Mn]+83.16[wt. % Si]+354.2[wt. % N]+17.402d(−1/2)
However, this is based primarily upon the effect of mixed crystal hardening which, according to this equation, is weaker for manganese than for silicon. However, as mentioned above, manganese simultaneously increases hardenability considerably, as a result of which the proportion of a strength-increasing second phase increases significantly in the case of multi-phase steels. Therefore, the addition of 0.1% silicon is to be equated in a first approximation to the addition of 0.1% of manganese in terms of the increase in strength. For a steel of the composition according to the invention and an annealing procedure which includes the time-temperature parameters according to the invention, the following relationship has been produced on an empirical basis for the yield strength and the tensile strength (TS):
YS (MPa)=185+147.9 [wt. % Si]+161.1 [wt. % Mn]
TS (MPa)=574+189.4 [wt. % Si]+174.1 [wt. % Mn]
In comparison with 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.
On the one hand, chromium (Cr) in dissolved form and even in small quantities can considerably increase the hardenability of steel. On the other hand, with corresponding temperature control 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.
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.
Moreover, 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.
Therefore, 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.
Therefore, the content of molybdenum is set between more than 0.200 to 0.300 wt. %. For reasons relating to cost, the Mo content is advantageously set to a range between more than 0.200 and 0.250 wt. %.
As a compromise between the required mechanical properties and hot-dip capability it 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.
In conjunction with oxygen, 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.
Nickel (Ni): In conjunction with oxygen, nickel can form, at the grain boundaries, noxious oxides which can produce negative effects particularly for hot-deformation processes. Therefore, the content of nickel 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 (Al) 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.
Aluminium nitride is not precipitated if titanium is present in sufficient quantities. Titanium nitrides have a lower enthalpy of formation and are formed at higher temperatures.
In the dissolved state, 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. The reason for this approach resides in aluminium being slightly less critical for the galvanisation reaction than silicon.
Therefore, 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 (Nb): 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.
Titanium (Ti): By reason of its high affinity to nitrogen, titanium is primarily precipitated as TiN during solidification. Moreover, it appears together with niobium as a mixed carbide. TiN is highly significant for the grain size stability in the pusher-type furnace. The precipitations have a high level of temperature stability so that, in contrast to the mixed carbides, they are present at 1200° C. predominantly as particles which inhibit the grain growth. Titanium also has a delaying effect upon the recrystallisation during hot rolling but is less effective in this regard than niobium. Titanium functions by means of precipitation hardening. The larger TiN particles are less effective than the more finely distributed mixed carbides. The best efficacy is achieved in the range of 0.005 to 0.060 wt. % titanium, therefore this represents the alloy span according to the invention. For this purpose, contents of 0.025 to 0.045 wt. % have proven to be advantageous.
Boron (B): 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 (
Moreover, boron is an effective element for increasing hardenability which becomes effective even in very small quantities. The martensite starting temperature remains unaffected thereby. 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. However, excessively high contents of boron are hazardous because iron boride can form which has a negative effect upon hardenability, deformability and toughness of the material. Boron also tends to form oxides or mixed oxides when annealing is performed during the continuous hot-dip coating procedure, which impair the galvanisation quality. The above-stated measures for setting the furnace regions during continuous hot-dip coating serve to reduce the formation of oxides on the steel surface.
For the aforementioned reasons, 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. On the other hand, by means of targeted alloying of nitrogen in conjunction with the microalloy elements 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.
According to the invention, the N content is therefore fixed to values of ≥0.0020 to ≤0.0120 wt. %.
It has proven to be advantageous for maintaining the required properties of the steel if the nitrogen is added in dependence upon the sum of Ti+Nb+B.
In the case of a total content of Ti+Nb+B of ≥0.010 to ≤0.070 wt. %, the content of nitrogen must be kept to values of ≥0.0020 to ≤0.0090 wt. %. For a total content of Ti+Nb+B of >0.070 wt. %, nitrogen contents of ≥0.0040 to ≤0.0120 wt. % have proven to be advantageous.
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.
During the interaction of the microalloy elements niobium and titanium with boron, 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.
Furthermore, maximum contents of ≤0.365 wt. % have proven to be successful as total contents of Ti+Nb+V+Mo+B for the aforementioned reasons.
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.
For the stated reasons, a Ca content of ≥0.0010 to ≤0.0060 wt. % is set. Contents of ≤0.0030 wt. % have proven to be advantageous.
Tests conducted using the steel according to the invention have revealed that with inter-critical annealing between Ac1 and Ac3 or with austenitising annealing over Ac3 with concluding controlled cooling, a dual-phase steel having a minimum tensile strength of 980 MPa in a thickness of 0.50 to 3.00 mm can be produced which is characterised by an adequate tolerance to process fluctuations.
Therefore, a clearly widened process window is provided for the alloy composition according to the invention in comparison with known alloy concepts.
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 tests also revealed that the set microstructure proportions are retained after inter-critical annealing between Ac1 and Ac3 or austenitising annealing over Ac3 with subsequent controlled cooling even after a further process step of hot-dip finishing at temperatures between 400 to 470° C. e.g. with zinc or zinc-magnesium.
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).
Moreover, 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. For instance, 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.
In order to ensure the cold-rollability of a hot strip produced from the steel according to the invention, 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.
As part of the further processing of the steel strip according to the invention, it is thus possible to produce a hardened component e.g. for the automotive industry.
In this case, 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.
In an advantageous manner, 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.
During hardening, 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. At the same time, the hardness of the component is reduced such that the required toughness values are achieved.
Further features, advantages and details of the invention will be apparent from the following description of exemplified embodiments illustrated in a drawing.
It is shown in:
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.
By means of the different temperature controls according to the invention within said range, mutually different characteristic values and/or also different hole expansion results and bending angles are produced. Differences in principle are thus the temperature-time parameters during the heat treatment and the following cooling.
Method 1 (
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.
Method 2 (
Method 3 (
Method 3 corresponds e.g. to a process being carried out in a continuous annealing installation, as depicted in
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.
For industrial manufacture for the hot-dip galvanising according to method 3 and as shown in
A steel according to the invention with 0.104% C; 0.487% Si; 2.248% Mn; 0.016% P; 0.0009% S; 0.0052% N; 0.043 Al; 0.321% Cr; 0.214% Mo; 0.0278% Ti; 0.0374% Nb; 0.0007% B; 0.0016% Ca according to method 3 corresponding to
The yield strength ratio Re/Rm in the transverse direction was 65%.
The material characteristic values transversely to the rolling direction would correspond e.g. to a HC660XD.
The yield strength ratio Re/Rm in the longitudinal direction was 63%.
The material characteristic values longitudinally to the rolling direction would correspond e.g. to a CR590Y980T-DP.
A steel according to the invention with 0.104% C; 0.495% Si; 2.226% Mn; 0.011% P; 0.0008% S; 0.0048% N; 0.052 Al; 0.317% Cr; 0.212% Mo; 0.0381% Ti; 0.0361% Nb; 0.0008% B; 0.0019% Ca according to method 3 corresponding to
The yield strength ratio Re/Rm in the transverse direction was 66%.
The material characteristic values transversely to the rolling direction would correspond e.g. to a HC660XD.
The yield strength ratio Re/Rm in the longitudinal direction was 64%.
The material characteristic values longitudinally to the rolling direction would correspond e.g. to a CR590Y980T-DP.
A steel according to the invention with 0.102% C; 0.488% Si; 2.170% Mn; 0.012% P; 0.0008% S; 0.0048% N; 0.041 Al; 0.333% Cr; 0.218% Mo; 0.0351% Ti; 0.0345% Nb; 0.0009% B; 0.0018% Ca according to method 3 corresponding to
The yield strength ratio Re/Rm in the transverse direction was 68%.
The material characteristic values transversely to the rolling direction would correspond e.g. to a HC660XD.
The yield strength ratio Re/Rm in the longitudinal direction was 67%.
The material characteristic values longitudinally to the rolling direction would correspond e.g. to a CR590Y980T-DP.
A steel according to the invention with 0.103% C; 0.493% Si; 2.244% Mn; 0.013% P; 0.0009% S; 0.0049% N; 0.042 Al; 0.332% Cr; 0.226% Mo; 0.0315% Ti; 0.0346% Nb; 0.0007% B; 0.0013% Ca according to method 3 corresponding to
The yield strength ratio Re/Rm in the transverse direction was 60%.
The material characteristic values transversely to the rolling direction would correspond e.g. to a HC660XD.
The yield strength ratio Re/Rm in the longitudinal direction was 60%.
The material characteristic values longitudinally to the rolling direction would correspond e.g. to a CR590Y980T-DP.
A steel according to the invention with 0.100% C; 0.467% Si; 2.169% Mn; 0.011% P; 0.0008% S; 0.0050% N; 0.043 Al; 0.340% Cr; 0.223% Mo; 0.0293% Ti; 0.0387% Nb; 0.0009% B; 0.0019% Ca according to method 3 corresponding to
The yield strength ratio Re/Rm in the transverse direction was 67%.
The material characteristic values transversely to the rolling direction would correspond e.g. to a HC660XD.
The yield strength ratio Re/Rm in the longitudinal direction was 68%.
The material characteristic values longitudinally to the rolling direction would correspond e.g. to a CR590Y980T-DP.
The invention has been described above with the aid of steel sheets with an end thickness to be achieved of 0.50 mm to 3.00 mm. It is also possible, if required, to produce end thicknesses in the range of 0.10 mm to 4.00 mm.
Claims
1.-37. (canceled)
38. An ultra high strength multi-phase steel having a minimum tensile strength of 980 MPa, said multi-phase steel comprising in wt. %: C ≥0.075 to ≤0.115 Si ≥0.400 to ≤0.500 Mn ≥1.900 to ≤2.350 Cr ≥0.250 to ≤0.400 Al ≥0.005 to ≤0.060 N ≥0.0020 to ≤0.0120 S ≤0.0020 Nb ≥0.005 to ≤0.060 Ti ≥0.005 to ≤0.060 B >0.0005 to ≤0.0010 Mo >0.200 to ≤0.300 Ca >0.0010 to ≤0.0060 Cu ≤0.050 Ni ≤0.050 the remainder being iron, including typical steel-associated, smelting-related impurities, wherein a total content of Mn+Si+Cr is ≥2.500 to ≤3.250 wt. % with regard to a widest possible processing window during annealing, in particular continuous annealing, of cold strips of said steel.
39. The ultra high strength multi-phase steel of claim 38, wherein the total content of Mn+Si+Cr is ≥2.500 to ≤3.100 wt. % or ≥2.650 to ≤3.150 wt. % or ≥2.700 to ≤3.250 wt. %.
40. The ultra high strength multi-phase steel of claim 38, wherein a B content is ≤0.0009 wt. %, in particular the B content is ≥0.0006 to ≤0.0009 wt. %.
41. The ultra high strength multi-phase steel of claim 38, wherein a C content is ≤0.100 wt. % and a carbon equivalent CEV (IIW) is ≤0.62%, or a C content is ≤0.105 wt. % and a carbon equivalent CEV(IIW) is ≤0.64%, or a C content is ≤0.115 wt. % and a carbon equivalent CEV(IIW) is ≤0.66%.
42. The ultra high strength multi-phase steel of claim 38, wherein a Mn content is ≥1.900 to ≤2.350 wt. %.
43. The ultra high strength multi-phase steel of claim 42, wherein the Mn content is ≥1.900 to ≤2.200 wt. %, or the Mn content is ≥2.050 to ≤2.250 wt. %, or the Mn content is ≥2.100 to ≤2.350 wt. %.
44. The ultra high strength multi-phase steel of claim 38, wherein a N content is ≥0.0020 to ≤0.0090 wt. %, when a sum of Ti+Nb+B is ≥0.010 to ≤0.070 wt. %, or a N content is ≥0.0040 to ≤0.0120 wt. %, when a sum of Ti+Nb+B is >0.070 wt. %
45. The ultra high strength multi-phase steel of claim 38, wherein a S content is ≤0.0015 wt. %, in particular the S content is ≤0.0010 wt. %.
46. The ultra high strength multi-phase steel of claim 38, wherein a Mo content is ≤0.250 wt. %.
47. The ultra high strength multi-phase steel of claim 38, wherein a Ti content is ≥0.025 to ≤0.045 wt. %.
48. The ultra high strength multi-phase steel of claim 38, wherein a Nb content is ≥0.025 to ≤0.045 wt. %.
49. The ultra high strength multi-phase steel of claim 38, wherein a sum of Nb+Ti is ≤0.100 wt. %, in particular the sum of Nb+Ti is ≤0.090 wt. %.
50. The ultra high strength multi-phase steel of claim 38, wherein a sum of Cr+Mo is ≤0.650 wt. %.
51. The ultra high strength multi-phase steel of claim 38, wherein a sum Ti+Nb+B is ≤0.102 wt. %, in particular the sum of Ti+Nb+B is ≤0.092 wt. %.
52. The ultra high strength multi-phase steel of claim 38, wherein a sum of Ti+Nb+B+Mo+V is ≤0.365 wt. %.
53. The ultra high strength multi-phase steel of claim 38, wherein a Ca content is ≤0.0030 wt. %.
54. The ultra high strength multi-phase steel of claim 38, wherein additions of Si and Mn with respect to a strength property to be achieved are exchangeable according to the following relationships:
- YS (MPa)=185+147.9 [% Si]+161.1 [% Mn]
- TS (MPa)=574+189.4 [% Si]+174.1 [% Mn].
55. A method, comprising:
- producing a hot strip from a multi-phase steel as set forth in claim 38;
- cold-rolling the hot strip to produce a steel strip with a desired end thickness; and
- subsequently annealing, in particular continuously annealing, the steel strip.
56. The method of claim 55, wherein steel strips are cold-rolled in a wide range of degrees of thinning by rolling of 10% to 70% with the desired end thickness by proceeding from a selected master hot strip having a specific thickness or selected hot strips having different thicknesses.
57. The method of claim 55, wherein a chemical composition of the multi-phase steel is selected in dependence upon the desired end thickness of the steel strip.
58. The method of claim 57, wherein 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 in dependence upon the end thickness 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 over 1.00 mm to 2.00 mm inclusive: sum of Mn+Si+Cr≥2.650 to ≤3.150 wt. %,
- end thickness over 2.00 mm to 3.00 mm inclusive: sum of Mn+Si+Cr≥2.700 to ≤3.250 wt. %.
59. The method of claim 57, wherein the chemical composition of the multi-phase steel is selected in dependence upon the end thickness as follows:
- end thickness 0.50 mm to 1.00 mm inclusive: C content ≤0.100% and carbon equivalent CEV (IIW)≤0.62%,
- end thickness over 1.00 mm to 2.00 mm inclusive: C content ≤0.105% and the carbon equivalent CEV (IIW)≤0.64%,
- end thickness over 2.00 mm to 3.00 mm inclusive: C content ≤0.115% and the carbon equivalent CEV (IIW)≤0.66%.
60. The method of claim 57, wherein the chemical composition of the multi-phase steel is selected in dependence upon the end thickness as follows:
- end thickness 0.50 mm to 1.00 mm inclusive: Mn content ≥1.900 to ≤2.200%,
- end thickness over 1.00 mm to 2.00 mm inclusive: Mn content ≥2.050 to ≤2.250%,
- end thickness over 2.00 mm to 3.00 mm inclusive: Mn content ≥2.100 to ≤2.350%.
61. The method of claim 55, wherein in order to produce a required microstructure, the steel strip cold-rolled to the end thickness is heated during the continuous annealing to a temperature in a range of approximately 700 to 950° C., and the annealed steel strip is subsequently cooled from the annealing temperature at a cooling rate between approximately 15 and 100° C./s to a first intermediate temperature of approximately 300 to 500° C., then at a cooling rate between approximately 15 and 100° C./s to a second intermediate temperature of approximately 160 to 250° C., then the steel strip is cooled in air at a cooling rate of approximately 2 to 30° C./s until room temperature is reached or at a cooling rate between approximately 15 and 100° C./s from the first intermediate temperature to room temperature.
62. The method of claim 55, wherein in order to produce a required microstructure, the steel strip cold-rolled to the end thickness is heated during the continuous annealing to a temperature in a range of approximately 700 to 950° C. and, during a hot-dip finishing procedure after the heating and subsequent cooling to a temperature of approximately 400 and 470° C., cooling is stopped prior to entry into a melting bath and, after undergoing the hot-dip finishing procedure, cooling is continued at a cooling rate between approximately 15 and 100° C./s to an intermediate temperature of approximately 200 to 250° C., and then the steel strip is cooled in air at a cooling rate of approximately 2 to 30° C./s until room temperature is reached.
63. The method of claim 55, wherein in order to produce a required microstructure, the steel strip cold-rolled to the end thickness is heated during the continuous annealing to a temperature in a range of approximately 700 to 950° C. and, during a hot-dip finishing procedure after the heating and subsequent cooling to an intermediate temperature of approximately 200 to 250° C. prior to entry into a melting bath, the temperature is maintained for approximately 1 to 20 s and then the steel strip is heated to a temperature of approximately 400 to 470° C. and, after undergoing the hot-dip finishing procedure, cooling at a cooling rate between approximately 15 and 100° C./s to an intermediate temperature of approximately 200 to 250° C. and subsequently cooling in air at a cooling rate of approximately 2 to 30° C./s to room temperature.
64. The method of claim 55, wherein, during the continuous annealing, an oxidation potential during annealing with an installation configuration comprised of a direct fired furnace region (NOF) and a radiant tube furnace (RTF) is increased by a CO content in the NOF of less than 4 vol. %, wherein in the RTF an oxygen partial pressure of a furnace atmosphere which is reducing for iron is set according to the following equation, wherein Si, Mn, Cr and B designate corresponding alloy proportions in the steel in wt. % and pO2 designates the oxygen partial pressure in mbar, and in order to avoid oxidation of the strip directly prior to dipping in a melting bath a dew point of the gas atmosphere is set at −30° C. or below.
- −18>Log pO2≥−5*Si−0.3−2.2*Mn−0.45−0.1*Cr−0.4−12.5*(−ln B)0.25
65. The method of claim 55, wherein during annealing only with a radiant tube furnace an oxygen partial pressure of a furnace atmosphere satisfies the following equation, wherein Si, Mn, Cr and B designate corresponding alloy proportions in the steel in wt. % and pO2 designates the oxygen partial pressure in mbar, and in order to avoid oxidation of the strip directly prior to dipping in a melting bath a dew point of the gas atmosphere is set at −30° C. or below.
- −12>Log pO2≥−5*Si−0.25−3*Mn−0.5−0.1*Cr−0.5−7*(−ln B)0.5
66. The method of claim 55, wherein when strips of different thickness are involved during continuous annealing comparable microstructure states and mechanical characteristic values of the strips are set by adapting a rate of installation throughput rate during heat treatment.
67. The method of claim 55, wherein the steel strip is temper-rolled subsequent to the annealing or hot-dip finishing procedure.
68. The method of claim 55, wherein the steel strip is stretch-bend-straightened subsequent to the annealing or hot-dip finishing procedure.
69. The method of claim 55, further comprising:
- cutting a blank from the steel strip;
- heating the blank to a temperature above Ac3;
- deforming the heated blank to produce a component; and
- hardening the component in a tool or in air.
70. The method of claim 55, wherein the steel strip is part of a sub-peritectic grade chain in the absence of vanadium as a microalloy element.
71. A steel strip produced by a method as set forth in claim 55, said steel strip comprising a minimum hole expansion according to ISO 16630 of 20%, in particular of 25%.
72. The steel strip of claim 71, further comprising a minimum bending angle according to VDA 238-100 of 70° in a longitudinal direction or transverse direction, in particular of 85°.
73. The steel strip of claim 71, further comprising a minimum product value Rm×α (tensile strength×bending angle according to VDA 238-100) of 100000 MPa°, in particular of 120000 MPa°.
74. The steel strip of claim 71, further comprising a delayed fracture free state for at least 6 months thus meeting a requirement of SEP 1970 for hole pull and hoop test pieces.
Type: Application
Filed: Jul 7, 2016
Publication Date: May 2, 2019
Applicant: Salzgitter Fachstahl GmbH (Salzgitter)
Inventors: THOMAS SCHULZ (Salzgitter), SASCHA KLUGE (Salzgitter), CHRISTIAN MEYER (Ribbesbüttel), PETER MATTHIES (Clausthal-Zellerfeld)
Application Number: 15/743,229
