US20200263283A1

US20200263283A1 - Ultrahigh strength multiphase steel and method for producing a steel strip from said multiphase steel

Info

Publication number: US20200263283A1
Application number: US16/652,933
Authority: US
Inventors: Thomas Schulz
Original assignee: Salzgitter Flachstahl GmbH
Current assignee: Salzgitter Flachstahl GmbH
Priority date: 2017-10-06
Filing date: 2018-09-27
Publication date: 2020-08-20
Also published as: CN111247258B; EP3692178A1; DE102017123236A1; WO2019068560A1; KR20200063167A; RU2742998C1; CN111247258A; EP3692178B1

Abstract

The invention relates to an ultrahigh strength multiphase steel having a minimum tensile strength of 980 MPa containing (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≥0.010 to ≤0.060; N≥0.0020 to ≤0.0120; P≤0.020; S≤0.0020; Ti≥0.005 to ≤0.060; Nb≥0.005 to ≤0.060; V≥0.005 to ≤0.020; B≥0.0005 to ≤0.0010; Mo≥0.200 to ≤0.300; Ca≥0.0010 to ≤0.0060; Cu≤0.050; Ni≤0.050; Sn≤0.040; H≤0.0010; and residual iron, including customary steel-accompanying smelting-related impurities, wherein the total content of Mn+Si+Cr is ≥1.750 to ≤2.250 wt. % 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.

Description

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 production and excellent material properties for subsequent processing, in particular for lightweight vehicle construction according to the preamble of claim 1. Advantageous developments are described in dependent claims 2 to 24.
The invention also relates to a method for producing steel strips from such a steel according to claim 25 and steel strips produced therewith according to claim 38.
In particular, the invention relates to steels having a tensile strength in the region of at least 980 MPa, in the non-quenched state, for producing components which have improved deformability, such as e.g. with regard to hole expansion and improved joining suitability, such as e.g. welding properties.
The fiercely competitive car market means that producers are constantly forced to find solutions for reducing fleet fuel consumption and CO2 exhaust emissions 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 favorable possible behavior of the individual components in the event of high static and dynamic loading during operation and also in the event of a crash.
The steel suppliers take the aforementioned problem into account by providing ultra high strength steels. Furthermore, by providing ultra high strength steels having a smaller sheet thickness, the weight of the vehicle components can be reduced whilst the behavior 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 high component requirements according to toughness, edge crack insensitivity, improved bending angle and bending radius, 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. stamping and deforming, optional thermal quenching with subsequent optional tempering, welding and/or surface post-treatment, such as phosphatizing and cathodic dip coating, 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.
Improved joining suitability, e.g. in the form of better general welding capability, as well as a larger usable welding area for resistance spot welding and improved failure behavior of the weld seam (fracture pattern) under mechanical stress, and sufficient resistant to delayed hydrogen embrittlement (i.e. delayed fracture free) are also required to an increasing extent. The same applies to the welding suitability of ultra high strength steels in the production of pipes which are produced e.g. by means of the high-frequency induction welding method (HFI).
The hole expansion capability is a material property which describes the resistance of the material to crack initiation and crack propagation in deformation operations in regions close to the edge, such as e.g. during plunging.
The hole expansion test is regulated e.g. in the ISO 16630 standard. According to this, prefabricated holes which are stamped e.g. into a metal sheet are expanded by means of a mandrel. The measurement variable is the change in hole diameter, related to the initial diameter, at which the first crack through the metal sheet occurs at the edge of the hole.
Improved edge crack insensitivity signifies an increased deformation capability of the sheet edges and can be described by an increased hole expansion capability. These circumstances are also known as the synonyms “Low Edge Crack” (LEC) or by “High Hole Expansion” (HHE) and Xpand®.
The bending angle describes a material property which allows conclusions to be drawn in respect of the material behavior in deformation operations with dominant bending proportions (e.g. during folding) or even in the event of crash loadings. Therefore, increased bending angles increase passenger compartment safety.
The determination of the bending angle (a) is regulated e.g. by means of the plate bending test in the standard VDA 238-100.
The aforementioned properties are important for components which have a very complex formation.
Improved welding capability is known to be achieved inter alia by means of a reduced carbon equivalent.
Terms such as “underperitectic” (UP) or the already known “Low Carbon Equivalent” (LCE) are used as synonyms for this. The carbon content is typically less than 0.120 wt. %.
Furthermore, the failure behavior or the fracture pattern of the weld seam can be improved by significant addition by alloying with microalloy elements, in the case of low-carbon steels with a reduced carbon equivalent.
High-strength components must have sufficient resistance to material embrittlement with respect to hydrogen.
The resistance test of Advanced High Strength Steels (AHSS) for automotive applications against production-related, hydrogen-induced brittle fracture is regulated in SEP1970 and is tested using the yoke test piece and hole pull test piece.
In vehicle construction, dual-phase steels are being increasingly used, said steels consisting of a ferritic basic microstructure, into which a martensitic second phase is incorporated. It has been found that in the case of low-carbon, micro-alloyed steels, proportions of further phases, such as bainite and residual austenite have an advantageous effect e.g. upon the hole expansion behavior, the bending behavior and the hydrogen-inducted brittle fracture behavior. In this case, bainite can be present in different manifestations, such as e.g. upper and lower bainite.
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.
The combination of properties required of the steel material ultimately represents a component-specific compromise of individual properties. However, these properties are often no longer adequate in the case of ever more complex component geometries.
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 structural compositions. Complex-phase steels are, according to EN 10346, steels which contain small proportions of martensite, residual austenite and/or perlite in a ferritic/bainitic basic microstructure, wherein strong grain refinement is produced by means of delayed recrystallization or by precipitation of microalloy elements.
These complex-phase steels have higher yield strengths, a greater yield strength ratio, lower cold solidification and a higher hole expansion capability.
Ferritic-bainitic steels are, according to EN 10346, steels which contain bainite or solidified bainite in a matrix of ferrite and/or solidified ferrite. The strength of the matrix is produced by means of a high dislocation density, by means of grain refinement and the precipitation of microalloy elements.
Dual-phase steels are, according to EN 10346, steels which have a ferritic basic microstructure, in which a martensitic second phase is incorporated in the shape of an island, occasionally with proportions of bainite as the second phase. In the case of high tensile strength, dual-phase steels demonstrate a low yield strength ratio and strong cold solidification.
TRIP steels are, according to EN 10346, steels which have a predominantly ferritic basic microstructure, in which bainite and residual austenite is incorporated, which can convert into martensite during deformation (TRIP effect). Owing to its intense cold solidification, the steel achieves high values for uniform elongation and tensile strength. High component strengths can be achieved in conjunction with the bake-hardening effect. These steels are suitable both for stretch-drawing and deep-drawing. However, during material deformation, higher sheet holder forces and press forces are required. A comparatively strong resilience is to be taken into account.
The high-strength steels comprising a single-phase microstructure include e.g. bainitic and martensitic steels.
Bainitic steels are, according to EN 10346, steels which are characterized by a very high yield strength and tensile strength with sufficiently high elongation for cold-forming processes. An effective welding capability is provided by reason of the chemical composition. The microstructure consists typically of bainite. The microstructure can contain in isolation small proportions of other phases, such as e.g. martensite and ferrite.
Martensitic steels are, according to EN 10346, steels which, by reason of thermomechanical rolling, contain small proportions of ferrite and/or bainite in a basic microstructure of martensite. This steel grade is characterized by a very high yield strength and tensile strength with sufficiently high elongation for cold-forming processes. Within the group of multi-phase steels, the martensitic steels have the highest tensile strength values. The suitability for deep-drawing is limited. The martensitic steels are suitable predominantly for bending deformation methods, such as roll-forming.
Heat-treatment steels are, according to EN 10083, steels which acquire a high tensile strength and fatigue strength by means of quenching (=hardening and tempering). If the cooling during hardening in air produces bainite or martensite, the method is called “air-hardening”. The strength/tensile strength ratio can be influenced in a targeted manner by tempering which is effected after hardening.
These steels are currently being used in structural components, chassis components and crash-relevant components, and as flexibly cold-rolled strips.
A significant weight reduction through the loading-adapted selection of the sheet thickness over the component length is made possible by this Tailor Rolled Blank lightweight construction technology (TRB®).
However, with currently known alloys and available continuous annealing installations for greatly varying sheet thicknesses, the production of TRB® s with multi-phase microstructures is not possible without limitations, such as e.g. for the 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 recrystallizing manner to produce a 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 aforementioned properties are significantly influenced e.g. by the steel compositions, the process parameters during hot rolling, the process parameters during acid-cleaning (e.g. stretch-bend-straightening) and the process parameters during cold-rolling even prior to continuous annealing.
The steel composition is fixed by analysis regulations which define MIN and MAX ranges.
The process parameters during hot-rolling, such as e.g. the standard slab thickness, slab lying time, slab output temperature, pass plan during pre-strip rolling, standard pre-strip thickness, entry temperature into the hot strip line, pass plan during hot-rolling, end rolling temperature, hot strip cooling pattern, reeling temperature, are fixed depending upon the multi-phase steel to be produced.
During acid-cleaning, optional stretch-bend-straightening (stretch-forming) influences the subsequent process step.
During cold-rolling, the hot strip thickness for producing a cold-rolling thickness is already fixed by a standard degree of thinning by rolling at the time of order conversion into the technical specifications (process parameters).
The thickness of the pre-strip during the hot-rolling process describes the starting thickness prior to entry into the multi-frame hot strip line, wherein the pre-strip has been manufactured in a reversing manner, in a plurality of passes (runs) from one slab having a defined standard thickness.
Typical slab thicknesses are between 250 mm and 300 mm (standard 250 mm, considered further here), the pre-strip thicknesses in the case of multi-phase steels typically range between 40 mm to 60 mm.
Typically, the pre-strip thicknesses for the subsequent hot-rolling are relatively constant, depending upon the material composition, e.g. at 45 mm (referred to here as standard).
Values below or above this produce changed technological hot strip characteristic values, such as tensile strength and yield strength which, in turn, thereby influence the subsequent deformation during cold-rolling, like cold solidification behavior.
In order to achieve the final technological fine sheet characteristic values required by the standards, according to the prior art a material-dependent pre-strip thickness is fixed in the case of continuous annealing treatment in order to ensure normal recrystallization. In the case of classic steels, values exceeding or falling short of this influence the final technological characteristic values to such an extent that considerable batch fluctuations (scatter range) can occur.
The degree of thinning by 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.
Typically, the degrees of thinning by cold-rolling are relatively constant, they are up to approximately 40% in the case of thicker cold strips of more than 2 mm and are up to approximately 60% in the case of cold strips of up to 1 mm in thickness.
In order to achieve the technological characteristic values required by the standards, according to the prior art, on average a degree of thinning by cold-rolling of approximately 50% is required in the case of continuous annealing treatment in order to ensure normal recrystallization. In the case of classic steels, values exceeding or falling short of this 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 recrystallization temperature, in order to set a corresponding dislocation density for the recrystallization annealing.
If the degree of thinning by cold-rolling is too low (even in local regions), the critical threshold for recrystallization cannot be overcome and so a fine-grain and relatively uniform microstructure cannot be achieved. After recrystallization, 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 galvanized, the annealing treatment is typically performed in a continuous hot-galvanizing installation, in which the heat treatment or annealing and the downstream galvanizing 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 decisive process parameter for material having a relatively constant degree of thinning by 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 is 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 pre-strip thicknesses or excessively low or excessively high degrees of thinning by 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-optimized components consisting of a cold strip but also consisting of a hot 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 pre-strip thicknesses or variable degrees of thinning by cold-rolling cannot be achieved in practical terms with the known alloy concepts during continuous annealing. The degree of thinning by cold-rolling which is required for recrystallization 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 characterized by an excessively narrow process window and therefore are unsuitable in particular for cold strip manufacture with variable pre-strip thicknesses and variable degrees of thinning by 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 either variable pre-strip thicknesses or variable degrees of thinning by cold-rolling with a single hot strip thickness (master hot strip thickness) whilst producing uniform material properties.
Laid-open document DE 10 2015 111 177 A1 discloses an ultra high strength multi-phase steel having minimum tensile strengths of 980 MPa which already has a very wide process window for the continuous annealing of hot or cold strips, and also e.g. with a single hot strip thickness (master hot strip thickness) thus achieving variable degrees of thinning by cold-rolling, continuously annealed cold strips having different thicknesses and with uniform material properties can be achieved.
Laid-open document DE 10 2014 017 274 A1 discloses an ultra high strength, air-hardenable multi-phase steel having minimum tensile strengths, in the non-air-hardened state, of 950 MPa which already has a very wide process window for the continuous annealing of hot or cold strips, and also e.g. with a single hot strip thickness (master hot strip thickness) thus achieving variable degrees of thinning by cold-rolling, continuously annealed cold strips having different thicknesses and with uniform material properties can be achieved and are suitable for the subsequent air-hardening process.
The aim of achieving the resulting mechanical-technological properties in a narrow region over the length and width of the strip by the controlled setting of the volume proportions of the microstructure components has the greatest priority and is only possible by means of an increased process window. The known alloy concepts are characterized by a process window which is too narrow, and therefore is unsuitable for solving the present problem, in particular in the case of flexibly rolled strips. With the known alloy concepts it is currently possible to produce only steels of one strength class with defined cross-sectional regions (strip thickness and strip width) and so changed alloy concepts are required for different strength classes and/or cross-sectional regions.
Steel production is witnessing a trend for reducing the carbon equivalent in order to achieve improved cold processing (cold-rolling, cold-forming) and better usage properties.
However, welding suitability—characterized inter alia by the carbon equivalent—is an important evaluation variable.
For example, in the following carbon equivalents
CEV(IIW)=C+Mn/6+(Cu+Ni)/15+(Cr+Mo+V)/5
CET=C+(Mn+Mo)/10+(Cr+Cu)/20+Ni/40
PCM=C+(Mn+Cu+Cr)/20+Ni/60+Mo/15+V/10+5 B
the characteristic standard elements, such as carbon and manganese, as well as chromium or molybdenum and vanadium are taken into consideration (contents in wt. %).
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 as well as the 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 behavior. With correspondingly optimized steel grade adaptations, the steel user expects higher values than with the standard material. However, welding suitability, characterized by the carbon equivalent, also comes further into focus.
The automotive industry is increasingly demanding grades of steel which have requirements in terms of the ratio of yield strength (Re) or elasticity limit (Rp0.2) to tensile strength which differ considerably depending upon the application. This leads to steel developments with a comparatively large yield strength interval with a standard 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 characterized 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 favorable 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 micro-alloys individually or in combinations, and in the material-characterizing properties, such as hole expansion and reduced carbon equivalent, etc.
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.
Thickness ranges below 0.50 and above 3.00 mm are feasible.
Overall, in the case of the known steel grades there is the problem that with regard to the required minimum degree of thinning by rolling during cold-rolling for a complete recrystallization after continuous annealing, at a given pre-strip thickness for producing a master hot strip thickness after hot-rolling there is no longer any manufacturing flexibility (see FIG. 1, process steps 6, 8 and 9 are then necessary) with regard to different cold strip thicknesses to be achieved. In particular, it is not possible to produce different cold strip thicknesses in the case of a constant master hot strip thickness with comparable material properties on the produced cold strip by reason of a process window which is too small. Moreover, the specification of a constant pre-strip thickness for producing a specified constant master hot strip thickness restricts the manufacturing flexibility.
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 different pre-strip thicknesses, 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). Moreover, instead of constant pre-strip thicknesses variable pre-strip thicknesses can be used prior to the hot-rolling.
In this case, the most uniform possible cold strip material properties are to be achieved independently of the set pre-strip thickness and the set degree of cold-rolling.
Moreover, the process window for the annealing, in particular continuous annealing, of steel strips which are cold-rolled to an end thickness is to be extended such that, in addition to strips having different cross-sections (cross-sectional jump), it is also possible to produce steel strips having a thickness (TRB®) which vanes 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. %:

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≥0.010 to ≤0.060

N≥0.0020 to ≤0.0120

P≤0.020

S≤0.0020

Ti≥0.005 to ≤0.060

Nb≥0.005 to ≤0.060

V≥0.005 to ≤0.020

B≥0.0005 to ≤0.0010

Mo≥0.200 to ≤0.300

Ca≥0.0010 to ≤0.0060

Cu≤0.050

Ni≤0.050

Sn≤0.040

H≤0.0010

with the remainder being iron, including typical steel-associated, smelting-related impurities, wherein the total content of Mn-Si+Cr is ≥1.750 wt. % to ≤2.250 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.
In mathematical terms, this is to mean that the content indications of Mn and Cr are added and that of Si is subtracted and the total (result) thus obtained is to be greater than or equal to 1.750 and less than or equal to 2.250 wt. %. The same applies to the further corresponding total contents.
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 pre-strip thickness prior to hot-rolling, as well as variable degrees of thinning by cold-rolling during cold-rolling. By means of variable pre-strip thicknesses, the cold-rolling process can be positively influenced by virtue of the fact that the steps of soft annealing of hot strip are performed prior to cold-rolling, double cold-rolling, soft annealing of the cold-rolled strip prior to the next cold-rolling step, without negative consequences on the production of the above-described master hot strip thickness or master cold strip thickness.
What is significant for this is a selected, closely maintained alloy composition with the emphasis on a restricted and cold strip thickness-dependent chromium content which has proven to be very positive for achieving uniform material properties with different pre-strip thicknesses, as well as different degrees of thinning by 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 and a standard pre-strip thickness is necessary can be disregarded to the extent that a selected pre-strip thickness and 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.
Therefore, a pre-strip can be produced from the multi-phase steel in the state of a slab, said pre-strip subsequently being hot-rolled with the hot strip thickness to be achieved.
It is also possible, proceeding from a previously fixed slab thickness of e.g. 250 mm and a previously selected pre-strip having a defined but variable thickness, to hot-roll hot strips with the same thickness with degrees of thinning by rolling of 72% to 87% with the end thickness to be achieved.
In an advantageous manner, in the case of steel strips of different thickness during continuous annealing comparable microstructure states and mechanical characteristic values of the strips can be set by adapting the installation throughput rate during the heat treatment.
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 a 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 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 multi-phase 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 varying pre-strip thickness, 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 cold-rolling of 10% to 70%, steel 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 cold 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 to 3.00 mm and the chemical composition of the multi-phase steel is selected as follows in dependence upon the end thickness to be achieved, even if variable pre-strip thicknesses are used.
In relation to the possible use of variable pre-strip thicknesses, it proved to be particularly advantageous that the total content of Mn-Si+Cr is selected as follows in dependence upon the end thickness of the cold strip to be achieved:
end thickness 0.50 to 1.00 mm inclusive:

sum of Mn-Si+Cr≥1.750 wt. % to ≤2.030 wt. %,

end thickness over 1.00 to 2.00 mm inclusive:

sum of Mn-Si+Cr≥1.940 wt. % to ≤2.110 wt. %

end thickness over 2.00 to 3.00 mm inclusive:

sum of Mn-Si+Cr≥2.020 wt. % to ≤2.220 wt. %.

Furthermore, it proved to be advantageous that the total content of Mn-Si+Cr+Mo is selected as follows in dependence upon the end thickness of the cold strip to be achieved:
end thickness 0.50 to 1.00 mm inclusive:

sum of Mn-Si+Cr+Mo≥1.950 wt. % to ≤2.280 wt. %,

end thickness over 1.00 to 2.00 mm inclusive:

sum of Mn-Si+Cr+Mo≥2.140 wt. % to ≤2.360 wt. %,

end thickness over 2.00 mm to 3.00 mm inclusive:

sum of Mn-Si+Cr+Mo≥2.220 wt. % to ≤2.470 wt. %.

Therefore, the end thickness of the steel strip to be achieved is associated with the alloy composition of the pre-strip or hot strip produced from the multi-phase steel.
It has also proven to be advantageous that the carbon equivalent CEV (IIW) is selected as follows in dependence upon the end thickness of the cold strip to be achieved:
end thickness 0.50 mm to 1.00 mm inclusive:
C content ≤0.100 wt. % and carbon equivalent CEV (IIW) ≤0.62%,
end thickness over 1.00 mm to 2.00 mm inclusive:
C content ≤0.105 wt. % and the carbon equivalent CEV (IIW) ≤0.64%,
end thickness over 2.00 mm to 3.00 mm inclusive:
C content ≤0.115 wt. % and the carbon equivalent CEV (IIW) ≤0.66%.
It has also proven to be advantageous that the Mn content is selected as follows in dependence upon the end thickness of the cold strip to be achieved:
end thickness 0.50 to 1.00 mm inclusive:
Mn content ≥1.900 wt. % to ≤2.200 wt. %,
end thickness over 1.00 to 2.00 mm inclusive:
Mn content ≥2.050 wt. % to ≤2.250 wt. %,
end thickness over 2.00 mm to 3.00 mm inclusive:
Mn content ≥2.100 wt. % to ≤2.350 wt. %.
In relation to the use of variable pre-strip thicknesses, it proved to be particularly advantageous that the Cr content and the carbon equivalent CEV (1W) are selected as follows in dependence upon the end thickness of the cold strip to be achieved:
end thickness 0.50 to 1.00 mm inclusive:
Cr content ≥0.260 wt. % to ≤0.330 wt. %
and carbon equivalent CEV (IIW) ≤0.62%,
end thickness over 1.00 to 2.00 mm inclusive:
Cr content ≥0.290 wt. % to ≤0.360 wt. %
and the carbon equivalent CEV (IIW) ≤0.64%,
end thickness over 2.00 to 3.00 mm inclusive:
Cr content ≥0.320 wt. % to ≤0.370 wt. %
and the carbon equivalent CEV (IIW)≤0.66%.
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 cold-rolling.
If, according to the invention, 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-optimized components from this material by deformation technology.
The material produced can be produced as a cold strip via a hot-dip galvanizing 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, FIG. 8a ). Alternatively, 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.
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 then performed at a cooling rate of approximately 2 to 30° C./s (see method 2, FIG. 8b ).
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 re-heating 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. 8c ).
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 boron in narrow regions 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, and the elements carbon, chromium, molybdenum and boron also act in a similar manner. 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 content 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 favorable 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 behavior occurs in the event of hydrogen embrittlement or in the event of hydrogen-induced tension crack corrosion. Even in the case of the delayed crack, the 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.
For the reasons stated above, the hydrogen content in the steel according to the invention is limited to ≤0.0010 wt. % (10 ppm) or advantageously to ≤0.0008 wt. %, optimally to ≤0.0005 wt. %.
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. Phosphorus increases the hardness by means of mixed crystal hardening and improves the hardenability. 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 behavior 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 behavior, 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 α-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 to 2.00 mm inclusive (C≤0.105 wt. %)
end thickness over 2.00 to 3.00 mm inclusive (C≤0.115 wt. %)
Furthermore, it is advantageous to adhere to a band thickness-dependent differentiation of the carbon content in combination with the carbon equivalent CEV(IIW).
End thickness 0.50 to 1.00 mm inclusive (C≤0,100 wt %) with a carbon equivalent CEV(IIW) of ≤0.62,
End thickness over 1.00 to 2.00 mm inclusive (C≤0.105 wt. %) with a carbon equivalent CEV(IIW) of ≤0.64%,
End thickness over 2.00 to 3.00 mm inclusive (C≤0.115 wt. %) with a carbon equivalent CEV(IIW) of ≤0.66%.
During casting, silicon (Si) binds oxygen and is therefore used for killing 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 e.g. hole expansion and bending capability.
In a manner characteristic of the material, the addition of silicon produces a strong mixed crystal hardening. 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 stabilized. Since silicon hinders the formation of carbide, the austenite is additionally stabilized. Therefore, during the accelerated cooling the formation of bainite can be suppressed in favor 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 and so 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 of 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 galvanizing 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-galvanized by reason of the very poor wettability of the steel surface with the liquid zinc.
In addition to the recrystallization 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 oxidizing, 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-galvanized 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 recrystallization 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 FIG. 8c ), the internal oxidation of the alloy elements can be influenced in a targeted manner by setting the oxygen partial pressure of the furnace atmosphere (N₂—H₂-protective gas atmosphere). The set oxygen partial pressure must satisfy the following equation, wherein the furnace temperature is between 700 and 950° C.
−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-oxidizing furnace (NOF) and a downstream radiant tube furnace (see method 2 in FIG. 8b ), 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. %.
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 oxidized 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 pO₂defines the oxygen partial pressure in mbar.
In the transition region between the furnace→zinc pot (muzzle), 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 the hot-dip finishing procedure (in this case e.g. hot-dip galvanizing), the method route selected is that involving continuous annealing with subsequent electrolytic galvanizing (see method 1 in FIG. 8a ), it is not necessary to implement any particular measures in order to ensure the galvanizing capability. It is known that the galvanization 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 galvanizing, 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 galvanization, 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.
In order to ensure the widest possible process window during annealing and sufficient galvanizing 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 the noxious sulphur into manganese sulphides. Moreover, by means of mixed crystal hardening, manganese increases the strength of the ferrite and shifts the α-/γ-conversion towards lower temperatures.
A main reason for adding manganese by alloying 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 wt. % 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 wt. % 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 wt. % to ≤2.250 wt. % and in the case of end thicknesses of 2.00 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 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]+17.402 d^(−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% 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)=185.7+147.9 [wt. % Si]+161.1 [wt. % Mn]
TS (MPa)=574.8+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 and so 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 in order 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 wt. % to 0.400 wt. %.
In order to achieve the required minimum strengths, it is advantageous to adhere to a band thickness-dependent differentiation of the chromium content, in particular for processing with variable pre-strip thicknesses.
In the case of an end thickness of 0.50 to 1.00 mm inclusive, the chromium content is preferably in a range between ≥0.260 wt. % to ≤0.330 wt. %, in the case of end thicknesses of 1.00 to 2.00 mm inclusive, said chromium content is in a range between a 0.290 wt. % to ≤0.360 wt. % and in the case of end thicknesses of 2.00 to 3.00 mm inclusive, it is in a range between a 0.320 wt. % to ≤0.370 wt. %.
In order to achieve the required minimum strengths, it is advantageous to adhere to a band thickness-dependent differentiation of the chromium content in combination with the carbon equivalent CEV(IIW), in this case in particular also for processing with variable pre-strip thicknesses.
In the case of an end thickness of 0.50 to 1.00 mm inclusive, the chromium content is preferably in a range between ≥0.260 wt. % to ≤0.330 wt. %, in the case of a carbon equivalent CEV(IIW) of ≤0.62%, in the case of end thicknesses of 1.00 to 2.00 mm inclusive, said chromium content is in a range between ≥0.290 wt. % to ≤0.360 wt. % in the case of a carbon equivalent CEV(IIW) of ≤0.66% and in the case of end thicknesses of 2.00 to 3.00 mm inclusive, it is in a range between ≥0.320 wt. % to ≤0.370 wt. % in the case of a carbon equivalent CEV(IIW) of ≤0.66%.
Chromium contents between ≥0.250 wt. % to <0.370 wt. % can be used in the case of end thicknesses of less than 0.50 mm and chromium contents between >0.370 wt. % to ≤0.400 wt. % can be used in the case of end thicknesses of greater than 3.00 mm.
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 and so 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 wt. % to 0.300 wt. %. For reasons relating to cost, the Mo content is advantageously set to a range between more than 0.200 wt. % to 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. %.
In order to achieve the required mechanical characteristic values, primarily the minimum tensile strength, it is advantageous to adhere to the total content of manganese, silicon and chromium, via a total formula Mn-Si+Cr, wherein this is to be limited between ≥1.750 wt. % to ≤2.250 wt. %, in particular for processing with variable pre-strip thicknesses.
In order to achieve the required mechanical characteristic values, primarily the minimum tensile strength, it has proven to be advantageous to fix the total content of manganese, silicon, chromium and molybdenum via a total formula Mn-Si+Cr+Mo, wherein this is to be limited between ≥1.950 wt. % to ≤2.500 wt. %, in particular for processing strips with variable pre-strip thicknesses.
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): in the case of the present alloy concept, the content of vanadium of ≥0.005 wt. % to ≤0.020 wt. % is fixed, optimally limited to ≥0.005 wt. % to ≤0.015 wt. %.
Tin (Sn): since, in the case of the present alloy concept, an addition of tin is not necessary, the content of tin is fixed to ≤0.040 wt. % and therefore limited to unavoidable quantities associated with steel.
Aluminium (Al) is generally added to the steel by alloying in order to bind the oxygen and nitrogen dissolved in the iron. Oxygen and nitrogen are thus converted 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 galvanization reaction than silicon.
Therefore, the aluminium content is limited to 0.010 wt. % to a maximum of 0.060 wt. % and is added for the purpose of killing the steel.
Niobium (Nb): niobium acts differently in steel. During hot-rolling on the production line, it delays the recrystallization 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 recrystallization. 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 galvanizing. Above 0.060 wt. % niobium, no additional effect is to be expected. Contents of 0.025 wt. % 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 and so, 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 recrystallization 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 wt. % to 0.060 wt. % titanium, therefore this represents the alloy span according to the invention. For this purpose, contents of 0.025 wt. % 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 cold-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 cold-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 cold-rolling, defined characteristic value ranges after the process steps (FIG. 8a, 8b or 8 c) also with the material with variable degrees of thinning by cold-rolling on the basis of a master hot strip thickness or on the basis of a master cold strip thickness.
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 galvanization 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 wt. % to 0.0010 wt. %, advantageously to values ≤0.0009 wt. % 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 addition by 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 a 0.0020 wt. % to ≤0.0120 wt. %.
It has been demonstrated to be advantageous to fix a total in the case of the contents of hydrogen and nitrogen, wherein an optimum for H+N is between ≥0.0025 wt. % to ≤0.0130 wt. %.
It has proven to be advantageous for maintaining the required properties of the steel if the nitrogen is added in dependence upon the total of Ti+Nb+B.
In the case of a total content of T+Nb+B of a 0.010 wt. % to ≤0.080 wt. %, the content of nitrogen must be kept to values of ≥0.0020 wt. % to ≤0.0090 wt. %. For a total content of Ti+Nb+B of ≥0.050 wt. %, nitrogen contents of a 0.0040 wt. % 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 0.005 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 wt. % to ≤0.0060 wt. % is set. Contents of ≤0.0030 wt. % have proven to be advantageous.
Tests with variable pre-strip thicknesses conducted using the steel according to the invention have revealed that with inter-critical annealing between Ac1 and Ac3 or with austenising annealing over Ac3 with concluding controlled cooling, a dual-phase steel having a minimum tensile strength of 980 MPa in a thickness of 1.50 mm, starting from a master hot strip of 2.30 mm was produced, but can also be produced in the thickness range of 0.50 to 3.00 mm which is characterized 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 achieved 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 characterized during further processing by high edge crack insensitivity and a high bending angle.
In an advantageous manner, it is thus possible to produce steel strips which have a minimum product value Rm×α (tensile strength×[bending angle according to VDA 238-100]) of 100000 MPa×°, in particular of 120000 MPa×°.
Moreover, the steel strips according to the invention have a delayed fracture free-state for at least 6 months whilst meet the requirements according to SEP 1970 for hole pull and hoop test pieces after provision by the steel producer,
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 minimized.
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 which is then heated to a temperature above Ac3. The heated blank is deformed into a component and is then hardened in a deforming tool or in air with optional subsequent tempering.
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.

In the figures:

FIG. 1 shows (schematically) the process chain for the production of a strip consisting of the steel according to the invention,

FIG. 2 shows (schematically) the time-temperature curve of the process steps of hot-rolling and cold-rolling and continuous annealing (with optional hot-dip finishing), as well as component manufacturing, optional quenching (air-hardening) and optional tempering by way of example for the steel according to the invention,

FIG. 3 shows the chemical composition (examples 1 to 4) of the steel according to the invention,

FIG. 4a shows mechanical characteristic values (transversely to the rolling direction) of the steel according to the invention in the hot-rolled state (HR),

FIG. 4b shows mechanical characteristic values (transversely to the rolling direction) of the steel according to the invention in the cold-rolled state (CR),

FIG. 5a shows solidification behavior during cold-rolling of the steel according to the invention, characteristic values transversely to the rolling direction,

FIG. 5b shows solidification behavior during cold-rolling of the steel according to the invention, cold flow curve,

FIG. 6a shows mechanical characteristic values (transversely to the rolling direction) of the steel according to the invention in the fine sheet state (HDG),

FIG. 6b shows results of the hole expansion tests according to ISO 16630 and of the plate bending test according to VDA 238-100 on the steel according to the invention, in the fine sheet state (HDG),

FIG. 7a shows mechanical characteristic values (transversely to the rolling direction) of the steel according to the invention in the state HR, CR and HDG; example 1 (pre-strip thickness 40 mm),

FIG. 7b shows mechanical characteristic values (transversely to the rolling direction) of the steel according to the invention in the state HR, CR and HDG; example 2 (pre-strip thickness 45 mm),

FIG. 7c shows mechanical characteristic values (transversely to the rolling direction) of the steel according to the invention in the state HR, CR and HDG; example 3 (pre-strip thickness 50 mm),

FIG. 7d shows mechanical characteristic values (transversely to the rolling direction) of the steel according to the invention in the state HR, CR and HDG; example 4 (pre-strip thickness 55 mm),

FIG. 7e shows mechanical characteristic values (transversely to the rolling direction) of the steel according to the invention in the state HR, CR and HDG as an overview,

FIG. 8a shows method 1, temperature-time curves (annealing variants shown schematically),

FIG. 8b shows method 2, temperature-time curves (annealing variants shown schematically),

FIG. 8c shows method 3, temperature-time curves (annealing variants shown schematically).

FIG. 1 schematically illustrates the process chain for the production of a strip from the steel according to the invention. The process routes possible in the invention are shown. Up until acid-cleaning, the process route is the same for all steels according to the invention, thereafter, deviating process routes are followed depending on the desired results. For example, after acid-cleaning, the acid-cleaned hot strip can be cold-rolled and hot-dip finished with different degrees of thinning by rolling. A soft annealed hot strip or soft annealed cold strip can also be cold-rolled and hot-dip finished.
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 galvanizing. From the optionally coated material a complex component can now be produced. Subsequently, a quenching process can optionally take place, such as e.g. air-hardening where the heat-treated component is cooled in the air. Optionally, a tempering step can conclude the thermal treatment of the component.
FIG. 2 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. The time- and temperature-dependent conversion for the hot rolling process and also for a heat treatment after cold-rolling, component manufacture, as well as optional quenching with optional tempering are shown.
FIG. 3 shows in examples 1 to 4 which come from a melt in order to exclude the analytical influence in this case, the alloy compositions of the steel according to the invention, depending upon the produced pre-strip thickness. Cold strips having a cold strip desired thickness of 1.50 mm were produced from a hot strip desired thickness of 2.30 mm. Depending upon the pre-strip thickness, to be produced, prior to hot-rolling, example 1 shows the alloy composition for a pre-strip thickness of 40 mm, example 2 for a pre-strip thickness of 45 mm, example 3 for a pre-strip thickness of 50, example 4 for a pre-strip having a thickness of 55 mm.
FIG. 4 shows the mechanical characteristic values (transversely to the rolling direction) of the steel according to the invention in the hot-rolled state (HR, Hot Rolled) in FIG. 4a and in the cold-rolled state (CR, Cold Rolled) in FIG. 4 b.
FIG. 5 shows the solidification behavior, via the mechanical characteristic values transversely to the rolling direction, during cold-rolling of the steel according to the invention, in a table in FIG. 5a and in a graph as a cold flow curve in FIG. 5 b.
FIG. 6 shows the mechanical characteristic values (transversely to the rolling direction) of the steel according to the invention in the fine sheet state (HDG, Hot Dipped Galvanized) in FIG. 6a and the results of the hole expansion tests according to ISO 16630 and of the plate bending test according to VDA 238-100 in the fine sheet state (HDG) longitudinally and transversely to the rolling direction, as well as the corresponding products with the tensile strength, in FIG. 6 b.
FIG. 7 shows the mechanical characteristic values (transversely to the rolling direction) of the steel according to the invention in the state HR, CR and HDG using a pre-strip thickness of 40 mm in FIG. 7 a, 45 mm in FIG. 7 b, 50 mm in FIG. 7 c, 55 mm in FIG. 7d and in FIG. 7e as a summarizing graphical overview.
FIG. 8 schematically illustrates three variations of the temperature-time curves according to the invention in the case of annealing treatment and cooling and of austenitisation conditions which differ in each case.
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. Principal differences are thus the temperature-time parameters during the heat treatment and the following cooling.
Method 1 (FIG. 8a ) presents the annealing and cooling of the produced steel strip cold-rolled to the end thickness in a continuous annealing installation. Firstly, 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. 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.
Method 2 (FIG. 8b ) 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 galvanizing with combined direct fired furnace and radiant tube furnace, as depicted in FIG. 8 b.
Method 3 (FIG. 8c ) 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.). 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. 8c . In addition, in this case, by means of an induction furnace, reheating of the steel is optionally achieved directly prior to the zinc bath.
The decreases from the slab with respect to the pre-strip vary in the subsequent examples from 78% to 84% for subsequent hot-rolling to a hot strip thickness of 2.30 mm with corresponding decreases of 94% to 96%. In a single cold-rolling step, the cold strip desired thickness of 1.50 mm is achieved with a degree of thinning by cold-rolling of 35%. It is impressively shown that both for very low pre-strip thicknesses and also for greater pre-strip thicknesses, and the range therebetween, relatively uniform values provided with a conventional fluctuation range are achieved for the tensile strength and yield strength, transversely to the rolling direction. The steel according to the invention similarly permits the use of a master hot strip thickness with varying degrees of thinning by cold-rolling, as well as the use of master cold strip thicknesses without influencing the previous fact.
By way of example, for industrial manufacturing for the hot-dip galvanizing (HDG) according to method 3 as shown in FIG. 8c , the following examples form part of so-called feasibility tests which are Intended to prove that the variable pre-strip thicknesses can significantly influence the cold-rollability, such as the necessary rolling forces, without the higher hot strip strength (HR), and higher cold strip strength (CR), with a decreasing pre-strip thickness, leading to considerable fluctuations in the fine sheet (HDG):

EXAMPLE 1

(1.50 mm cold strip from 2.30 mm master hot strip and a pre-strip thickness of 40 mm)
Alloy composition in wt. %. A steel according to the invention comprising 0.104% C; 0.443% Si; 2.178% Mn; 0.012% P; 0.0004% S; 0.0045% N; 0.038 Al; 0.330% Cr; 0.208% Mo; 0.0372% Ti; 0.0332% Nb; 0.007% V; 0.0006% B; 0.0020% Ca; 0.027% Cu; 0.047% Ni; 0.008% Sn; 0.00038% H according to method 3 corresponding to FIG. 8c hot-dip finished, the slab material of 250 mm was rolled prior to the hot-rolling in the pre-line into a pre-strip of 40 mm in a reversing manner with a percentage decrease of 84% and subsequently hot-rolled in the hot wide strip line at a desired end rolling temperature of 910° C. with a decrease of 94% and reeled at a desired reeling temperature of 650° C. with a master hot strip thickness of 2.30 mm and cold rolled after acid-cleaning without additional heat treatment (such as e.g. batch-type annealing) to 1.50 mm in one pass (degree of thinning by cold-rolling 35%).
Fine Sheet State (HDG)
The yield strength ratio Re/Rm in the transverse direction was 66%.
elasticity limit (Rp0.2) 706 MPa

tensile strength (Rm) 1071 MPa

elongation at fracture (A80) 10.9%

bake-hardening-index (BH2) 492 MPa

hole expansion ratio according to ISO 16630 39%

bending angle according to VDA 238-100 121°/112°

(longitudinal, transverse)

The material characteristic values transversely to the rolling direction would correspond e.g. to a HC660XD.
Initial State (HR)
The yield strength ratio Re/Rm in the transverse direction was 77%.


yield strength (Re)	826	MPa
tensile strength (Rm)	1070	MPa

	elongation at fracture (A80)	10.0%

Intermediate State (CR) in the Transverse Direction


	yield strength (Re)	1246 MPa
	tensile strength (Rm)	1305 MPa
	elongation at fracture (A80)	2.0%

EXAMPLE 2

(1.50 mm cold strip from 2.30 mm master hot strip and a pre-strip thickness of 45 mm)
Alloy composition in wt. %
A steel according to the invention comprising 0.104% C; 0.443% Si; 2.178% Mn; 0.012% P; 0.0004% S; 0.0045% N; 0.038 Al; 0.330% Cr; 0.208% Mo; 0.0344% Ti; 0.0372% Nb; 0.007% V; 0.0006% B; 0.0020% Ca; 0.027% Cu; 0.047% Ni; 0.008% Sn; 0.00038% H according to method 3 corresponding to FIG. 8c hot-dip finished, the slab material of 250 mm was rolled prior to the hot-rolling in the pre-line into a pre-strip of 45 mm in a reversing manner with a percentage decrease of 82% and subsequently hot-rolled in the hot wide strip line at a desired end rolling temperature of 910° C. with a decrease of 95% and reeled at a desired reeling temperature of 650° C. with a master hot strip thickness of 2.30 mm and cold rolled after acid-cleaning without additional heat treatment (such as e.g. batch-type annealing) to 1.50 mm in one pass (degree of thinning by cold-rolling 35%).
Fine Sheet State (HDG)
The yield strength ratio Re/Rm in the transverse direction was 67%.
elasticity limit (Rp0.2) 720 MPa

tensile strength (Rm) 1077 MPa

elongation at fracture (A80) 10.4%

bake-hardening-index (BH2) 51 MPa

hole expansion ratio according to ISO 16630 35%

bending angle according to VDA 238-100 128°/114°

(longitudinal, transverse)

The material characteristic values transversely to the rolling direction would correspond e.g. to a HC660XD.
Initial State (HR)
The yield strength ratio Re/Rm in the transverse direction was 70%.


yield strength (Re)	725	MPa
tensile strength (Rm)	1030	MPa

	elongation at fracture (A80)	10.2%

Intermediate State (CR) in the Transverse Direction


	yield strength (Re)	1224 MPa
	tensile strength (Rm)	1260 MPa
	elongation at fracture (A80)	1.5%

EXAMPLE 3

(1.50 mm cold strip from 2.30 mm master hot strip and a pre-strip thickness of 50 mm)
Alloy composition in wt. %
A steel according to the invention comprising 0.104% C; 0.443% Si; 2.178% Mn; 0.012% P; 0.0004% S; 0.0045% N; 0.038 Al; 0.330% Cr; 0.208% Mo; 0.0344% Ti; 0.0372% Nb; 0.007% V; 0.0006% B; 0.0020% Ca; 0.027% Cu; 0.047% Ni; 0.008% Sn; 0.00038% H according to method 3 corresponding to FIG. 8c hot-dip finished, the slab material of 250 mm was rolled prior to the hot-rolling in the pre-line into a pre-strip of 50 mm in a reversing manner with a percentage decrease of 80% and subsequently hot-rolled in the hot wide strip line at a desired end rolling temperature of 910° C. with a decrease of 96% and reeled at a desired reeling temperature of 650° C. with a master hot strip thickness of 2.30 mm and cold rolled after acid-cleaning without additional heat treatment (such as e.g. batch-type annealing) to 1.50 mm in one pass (degree of thinning by cold-rolling 35%).
Fine Sheet State (HDG)
The yield strength ratio Re/Rm in the transverse direction was 65%.
elasticity limit (Rp0.2) 704 MPa

tensile strength (Rm) 1084 MPa

elongation at fracture (A80) 10.4%

bake-hardening-index (BH2) 55 MPa

hole expansion ratio according to ISO 16630 38%

bending angle according to VDA 238-100 127°/115°

(longitudinal, transverse)

The material characteristic values transversely to the rolling direction would correspond e.g. to a HC660XD.
Initial State (HR)
The yield strength ratio Re/Rm in the transverse direction was 69%.


yield strength (Re)	695	MPa
tensile strength (Rm)	1010	MPa

	elongation at fracture (A80)	8.8%

Intermediate State (CR) in the Transverse Direction


	yield strength (Re)	1203 MPa
	tensile strength (Rm)	1255 MPa
	elongation at fracture (A80)	1.9%

EXAMPLE 4

(1.50 mm cold strip from 2.30 mm master hot strip and a pre-strip thickness of 55 mm)
Alloy composition in wt. %
A steel according to the invention comprising 0.104% C; 0.443% Si; 2.178% Mn; 0.012% P; 0.0004% S; 0.0045% N; 0.038 Al; 0.330% Cr; 0.208% Mo; 0.0344% TI; 0.0372% Nb; 0.007% V; 0.0006% B; 0.0020% Ca; 0.027% Cu; 0.047% Ni; 0.008% Sn; 0.00038% H according to method 3 corresponding to FIG. 8c hot-dip finished, the slab material of 250 mm was rolled prior to the hot-rolling in the pre-line into a pre-strip of 55 mm in a reversing manner with a percentage decrease of 78% and subsequently hot-rolled in the hot wide strip line at a desired end rolling temperature of 910° C. with a decrease of 96% and reeled at a desired reeling temperature of 650° C. with a master hot strip thickness of 2.30 mm and cold rolled after acid-cleaning without additional heat treatment (such as e.g. batch-type annealing) to 1.50 mm in one pass (degree of thinning by cold-rolling 35%).
Fine Sheet State (HDG)
The yield strength ratio Re/Rm in the transverse direction was 66%.


elasticity limit (Rp0.2)	708	MPa
tensile strength (Rm)	1077	MPa

elongation at fracture (A80)

10.4%

bake-hardening-index (BH2)

58

MPa

	hole expansion ratio according to ISO 16630	40%
	bending angle according to VDA 238-100	123°/111°
	(longitudinal, transverse)

Initial State (HR)
The yield strength ratio Re/Rm in the transverse direction was 70%.


	yield strength (Re)	679 MPa
	tensile strength (Rm)	967 MPa
	elongation at fracture (A80)	9.6%

Intermediate State (CR) in the Transverse Direction


	yield strength (Re)	1158 MPa
	tensile strength (Rm)	1230 MPa
	elongation at fracture (A80)	2.5%

CONCLUSION

It is not possible to see a significant influence of the pre-strip thickness on the mechanical characteristic values on the fine sheet (HDG).
This statement applies to the degree of thinning by cold-rolling of 35% used in the examples, but could also be applied without restriction to variable degrees of thinning by cold-rolling.
The invention has been described above with the aid of fine sheet steel sheets with an end thickness to be achieved of 1.50 mm in the thickness range of 0.50 to 3.00 mm. It is also possible, if required, to produce end thicknesses in the range of 0.10 to 4.00 mm.

Claims

1.-41. (canceled)

42. A method, comprising:

producing a pre-strip from a multi-phase steel in a state of a slab, with the multi-phase steel having a minimum tensile strength of 980 MPa in a non-quenched state containing (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≥0.010 to ≤0.060

N≥0.0020 to ≤0.0120

P≤0.020

S≤0.0020

Ti≥0.005 to ≤0.060

Nb≥0.005 to ≤0.060

V≥0.005 to ≤0.020

B≥0.0005 to ≤0.0010

Mo≥0.200 to ≤0.300

Ca≥0.0010 to ≤0.0060

Cu≤0.050

Ni≤0.050

Sn≤0.040

H≤0.0010,

with the remainder being iron, including typical steel-associated, smelting-related impurities, wherein the total content of Mn-Si+Cr is ≥1.750 wt. % to ≤2.250 wt. % with regard to a processing window which is as wide as possible during annealing of cold strips of this steel;

hot-rolling the pre-strip into a steel strip with a hot strip thickness to be achieved;

proceeding from a previously fixed slab thickness and a previously selected pre-strip having a defined but variable thickness, hot-rolling hot strips with a same thickness with a degree of thinning by rolling of 72% to 87% with end thickness to be achieved;

cold-rolling the hot strip into a cold strip with an end thickness to be achieved,

heating the steel strip cold-rolled to the end thickness during the continuous annealing to an annealing temperature in a range of approximately 700 to 950° C. to produce a required multi-phase microstructure;

cooling the annealed steel strip 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. followed by a cooling rate between approximately 15 and 100° C./s to a second intermediate temperature of approximately 160 to 250° C., and cooling the steel strip 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, or

cooling the annealed steel strip to a temperature of approximately 400 to 470° C. such that cooling is stopped prior to entry of the steel strip into a melting bath, then the steel strip undergoes a hot-dip finishing procedure, and after undergoing the hot-dip finishing procedure, continuing cooling at a cooling rate between approximately 15 and 100° C./s to an intermediate temperature of approximately 200 to 250° C., and cooling the steel strip in air at a cooling rate of approximately 2 to 30° C./s until room temperature is reached, or

cooling the annealed steel strip to an intermediate temperature of approximately 200 to 250° C. such that prior to entry of the steel strip into a melting bath, maintaining the steel strip at the intermediate temperature for approximately 1 to 20 s, then the steel strip is heated to a temperature of approximately 400 to 470° C., undergoes a hot-dip finishing procedure, and after undergoing the hot-dip finishing procedure, is cooled again at a cooling rate between approximately 15 and 100° C./s to an intermediate temperature of approximately 200 to 250° C., and subsequently cooled in air at a cooling rate of approximately 2 to 30° C./s to room temperature.

43. The method of claim 42, wherein the cold strip is continuously annealed.

44. The method of claim 43, wherein, proceeding from a selected hot strip having a specific thickness or selected hot strips having different thicknesses, cold strips with degrees of thinning by cold-rolling of 10% to 70% are produced with the end thickness to be achieved.

45. The method of claim 43, further comprising, during the continuous annealing, increasing an oxidation potential during annealing with an installation configuration comprised of a direct fired furnace region (NOF) and a radiant tube furnace (RTF) by a CO content in the NOF of less than 4 vol. %; and setting in the RTF an oxygen partial pressure of a furnace atmosphere, which is reducing for iron, in accordance with a following equation,

−18>Log pO₂≥-5*Si^−0.3-2.2*Mn^−0.45-0.1*Cr^−0.4-12.5*(−ln B)^0.25

wherein

Si, Mn, Cr and B designate corresponding alloy proportions in steel in wt. %,

pO₂designates the oxygen partial pressure in mbar, and

wherein in order to avoid oxidation of the steel strip directly prior to dipping in the melting bath a dew point of a gas atmosphere is set at −30° C. or below.

46. The method of claim 43, further comprising, during the continuous annealing, increasing an oxidation potential during annealing with an installation configuration comprised of only a radiant tube furnace (RFT) by setting in the RTF an oxygen partial pressure of a furnace atmosphere, which is reducing for iron, in accordance with a following equation,

−12>Log pO₂≥-5*Si^−0.25-3*Mn^−0.5-0.1*Cr^−0.5-7*(−ln B)^0.5

wherein Si, Mn, Cr and B designate corresponding alloy proportions in steel in wt. %,

pO₂designates the oxygen partial pressure in mbar, and

wherein in order to avoid oxidation of the steel strip directly prior to dipping in the melting bath the dew point of the gas atmosphere is set at −30° C. or below.

47. The method of claim 42, further comprising temper-rolling the steel strip after undergoing annealing or the hot-dip finishing procedure.

48. The method of claim 42, further comprising stretch-bending-straightening the steel strip after undergoing annealing or the hot-dip finishing procedure.

49. The method of claim 42, further comprising:

cutting a blank from the steel strip;

heating the blank to a temperature above Ac3;

deforming the heated blank into a component; and

hardening the component in a tool or in air.

50. A steel strip produced by a method as set forth in claim 42, said steel strip comprising a minimum hole expansion value according to ISO 16630 of at least 20%.

51. The steel strip of claim 50, wherein the minimum hole expansion value according to ISO 16630 is 25%.

52. The steel strip of claim 50, comprising a minimum bending angle according to VDA 238-100 of 70° in a longitudinal direction or transverse direction.

53. The steel strip of claim 50, comprising a minimum bending angle according to VDA 238-100 of 85° in a longitudinal direction or transverse direction.

54. The steel strip of claim 50, comprising a minimum product value Rm×α of 100000 MPa, wherein

Rm is a tensile strength, and

α is a bending angle according to VDA 238-100.

55. The steel strip of claim 50, comprising a minimum product value Rm×α of 120000 MPa, wherein

Rm is a tensile strength, and

α is a bending angle according to VDA 238-100.

56. The steel strip of claim 50, comprising a delayed fracture free state for at least 6 months thus meeting the requirements of SEP 1970 for hole pull and hoop test pieces.

57. A steel strip, 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≥0.010 to ≤0.060

N≥0.0020 to ≤0.0120

P≤0.020

S≤0.0020

Ti≥0.005 to ≤0.060

Nb≥0.005 to ≤0.060

V≥0.005 to ≤0.020

B≥0.0005 to ≤0.0010

Mo≥0.200 to ≤0.300

Ca≥0.0010 to ≤0.0060

Cu≤0.050

Ni≤0.050

Sn≤0.040

H≤0.0010,

with the remainder being iron, including typical steel-associated, smelting-related impurities, wherein a total content of Mn-Si+Cr is ≥1.750 wt. % to ≤2.250 wt. %,

said steel strip comprising a minimum hole expansion value according to ISO 16630 of at least 20%.