WO2019068560A1

WO2019068560A1 - Ultrahigh strength multiphase steel and method for producing a steel strip from said multiphase steel

Info

Publication number: WO2019068560A1
Application number: PCT/EP2018/076307
Authority: WO
Inventors: Thomas Schulz
Original assignee: Salzgitter Flachstahl Gmbh
Priority date: 2017-10-06
Filing date: 2018-09-27
Publication date: 2019-04-11
Also published as: RU2742998C1; DE102017123236A1; KR20200063167A; CN111247258A; US20200263283A1; CN111247258B; EP3692178A1; 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; AI ≥ 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

Highest strength multi-phase steel and process for producing a steel strip from this multi-phase steel

The invention relates to a high-strength multi-phase steel with dual-phase structure or complex phase structure and small amounts of retained austenite with

improved manufacturing properties and excellent material properties in the subsequent processing, in particular for the lightweight vehicle construction, according to the preamble of claim 1. Advantageous developments are

Subject matter of the subclaims 2 to 24.

The invention further relates to a method for producing steel strips from such a steel according to claim 25 and to steel strips produced therewith according to claim 38. More particularly, the invention relates to steels with a tensile strength in the range of at least 980 MPa, in the un-tempered state, for the production of components which have an improved formability, such as with regard to a hole widening and improved joining suitability, such as, for example, welding properties.

The hotly contested automotive market is forcing manufacturers to steadily reduce solutions to reduce fleet fuel consumption and CO 2 emissions

Maintaining the utmost comfort and occupant protection. On the one hand, the weight saving of all vehicle components plays a decisive role, on the other hand, however, the best possible behavior of the individual

Components with high static and dynamic load during operation as well as in the event of a crash.

The steel suppliers contribute by providing high strength steels. Task invoice. In addition, by providing

high-strength steels with lower sheet thickness, the weight of the vehicle components are reduced with the same and possibly even improved component behavior.

In addition to the required weight reduction, these newly developed steels must meet the high material requirements with regard to yield strength, tensile strength and Elongation and bake-hardening index are sufficient, as well as the high

Component requirements for toughness, edge crack resistance, improved bending angle and bending radius, energy absorption and defined solidification on the work-hardening effect and the bake-hardening effect have.

In addition, good processability must be ensured. This concerns both the processes at the automobile manufacturer, for example punching and forming, optional thermal tempering with subsequent optional tempering,

Welding and / or surface post-treatment, such as phosphating and cathodic dip painting, as well as the manufacturing processes of the primary supplier, such as surface finishing by metallic or organic coating.

Increasingly, improved joinability is also being demanded, for example, in terms of better general weldability, such as a larger usable weld area in resistance spot welding and improved weld failure (fracture pattern) under mechanical stress, as well as sufficient resistance to delayed hydrogen embrittlement (i.e., delayed fracture free). The same applies to the weldability of high-strength steels in the production of pipes, which are produced for example by means of the high-frequency induction welding (HFI).

Hole expanding capability is a material property that describes the resistance of the material to crack initiation and crack propagation during forming operations in near edge areas, such as collaring.

The Lochaufweiteversuch is normatively regulated, for example, in ISO 16630. Thereafter, prefabricated, for example punched in a sheet holes are widened by means of a mandrel. The measured quantity is the change in the diameter of the hole in relation to the initial diameter at which the first crack occurs through the sheet at the edge of the hole.

An improved edge crack resistance means an increased formability of the sheet edges and can be described by an increased Lochaufweitvermögen. This fact is under the synonyms "Low Edge Crack" (LEC) or known as "High Hole Expansion" (HHE) and xpand®.

The bending angle describes a material property, which gives conclusions on the material behavior in forming operations with dominant bending components (for example, when folding) or in crash loads. Increased bending angles thus increase passenger compartment safety.

The determination of the bending angle (a) is normatively regulated, for example, via the platelet bending test in VDA 238-100.

The above properties are important for components that have a very complex shape.

Improved weldability is known i.a. achieved by a lowered carbon equivalent.

Synonyms such as "unterjDeritektisch" (UP) and the already known "Low Carbon Equivalent" (LCE) stand for this. The carbon content is usually less than 0.120 wt .-%.

Furthermore, the failure behavior or the fracture pattern of the weld can be improved by a clear alloying with micro-alloying elements, in low-carbon steels with lowered carbon equivalent. High-strength components must be sufficiently resistant to hydrogen

Have resistance to material embrittlement.

The testing of the durability of Advanced High Strength Steels (AHSS) for the automotive industry compared to production-induced hydrogen-induced brittle fractures is regulated in the SEP1970 and tested on the hanger sample and the sample size.

In vehicle construction, dual-phase steels are increasingly being used, which consist of a ferritic basic structure in which a martensitic second phase is incorporated. It has been found that in low-carbon, micro-alloyed steels shares further phases such as bainite and retained austenite advantageous for example the Lochaufweitverhalten, the bending behavior and the hydrogen-induced

Affect brittle fracture behavior. The bainite can here in different

Manifestations, such as upper and lower bainite, are present. The characteristic processing properties of the dual-phase steels, such as a very low yield ratio with simultaneously very high tensile strength, high work hardening and good cold workability, are well known.

The combination of properties required for the steel material ultimately represents a component-specific compromise of individual properties. However, these properties are often no longer sufficient with increasingly complex component geometries.

Increasingly, multi-phase steels are also used in the automotive industry, 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 amounts of martensite, retained austenite and / or pearlite in a ferritic / bainitic matrix, whereby a pronounced grain refining is effected by delayed recrystallization or by precipitation of micro-alloying elements.

These complex phase steels have higher yield strengths, a higher yield ratio, a lower strain hardening and a higher hole spreading capacity compared to dual phase steels. Ferritic-bainitic steels are according to EN 10346 steels containing bainite or solidified bainite in a matrix of ferrite and / or solidified ferrite. The strength of the matrix is brought about by a high dislocation density, grain refining and the excretion of micro-alloying elements. Dual-phase steels are, according to EN 10346, steels with a ferritic basic structure, in which a martensitic second phase is insular, occasionally also with proportions of bainite as second phase. At high tensile strength, dual phase steels exhibit a low yield ratio and high work hardening. TRIP steels are according to EN 10346 steels with a predominantly ferritic Basic structure in which bainite and retained austenite is embedded, which can transform to martensite during transformation (TRIP effect). Because of its strong work hardening, the steel achieves high levels of uniform elongation and

Tensile strenght. In conjunction with the bake hardening effect, high component strengths can be achieved. These steels are suitable both for stretch drawing and deep drawing. However, material conversion requires higher blankholder forces and press forces. A comparatively strong springback must be considered. The high-strength steels with single-phase structure include, for example, bainitic and martensitic steels.

Bainitic steels are according to EN 10346 steels, which are characterized by a very high

Yield strength and tensile strength at a sufficiently high elongation for cold forming processes. Due to the chemical composition a good weldability is given. The microstructure typically consists of bainite. Occasionally small fractions of other phases, such as, for example, martensite and ferrite, may be present in the microstructure. Martensitic steels are, according to EN 10346, steels which contain small amounts of ferrite and / or bainite in a matrix of martensite due to thermomechanical rolling. This steel grade is characterized by a very high yield strength and tensile strength at a sufficiently high elongation for cold forming processes. Within the group of multiphase steels, the martensitic steels have the highest tensile strength values. The suitability for thermoforming is limited. The martensitic steels are mainly suitable for bending forming processes, such as roll forming.

Compared to EN 10083, tempered steels are steels which have high tensile strength and fatigue strength due to tempering (= hardening and tempering). Performs the cooling during

Hardening in air to bainite or martensite, the process is called "air hardening." A tempering after hardening can have a specific effect on the strength / toughness ratio.

These steels are currently used in structural, chassis and crash-relevant components, as well as flexibly cold-rolled strips. The Tailor Rolled Blank Lightweight Technology (TRB®) enables a significant weight reduction due to the load-adapted choice of sheet thickness over the component length.

However, with today's known alloys and available continuous annealing systems for widely varying sheet thicknesses, the production of TRB®s is included

Multi-phase structure is not possible without restrictions, such as for the heat treatment before cold rolling. In areas with different sheet thicknesses, a homogeneous multi-phase microstructure in cold- as well as hot-rolled steel strips can not be set due to a temperature gradient occurring in the common process windows.

If thin sheets are to be produced, the cold-rolled steel strips are usually, for economic reasons, re-annealed in a continuous annealing process to form a thin sheet that can be readily formed.

Depending on the alloy composition and the strip cross-section, the process parameters, such as throughput speed, annealing temperatures and cooling rate, are set according to the required mechanical and technological properties with the necessary structure.

The above-mentioned properties are significantly influenced by, for example, the steel compositions, the process parameters during hot rolling, the process parameters during pickling (for example, the stretch bend density) and the process parameters during cold rolling even before the continuous annealing.

The steel composition is determined by analytical rules defining MIN and MAX ranges.

Depending on the multiphase steel to be produced, the process parameters during hot rolling, such as standard slab thickness, slab lay time, slab discharge temperature, pre-strip rolling pass schedule, standard pre-strip thickness, hot strip line entry temperature, hot rolling pass schedule, final roll temperature, hot strip cooling pattern, coiler temperature, are set. During pickling, an optional stretch bending (stretching) affects the subsequent process step. For cold rolling, the hot strip thickness for the presentation of a cold rolling thickness by a standard Kaltabwalzgrad already during the order conversion in the technical specifications (process parameters) are determined.

The thickness of the pre-strip in the hot rolling process, describes the initial thickness prior to entering the multi-stand hot strip mill, wherein the pre-strip was reversibly made in several passes (passes) from a slab with a defined standard thickness.

Typical slab thicknesses are between 250 mm and 300 mm (standard 250 mm, further considered here), the pre-strip thicknesses usually range between 40 mm to 60 mm for the multiphase steels.

Usually, the pre-strip thicknesses for the subsequent hot rolling are relatively constant, depending on the material composition, for example at 45 mm (called standard here).

Values below or above cause altered technological hot strip parameters, such as tensile strength and yield strength, which in turn are the subsequent

As a result, deformation during cold rolling is influenced by this, as is the work hardening behavior.

In order to achieve the final technological sheet characteristics demanded by the standards, a material-dependent pre-strip thickness is determined in the continuous annealing treatment to ensure proper recrystallization according to the prior art. Underruns or exceedances influence the final technological characteristics of classic steels to such an extent that considerable batch fluctuations (scattering range) can occur

The rolling degree during cold rolling (cold rolling degree) describes the percentage ratio of the difference between the hot strip starting thickness and the finished cold strip thickness based on the hot strip exit thickness.

Usually the Kaltabwalzgrade are relatively constant, they are in thicker cold tapes of about 2 mm up to about 40% and up to about 60% for cold tapes up to 1 mm thickness.

In order to achieve the technological characteristics required by the standards, a cold rolling degree of about 50% is required on average in continuous annealing to ensure proper recrystallization, according to the state of the art. Underruns or exceedances in conventional steels lead to fluctuating technological characteristics, as described in the TRB®'s.

To achieve a fine-grained structure after the continuous annealing process, it is known to set a minimum degree of cold rolling as a function of the recrystallization temperature in order to obtain a corresponding dislocation density for the

Adjust recrystallization annealing.

If the degree of cold rolling is too low (even in local areas), the critical threshold for recrystallization can not be overcome, so that a fine-grained and relatively uniform structure can not be achieved. Due to different particle sizes in the cold strip, even after recrystallization, different grain sizes appear in the final microstructure, which leads to characteristic fluctuations. Different sized grains may increase upon cooling from the oven temperature

transform different phase components and provide for a further inhomogeneity.

In order to achieve the respectively required microstructure, the cold strip is heated in a continuous annealing furnace to a temperature at which the required microstructure formation (for example dual or complex phase structure) is established during cooling.

If, due to high corrosion protection requirements, the surface of the cold strip is to be hot-dip galvanized, the annealing treatment is usually carried out in a continuous hot-dip galvanizing plant, in which the heat treatment or Annealing and the downstream galvanizing take place in a continuous process.

In the continuous annealing of hot-rolled or cold-rolled steel strips with, for example, EP 2 028 282 A1 and EP 2 031 081 A1 known alloy concepts for high-strength dual-phase steels with minimum tensile strengths of about 980 MPa, there is the problem that only a small process window for the annealing parameters is available. Thus, even with minimal cross-sectional changes (thickness, width) adjustments of the process parameters for achieving uniform mechanical properties are required.

With expanded process windows, the required strip properties are possible with the same process parameters even with larger cross-sectional changes of the strips to be annealed.

In addition to flexibly rolled strips with different sheet thicknesses over the strip length, this applies above all to strips with different thicknesses and / or different widths, which must be successively annealed. A homogeneous temperature distribution is difficult to achieve, especially at different thicknesses in the transition region from one belt to another. This can result in alloy compositions with too small process windows in the continuous annealing, for example, that the thinner strip is driven too slowly through the oven and thereby the productivity is lowered, or that the thicker strip is driven too fast through the oven and the required annealing temperature for the desired structure is not achieved. The consequences are increased rejects.

The decisive process parameter for material with a relatively constant degree of cold rolling is therefore the adjustment of the speed during continuous annealing, since the phase transformation takes place in a temperature- and time-dependent manner. ever

The less sensitive the steel is to the uniformity of the mechanical properties with changes in the temperature and time course during the continuous annealing, the greater the process window. Particularly serious is the problem of too narrow a process window in the annealing of cold strips, which have too low or too high pre-strip thicknesses or too low or too high Kaltabwalzgrade, as well as in the annealing of strips with over the tape length varying sheet thicknesses for

Production of load-optimized components from cold strip, but also from hot strip.

A method for producing a steel strip with different thickness over the strip length is described for example in DE 100 37 867 A1. When using the known alloy concepts for the group of multiphase steels, it is difficult to achieve uniform mechanical properties over the entire strip length of the strip due to the narrow process window already in the continuous annealing of different thickness tapes. Complex-phase steels also have an even narrower process window than dual-phase steels.

To set relatively homogeneous mechanical and technological properties of different cold strips with variable pre-strip thicknesses or variable Kaltabwalzgraden, is practically impossible to achieve with the known alloy concepts in continuous annealing. The necessary for the recrystallization annealing

Cold rolling degree leads to a very significant limitation in the flexibility of material production within the entire process chain. Already the final

Cold strip thickness determines the thickness of the hot strip and thus the hot strip production parameters. For flexibly rolled cold strips made of multiphase steels of known compositions, because of the too small process window areas with lower sheet thickness due to the conversion processes during cooling either too high strengths due to excessive martensite on or the areas with larger sheet thickness reach too low strengths due to low martensite.

Homogenous mechanical-technological properties over the strip length or width are practically impossible to achieve with the known alloy concepts in continuous annealing.

The known alloy concepts for multiphase steels are characterized by a too narrow process window and therefore especially for cold-rolled strip production variable pre-strip thicknesses and variable Kaltabwalzgraden, as well as for flexibly rolled strips, unsuitable.

German Offenlegungsschrift DE 10 2012 002 079 A1 discloses a high-strength, multi-phase steel with minimum tensile strengths of 950 MPa, which indeed already has a very wide process window for the continuous annealing of hot or cold strips, but it has been shown that even with this steel neither variable pre-strip thicknesses, nor variable Kaltabwalzgrade be achieved with a single hot strip thickness (Masterwarmbanddicke) under realization of uniform material properties.

German Offenlegungsschrift DE 10 2015 1 1 177 A1 discloses a high-strength multiphase steel with minimum tensile strengths of 980 MPa, which already has a very wide process window for continuous annealing of hot or cold strips, as well as, for example, a single hot strip thickness (master hot strip thickness) Variable Kaltabwalzgrade, pass annealed cold strips with different thicknesses and with uniform material properties can achieve.

German Offenlegungsschrift DE 10 2014 017 274 A1 discloses a high-strength air-hardenable multiphase steel with minimum tensile strengths in the non-air-cured state of 950 MPa, which already has a very wide process window for the continuous annealing of hot or cold strips, as well as, for example, with a single hot strip thickness (US Pat. Masterwarmbanddicke) under realization of variable Kaltabwalzgrade, pass annealed cold strips, with different thicknesses and with uniform material properties can be achieved and is suitable for the subsequent air hardening process.

The goal of achieving the resulting mechanical and technological properties in a narrow range over bandwidth and strip length by the controlled adjustment of the volume fractions of the structural components has top priority and is only possible through an enlarged process window. The known alloy concepts are characterized by too narrow a process window and therefore unsuitable for solving the present problem, in particular in flexibly rolled strips. With the known alloying concepts, only steels of a strength class with defined cross-sectional areas (strip thickness and strip width) can currently be represented, see above that for different strength classes and / or cross-sectional areas changed alloy concepts are necessary.

In steelmaking, there is a trend toward reducing the carbon equivalent to achieve improved cold working (cold rolling, cold working) and better performance.

But the weldability - characterized among other things by the carbon equivalent - is an important criterion.

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, and chromium or molybdenum and vanadium considered (contents in wt .-%).

The prior art is also that an increase in the strength by the quantitative increase of carbon and / or silicon and / or manganese and a

Increased strength over the microstructural settings and solid solution hardening (solid solution hardening) is achieved.

Increasing the quantity of the aforementioned elements, however, causes the material processing properties to deteriorate, for example during welding, forming and hot dipping.

However, steelmaking is showing a trend towards reducing the

Carbon and / or manganese content for improved cold working and performance.

An example is the hole expansion test to describe and quantify the

Edge crack behavior. With appropriately optimized steel grade adjustments, the steel user expects higher values than the standard material. But also the

Weldability, characterized by the carbon equivalent, continues to get into focus. The automotive industry is increasingly in demand for steel grades with significantly different requirements in terms of application, depending on the application

Ratio of yield strength (Re) or yield strength (Rp0.2) to tensile strength. This leads to steel developments with a comparatively large yield point interval at a normative tensile strength interval.

A low yield ratio (Re / Rm) is typical for a dual-phase steel and is used primarily for formability in drawing and deep drawing operations.

A higher yield ratio (Re / Rm), which is typical for complex phase steels, is also distinguished by resistance to edge cracks. This is due to the smaller differences in the strengths of each

Microstructure components lead back, which has a favorable effect on a homogeneous deformation in the area of the cutting edge.

The analytical landscape for achieving multiphase steels with minimum tensile strengths of 980 MPa is very diverse and shows very large alloy ranges in the strength-increasing elements carbon, manganese, phosphorus, aluminum and chromium and / or molybdenum, as well as in the addition of micro-alloys individually or in combinations, as well as in the material characterizing special properties, such as hole widening and lowered carbon equivalent, etc.

The range of dimensions is broad and lies in the thickness range from 0.50 to 3.00 mm, whereby the range between 0.80 to 2.10 mm is relevant in terms of quantity.

Thickness ranges below 0.50 and over 3.00 mm are conceivable.

Overall, there is the problem with the known steel grades that, in view of the required minimum rolling rates during cold rolling for a complete

Recrystallization after continuous annealing, at a given Vorbanddicke for producing a Masterwarmbanddicke after hot rolling no manufacturing flexibility (s.Figure 1, process steps 6,8 and 9 are then necessary) is more in terms of achievable different cold strip thicknesses. In particular, the production of different cold strip thicknesses at a constant master heat Thickness with comparable material properties on the produced cold strip due to a too small process window not possible. In addition, the specification of a constant pre-strip thickness for producing a predetermined constant master hot-strip thickness additionally limits the manufacturing flexibility.

The invention is therefore based on the object of specifying a new alloy concept for a high-strength multiphase steel, a method for producing a steel strip from this high-strength multiphase steel and a steel strip produced by this method, with which the process window for the continuous annealing of cold strips can be expanded that from different

Vorbanddicken, a predetermined hot strip thickness (Masterwarmbanddicke) different cold strip thicknesses or from different hot strip thicknesses a cold strip thickness (Masterkaltbanddicke) can be made. In addition, variable pre-strip thicknesses before hot rolling should be used instead of constant pre-strip thicknesses.

This should be achieved as uniform as possible cold strip material properties, regardless of the set pre-strip thickness and the set cold rolling.

In addition, the process window for the annealing, in particular continuous annealing, of cold rolled steel strip to be extended so that in addition to bands with different cross sections (jump in cross section) and steel bands over band length and possibly bandwidth varying thickness (TRB®) with the most homogeneous mechanical and technological Properties can be generated.

According to the teaching of the invention, this object is achieved by a high-strength multiphase steel having a minimum tensile strength of 980 MPa with the following contents in% by weight:

C> 0.075 to <0.1 15

Si> 0.400 to <0.500

Mn> 1, 900 to <2,350

Cr> 0.250 to <0.400

AI> 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 remainder of iron, including common steel-accompanying melting-related

Contaminants in which, with a view to the widest possible process window in the annealing, in particular continuous annealing, of cold tapes made of this steel, the total amount of Mn-Si + Cr> 1, 750 wt .-% to <2.250 wt .-%. Mathematically, this means that the content information of Mn and Cr is added and that of Si is subtracted and the sum thus obtained (result) greater than or equal to 1, 750 and less equal to 2.250 wt .-% should be. The same applies to the other corresponding total quantities. With the alloy concept according to the invention, the mechanical properties are reliably achieved in a narrow range for cold strips with variable pre-strip thickness before hot rolling, as well as variable cold rolling degrees during cold rolling. By variable pre-strip thicknesses, the cold rolling process can be positively influenced that the steps annealing of hot strip before cold rolling, double cold rolling, annealing the cold rolled strip before the next cold rolling step, without negative consequences on the production of the above-described Masterwarmbanddicke or Masterkaltbanddicke.

Decisive for this is a selected narrow alloy composition with a focus on a limited and cold-tape thickness-dependent chrome alloy. Content that has been found to be very positive for achieving uniform material properties at different pre-strip thicknesses, as well as different Kaltabwalzgraden. In addition, the representable mechanical properties in a narrow range over bandwidth and

Tape length achieved by the controlled adjustment of the volume fractions of the structural phases.

Furthermore, the previous manufacturing philosophy that the final cold rolled strip thickness (final thickness) determines the necessary hot strip thickness and a standard pre-strip thickness is necessary can be relied upon to require a selected pre-strip thickness and only a selected master hot strip thickness for different cold strip thicknesses. However, it is also advantageously possible to produce a cold strip thickness to be achieved from different hot strip thicknesses analogously. This significantly increases flexibility in manufacturing and also reduces production costs.

Thus, from the multiphase steel in the state of a slab a preliminary band can be generated, which subsequently with the hot strip thickness to be achieved

is hot rolled.

It is also possible, starting from a previously determined slab thickness of, for example, 250 mm and a previously selected pre-strip strip having a defined but variable thickness to hot roll hot strips of equal thickness with degrees of roll of 72% to 87% with the final thickness to be achieved.

Advantageously, it can be used with different thickness steel strips

Continuous annealing comparable microstructural conditions and mechanical characteristics of the bands by adjusting the system throughput speed in the course of

Heat treatment can be adjusted.

The steel according to the invention also offers the advantage of a significantly enlarged process window compared to the known steels. This results in increased process reliability in the continuous annealing of cold strip with multi-phase structure. Thus, for pass-annealed cold strips, more homogeneous mechanical and technological properties can be achieved for strips with variable degrees of cold rolling as well as in the Band or in the transition region of two bands even at different

Cross sections and otherwise the same process parameters are guaranteed.

According to the invention, a steel strip can be produced from the inventive multiphase steel in which a hot strip is produced from the multiphase steel, from the hot strip the steel strip is cold rolled with the final thickness to be achieved and then the steel strip is annealed, in particular continuously annealed.

The properties of the multiphase steel make it possible to cold-roll steel strips of the final thickness to be achieved, starting from a variable pre-strip thickness, a selected master hot strip with a particular thickness or selected hot strips of different thicknesses in a wide range of cold rolling degrees of 10% to 70%. Here, the chemical composition of the multi-phase steel is selected according to the invention depending on the final thickness of the cold strip to be achieved. It is thus possible, within selectable thickness graduations of the cold strip to be achieved, to produce a master cold strip with a uniform thickness from a master heat strip having a thickness corresponding cold strips with one or more end thicknesses or else from different hot strip thicknesses.

In order to achieve uniform mechanical properties, it has been found advantageous that the steel strip is cold rolled to a final thickness of 0.50 to 3.00 mm and, depending on the final thickness to be achieved, the chemical composition of the multiphase steel is chosen as follows, even if find variable Vorbanddicken application.

With regard to the possible use of variable pre-strip thicknesses, it turned out to be particularly advantageous that the sum amount of Mn-Si + Cr is chosen as a function of the final thickness of the cold strip to be obtained as follows:

Final thickness 0.50 to 1.00 mm inclusive:

Sum of Mn-Si + Cr> 1, 750 wt% to <2.030 wt%,

Final thickness over 1.00 to 2.00 mm inclusive:

Sum of Mn-Si + Cr> 1, 940 wt.% To <2, 1 10 wt.% Final thickness over 2.00 up to and including 3.00 mm:

Sum of Mn-Si + Cr> 2.020 wt% to <2.220 wt%.

Furthermore, it turned out to be advantageous that the sum amount of Mn-Si + Cr + Mo is selected as a function of the final thickness of the cold strip to be obtained as follows:

Final thickness 0.50 to 1.00 mm inclusive:

Sum of Mn-Si + Cr + Mo> 1.950% by weight to <2.280% by weight,

Final thickness over 1.00 to 2.00 mm inclusive:

Sum of Mn-Si + Cr + Mo> 2.140 wt% to <2.356 wt%

Final thickness over 2.00 mm up to and including 3.00 mm:

Sum of Mn-Si + Cr + Mo> 2.220 wt% to <2.470 wt%. The final thickness of the steel strip to be achieved is thus related to the alloy composition of the multiphase steel

Vorbandes or hot strip.

In addition, it has been found to be advantageous that the carbon equivalent CEV (NW) is chosen as follows, depending on the final thickness of the cold strip to be achieved:

Final thickness 0.50 mm up to and including 1.00 mm:

C content <0.100% by weight and carbon equivalent CEV (NW) <0.62%,

Final thickness over 1, 00 mm up to and including 2,00 mm:

C content <0.105% by weight and the carbon equivalent CEV (NW) <0.64%, final thickness over 2.00 mm up to and including 3.00 mm:

C content <0.1 15% by weight and the carbon equivalent CEV (NW) <0.66%.

It has also proven to be advantageous that the Mn content is selected as a function of the final thickness of the cold strip to be obtained as follows:

Final thickness 0.50 up to and including 1.00 mm:

Mn content> 1, 900% by weight to <2.200% by weight,

Final thickness over 1.00 to 2.00 mm inclusive:

Mn content> 2.050% by weight to <2.250% by weight, Final thickness over 2.00 mm up to and including 3.00 mm:

Mn content> 2,100% by weight to <2,350% by weight.

With regard to the use of variable pre-strip thicknesses, it turned out to be particularly advantageous that the Cr content and the carbon equivalent CEV (NW) be chosen as a function of the final thickness of the cold strip to be obtained as follows:

Final thickness 0.50 up to and including 1.00 mm:

Cr content> 0.260% by weight to <0.330% by weight

and carbon equivalent CEV (NW) <0.62%,

Final thickness over 1.00 to 2.00 mm inclusive:

Cr content> 0.290% by weight to <0.360% by weight

and the carbon equivalent CEV (NW) <0.64%,

Final thickness over 2.00 up to and including 3.00 mm:

Cr content> 0.320% by weight to <0.370% by weight

and the carbon equivalent CEV (NW) <0.66%.

This applies to the continuous annealing of successive belts with different belt cross-sections, as well as to belts with varying sheet thickness over belt length or belt width. For example, a processing of cold strips with variable Kaltabwalzgraden is possible.

Are inventively produced by continuous annealing highest strength

Cold strips produced from multiphase steel with varying sheet thicknesses, can be made of this material advantageous stress-optimized components forming technology.

The material produced can be produced as a cold strip via a hot-dip galvanizing line or a pure continuous annealing plant in the dressed and undressed and also in the heat-treated state (overaging) and in the stretched and unstretched state (stretch bending strains).

At the same time, it is possible to adjust the microstructural fractions by selective variation of the process parameters so that steels in different strength classes, for example with yield strengths between 550 MPa and 950 MPa, as well as Tensile strengths between 980 MPa and 1 140 MPa are displayed.

With the alloy composition according to the invention, steel strips can be produced by an intercritical annealing between Ac1 and Ac3 or in an austenitizing annealing over Ac3 with final controlled cooling, which leads to a dual or multi-phase structure.

Annealing temperatures of about 700 to 950 ° C have proved to be advantageous. Depending on the overall process (only continuous annealing or with additional

Hot dipping) there are according to the invention different approaches for a heat treatment.

In a continuous annealing plant without subsequent hot-dip finishing, the cold rolled steel strip is cooled to an intermediate temperature of about 160 to 250 ° C starting from the annealing temperature at a cooling rate of about 15 to 100 ° C / s. Optionally, it is possible to cool down in advance to a previous intermediate temperature of 300 to 500 ° C. at a cooling rate of approximately 15 to 100 ° C./s. The cooling to room temperature is finally carried out at a cooling rate of about 2 to 30 ° C / s (see method 1, Figure 8a).

Alternatively, it may be cooled to room temperature at a cooling rate between about 15 and 100 ° C / s from the intermediate temperature of 300 to 500 ° C.

In a heat treatment in the context of a hot dip refinement, there are two ways of temperature control. The cooling is stopped as described above before entering the molten bath and continued until after leaving the bath until reaching the intermediate temperature of about 200 to 250 ° C.

Depending on the molten bath temperature, this results in a holding temperature in the molten bath of about 400 to 470 ° C. The cooling to room temperature is again at a cooling rate of about 2 to 30 ° C / s (see method 2, Figure 8b).

The second variant of the temperature control in the hot dip finishing includes holding the temperature for about 1 to 20 seconds at the intermediate temperature of about 200 to 350 ° C and then reheating to the

Hot dip refinement required temperature of approx. 400 to 470 ° C. The ribbon is cooled to approx. 200 to 250 ° C after refining. The cooling to room temperature takes place again at a cooling rate of about 2 to 30 ° C / s (see method 3, Figure 8c). In the known dual-phase steels, besides carbon, manganese, chromium and silicon are also responsible for the transformation of austenite to martensite. Only the inventive combination of alloyed within the specified limits elements carbon, silicon, manganese, nitrogen, molybdenum and chromium and niobium, titanium and boron in tight areas on the one hand ensures the required mechanical properties such as minimum tensile strength of 980 MPa at the same time significantly widened process window in the continuous annealing ,

Material characteristic is also that the addition of manganese with increasing weight percent of the ferrite is shifted to longer times and lower temperatures during cooling, similar effect also the elements carbon, chromium, molybdenum and boron. The proportions of ferrite are thereby increased levels of Bainite more or less reduced depending on the process parameters. By setting a low carbon content of <0.1-15% by weight, the carbon equivalent can be reduced, thereby improving weldability and avoiding excessive weld hardening. In resistance spot welding, moreover, the electrode life can be significantly increased.

The effect of the elements in the alloy according to the invention is described in more detail below. Accompanying elements are unavoidable and are considered in the analysis concept with regard to their effect, if necessary. Accompanying elements are elements that are already present in iron ore, or

due to production, get into the steel. Because of their predominantly negative influences, they are usually undesirable. An attempt is made to remove them to a tolerable level or to convert them into more harmless forms. Hydrogen (H) can be the only element without creating lattice strains diffuse through the iron grid. This causes the hydrogen in the

Iron grating is relatively mobile and can be relatively easily absorbed during the processing of the steel. Hydrogen can only be taken up in atomic (ionic) form in the iron lattice.

Hydrogen has a strong embrittlement and preferably diffuses to energy-favorable sites (defects, grain boundaries, etc.). In this case, defects act as hydrogen traps and can significantly increase the residence time of the hydrogen in the material.

By recombination to molecular hydrogen, cold cracks can arise. This behavior occurs in hydrogen embrittlement or in hydrogen-induced stress corrosion cracking. Even with the delayed crack, the so-called

Delayed fracture, which occurs without external stress, is often called hydrogen as the causative reason. Therefore, the hydrogen content in the steel should be as low as possible.

For the aforementioned reasons, the hydrogen content in the steel according to the invention is limited to <0.0010% by weight (10 ppm) or advantageously to <0.0008% by weight, optimally to <0.0005% by weight.

A more uniform structure, the u.a. achieved by its widened process window, also reduces the susceptibility to hydrogen embrittlement.

Oxygen (O): In the molten state, the steel has a relatively high absorption capacity for gases. At room temperature, however, oxygen is only soluble in very small quantities. Similar to hydrogen, oxygen can only diffuse into the material in atomic form. Due to the strong embrittling effect and the negative effects on the aging resistance, as much as possible is attempted during production to reduce the oxygen content.

For the reduction of oxygen exist on the one hand procedural approaches such as a vacuum treatment and on the other analytical approaches. By adding certain alloying elements, the oxygen can be more harmless States are transferred. Thus, a binding of the oxygen in the course of a deoxidation of the steel with manganese, silicon and / or aluminum is usually common. However, the resulting oxides can cause negative properties as defects in the material.

For the above reasons, therefore, the oxygen content in the steel should be as low as possible.

Phosphorus (P) is a trace element from iron ore and is dissolved in the iron lattice as a substitution atom. Phosphorus increases hardness by solid solution strengthening and improves hardenability. However, it is generally attempted to lower the phosphorus content as much as possible, since it is highly prone to segregation, among other things due to its low solubility in the solidifying medium, and greatly reduces the toughness. Due to the addition of phosphorus at the grain boundaries, grain boundary fractures occur. In addition, phosphorus increases the transition temperature from tough to brittle behavior up to 300 ° C. During hot rolling, near-surface phosphorus oxides at the grain boundaries can lead to breakage cracks. However, in some steels, phosphorus is used as a micro-alloying element in small quantities (<0.1% by weight) due to its low cost and high strength enhancement, for example, in higher-strength IF (interstitial free) steels, bake hardening steels or some alloy concepts for dual-phase steels. The steel according to the invention differs from known analysis concepts, which use phosphorus as a mixed-crystal former, inter alia by the fact that phosphorus is not alloyed but is adjusted as low as possible.

For the aforementioned reasons, the phosphorus content in the steel according to the invention is limited to unavoidable amounts in steelmaking. Preferably, P should be <0.020 wt%.

Like phosphorus, sulfur (S) is bound as a trace element in iron ore. Sulfur is undesirable in steel (except free-cutting steels), as it tends to segregate and has a strong embrittlement. It is therefore attempted to minimize the content of sulfur in the melt, for example by a vacuum treatment, to reach. Furthermore, the existing sulfur is converted by adding manganese into the relatively harmless compound manganese sulfide (MnS). The manganese sulfides are often rolled in rows during the rolling process and act as nucleation sites for the transformation. This leads, especially in the case of diffusion-controlled transformation, to a line-shaped structure and can lead to impaired mechanical properties in the case of pronounced bristleness, for example to pronounced Martensitzeilen instead of distributed Martensitinseln, anisotropic

Material behavior, reduced elongation at break. For the aforementioned reasons, the sulfur content of the steel according to the invention is limited to <0.0020% by weight or advantageously to <0.0015% by weight, optimally to <0.0010% by weight.

Alloying elements are usually added to the steel in order to specifically influence certain properties. An alloying element in different steels can influence different properties. The effect generally depends strongly on the amount and the solution state in the material. The connections can therefore be quite varied and complex. In the following, the effect of the alloying elements will be discussed in greater detail.

Carbon (C) is considered the most important alloying element in steel. Through its targeted introduction of up to 2.06 wt .-% iron is only for steel. Often the carbon content is drastically lowered during steelmaking. at

Dual-phase steels for a continuous hot-dip finishing is its proportion according to EN 10346 or VDA 239-100 maximum 0.230 wt .-%, a minimum value is not specified. Due to its comparatively small atomic radius, carbon is interstitially dissolved in the iron lattice. The solubility is 0.02% maximum in α-iron and 2.06% maximum in iron. Carbon in solute significantly increases the hardenability of steel and is therefore essential for the formation of a sufficient amount of martensite. However, excessive carbon contents increase the hardness difference between ferrite and martensite and limit weldability. In order to meet the requirements for, for example, high hole widening and bending angles as well as improved weldability, the steel according to the invention contains

Carbon contents of <0.1 15 wt .-%.

Due to the different solubility of the carbon in the phases pronounced diffusion processes in the phase transformation are necessary, which can lead to very different kinetic conditions. In addition, carbon increases the thermodynamic stability of austenite, which is shown in the phase diagram in an extension of the austenite area to lower temperatures. As the constrained carbon content in martensite increases, the lattice distortions and, associated therewith, the strength of the diffusion-free phase are increased.

Carbon also forms carbides. An almost in every steel occurring

Structural phase is the cementite (FesC). However, significantly harder special carbides with other metals such as chromium, titanium, niobium but also vanadium can form. Not only the species but also the distribution and size of the precipitates is of crucial importance for the resulting increase in strength. On the one hand, to ensure sufficient strength and, on the other hand, good weldability, improved hole widening, improved bending angle and sufficient resistance to hydrogen-induced cracking (delayed fracture free), the minimum C content is set at 0.075% by weight and the maximum C- Content determined to 0.1 15 wt .-%, advantageous are contents with a cross-sectional differentiation, such as:

Final thickness 0.50 mm up to and including 1.00 mm (C <0.100% by weight)

Final thickness over 1.00 to 2.00 mm inclusive (C <0.105 wt.%)

Final thickness over 2.00 to 3.00 mm inclusive (C <0.1 15 wt .-%). Furthermore, it is advantageous to comply with a band-thickness-dependent differentiation of the carbon content in combination with the carbon equivalent CEV (IIW).

Final thickness 0.50 to and including 1.00 mm (C <0.100 wt.%) With a carbon equivalent CEV (IIW) of <0.62, Final thickness above 1.00 to 2.00 mm inclusive (C <0.105 wt .-%) at a

Carbon equivalent CEV (IIW) <0.64%,

Final thickness over 2.00 to 3.00 mm inclusive (C <0.1 15 wt .-%) at a

Carbon equivalent CEV (IIW) of <0.66%.

Silicon (Si) binds oxygen during casting and is therefore used to calm down during the deoxidation of the steel. Important for the later steel properties is that the Seigerungskoeffizient is significantly lower than, for example, that of manganese (0.16 compared to 0.87). Seingings generally result in a line arrangement of the structural components that degrade the forming properties, such as hole widening and bending capability.

In terms of material characteristics, the addition of silicon causes strong solid solution hardening. Approximately, addition of 0.1% silicon causes an increase in tensile strength of about 10 MPa, with elongation only slightly deteriorating when added up to 2.2% silicon. This was investigated for different sheet thicknesses and annealing temperatures. The increase from 0.2 to 0.5% silicon caused an increase in strength of about 20 MPa in the yield strength and about 70 MPa in the tensile strength. The elongation at break decreases by about 2%. The latter is partly due 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 reduce the ductility as brittle phases, which in turn improves the formability. Due to the low strength-increasing effect of silicon within the range of the steel according to the invention, the basis for a broad process window is created.

Another important effect is that silicon shifts the formation of ferrite to shorter times and temperatures, thus allowing the formation of sufficient ferrite before quenching. Hot rolling thereby provides a basis for improved cold rollability. In hot dipping, the accelerated ferrite formation enriches the austenite with carbon and stabilizes it. Since silicon hinders carbide formation, the austenite is additionally stabilized. Thus, the accelerated cooling can suppress the formation of bainite in favor of martensite. The addition of silicon in the range according to the invention has led to further surprising effects described below. For example, the carbide formation delay described above could be due to aluminum

be brought about. However, aluminum forms stable nitrides, so that there is insufficient nitrogen available for the formation of carbonitrides with micro-alloying elements. By alloying with silicon, this problem does not exist because silicon forms neither carbides nor nitrides. Thus, silicon has an indirect positive effect on precipitation formation by microalloys, which in turn has a positive effect on the strength of the material. Since the increase in the transformation temperatures by silicon tends to favor grain coarsening, micro-alloying with niobium, titanium and boron is particularly expedient, as is the targeted adjustment of the nitrogen content in the steel according to the invention.

When hot rolling, it is known to occur in higher silicon-alloyed steels to form strongly adhering red scale and increased risk of Zundereinwalzungen what influence on the subsequent Beizings and the

Can have pickling productivity. This effect could not be detected in the steel according to the invention with 0.400 to 0.500% silicon, if the pickling is advantageously carried out with hydrochloric acid instead of sulfuric acid.

Regarding the galvanizability of silicon-containing steels, i.a. in DE 196 10 675 C1 stated that steels with up to 0.800 wt .-% silicon or up to 2,000 wt .-% silicon are not hot dip galvanized due to the very poor wettability of the steel surface with the liquid zinc.

In addition to the recrystallization of the hard cold-rolled strip, the atmospheric conditions during the annealing treatment in a continuous hot-dip coating plant cause a reduction of iron oxide, which can form, for example, during cold rolling or as a result of storage at room temperature on the surface. For oxygen-affinity alloy components, such as

Silicon, manganese, chromium, boron, the gas atmosphere is oxidizing with the result that segregation and selective oxidation of these elements can occur. The selective oxidation can take place both externally, that is on the substrate surface, and internally within the metallic matrix. It is known in particular that silicon diffuses during the annealing to the surface and forms oxides on the steel surface alone or together with manganese. These oxides can prevent contact between the substrate and the melt and prevent or worsen the wetting reaction. As a result, undiluted spots, so-called "bare spots", or even large areas without coating may occur

Wetting reaction with the result of insufficient inhibiting layer formation, the adhesion of the zinc or zinc alloy layer on the steel substrate are reduced.

Contrary to this general knowledge was surprisingly found in experiments that can be achieved only by a suitable Ofenfahrweise during recrystallization and when passing through the hot dip a good Schmelztauchveredelung the steel strip and a good adhesion of the coating.

For this purpose, it must first be ensured that the strip surface is free from scale residues, pickling or rolling oil or other dirt particles by a chemical-mechanical or thermal-hydro-mechanical pre-cleaning. In order to prevent silicon oxides from reaching the strip surface, further methods are to be taken which promote the internal oxidation of the alloying elements below the surface of the material. Depending on the system configuration, different measures are used here. In a plant configuration in which the annealing process step is performed exclusively in a radiant tube furnace (RTF) (see process 3 in FIG. 8c), the internal oxidation of the alloying elements can be targeted by adjusting the oxygen partial pressure of the furnace atmosphere (N2-H2 shielding gas atmosphere) to be influenced. The set oxygen partial pressure must satisfy the following equation, with the furnace temperature between 700 and 950 ° C.

-12> Log p0 ₂ > -5 ^* Sr ⁰ ^'25 - 3 ^* Mn- °' ⁵ -0, rCr ° ' ⁵ -7 ^* (-ln Bf ⁵

Here, Si, Mn, Cr, B denote the corresponding alloying proportions in the steel in wt .-% and p02 the oxygen partial pressure in mbar. In a plant configuration in which the furnace area consists of a combination of a direct fired furnace (DFF) and a subsequent radiant tube furnace (see process 2 in Figure 8b), selective oxidation can be used also influence the alloying elements via the gas atmospheres of the furnace areas.

The combustion reaction in the NOF can be used to adjust the oxygen partial pressure and thus the oxidation potential for iron and the alloying elements. This should be adjusted so that the oxidation of the alloying elements takes place internally below the steel surface and, if necessary, a thin iron oxide layer is formed on the steel surface after passing through the NOF region. This is achieved, for example, by reducing the CO value below 4% by volume. In the subsequent radiant tube furnace, the optionally formed iron oxide layer is reduced under Isb-F protective gas atmosphere and likewise the alloying elements are further internally oxidized. The set oxygen partial pressure in this furnace area must satisfy the following equation, with the furnace temperature between 700 and 950 ° C.

-18> log p0 ₂ > -5 ^* Si ^"3 - 2,2 ^* Mn- ° ^'45 -0, rCr ⁴ -12,5 ^* (-ln Bf ²⁵

Here, Si, Mn, Cr, B denote the corresponding alloying proportions in the steel in wt .-% and p0 ₂ the oxygen partial pressure in mbar.

In the transition zone between furnace -> Zinkpott (proboscis) is the dew point of the gas atmosphere (Isb-F-Schutzgasatmosphäre) and thus adjust the oxygen partial pressure so that oxidation of the tape before immersion in the

Melt bath is avoided. Dew points in the range from -30 to -40 ° C have proven to be advantageous.

The measures described above in the furnace area of the continuous hot-dip coating installation prevent the superficial formation of oxides and achieve a uniform, good wettability of the strip surface with the liquid melt. If instead of the hot dipping refinement (here eg the hot dip galvanizing) the

Process route via a continuous annealing with subsequent electrolytic galvanizing selected (see method 1 in Figure 8a), are not special

Precautions necessary to ensure galvanizability. It is known that the galvanizing of higher-alloyed steels is much easier to realize by electrolytic deposition than by continuous hot-dip processes. In electrolytic galvanizing, pure zinc is deposited directly on the strip surface. In order not to hinder the electron flow between the steel strip and the zinc ions and thus the zinc plating, it must be ensured that no surface-covering oxide layer is present on the strip surface. This condition is usually ensured by a standard reducing atmosphere during annealing and pre-cleaning prior to electrolysis. In order to have the widest possible process window in the annealing and a sufficient

To ensure galvanizability, the minimum silicon content is set at 0.400% by weight and the maximum silicon content at 0.500% by weight.

Manganese (Mn) is added to almost all steels for desulfurization to convert the harmful sulfur into manganese sulphides. In addition, manganese increases the strength of the ferrite by solid solution strengthening and shifts the α / γ conversion to lower temperatures.

One of the main reasons for adding manganese to multiphase steels, such as dual-phase steels, is the significant improvement in hardenability. Due to the diffusion hindrance, the pearlite and bainite transformation is shifted to longer times and the martensite start temperature is lowered.

At the same time, however, the addition of manganese increases the hardness ratio between martensite and ferrite. In addition, the line of the structure is reinforced. A high hardness difference between the phases and the formation of Martensitzeilen result in a lower Lochaufweitvermögen, which is equivalent to an increased edge crack sensitivity. Like silicon, manganese tends to form oxides on the steel surface during the annealing treatment. Depending on the annealing parameters and the contents of other alloying elements (especially silicon and aluminum), manganese oxides (for example MnO) and / or Mn mixed oxides (for example Mn2SiC> 4) can occur. However, with a low Si / Mn or Al / Mn ratio, manganese is considered to be less critical, since giobuare oxides rather than oxide films are formed. However, high levels of manganese can negatively affect the appearance of the zinc layer and zinc adhesion. By the above measures for adjusting the furnace areas in the continuous hot dip coating, the formation of Mn oxides or Mn mixed oxides is reduced on the steel surface after annealing.

The manganese content is determined for the reasons mentioned to 1, 900 wt .-% to 2.350 wt .-%. In order to achieve the required minimum strengths, it is advantageous to adhere to a band-thickness-dependent differentiation of the manganese content.

At a final thickness of 0.50 to 1.00 mm inclusive, the manganese content is preferably in a range between> 1.900% by weight and <2.200% by weight, with final thicknesses of 1.00 to 2.00 inclusive mm between> 2.050 wt .-% to <2.250 wt .-% and at final thicknesses of 2.00 to 3.00 mm inclusive between> 2.100 wt .-% to <2.350 wt .-%.

Another peculiarity of the invention is that the variation of the manganese content can be compensated by simultaneously changing the silicon content. The increase in strength (here the yield strength, YS) by manganese and silicon is generally well described 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 on the effect of solid solution hardening, which is weaker for manganese according to this equation than for silicon. At the same time, however, manganese, as mentioned above, significantly increases the hardenability, which significantly increases the proportion of strength-increasing second phase in multiphase steels. Therefore, the addition of 0.1% silicon in a first approximation is equivalent to the addition of 0.1% manganese in terms of strength enhancement. For a steel of the composition according to the invention and an annealing which includes the time-temperature parameters according to the invention, the following relationship for yield strength and tensile strength (TS) has been established empirically:

YS (MPa) = 185.7 + 147.9 [wt% Si] + 161, 1 [wt% Mn]

TS (MPa) = 574.8 + 189.4 [wt% Si] + 174.1 [wt% Mn]

Compared to the Pickering equation, the coefficients of manganese and silicon are approximately the same for both the yield strength and the tensile strength, which gives the possibility of substitution of manganese by silicon. Chromium (Cr) on the one hand in dissolved form in small quantities

Significantly increase the hardenability of steel. On the other hand, chromium causes particle hardening with appropriate temperature control in the form of chromium carbides. The associated increase in the number of seed sites with simultaneously reduced content of carbon leads to a reduction in the hardenability.

In dual-phase steels, the addition of chromium mainly results in the

Hardenability improved. In the dissolved state, chromium shifts the pearlite and bainite transformation to longer times and at the same time lowers the martensite start temperature.

Another important effect is that chromium increases the tempering resistance significantly, so that there is almost no loss of strength in the hot dip.

Chromium is also a carbide former. If chromium-iron mixed carbides are present, the austenitizing temperature must be high enough before curing to dissolve the chromium carbides. Otherwise, the increased germ count may lead to a deterioration of the hardenability.

Chromium also tends to form oxides on the steel surface during the annealing treatment, which may degrade the hot dipping quality. By the abovementioned measures for adjusting the furnace areas in continuous hot-dip coating reduce the formation of Cr oxides or Cr mixed oxides on the steel surface after annealing. The chromium content is therefore set at levels of 0.250 wt .-% to 0.400 wt .-%.

In order to achieve the required minimum strengths, it is advantageous to comply with a band-thickness-dependent differentiation of the chromium content, especially for processing with variable pre-strip thicknesses.

At a final 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 .-%, with final thicknesses of 1.00 to 2.00 mm inclusive between > 0.290 wt .-% to <0.360 wt .-% and at final thicknesses of 2.00 to 3.00 mm inclusive between> 0.320 wt .-% to <0.370 wt .-%.

In order to achieve the required minimum strengths, it is advantageous to differentiate the thickness of the banddischarge in combination with the

Carbon equivalent CEV (IIW), also here especially for the processing with variable pre-strip thicknesses.

With a final 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 .-% with a carbon equivalent CEV (IIW) of <0.62%, at final thicknesses of 1.00 to 2.00 mm including between 0.290 wt .-% to <0.360 wt .-% at a

Carbon equivalent CEV (IIW) of <0.66% and at final thicknesses of 2.00 to 3.00 mm including between> 0.320 wt .-% to <0.370 wt .-% with a carbon equivalent CEV (IIW) of <0.66 %.

Chromium contents between> 0.250% by weight and <0.370% by weight can be used with final thicknesses of small 0.50 mm and between> 0.370% by weight and <0.400% by weight greater than 3.00 mm. Molybdenum (Mo): The addition of molybdenum is similar to that of chromium and Manganese to improve hardenability. The pearlite and bainite transformation is postponed to longer times and the martensite start temperature is lowered. At the same time, molybdenum is a strong carbide former that gives rise to finely divided mixed carbides, including titanium. Molybdenum also increases the tempering resistance significantly, so that in the hot dip no strength losses are expected. Molybdenum also works by solid solution hardening, but is less effective than manganese and silicon.

The content of molybdenum is therefore set between more than 0.200 wt .-% to 0.300 wt .-%. For cost reasons, the Mo content is advantageously adjusted to a range between more than 0.200 wt .-% to 0.250 wt .-%.

As a compromise between the required mechanical properties and

Hot-dip dipability has proved to be advantageous for the alloy concept according to the invention a sum content of Mo + Cr of <0.650% by weight.

To achieve the required mechanical properties, especially the minimum tensile strength, it is advantageous to comply with the total content of manganese, silicon and chromium, a molecular formula Mn-Si + Cr, wherein this between> 1, 750 wt .-% to <2.250 wt .-%, especially for processing with variable pre-strip thicknesses.

The determination of the total content of manganese, silicon, chromium and molybdenum via a molecular formula Mn-Si + Cr + Mo has proved to be advantageous for achieving the required mechanical characteristics, in particular the minimum tensile strength, whereby between> 1.950% by weight. % to <2.500 wt .-% is limited, especially for the processing of strips with variable pre-strip thicknesses.

Copper (Cu): The addition of copper can increase the tensile strength and hardenability. In conjunction with nickel, chromium and phosphorus, copper can be a

form protective oxide layer on the surface, which can significantly reduce the corrosion rate.

When combined with oxygen, copper can form harmful oxides at the grain boundaries, which can be detrimental to hot working processes can. The content of copper is therefore fixed at <0.050% by weight and thus limited to quantities that are unavoidable in steel production.

Nickel (Ni): In combination with oxygen, nickel can form harmful oxides at the grain boundaries, which can cause negative effects, especially for hot forming processes. The content of nickel is therefore fixed at <0.050% by weight and thus limited to quantities that are unavoidable in steel production.

Vanadium (V): In the present alloy concept, the content of

Vanadium from> 0.005 wt .-% to <0.020 wt .-% set, optimally limited to> 0.005 wt .-% to <0.015 wt .-% limited.

Tin (Sn): Since addition of tin is not necessary with the present alloy concept, the content of tin is determined to <0.040% by weight, thus limiting unavoidable steel-accompanying amounts.

Aluminum (AI) is usually added to the steel to bind the dissolved oxygen in the iron and nitrogen. Oxygen and nitrogen become so in

Converted aluminum oxides and aluminum nitrides. These precipitations can cause a grain refining by increasing the germination sites and thus increase the toughness properties and strength values.

Aluminum nitride is not precipitated when titanium is present in sufficient quantities. Titanium nitrides have a lower formation enthalpy and are formed at higher temperatures.

When dissolved, aluminum such as silicon shifts ferrite formation to shorter times, allowing the formation of sufficient ferrite in dual phase steel. It also suppresses carbide formation, leading to a delayed transformation of austenite. For this reason, aluminum is also used as an alloying element in residual austenitic steels (TRIP steels) to substitute a part of the silicon. The reason for this approach is that aluminum is slightly less critical to the galvanizing reaction than silicon. The aluminum content is therefore based on 0.010 wt .-% to a maximum of 0.060 wt .-% limited and is added to calm the steel.

Niobium (Nb): Niobium has different effects in steel. During hot rolling in the finishing train, it retards recrystallization by forming finely divided precipitates, increasing the nucleation density and producing a finer grain after conversion. The proportion of dissolved niobium also works

rekristallisationshemmend. The excretions increase the strength of the final product. These can be carbides or carbonitrides. Often these are mixed carbides in which titanium is also incorporated. This effect begins at 0.005 wt .-% and is most evident from 0.010 wt .-% niobium. The precipitates also prevent grain growth during (partial) austenitization in the hot dip galvanizing. Above 0.060 wt.% Niobium, no additional effect is to be expected. Levels of 0.025 wt .-% to 0.045 wt .-% have been found to be advantageous.

Titanium (Ti): Due to its high affinity to nitrogen, titanium is primarily precipitated as TiN during solidification. In addition, it joins with niobium as

Mixed carbide on. TiN is of great importance for grain size stability in the blast furnace. The precipitates have a high temperature stability, so that they exist in contrast to the mixed carbides at 1200 ° C largely as particles that impede grain growth. Titanium also retards recrystallization during hot rolling, but is less effective than niobium. Titanium works by precipitation hardening. The larger TiN particles are less effective than the finely divided mixed carbides. The best effectiveness 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, contents of 0.025 wt .-% to 0.045 wt .-% have been found to be advantageous.

Boron (B): Boron is an extremely effective alloying agent for realizing variable degrees of cold rolling. In experiments, it has surprisingly been found that the very narrow range for the addition of boron according to the invention has a pronounced effect on the uniformity of the mechanical properties of the cold-rolled strips with variable degree of cold rolling produced in the subsequent processing. This pronounced effect first leads to the possibility of defining characteristic ranges according to the process variables instead of a relatively constant degree of cold rolling. Steps (Figures 8a, 8b and 8c) also for the material with variable Kaltabwalzgraden based on a Masterwarmbanddicke or based on a Masterkaltbanddicke adjust. In addition, boron is an effective hardening enhancer that is effective in very small quantities. The martensite start temperature remains unaffected. To be effective, boron must be in solid solution. Since it has a high affinity for nitrogen, the nitrogen must first be set, preferably by the stoichiometrically necessary amount of titanium. Due to its low solubility in iron, the dissolved boron prefers to attach to the

Austenite grain boundaries. There it partially forms Fe-B carbides, which are coherent and reduce the grain boundary energy. Both effects have a retarding effect on ferrite and pearlite formation and thus increase the hardenability of the steel. Excessive levels of boron, however, are detrimental as iron boride can form, adversely affecting the hardenability, formability and toughness of the material. Boron also tends to form oxides or mixed oxides during annealing during the continuous hot-dip coating, which deteriorate the quality of galvanizing. The above measures for adjusting the furnace areas in continuous hot dip coating reduce the formation of oxides on the steel surface.

For the aforementioned reasons, the boron content for the inventive

Alloy concept set 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% by weight

Nitrogen (N) can be both an alloying element and a companion element from steelmaking. Excessive levels of nitrogen cause an increase in strength associated with rapid loss of toughness and aging effects.

On the other hand, a fine grain hardening via titanium nitrides and niobium (karbo) nitrides can be achieved by a targeted addition of nitrogen in conjunction with the micro-alloying elements titanium and niobium. In addition, coarse grain formation upon re-heating before hot rolling is suppressed. According to the invention, the N content is therefore to values of> 0.0020 wt .-% to < 0.0120 wt .-% set.

Advantageously, the determination of a sum in the contents of hydrogen and nitrogen has been shown, with an optimum for H + N between> 0.0025 wt .-% to <0.0130 wt .-% is.

It has proved to be advantageous for compliance with the required properties of the steel when the nitrogen is added as a function of the sum of Ti + Nb + B.

With a total content of Ti + Nb + B of> 0.010% by weight to <0.080% by weight, the content of nitrogen should be kept to values of> 0.0020% by weight to <0.0090% by weight. For a sum content of Ti + Nb + B of> 0.050 wt .-% nitrogen contents of> 0.0040 wt .-% to <0.0120 wt .-% have proven to be advantageous.

For the sum amounts of niobium and titanium, contents of <0.100% by weight have been found to be advantageous and, owing to the principal interchangeability of niobium and titanium, to a minimum niobium content of 0.005% by weight and, for cost reasons, particularly advantageously <0.090% by weight. % proved.

In the interaction of the micro-alloying elements niobium and titanium with boron, sum amounts of <0.102% by weight have proved to be advantageous and particularly advantageous <0.092% by weight. Higher contents no longer have an improving effect according to the invention.

In addition, maximum contents of <0.365% by weight of Ti + Nb + V + Mo + B have proved to be the maximum amounts for the abovementioned reasons. Calcium (Ca): An addition of calcium in the form of calcium-silicon mixed compounds causes deoxidation during steelmaking and

Desulphurization of the molten phase. Thus, reaction products are transferred to the slag and the steel is cleaned. The increased purity leads according to the invention to better properties in the final product. For these reasons, a Ca content of> 0.0010 wt .-% to <0.0060 wt .-% is set. Levels of <0.0030 wt .-% have been found to be advantageous. In tests with variable pre-strip thicknesses carried out with the steel according to the invention, it was found that with an intercritical annealing between Ac1 and Ac3 or an austenitizing annealing over Ac3 with finally controlled cooling, a dual-phase steel with a minimum tensile strength of 980 MPa in a thickness of 1.50 mm, starting from a

Masterswarmband of 2.30 mm, was produced, but also in the thickness range 0.50 to 3.00 mm can be generated, which is characterized by a sufficient tolerance to process variations.

This is a significantly expanded process window for the alloy composition according to the invention in comparison to known alloy concepts.

The annealing temperatures for the dual-phase structure to be achieved are between about 700 and 950 ° C. for the steel according to the invention, so that depending on

Temperature range reaches a partially austenitic (two-phase area) or a fully austenitic structure (austenitic area).

The experiments also showed that the set microstructural fractions after the intercritical annealing between Ac1 and Ac3 and austenitizing annealing over Ac3 with subsequent controlled cooling even after another process step of the Schmelztauchveredelung at temperatures between 400 to 470 ° C, for example with zinc or zinc magnesium remain.

The continuous annealed and occasionally hot-dip refined material can be in the dressed (cold rolled) or undressed state and / or in the stretch bent or non-stretch bent state and also in the

heat treated condition (aging) are made.

The steel strips of the alloy composition according to the invention are also characterized in the further processing by a high Kantenrissunempfindlich- speed and a high bending angle. Advantageously, steel strips can thus be produced which have a minimum product value Rm x α (tensile strength x [bending angle according to VDA 238-100]) of 100,000 MPa x °, in particular of 120000 MPa x °.

In addition, the steel strips according to the invention have a delayed fracture-free state for at least 6 months, while meeting the requirements of SEP 1970 for hole pull and iron-on samples made available by the steel manufacturer. The very small differences in the characteristic of the steel strip along and across its rolling direction are advantageous for later material use. Thus, the cutting of blanks from a strip regardless of the rolling direction (for example, transversely, longitudinally and diagonally or at an angle to the rolling direction) take place and so the waste can be minimized.

To the cold rollability of a steel produced from the steel according to the invention

To ensure hot strip, the hot strip according to the invention with

Final rolling temperatures in the austenitic range above Ar3 and reel temperatures above the bainite start temperature generated.

In the course of further processing of the steel strip according to the invention, it is thus possible to produce a hardened component, for example for the automotive industry.

Here, a board is cut from a steel strip according to the invention, which is then heated to a temperature above Ac3. The heated board is formed into a component and then cured in a forming tool or in air, with optional subsequent tempering.

Advantageously, the steel according to the invention has the property that the hardening takes place already on cooling at still air, so that a separate cooling of the forming tool can be omitted.

During hardening, the structure of the steel is converted by heating into the austenitic region, preferably to temperatures above 950 ° C. under a protective gas atmosphere. During subsequent cooling in air or inert gas takes place Formation of a martensitic microstructure for a high-strength component.

A subsequent tempering makes it possible to reduce residual stresses in the hardened component. At the same time, the hardness of the component is reduced so that the required toughness values are achieved.

Further features, advantages and details of the invention will become apparent from the following description of exemplary embodiments illustrated in a drawing.

Show it:

FIG. 1 process chain (schematic) for the production of a strip from the steel according to the invention,

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

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

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

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

FIG. 5a hardening behavior during cold rolling of the steel according to the invention, characteristic values transverse to the rolling direction, FIG.

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

FIG. 6 a shows mechanical characteristics (transverse to the rolling direction) of the steel according to the invention, in the state of sheet metal (HDG),

FIG. 6b Results of the hole expansion tests according to ISO 16630 and of the

Platelet bending test according to VDA 238-100 on the steel according to the invention, in the thin sheet state (HDG),

Figure 7a mechanical characteristics (transverse to the rolling direction) of the steel according to the invention

in the state of HR, CR and HDG; Example 1 (pre-strip thickness 40 mm), Figure 7b mechanical characteristics (transverse to the rolling direction) of the steel according to the invention

in the state of HR, CR and HDG; Example 2 (pre-strip thickness 45 mm),

Figure 7c mechanical characteristics (transverse to the rolling direction) of the steel according to the invention

in the state of HR, CR and HDG; Example 3 (pre-strip thickness 50 mm),

Figure 7d mechanical characteristics (transverse to the rolling direction) of the invention

steel

in the state of HR, CR and HDG; Example 4 (pre-strip thickness 55 mm),

Figure 7e mechanical characteristics (transverse to the rolling direction) of the steel according to the invention

in the status HR, CR and HDG as overview,

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

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

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

Figure 1 shows schematically the process chain for the production of a strip of the steel according to the invention. Shown are the possible process routes in the invention. Until pickling, the process route is the same for all steels according to the invention, after which deviating process routes take place, depending on the desired results. For example, after pickling, the pickled hot strip may be cold rolled and hot dip refined with varying degrees of rolling. Also soft annealed hot strip or annealed cold strip can be cold rolled and

be melt-dip refined.

Material can also be optionally processed without hot dip finishing, ie only in the context of continuous annealing with and without subsequent electrolytic galvanizing. From the optionally coated material, a complex component can now be produced. Following this can optionally be

Annealing process, such as the air hardening, where the heat-treated component is cooled in air. Optionally, a tempering stage can complete the thermal treatment of the component.

FIG. 2 shows schematically the time-temperature profile of the process steps

Hot rolling and continuous annealing of strips of the invention Alloy composition. Shown are the time- and temperature-dependent transformation for the hot rolling process as well as for post-cold-rolled heat treatment, component fabrication and optional tempering with optional tempering.

FIG. 3 shows in examples 1 to 4, which originate from a melt, in order to exclude the analytical influence, the alloy compositions of the steel according to the invention, depending on the produced pre-strip thickness. From a hot strip nominal thickness of 2.30 mm, cold strips with a cold strip nominal thickness of 1.50 mm were produced. Depending on the pre-strip thickness to be produced before hot rolling, Example 1 shows the alloy composition for a

Vorbanddicke 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 Vorband with a thickness of 55 mm.

FIG. 4 shows the mechanical characteristic values (transverse 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. 4b. FIG. 5 shows the solidification behavior, via the mechanical characteristics transverse to the rolling direction, during cold rolling of the steel according to the invention, in tabular form in FIG. 5a and graphically as cold flow curve in FIG. 5b.

FIG. 6 shows the mechanical characteristics (transverse to the rolling direction) of the steel according to the invention in the thin sheet state (HDG, Hot Dipped Galvanized) in FIG. 6a and the results of the hole expansion tests according to ISO 16630 and the platelet bending test according to VDA 238-100 in the thin sheet state (HDG). along and across the rolling direction, as well as the corresponding products with the tensile strength, in Figure 6b. FIG. 7 shows the mechanical characteristic values (transverse 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. 7a, 45 mm in FIG. 7b, 50 mm in FIG. 7c, 55 mm in FIG. 7d and FIG in Figure 7e as a summary graphical overview. FIG. 8 schematically shows three variants of the temperature Time courses during the annealing treatment and cooling and in each case different austenitizing conditions.

As a result of the different temperature guides according to the invention within the specified span, mutually different characteristic values or also different hole widening results and bending angles result. The basic differences are the temperature-time parameters during the heat treatment and the subsequent cooling. Process 1 (FIG. 8a) shows the annealing and cooling of the steel strip produced and cold-rolled to final thickness in a continuous annealing plant. First, the tape 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 about 15 and 100 ° C / s up to an intermediate temperature (ZT) of about 200 to 250 ° C. On the representation of a second intermediate temperature (about 300 to 500 ° C) is omitted in this schematic representation.

Subsequently, the steel strip is cooled at a cooling rate between about 2 and 30 ° C / s until reaching room temperature (RT) in air or the

Cooling at a cooling rate between about 15 and 100 ° C / s is maintained to room temperature.

The process 2 (Figure 8b) shows the process according to method 1, but the cooling of the steel strip for the purpose of a hot dip finishing briefly interrupted when passing through the hot dipping vessel, then the cooling at a cooling rate between about 15 and 100 ° C / s up to an intermediate temperature of about 200 to 250 ° C continue. Subsequently, the steel strip is cooled at a cooling rate between about 2 and 30 ° C / s until it reaches room temperature in air. Process 2 corresponds to annealing, for example hot dip galvanizing with a combined direct fired furnace and radiant tube furnace, as described in FIG. 8b.

The method 3 (Figure 8c) also shows the process according to method 1 in a hot dipping refinement, however, the cooling of the steel strip by a short Break (about 1 to 20 s) at an intermediate temperature in the range of about 200 to 400 ° C interrupted and heated up to the temperature (ST), which is necessary for hot dip refining (about 400 to 470 ° C), reheated. Subsequently, the steel strip is again cooled to an intermediate temperature of about 200 to 250 ° C. With a cooling rate of about 2 and 30 ° C / s takes place until reaching the

Room temperature in air the final cooling of the steel strip.

The method 3 corresponds for example to a process control in a continuous annealing plant, as described in Figure 8c. In addition, here by means of

Induction furnace, a reheating of the steel can be achieved optionally directly in front of the zinc bath.

The decreases from slab to sliver vary in the examples below from 78% to 84% for subsequent hot rolling to one

Hot strip thickness of 2.30 mm with corresponding decreases of 94% to 96%. In a single cold rolling step, with a Kaltabwalzgrad of 35%

Cold strip nominal thickness of 1.50 mm realized. It is impressively shown that both for very low pre-strip thicknesses and higher pre-strip thicknesses, and the area between them, relatively uniform values of the tensile strength and yield strength which are provided with the usual fluctuation range are obtained, transverse to the rolling direction. The steel according to the invention equally allows the use of master hot strip thickness with varying degrees of cold rolling, as well as the use of master cold strip thicknesses, without affecting the previous fact. For industrial production for hot dip galvanizing (HDG) according to method 3 of FIG. 8c, the following examples are exemplary within the framework 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 causing any problems the higher hot strip strength (HR) and higher cold strip strength (CR), with decreasing pre-strip thickness, would lead to significant fluctuations in the thin sheet (HDG):

example 1

(1.50 mm cold strip of 2.30 mm master warming tape and a pre-strip thickness of 40 mm) Alloy composition in wt .-%. An inventive steel with 0.104% C; 0.443% Si; 2.178% Mn; 0.012% P; 0.0004% S; 0.0045% N; 0.038 AI; 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 according to FIG. 8c, hot-dip refined, the slab material of 250 mm was hot rolled in the roughing mill to a pre-strip of 40 mm reversing rolled with a percentage decrease of 84% and then in the hot strip mill at a final target roll temperature of 910 Hot rolled at a reduction of 94% and unwound at a coiler temperature of 650 ° C with a master heat-treated thickness of 2.30 mm and after pickling without additional

Heat treatment (such as bell annealing) to 1, 50 mm in one

Continuously cold rolled (cold rolling degree 35%).

Sheet condition (HDG)

The yield ratio Re / Rm in the transverse direction was 66%.

Yield strength (Rp0.2) 706 MPa

Tensile strength (Rm) 1071 MPa

Elongation at break (A80) 10.9%

Bake hardening index (BH2) 492 MPa

- hole expansion ratio according to ISO 1663039%

Bending angle according to VDA 238-100 (longitudinal, transverse) 12171 12 °

The material characteristics transverse to the rolling direction would correspond for example to a HC660XD. Initial state (HR)

The yield ratio Re / Rm in the transverse direction was 77%.

Yield strength (Re) 826 MPa

Tensile strength (Rm) 1070 MPa

Elongation at break (A80) 10.0%

Intermediate state (CR) in the transverse direction

Yield strength (Re) 1246 MPa

Tensile strength (Rm) 1305 MPa

Elongation at break (A80) 2.0% Example 2

(1.50 mm cold strip of 2.30 mm master warming tape and a pre-strip thickness of 45 mm)

Alloy composition in% by weight

An inventive steel with 0.104% C; 0.443% Si; 2.178% Mn; 0.012% P; 0.0004% S; 0.0045% N; 0.038 AI; 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 process 3 according to FIG. 8c, hot-dip refined, the slab material of 250 mm was reversibly rolled before hot rolling in the roughing train to a preliminary strip of 45 mm with a percentage decrease of 82% and then in the hot strip mill at a final rolling target temperature from 910 ° C with a decrease of 95% hot rolled and at a reel target temperature of 650 ° C with a master hot strip thickness of 2.30 mm and cold after pickling without additional heat treatment (such as bell annealing) to 1.50 mm in one pass (Cold rolling degree 35%).

Sheet condition (HDG)

The yield ratio Re / Rm in the transverse direction was 67%.

- Elongation limit (Rp0.2) 720 MPa

Tensile strength (Rm) 1077 MPa

Elongation at break (A80) 10.4%

Bake hardening index (BH2) 51 MPa

Hole expansion ratio to ISO 1663035%

- Bending angle according to VDA 238-100 (longitudinal, transverse) 12871 14 °

The material characteristics transverse to the rolling direction would correspond for example to a HC660XD.

Initial state (HR)

The yield ratio Re / Rm in the transverse direction was 70%.

Yield strength (Re) 725 MPa

Tensile strength (Rm) 1030 MPa

Elongation at Break (A80) 10.2% Transverse Intermediate (CR) Yield strength (Re) 1224 MPa

Tensile strength (Rm) 1260 MPa

Elongation at break (A80) 1, 5% Example 3

(1.50 mm cold strip of 2.30 mm master warming tape and a pre-strip thickness of 50 mm)

% By weight alloy composition A 0.104% C steel of the invention; 0.443% Si; 2.178% Mn; 0.012% P;

0.0004% S; 0.0045% N; 0.038 AI; 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 after process 3 according to FIG. 8c, hot-dip refined, the 250 mm slab material was reversibly rolled in a roughing mill in the roughing train to a 50 mm pitch of 80% percent reduction and then in the hot strip mill at 910 final rolling target temperature ° C with a decrease of 96% hot rolled and at a reel target temperature of 650 ° C with a master hot strip thickness of 2.30 mm and after pickling without additional heat treatment (such as bell annealing) to 1.50 mm cold rolled in one pass (Kaltabwalzgrad 35%).

Sheet condition (HDG)

The yield ratio Re / Rm in the transverse direction was 65%.

Yield point (Rp0.2) 704 MPa

Tensile strength (Rm) 1084 MPa

Elongation at break (A80) 10.4%

Bake hardening index (BH2) 55 MPa

Hole expansion ratio to ISO 1663038%

Bending angle according to VDA 238-100 (longitudinal, transverse) 12771 15 °

Initial state (HR)

The yield ratio Re / Rm in the transverse direction was 69%.

Yield strength (Re) 695 MPa Tensile strength (Rm) 1010 MPa

Elongation at break (A80) 8.8%

Intermediate state (CR) in the transverse direction

Yield strength (Re) 1203 MPa

Tensile strength (Rm) 1255 MPa

Elongation at break (A80) 1, 9%

Example 4

(1.50 mm cold strip of 2.30 mm master warming tape and a pre-strip thickness of 55 mm)

Alloy composition in% by weight

An inventive steel with 0.104% C; 0.443% Si; 2.178% Mn; 0.012% P; 0.0004% S; 0.0045% N; 0.038 AI; 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 according to FIG. 8c, hot-dip refined, the slab material of 250 mm was hot rolled in the roughing mill to a pre-strip of 55 mm reversing rolled with a percentage decrease of 78% and then in the hot strip mill at a final rolling target temperature of

910 ° C with a decrease of 96% hot rolled and coiled at a reel target temperature of 650 ° C with a master hot strip thickness of 2.30 mm and after pickling without additional heat treatment (such as bell annealing) to 1.50 mm in one pass cold rolled ( Kaltabwalzgrad 35%).

Sheet condition (HDG)

The yield ratio Re / Rm in the transverse direction was 66%.

Yield point (Rp0.2) 708 MPa

Tensile strength (Rm) 1077 MPa

- Elongation at break (A80) 10.4%

Bake hardening index (BH2) 58 MPa

Hole expansion ratio according to ISO 1663040%

Bending angle according to VDA 238-100 (longitudinal, transverse) 12371 1 1 °

The yield ratio Re / Rm in the transverse direction was 70%.

Yield strength (Re) 679 MPa

Tensile strength (Rm) 967 MPa

Elongation at break (A80) 9.6%

Intermediate state (CR) in the transverse direction

Yield strength (Re) 1 158 MPa

Tensile strength (Rm) 1230 MPa

Elongation at break (A80) 2.5%

Conclusion:

A significant influence of the pre-strip thickness on the mechanical characteristics of the thin sheet (HDG) is not discernible.

This statement applies to the cold rolling rate of 35% used in the examples, but would also be fully transferable to variable cold rolling degrees.

Above, the invention has been described with reference to sheet steel plates with a final thickness of 1, 50 mm to be achieved in the thickness range 0.50 to 3.00 mm. It is also possible, if necessary, to produce final thicknesses in the range of 0.10 to 4.00 mm.

Claims

claims

1 . Highest strength multiphase steel having a minimum tensile strength of 980 MPa (in% by weight): c> 0.075 to <0.1 15

Si> 0.400 to <0.500

Mn> 1, 900 to <2,350

Cr> 0.250 to <0.400

AI> 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

Contaminants in which, with a view to the widest possible process window in the annealing, in particular continuous annealing, of cold tapes made of this steel, the total amount of Mn-Si + Cr> 1, 750 wt .-% to <2.250 wt .-%.

2. High-strength multiphase steel according to claim 1, characterized in that the sum amount of Mn-Si + Cr depending on the final thickness of the cold strip to be achieved has the following contents:

Final thickness 0.50 to 1.00 mm inclusive:

Sum of Mn-Si + Cr> 1, 750 wt% to <2.030 wt%,

Final thickness over 1.00 to 2.00 mm inclusive: Sum of Mn-Si + Cr> 1.940 wt.% To <2.1. 10

Final thickness over 2.00 up to and including 3.00 mm:

Sum of Mn-Si + Cr> 2.020 wt% to <2.220

3. Highest strength multiphase steel according to claim 1, characterized in that the sum amount of Mn-Si + Cr + Mo depending on the final thickness of the cold strip to be achieved has the following contents:

Final thickness 0.50 to 1.00 mm inclusive:

Sum of Mn-Si + Cr + Mo> 1.950% by weight to <2.280% by weight,

Final thickness over 1.00 to 2.00 mm inclusive:

Sum of Mn-Si + Cr + Mo> 2.140 wt% to <2.356 wt%

Final thickness over 2.00 mm up to and including 3.00 mm:

Sum of Mn-Si + Cr + Mo> 2.220 wt% to <2.470 wt%.

4. Highest strength multiphase steel according to one of claims 1 to 3, characterized in that

at a final thickness of the cold strip of 0.50 mm to 1.00 mm including the C content <0.100 wt .-%,

at a final thickness of the cold strip of more than 1.00 to 2.00 mm including the C content <0, 105 wt .-% and

at a final thickness of the cold strip of more than 2.00 to 3.00 mm including the C content <0.1 15 wt .-% is.

5. High-strength multi-phase steel according to one of claims 1 to 4, characterized in that the and carbon equivalent CEV (NW) depending on the final thickness of the cold strip to be achieved has the following contents:

Final thickness 0.50 mm up to and including 1.00 mm:

C content <0.100% by weight and carbon equivalent CEV (NW) <0.62%,

Final thickness over 1, 00 mm up to and including 2,00 mm:

C content <0.105% by weight and the carbon equivalent CEV (NW) <0.64%,

Final thickness over 2.00 mm up to and including 3.00 mm:

C content <0.1 15% by weight and the carbon equivalent CEV (NW) <0.66%.

6. Highest strength multiphase steel according to one of claims 1 to 5, characterized in that the Mn content depending on the final thickness to be achieved the cold strip has the following contents:

Final thickness 0.50 up to and including 1.00 mm:

Mn content> 1, 900% by weight to <2.200% by weight,

Final thickness over 1.00 to 2.00 mm inclusive:

Mn content> 2.050% by weight to <2.250% by weight,

Final thickness over 2.00 mm up to and including 3.00 mm:

Mn content> 2,100% by weight to <2,350% by weight.

7. High-strength multiphase steel according to one of claims 1 to 6, characterized in that the B content <0.0009 wt .-%, in particular> 0.0006% by weight to <0.0009 wt .-% is.

8. High-strength multiphase steel according to one of claims 1 to 7, characterized in that the sum of Nb + Ti <0.100 wt .-%, in particular the sum Nb + Ti <0.090 wt .-% is.

9. High-strength multiphase steel according to one of claims 1 to 8, characterized in that the sum of Ti + Nb + B <0.102 wt .-%, in particular <0.092 wt .-% is.

10. Highest strength multiphase steel according to claim 9, characterized in that at a sum of Ti + Nb + B of> 0.010 to <0.080 wt .-%, the N content> 0.0020 to <0.0090 wt .-% is or in particular at a sum of Ti + Nb + B of> 0.050 wt .-%,

the N content is> 0.0040 to <0.0120% by weight.

1 1. High-strength multiphase steel according to one of claims 1 to 10, characterized in that the S content is <0.0015 wt .-%,

in particular, the S content is <0.0010 wt .-%.

12. Highest strength multiphase steel according to one of claims 1 to 1 1, characterized in that the Mo content more than 0.200 wt .-% to 0.300 wt .-%, advantageously between more than 0.200 wt .-% and <0.250 wt. % is.

13. Highest strength multi-phase steel according to one of claims 1 to 12, characterized characterized in that the Ti content is> 0.025 wt .-% to <0.045 wt .-%.

14. High-strength multiphase steel according to one of claims 1 to 13, characterized in that the Nb content is> 0.025 wt .-% to <0.045 wt .-%.

15. High-strength multiphase steel according to one of claims 1 to 14, characterized in that the V content> 0.005 wt .-% to <0.020 wt .-%,

optimally> 0.005 wt .-% to <0.015 wt .-% is.

16. High-strength multiphase steel according to one of claims 1 to 15, characterized in that the sum Cr + Mo <0.650 wt .-% is.

17. High-strength multiphase steel according to one of claims 1 to 16, characterized in that the sum of Ti + Nb + B + Mo + V <0.365 wt .-% is.

18. High-strength multiphase steel according to one of claims 1 to 17, characterized in that the Ca content is <0.0030 wt .-%.

19. High-strength multi-phase steel according to one of claims 1 to 18, characterized in that the H content is <0.00050 wt .-%.

20. High-strength multiphase steel according to one of claims 1 to 19, characterized in that the sum of H + N> 0.0025 wt .-% to <0.0130 wt .-% is.

21. High-strength multiphase steel according to one of claims 1 to 20, characterized in that at a Cr content of> 0.260 wt .-% to <0.330 wt .-%, the carbon equivalent CEV (IIW) <0.62%, in a Cr Content> 0.290% by weight to <0.360% by weight of the carbon equivalent CEV (IIW) <0.64% or with a Cr content> 0.320% by weight to <0.370% by weight of the carbon equivalent CEV (IIW) <0.66%.

22. Highest strength multiphase steel according to claim 21, characterized in that the Cr content and the maximum carbon equivalent CEV (IIW) is chosen as a function of the final thickness of the cold strip to be obtained as follows: Final thickness 0.50 up to and including 1.00 mm:

Cr content> 0.260% by weight to <0.330% by weight, carbon equivalent CEV (NW) <0.62%,

Final thickness over 1.00 to 2.00 mm inclusive:

Cr content> 0.290% by weight to <0.360% by weight, carbon equivalent CEV (NW) <0.64%,

Final thickness over 2.00 up to and including 3.00 mm:

Cr content> 0.320 wt% to <0.370 wt%, carbon equivalent CEV (NW) <0.66%.

23. Highest strength multiphase steel according to one of claims 1 to 22, characterized in that the additions of Si and Mn in view of the strength properties to be achieved according to the relationships

YS (MPa) = 185.7 + 147.9 [% Si] + 161, 1 [% Mn]

TS (MPa) = 574.8 + 189.4 [% Si] + 174.1 [% Mn]

are interchangeable.

24. Highest strength multiphase steel according to one of claims 1 to 23, characterized in that the steel is curable by means of air cooling.

25. A method for producing a steel strip from a multiphase steel according to one of claims 1 to 24, characterized in that from the multiphase steel in the state of a slab a Vorband is produced, then from the pre-strip the steel strip is hot rolled with the desired hot strip thickness.

26. The method according to claim 25, characterized in that, starting from a previously determined slab thickness and a pre-selected pre-strip with a defined but variable thickness, hot strips with the same thickness

Rolling degrees of 72% to 87% are hot rolled with the final thickness to be achieved.

27. The method according to claims 25 and 26, characterized in that a hot strip is produced from the hot strip, a cold strip is cold rolled with the final thickness to be achieved and then annealed the cold strip, in particular is annealed.

28. The method according to claim 27, characterized in that, starting from a selected master warming band with a certain thickness, or selected hot strips with different thicknesses, cold strips are produced with Kaltabwalz- degrees of 10% to 70% with the final thickness to be achieved.

29. The method according to any one of claims 25 to 28, characterized in that to produce the required multi-phase structure, the cold rolled to final thickness steel strip during the continuous annealing to a temperature in the range of about 700 to 950 ° C heated and that the annealed steel strip then from the annealing temperature with a cooling rate between about 15 and 100 ° C / s up to a first intermediate temperature of about 300 to 500 ° C, following with a cooling rate between about 15 and 100 ° C / s up to a second intermediate temperature is cooled from about 160 to 250 ° C, then the steel strip with a

Cooling rate of about 2 to 30 ° C / s is cooled to room temperature in air or cooled at a cooling rate between about 15 and 100 ° C / s from the first intermediate temperature to room temperature.

30. The method according to any one of claims 25 to 28, characterized in that to produce the required multi-phase structure, the cold rolled to final thickness steel strip during the continuous annealing to a temperature in the range of about 700 to 950 ° C, then heated to a temperature of approx Cooled to 400 to 470 ° C, wherein the cooling stopped before entering the molten bath, then melt-refined and after the Schmelztauchveredelung the cooling at a cooling rate between about 15 and 100 ° C / s up to an intermediate temperature of about 200 to 250 ° C and then the steel strip with a

Cooling rate of about 2 and 30 ° C / s is cooled until reaching room temperature in air.

31. Method according to one of claims 25 to 28, characterized in that to produce the required structure, the cold rolled to final thickness steel strip during the continuous annealing to a temperature in the range of about 700 to 950 ° C heated, then to an intermediate temperature of about 200 to Cooled to 250 ° C and held before entering the molten bath, the temperature for about 1 to 20 s and then the steel strip reheated to a temperature of about 400 to 470 ° C, then melt-refined and after the Schmelztauchveredlung a new cooling with a cooling rate between about 15 and 100 ° C / s up to an intermediate temperature of about 200 to 250 ° C and then cooled at a cooling rate of about 2 and 30 ° C / s in air to room temperature.

32. The method according to any one of claims 25 to 31, characterized in that in the continuous annealing, the oxidation potential in an annealing with a system configuration consisting of directly fired furnace area (NOF) and a radiant tube furnace (RTF) by a CO content in the NOF of below 4 vol .-% is increased, wherein in the RTF, the oxygen partial pressure of the iron-reducing furnace atmosphere is set according to the following equation,

where Si, Mn, Cr, B denote the corresponding alloying proportions in the steel in wt .-% and PO2 the oxygen partial pressure in mbar and to avoid

Oxidation of the band just before immersion in the molten bath, the dew point of the gas atmosphere at -30 ° C or below is set.

33. The method according to claim 32, characterized in that, in the case of an annealing, the oxygen partial pressure of the furnace atmosphere satisfies the following equation only with a jet tube furnace,

where Si, Mn, Cr, B denote the corresponding alloying proportions in the steel in wt .-% and p02 the oxygen partial pressure in mbar and to avoid the

34. The method according to any one of claims 25 to 33, characterized in that at different thickness steel strips during continuous annealing comparable

Microstructure states and mechanical characteristics of the belts can be set by adjusting the system throughput speed in the course of the heat treatment.

35. The method according to claim 27, characterized in that the steel strip is trained after the annealing or hot dip finishing.

36. The method according to claim 27, characterized in that the steel strip is stretch bend-rectified following the annealing or hot dip finishing.

37. The method according to at least one of claims 25 to 36, characterized in that from the steel strip, a circuit board is cut, which then heated to a temperature above Ac3, the heated board is formed into a component and then cured in the tool or in air.

38. Steel strip produced by the method according to one of claims 25 to 37, having a minimum hole expansion value according to ISO 16630 of 20%, in particular of 25%.

39. Steel strip according to claim 38, comprising a minimum bending angle according to VDA 238-100 of 70 ° in the longitudinal or transverse direction, in particular of 85 °.

40. Steel strip according to claim 38 or 39, comprising a minimum product value Rm x α (tensile strength x [bending angle according to VDA 238-100]) of 100000 MPa °, in particular of 120000 MPa °.

41. Steel strip according to one of claims 38 to 40, having a delayed fracture free state for at least 6 months, meeting the requirements of SEP 1970 for Lochzug- and ironing sample.