Patent application title: HIGH STRENGTH MULTI-PHASE STEEL, AND METHOD FOR PRODUCING A STRIP FROM SAID STEEL
Thomas Schulz (Salzgitter, DE)
Marion Bechthold (Braunschweig, DE)
Norbert Kwiaton (Salzgitter, DE)
Alexander Georgiew (Wolfenbüttel, DE)
IPC8 Class: AC21D952FI
Class name: Chromium containing, but less than 9 percent molybdenum containing titanium, zirconium or niobium containing
Publication date: 2015-11-26
Patent application number: 20150337408
A high-strength multiphase steel with minimum tensile strengths of 580
MPa, preferably having a dual-phase structure for a cold-rolled or
hot-rolled steel strip with improved forming properties, particularly for
lightweight vehicle construction contains the elements (contents in
mass-%): C 0.075 to ≦0.105; Si 0.600 to ≦0.800; Mn 1.000 to
≦2.250; Cr 0.280 to ≦0.480; Al 0.010 to ≦0.060;
P≦0.020; N≦0.0100; S≦0.0150, remainder iron,
including typical steel-accompanying elements not mentioned above, which
are impurities introduced by smelting, with the condition that the Mn
content is preferably ≦1.500% for strip thicknesses up to 1 mm,
the Mn content is preferably ≦1.750% for strip thicknesses of 1 to
2 mm, and the Mn content is preferably ≧1.500% for strip
thicknesses ≧2 mm
17. A high strength multiphase steel with minimal strengths of 580 MPa preferably with dual phase microstructure for a cold or hot rolled steel strip having improved forming properties, in particular for the vehicle lightweight construction, composed of the following elements weight %: C 0.075 to ≦0.105 Si 0.600 to ≦0.800 Mn 1.000 to ≦2.250 Cr 0.280 to ≦0.480 Al 0.010 to ≦0.060 P≦0.020 N≦0.0100 S≦0.0150 remainder iron including usual steel accompanying elements not mentioned above which constitute smelting related impurities, with the proviso that at strip thicknesses up to 1 mm the Mn content is preferably ≦1.500%, that at strip thicknesses from 1 mm to 2 mm the Mn content is preferably ≦1.750%, and that at strip thicknesses 2 mm the Mn content is preferably ≧1.500%.
18. The steel of claim 1, wherein at strip thicknesses up to 1 mm a sum of Mn+Si+Cr is ≧1.88.ltoreq.2.60%.
19. The steel of claim 17, wherein at strip thicknesses of 1.00-2.00 mm the sum of Mn+Si+Cr is ≧2.2.ltoreq.3.00%
20. The steel of claim 17, wherein at strip thicknesses of ≧2.00 the sum of Mn+Si+Cr is ≧2.50.ltoreq.3.53%.
21. The steel of claim 17, wherein the N content is ≦0.0090%.
22. The Steel according of claim 17, wherein the N content is ≦0.0080%.
23. The steel of claim 17, wherein the S content is ≦0.0050%.
24. The steel of the claim 17, wherein the S content is ≦0.0030%.
25. A method for producing a cold or hot rolled steel strip from the steel of claim 17, comprising: heating the cold or hot rolled steel strip during a continuous annealing to an annealing temperature in the range of about 700 to 950.degree. C.; cooling the steel strip is from the annealing temperature to a first intermediate temperature of about 300 to 500.degree. C. at a cooling rate between about 15 and 100.degree. C./s; cooling the cold or hot rolled steel strip to a second intermediate temperature of about 200 to 250.degree. C. with a cooling rate between about 15 to 100.degree. C./s; and cooling the steel strip at air with a cooling rate of about 2 to 30.degree. C./s until reaching room temperature or maintaining the cooling with a cooling rate between about 15 to 100.degree. C./s from the first intermediate temperature to room temperature, wherein a dual phase microstructure is generated during the continuous annealing.
26. The method of claim 25, further comprising hot dip galvanizing the strip in a hot dip bath, wherein subsequent to the heating and subsequent cooling the cooling is halted prior to entering into the hot dip bath, and after the hot dip galvanizing the cooling is continued with a cooling rate between about 15 and 100.degree. C./s until reaching an intermediate temperature of about 200 to 250.degree. C., and subsequently the steel strip is cooled on air with a cooling rate between about 2 and 30.degree. C./s until reaching room temperature.
27. The method of claim 25, further comprising hot dip galvanizing the steel strip in a hot dip bath, wherein after the heating and subsequent cooling to the intermediate temperature of about 200 to 250.degree. C. and prior to entering the hot dip bath the temperature is held for about 1 to 20 s and subsequently the steel strip is reheated to the temperature of about 420 to 470.degree. C. and after the hot dip galvanizing the steel strip is cooled until reaching the intermediate temperature of about 200 to 250.degree. C. with a cooling rate between about 15 and 100.degree. C./s, and subsequently the steel strip is cooled on air with a cooling rate of about 2 and 30.degree. C./s until reaching room temperature.
28. The method of the claim 25, wherein the heating step is carried out in a plant configuration composed of a directly fired furnace region and a radiant-tube furnace, the method further comprising increasing an oxidation potential in the annealing by setting a CO-content below 4%, setting an atmosphere of the furnace to be reducing and setting a dew point at -30.degree. C. or below -30.degree. C. so as to avoid oxidation of the strip prior to immersion into the hot dip bath.
29. The method of claim 25, wherein the annealing is carried out by solely utilizing a radiant-tube furnace, wherein a dew the dew point in the furnace atmosphere is -30.degree. C. or above -30.degree. C.
30. The method of claim 29, wherein the dew point in the furnace atmosphere is -25.degree. C. or -20.degree. C.
31. The method of claim 25, further comprising adjusting a throughput speed as a function of varying thicknesses of individual steel strips during the heat treatment thereby establishing comparable microstructure states and mechanical characteristic values in the individual steel strips.
32. The method of claim 25, further comprising subsequent to the heat treatment skin passing the steel strip.
33. The method of claim 25, further comprising aligning the steels strip stretch bending subsequent to the heat treatment.
 The invention relates to a high strength multiphase steel according
to the preamble of claim 1.
 The invention also relates to a method for producing a hot and/or cold rolled strip from such a steel according to patent claim 9.
 The invention relates in particular to steels with tensile strengths in the range of from 580 MPa to maximally 700 MPa with low yield ultimate ratio of below 66% for producing components, which have excellent formability and improved welding properties.
 The hotly contested automobile market forces manufacturers to constantly seek solutions to lower the fleet consumption while at the same time maintaining a highest possible comfort and occupant protection. Hereby weight saving of all vehicle components plays an important role but also a best possible behavior of the individual components under conditions of high static or dynamic stress during use and also in the event of a crash. Pre-material suppliers seek to account for this requirement by providing high-strength and ultra-high-strength steels with thin sheet thickness to reduce the weight of the vehicle components while at the same time improving forming and component properties during manufacture and during use.
 High-strength and ultra-high-strength steels enable more lightweight vehicle components which as a consequence leads to reduced fuel consumption and reduced pollution due to the reduced CO2 proportion.
 These steels therefore have to meet relatively high demands regarding their strength and ductility, energy absorption capacity and during processing, such as for example during pickling, hot or cold forming, welding and/or metallic veredelung (organic coating or varnishing.
 Newly developed steels thus must met the demands on weight reduction, the increasing material demands on ultimate yield strength, strain hardening behavior and elongation a fracture at good formability, as well as the demands on the component of high tenacity, border crack resistance, energy absorption and hardenabiltiy via the work hardening effect and the bake hardening effect, but also improved suitability for joining in the form of improved weldabiltiy,
 Improved edge crack resistance means increased hole expansion capacity during forming and is known under synonymous terms such as low edge crack (LEC) or high hole expansion (HHE)
 Improved weldability is among other things achieved by a lowered carbon equivalent. Synonymous terms for this are under-peritectic (UP) or the already known low carbon equivalent (LCE).
 The steel according to the invention is also associated with the goal to reduce the thickness of the steels used in automobile construction (for example micro-alloyed LA-steels or LAD-steels) in order to save weight.
 For such a reduction of the thickness of the steel sheet a high strength steel for example dual phase steel (DP) has to be used in order to ensure sufficient strength of the motor vehicle components.
 In vehicle construction therefore dual phase steels are increasingly used, which consist of a ferritic basic structure in which a martensitic second phase and possibly a further phase with bainite and residual austenite is integrated. Bainite can be present in different forms.
 The specific material properties of the dual phase steels that determine the steel types such as a low yield ultimate ratio at very high tensile strength, a strong cold hardening and a good cold fomability are well known.
 Increasingly multiphase steels are also used such as complex-phase steels, ferritic-bainitic steels, bainitic steels, martensitic steels, TRIP-steels and the above described dual phase steels which are characterized by different microstructure compositions as described in EN 10346 (Norm of the European steel manufacturers) or VDA239-100 (Norm of an automobile steel user) and are described here in the following.
 Complex phase steels are steels that contain small proportions of martensite, residual austenite and/or perlite in a ferritic/bainitic basic structure, wherein an extreme grain refinement is caused by a delayed re-crystallization or by precipitations of micro alloy elements.
 Ferritic bainitic steels are steels which contain bainite or strain hardened bainite in a matrix of ferrite and/or strain hardened ferrite. The strength of the matrix is caused by a high dislocation density, by grain refinement and by the precipitation of micro alloy elements.
 Bainitic steels are steels that are characterized by a very high yield strength and tensile strength at a sufficiently high expansion for cold forming processes. The chemical composition results in a good weldability. The microstructure typically consists of bainite. In some cases small proportions of other phases such as marteniste and ferrite can be contained.
 Martensitic steels are steels that as a result of thermo mechanical rolling contain small proportions of ferrite and/or bainite in a basic structure of martensite. The steel type is characterized by a very high yield strength and tensile strength at sufficiently high expansion for cold forming processes. Within the group of multi-phase steels the martenisititc steels have the highest tensile strength values.
 Dual phase steels are steels are steels with a ferritic basic microstructure in which a martenisitc second phase is incorporated island like, possibly also with bainite as second phase. At high tensile strength dual phase steels exhibit a low yield ultimate ratio and a strong strain hardening.
 TRIP steels are steels with a predominantly ferritic basic microstructure in which residual austenite is incorporated which can transform into marteniste during deformation (TRIP effect). Due to its strong strain hardening the steel achieves high values of the uniform elongation and tensile strength.
 These multiphase steels are used in structural components, chassis and crash relevant components as steel plates, Tailored Blanks (welded steel plates) as well as flexibly cold tolled strips, so called TRBs.
 The Tailor Rolled Blank lightweight construction technology (TRB®) enables a significant weight reduction as a result of the load adjusted selection of sheet thickness over the length of the component
 In the continuous annealing system, a special heat treatment occurs in which the steel is provided with its low yield strength by relatively soft components such as ferrite or bainitic ferrite and is provided with is strength by its hard components such as martensite or carbon rich bainite.
 For economic reasons cold rolled high strength to ultra-high strength steel strips are usually subjected to recrystallizing annealing in the continuous annealing process to generate well formable steel sheet. Depending on the alloy composition and the strip cross section, the process parameters such as throughput speed, annealing temperature and cooling rate (cooling profile), are adjusted corresponding to the demanded mechanical-technological properties by way of the microstructure required therefore.
 For establishing the dual-phase microstructure, the unpickled or pickled hot strip in typical thicknesses between 1.50 mm to 4.00 mm, or cold strip in typical thicknesses of 0.50 mm to 3.00 mm, is heated in the continuous annealing furnace to such a temperature that the required microstructure forms during the cooling. The same applies for configuring a steel with complex phase microstructure, martensitic, ferritc-bainitic and also purely bainitic microstructure.
 Especially in the case of different thicknesses in the transition region of one strip to another, a constant temperature is difficult to achieve. In the case of alloy compositions with too narrow process windows this can lead to the fact that for example the thinner strip is either moved through the furnace too slowly, thereby lowering productivity, or that the thicker strip is moved through the furnace too fast and the required annealing temperatures and cooling gradients for achieving the desired microstructure are not reached. The result of this is increased waste.
 Widened process windows are necessary in order to enable the demanded strip properties at same process parameters also in the case of greater cross sectional changes of the strips to be annealed.
 The problem of a too narrow process window is especially pronounced in the annealing treatment when stress-optimized components made of hot or cold strip are to be produced, which have sheet thicknesses that vary across the strip length and strip width (for example as a result of flexible rolling).
 However, when strongly varying sheet thicknesses are involved, production of TRB®s with multiphase microstructure employing the presently known alloys and available continuous annealing systems requires increased efforts, for example an additional heat treatment prior to the cold rolling. In regions of different sheet thickness, i.e., when varying degrees of rolling reduction are present, a homogenous multiphase microstructure cannot be established in cold rolled and hot rolled steel strips due to the temperature difference in the conventional process windows.
 A method for producing a steel strip with different thickness across the strip length is for example described in DE 100 37 867 A1.
 When high demands on corrosion protection require the surface of the hot or cold strip to be hot dip galvanized, the annealing is usually carried out in a continuous annealing furnace arranged upstream of the hot dip galvanizing bath.
 Also in the case of hot strip the demanded microstructure is only established in the annealing in the continuous furnace depending on the alloy concept, in order to realize the demanded mechanical properties.
 Deciding process parameters are thus the adjustment of the annealing temperatures and the speed as well as the cooling rate (cooling gradient) in the continuous annealing because the phase transformation is temperature and time dependent. Thus, the less sensitive the steel is regarding the uniformity of the mechanical properties when temperature and time course change during the continuous annealing, the greater is the process window.
 In the continuous annealing of hot and cold rolled steel strips of different thickness utilizing for example the alloy concept for a multiphase steel known from DE 196 10 675 C1, the problem arises that although the alloy composition tested there satisfies the demanded mechanical properties, only a narrow process window is available of the annealing parameters to enable establishment of uniform mechanical properties across the strip length without having to adjust the process parameters
 A further disadvantage is that the very high Al-contents of 0.4-2.5% adversely affect steel production via conventional strip casting due to the risk of clogging during the ladle treatment and the uncontrolled reaction with the casting powder during solidification.
 When using the known alloy concepts for the group of multiphase steels, the narrow process window makes it already difficult during the continuous annealing of strips with different thicknesses to establish uniform mechanical properties over the entire length and width of the strip.
 In the case of flexibly rolled cold strip made of multiphase steels of known compositions, the too narrow process window either causes the regions with lower sheet thickness to have excessive strengths resulting from excessive martensite proportions due to the transformation processes during the cooling, or the regions with greater sheet thickness achieve insufficient strengths as a result of insufficient martensite proportions. Homogenous mechanical-technological properties across the strip length or width can practically not be achieved with the known alloy concepts in the continuous annealing.
 The goal to achieve the resulting mechanical-technological properties in a narrow region across the strip width and strip length through controlled adjustment of the volume proportions of the microstructure phases has highest priority and is therefore only possible through a widened process window. The known alloy concepts for multiphase steels are characterized by a too narrow process window and are therefore not suited for solving the present problem, in particular in the case of flexibly rolled strips. With the alloy concepts known to date only steels of a strength class with defined cross sectional regions (sheet thickness and strip width) can be produced, hence requiring different alloy concepts for different strength classes or cross sectional ranges.
 The state of the art is to increase the strength by increasing the amount of carbon and/or silicone and/or manganese and via the microstructure adjustment and solid solution strengthening (solid solution hardening).
 However, increasing the amounts of the aforementioned elements, increasingly worsens the material processing properties for example during welding, forming and hot dip coating, but also the industrial production in all process steps such as steel production, hot rolling picking, cold rolling and heat treatment with/vvithout hot dip coating places increased demands in the individual plants.
 On the other hand, there is also a trend in the steel production to reduce the carbon and/or manganese content in order to achieve a better cold processability and better performance properties.
 For describing and quantifying the cold processing, in particular the edge crack behavior, the hole expansion test according to ISO 11630 is used as one of multiple possible test methods.
 But also suitability for welding, characterized by the carbon equivalent has become an increased focus.
 For example in the following carbon equivalents (see FIG. 3)
 the characteristic standard elements such as Carbon and Manganese as well as Chromium or Molybdenum and Vanadium are taken into account.
 Silicone plays a minor role for calculating the carbon equivalent. This is of deciding importance with regard to the invention. The lowering of the carbon equivalent through lower contents of Carbon and Manganese is to be compensated by increasing the silicone content. Thus the edge crack resistance and welding suitability are improved at same strengths.
 A low yield strength ratio (Re/Rm) is typical for a dual-phase steel and serves in particular for the formability in stretching and deep drawing processes. This provides the constructor with information regarding the distance between ensuing plastic deformation and failing of the material at quasi static load. Correspondingly lower yield strength ratios represent a greater safety margin for the component failure.
 A higher yield strength ratio (Re/Rm), as it is typical for complex-phase steels, is also characterized by a resistance against edge cracks. This can be attributed to the smaller differences in the strengths of the individual microstructure components, which has a positive effect on a homogenous deformation in the region of the cutting edge.
 The analytical landscape for achieving multiphase steels with minimal strengths of 580 MPa has become more diverse and reveals very broad alloy ranges regarding the strength-promoting elements carbon, silicone, manganese, phosphorous, aluminum and chromium and/or molybdenum as well as regarding the addition of micro-alloys such as titanium and vanadium and regarding the material characterizing properties.
 The spectrum of dimensions is broad and lies in the thickness range of 0.50 to 4.00 mm. Predominantly strips up to about 1850 mm are used but also slit strip dimensions, which are generated by longitudinally separating the strips. Sheets or plates are generated by transverse separation of the strips.
 The invention is therefore based on the object to set forth a new alloy concept for a high strength multi-phase steel with a minimal tensile strength of 580 MPa to 700 MPa longitudinally and transversely to the rolling direction, preferably with dual-phase microstructure and a yield strength ratio of less than 66% with which the process window for the continuous annealing of hot and cold rolled strips can be widened so that beside strips with different cross sections also steel strips with thicknesses that vary over the strip length or strip width and the correspondingly varying cold rolling reduction degrees can be generated with mechanical technological properties that are as homogenous as possible. In addition a method for producing a strip made of this steel is set forth.
 According to the teaching of the invention this object is solved by a steel with the following contents in weight %:
C 0.075 to ≦0.105
Si 0.600 to ≦0.800
Mn 1.000 to ≦2.250
Cr 0.280 to ≦0.480
Al 0.010 to ≦0.060
 remainder iron including usual steel accompanying elements not mentioned above.
 The steel according to the invention is very well suited for a hot dip galvanizing and has a significantly widened process window compared to the known steels. This results in an increased process reliability during continuous annealing of cold and hot strip with dual-phase or multiphase microstructure. Thus more homogenous mechanical-technological properties can be ensured in the strip for continuously annealed hot or cold strips also in the case of different cross sections and otherwise same process parameters.
 This applies for the continuous annealing of subsequent strips with different strip cross sections as wells as for strips with varying strip thickness and strip length or strip width. This enables for example processing in selected thickness ranges (such as a strip thickness of smaller than 1 mm, a strip thickness of 1 to 2 mm and a strip thickness of 2 to 4 mm).
 When high-strength, hot strips or cold strips made of multiphase steel with varying sheet thicknesses are produced according to the invention in the continuous annealing method, stress-optimized components can advantageously be produced from this material by forming.
 The produced material can be produced as cold strip and also as hot strip via a hot dip galvanizing line or a pure continuous annealing line in the skin passed or non skin passed state and also in the heat treated state (intermediate annealing).
 With the alloy composition according to the invention steel strips can be produced by an intercritical annealing between Ac1 and Ac3 or an austenitic annealing above Ac3 with final controlled cooling, which leads to a dual phase or multiphase microstructure.
 Annealing temperatures of 700° C. to 950° C. have proven advantageous. Depending on the overall process there are different approaches for realizing the heat treatment.
 In a continuous annealing system without subsequent hot dip coating, the strip is cooled starting from the annealing temperature to an intermediate temperature of about 200 to 250° C. with a cooling rate of about 15 to 100° C./s. Optimally, cooling to a previous intermediate temperature of 300 to 500° C. can occur beforehand with a cooling rate of 15 to 100° C./s. Finally cooling to room temperature occurs with a cooling rate of about 2 to 30° C.
 In a heat treatment within the framework of a hot dip coating two possible temperature profiles exist. The cooling as described above is halted prior to entry into the dip bath and is only continued after emergence from the bath until reaching the intermediate temperature of about 200 to 250° C. Depending on the dip bath temperature, a holding temperature of about 420 to 470° C. results in this case. The cooling to room temperature occurs again with a cooling rate of 2 to 30° C./s.
 The second variant of the temperature profile in the hot dip coating includes holding the temperature for about 1 to 20 s at the intermediate temperature of 200 to 250° C. and subsequent reheating to the temperature of 420 to 470° C. required for the hot dip coating. After the hot dip coating the strip is cooled again to 200 to 250° C. The cooling to room temperature occurs again with a cooling rate of 2 to 30° C./s.
 Beside manganese, chromium and silicone, carbon is responsible for the transformation of austenite to martensite in classical dual-phase steels.
 Only the combination according to the invention of the added elements carbon, silicone, manganese and chromium ensures on one hand the demanded mechanical properties of minimal tensile strengths of 580 MPa and yield strength ratios of below 66% at simultaneous significantly widened process window in the continuous annealing.
 It is characteristic for the material that the addition of silicone causes a strong solid solution hardening. On estimate addition of 0.1% silicone causes an increase of the tensile strength by about 10 MPa, wherein up to 2.2% silicone only slightly slightly negatively affects the stretching. Among other things the latter is due to the fact that silicone lowers the solubility of carbon and ferrite which causes the ferrite to become softer which in turn improves the formability. In addition silicone prevents formation of carbides, which as brittle phases lower ductility. The low strength increasing effect of silicone within the range of the steel according to the invention creates the basis for a wide process window.
 Characteristic for the material is also that increasing weight percentages of added manganese causes shifting of the ferrite region toward longer times and lower temperatures during cooling. The proportions of ferrite are hereby reduced to a lesser or stronger degree by increased proportions of bainite depending on the process parameters.
 The basis for achieving the wide process window is the micro-alloying according to the invention of exclusively with niobium, while taking into account the above mentioned classical composition of carbon/silicone/manganese/chromium with a manganese content which is stepped and defined according to the strip thickness.
 Because the speed in the continuous annealing system is reduced with increasing cross section or strip thickness at same width, i.e., the time available for transformation increases, manganese has to take over this task in order to establish similar microstructure proportions over the selected thickness range (for example 0.5 to 4.0 mm) and to shift the phase transformations correspondingly, as schematically shown in FIG. 6 in the variants 1, 2 and 3.
 Tests have surprisingly shown that in particular the addition of silicone in amounts of 0.600-0.800 enables a wide process window for a great range of dimensions and to achieve the demanded tensile strength of at least 580 MPa for hot strip and at least 590 MPA for cold rerolled hot strip and cold strip.
 By setting a low carbon content of s 0.105% the carbon equivalent can be reduced which improves the weldability and excessive hardening is avoided. In addition the service life of the electrode in the resistance spot welding can be significantly increased.
 In the following the effect of the elements in the alloy according to the invention is described in more detail. The multiphase steels typically have a chemical composition in which alloy components are combined with and without micro-alloying elements. Accompanying elements are unavoidable and are taken into account regarding their effect when necessary.
 Accompanying elements are elements, which are already present in the iron ore or enter the steel due to manufacturing. Due to their predominantly negative effect they are usually undesired. It is sought to remove them to a tolerable content or to convert them into less deleterious forms.
 Hydrogen (H) is the only element capable of diffusing through the iron lattice without generating lattice tensions. As a result hydrogen is relatively mobile in the iron lattice and can be taken up relatively easily during manufacturing. Hydrogen can thereby only be taken up into the iron lattice in atomic (ionic) form.
 Hydrogen has a strong embrittling effect and diffuses preferably to energetically favorable sites (defects, grain boundaries etc.). The defects act as hydrogen traps and can significantly increase the retention time of the hydrogen in the material.
 The recombination to molecular hydrogen can lead to cold cracks. This behavior occurs in the hydrogen embrittlement or in hydrogen-induced stress corrosion. Hydrogen is also often named as the cause for the so-called delayed fracture, which occurs without external tensions.
 A more uniform microstructure, which is achieved with the steel according to the invention among other things by virtue of its widened process window, lowers the sensitivity against hydrogen embrittlement.
 Therefore the hydrogen content in the steel should be as low as possible.
 Oxygen (O): in the molten state, steel has a relatively great capacity for absorbing gases, however, at room temperature oxygen is only soluble in very low amounts. Analogous to hydrogen, oxygen can only diffuse into the material in atomic form. Due to the strongly embrittling effect and the negative effect on the ageing resistance, oxygen content is sought to be reduced during production as much as possible.
 For decreasing oxygen, production methods such as a vacuum treatment exist on one hand, and on the other hand analytical approaches. Oxygen can be converted into harmless states by adding certain alloy elements. Thus binding of oxygen via manganese, silicone and/or aluminum is common. However, the oxide produced thereby can cause negative properties in the material in the form of defects. On the other hand a fine precipitation of aluminum oxides can lead to a grain refinement.
 For the above stated reasons the oxygen content in the steel should be as low as possible.
 Nitrogen (N): is also an accompanying element in steel production. Steels with free nitrogen are prone to a strong ageing effect. Nitrogen already diffuses at low temperatures at dislocations and blocks the same. As a result it causes a strength increase associated with a fast loss of tenacity. Nitrogen can be bound in the form of nitrides by adding aluminum or titanium.
 For the foregoing reasons the nitrogen content is limited to ≦0.0100%, ≦0.0090% or optimally to ≦0.0080% or to unavoidable amounts during steel production.
 Sulfur (S): like phosphorous is bound as trace element in the iron ore. It is undesired in the steel (exception automate steels) because of its strong tendency for segregation and embrittling effect. It is therefore sought to achieve as low amounts of sulfur as possible in the metal (for example by a deep vacuum treatment). Further the present sulfur is converted into the relatively harmless compound manganese sulfide (MnS).
 The manganese sulfides are often rolled out band-like during rolling and function as germination sites for the transformation. Especially in the case of diffusion controlled transformation this leads to a microstructure that is configured band-like and can lead to decreased mechanical properties in the case of strongly pronounced banding (for example pronounced martensite bands instead of distributed martensite islands, anisotropic material behavior, reduced elongation at fracture).
 For the foregoing reasons the sulfur content is limited to ≦0.0150%, advantageously to ≦0.0050 or to unavoidable amounts during steel production.
 Phosphorous (P) is a trace element from the iron ore and is solubilized in the iron lattice as substitution atom. As a result of the solid solution strengthening phosphorous increases strength and improves hardenability.
 However, it is usually sought to lower the phosphorous content as far as possible because among other things due to its slow diffusion speed it has a strong tendency to segregation and strongly lowers tenacity. Deposition of phosphorus at the grain boundaries can lead to grain boundary cracks. In addition phosphorous increases the transition temperature from tenacious to brittle behavior by up to 300° C. During hot rolling, surface-proximate phosphorous oxides can lead to separation at the grain boundaries.
 However, due to the low costs and the high strength increase, phosphorous is used in some steels in low amounts (<0.1%) as micro-alloying element. For example in high strength steels (interstitial free) or also in some alloying concepts for dual-phase steels.
 For the aforementioned reasons phosphorous is limited to ≦0.020% or to unavoidable amounts during steel production.
 Alloying elements are usually added to the steel in order to influence properties in a targeted manner. An alloying element can influence different properties in different steels. The effect generally depends strongly on the amount and the solubility state in the material.
 The interrelations can thus be very diverse and complex. In the following the effect of the alloying elements is described in more detail.
 Carbon (C): counts as the most important alloy element in the steel. Its targeted introduction to up to 2.06% turns iron into steel in the first place. Oftentimes the carbon content is drastically lowered during steel production. In dual-phase steels for a continuous hot dip coating its content is maximally 0.230%, a minimal value is not given.
 Due to its relatively small atomic radius carbon is dissolved interstitially in the iron lattice. The solubility in the α-iron is maximally 0.02% and in the γ-iron maximally 2.06%. In solubilized form carbon increases the hardenability of steel significantly and is thus indispensible for the formation of sufficient amounts of martensite. Excessive carbon contents however increase the hardness difference between ferrite and martensite and limit weldability.
 In order to satisfy the demands on a high hole expansion, the steel according to the invention is under-peritectic.
 As a result of the different solubility, pronounced diffusion processes are necessary in the phase transformation, which can lead to very different kinetic conditions. In addition carbon increases the thermodynamic stability of the austenite, which becomes apparent in the phase diagram as a widening of the austenite region toward lower temperatures. With increasing force-solubilized carbon content in the martensite the lattice distortions increase and associated with this the strength of the non-diffusively generated phase.
 In addition carbon is required for the formation of carbides. A representative is zementite (Fe3C), which is present in almost every steel. However, significantly harder special carbides can form with other metals such as chromium, titanium, niobium and vanadium. Not only the type but also the distribution and size of the precipitations is of deciding importance for the resulting strength increase. In order to ensure a sufficient strength on one hand and a good weldability on the other hand, the minimal C-content is set to 0.075% and the maximal C-content to 0.105%.
 Silicone (Si): binds oxygen during casting and thus lowers segregation and contaminations in the steel. The segregation coefficient is significantly lower than for example in the case of manganese (0.16 compared to 0.87)
 A further important effect is that silicone shifts the formation of ferrite toward shorter times and thus enables the generation of sufficient amounts of ferrite prior to the quenching. During hot rolling this thus creates a basis for an improved cold rollability. The accelerated ferrite formation causes enrichment of the austenite with carbon during hot dip galvanizing and is thus stabilized the accelerated ferrite formation. Because silicone inhibits carbide formation the austenite is additionally stabilized. Thus the formation of bainite can be suppressed in the accelerated cooling in favor of martensite.
 In the finished product silicone in addition increases the strength and the yield ultimate ratio of the ferrite as a result of solid solution hardening at only slightly decreased elongation at fracture. In contrast to carbon and manganese silicone reduces the difference in hardness between the microstructure components ferrite and martensite because it increases the solubility for carbon in the ferrite.
 In addition silicone leads to a less banded arrangement of the microstructure components. A small hardness difference and a band free microstructure have a positive effect on the processability, in particular n the hole expansion.
 As is known, in steels with high-silicone alloyed steels strongly adhering red scale is thought to form and a higher risk of rolled-in scale arises during hot rolling, which can have an influence on the subsequent pickling result and the pickling productivity. This effect could not be detected in the steel according to the invention with 0.600% to 0.800% silicone when the pickling was advantageously performed with hydrochloric acid instead with sulfuric acid.
 With regard to the galvanization capacity of silicone containing steels, DE 196 10 675 C1 describes that steels with up to 0.800% silicone or up to 2.000% silicone are not hot dip coatable due to the very poor wettability of the steel surface with the liquid zinc.
 It is known in the state of the art that in the continuous galvanizing, silicone can diffuse during the annealing to the surface and by itself or together with manganese form film-like oxides. These oxides adversely affect the galvanization by impairing the galvanization reaction (solubilization of iron and formation of inhibition layer) during dipping of the steel strip into the zinc melt. This manifests itself in a poor zinc adhesion and un-galvanized regions.
 Against this general knowledge in the state of the art it was unexpectedly found in the tests that solely by suitably operating the furnace during recrystallizing annealing and during passage through the Zinc bath a good galvanizabiltiy of the steel strip and a good Zinc adhesion can be achieved.
 For this purpose, the surface first has to be freed of scale remnants, rolling oil or other dirt particles by a chemical or thermal-hydro-mechanical pre-cleaning. In order to prevent silicone oxides from reaching the surface, measures also have to be taken to promote the inner oxidation of the alloy elements below the surface of the material. Depending on the configuration of the plant, different measures are used for this purpose.
 In a configuration of a plant which combines a directly fired furnace (NOF=non oxidizing furnace) and a radiant tube furnace (RTF--radiant tube furnace) (see method 2 in FIGS. 2, 4, 8) the oxidation reduction method is used. Hereby a thin iron oxide layer is formed in the NOF by increasing the oxidation potential (for example by reducing the CO values to below 4%). Subsequently the alloy elements oxidize on the iron-/iron oxide boundary. In the subsequent RTF the iron oxide layer is reduced under protective atmosphere (N2/H2).
 Because silicone (as well as Manganese and Chromium) are less noble than iron, the furnace atmosphere may be reducing for iron but oxidizing for the alloy elements. The oxides of the alloy elements remain in the RTF in the former boundary iron/iron oxide, i.e., significantly below the steel strip surface (inner oxidation).
 Thus the surface spreading of oxides is prevented and a uniform wettability of the strip surface is achieved. At the same time the dew point in the region of the transition zone furnace→zinc pot (tuyere snout) is to be selected (preferably below -30° C.) so that an oxidation of the strip prior to dipping into the zinc bath is avoided. Advantageous are dew point of -35 or -40° C.
 In a configuration of a plant in which annealing is carried out only with a radiant-tube furnace (see method 3 in FIGS. 2, 4, 8) the inner oxidation can be promoted by a slight increase of the oxygen content in the furnace atmosphere. This is achieved by controlling the dew point (preferably >30° C. advantageously -25 or -20). As a result of the higher partial pressure of the oxygen increased diffusion of oxygen into the steel strip can take place and oxidize the alloy elements. On the other hand if only little oxygen is present in the atmosphere, the less noble elements diffuse to the strip surface where they form unwettable oxides. Also in this case the oxidation of the iron in the RTF and in the region of the tuyere snout should be avoided.
 When instead of the hot dip galvanizing the method route of the continuous annealing with subsequent electrolytic galvanizing is selected, no special measures are required to ensure the galvanizing ability. It is known that the galvanizing of higher alloyed steels is significantly easier to realize by electrolyte galvanizing than by continuous hot dip galvanizing. In electrolytic galvanizing pure zinc is deposited directly n the strip surface. In this case it only has to be ensured that no surface covering oxide layer is present on the strip surface in order to not inhibit migration of the electrons to the zinc ions. This condition is usually ensured by a standard reducing atmosphere during the annealing and a pre-cleaning prior to electrolysis.
 In order to ensure process window during the annealing that is as wide as possible and a sufficient galvanizing capacity the minimal Si-content is set to 0.600% and the maximal silicone content to 0.800%.
 Manganese (Mn) is added to almost every steel for de sulfurization in order to convert the deleterious sulfur into manganese sulfides. In addition as a result of solid solution strengthening, manganese increases the strength of the ferrite and shifts the α/γ transformation toward lower temperatures.
 A main reason for adding manganese in dual-phase steel is the significant improvement of the hardness penetration. Due to the diffusion impairment the perlite and bainite transformation is shifted toward longer times and the martensite start temperature is lowered.
 At the same time however, addition of manganese increases the hardness ration between martensite and ferrite. In addition the microstructure banding is increased. A high hardness difference between the phases and the formation of martensite bands results in a lower hole expansion capacity, which has an adverse affect on the edge crack resistance.
 Like silicone, manganese tends to form oxides on the steel surface during the annealing treatment. Depending on the annealing parameters and the content of other alloy elements (in particular Silicone and Aluminum), manganese oxides (for example MnO) and/or Mn mixed oxides (for example Mn2SiO4) may form. However, manganese is less critical at a low Si/Mn or Al/Mn ratio because rather globular oxides instead of oxide films form. Nevertheless high Manganese contents may negatively influence the appearance of the zinc layer and the zinc adhesion.
 For the stated reasons the Mn-content is set to 1.000 to 2.250%.
 For achieving the demanded minimal strengths the utilized it is manganese contents are advantageously differentiated based on the cross section. For a strip thickness range of <1.0 mm the Manganese content is preferably ≦1.50%, at strip thicknesses of 1 to 2 mm at ≦1.75%, and at strip thicknesses of >2 mm at ≦1.50%.
 Chromium (Cr) in solubilized form chromium already in small amounts significantly increase the hardenability of steel. On the other hand Chromium causes a precipitation hardening at corresponding temperature profile in the form of chromium carbides. The increase in dual-phase steels the addition of chromium mainly improves the hardness penetration. In the solubilized form chromium shifts the perlite and bainite transformation toward longer times and thereby at the same time lowers the martensite start temperature.
 In dual phase steels addition of chromium mainly improves the hardness penetration. Chromium shifts in the solubilized state the perlite and bainite transformation toward longer times and at the same time lowers the martensite start temperature.
 A further important effect is that chromium significantly increases the tempering resistance so that almost no strength losses occur in the zinc dip bath.
 In addition chromium is a carbide former. When chromium is present in the carbide form the austenizing temperature has to be selected high enough prior to the hardening in order to solubilize the chromium carbides. Otherwise the increased number of nuclei may lead to an impairment of the hardness penetration.
 Chromium also tends to form oxides on the steel surface during the annealing treatment, which may negatively affect the galvanization quality.
 The Cr content is therefore set to values of 0.280 to 0.480%.
 For meeting the demanded mechanical properties, the sum amount of Mn+Si+Cr is also to be selected depending on the sheet thickness. Advantageous at sheet thicknesses of ≦1 mm is a sum amount of ≧1.88 to ≦2.60%, at sheet thicknesses of 1 to 2 mm a sum amount of ≧2.20 to ≦3.00% and at sheet thicknesses ≧2 mm a sum amount of ≧2.50 to ≦3.53%.
 Molybdenum (Mo): similar to chromium, addition of molybdenum improves the hardenability. The perlite and baininte transformation is shifted toward longer times and the martensite start temperature is lowered.
 Molybdenum also significantly increases the tempering resistance so that no strength losses are to be expected in the zinc bath and causes an increase in strength of the ferrite as a result of solid solution strengthening.
 For reasons of cost Mo is therefore not added in the strength range of at least 580 MPa. The content of molybdenum is limited to unavoidable steel accompanying amounts.
 Copper (Cu): the addition of copper can increase the tensile strength and the hardness penetration. In connection with nickel, chromium and phosphorous copper can form a protective oxide layer on the surface, which significantly reduces the corrosion rate.
 In connection with oxygen copper can form deleterious oxides at the grain boundaries, which can have negative consequences in particular for hot forming processes. The copper content is therefore limited to amounts that are unavoidable during steel production.
 The contents of other alloy elements such as Nickel (Ni) or Tin (Sn) are limited to amounts that are unavoidable during the steel production.
 Micro-alloying elements are usually only added in very low amounts (<0.1%). In contrast to the alloying elements they are effective mainly through forming precipitations however they can also influence the properties in the solubilized state. In spite of the low added amounts, the micro-alloying elements strongly influence the production conditions such as processing and final properties.
 Aluminum (Al) is usually added to the steel in order to bind oxygen and nitrogen solubilized in the iron. In this way, oxygen is converted into aluminum oxides and aluminum nitrides. These precipitations can cause a grain refinement via increasing the nucleation sites and thus increase the tenacity and strength values.
 Aluminum nitride is not precipitated when Titanium is present in sufficient amounts. Titanium nitrides have a lower formation enthalpy and are formed at higher temperatures.
 In the solubilized state aluminum, like silicone, shifts the ferrite formation toward shorter times and thus enables the formation of sufficient amounts of ferrite in the dual-phase steel. In addition it suppresses the carbide formation and leads thus to a delayed transformation of the austenite. For this reason Al is also used as alloy element in residual austenite steels in order to substitute for a portion of the silicone by aluminum. The reason for this approach is that Al is less critical for the galvanization reaction than silicone.
 The Al-content is therefore limited to 0.01 to maximally 0.060% and is added for deoxidizing the steel.
 Niobium (Nb): For reasons of cost Niobium is not added and contents are limited to unavoidable amounts during steel production.
 Titanium (Ti): because in the present alloy concept addition of titanium is not required, the content of titanium is limited to unavoidable steel accompanying amounts.
 Vanadium (V): because in the present alloy concept addition of vanadium is not required, the content of vanadium is limited to unavoidable steel accompanying amounts.
 Boron (B): because in the present alloy concept addition of boron is not required, the content of boron is limited to unavoidable steel accompanying amounts.
 Tests conducted with the steel according to the invention have shown that with the present alloy concept a dual-phase steel with a minimal tensile strength of 580 MPa can be produced in a thickness range of 0.50 to 4.00 mm which is characterized by a sufficient tolerance with regard to process fluctuations by an inter-critical annealing between Ac1 and Ac3 or an austenizing annealing above Ac3 with final controlled cooling.
 With this, a significantly widened process window is established for the alloy composition according to the invention compared to known alloy concepts.
 The annealing temperatures for the dual-phase microstructure to be achieved are between about 700 and 950° C.; depending on the temperature range this achieves a partially austenitic (dual-phase region) microstructure or a fully austenitic microstructure (austenitic region).
 The tests show that the established microstructure proportions after the inter-critical annealing between AC1 and AC3 or the austenizing annealing above AC3 with subsequent controlled cooling are maintained also after a further process step "hot dip galvanizing" at temperatures between 420 to 470° C. for example in the case of zinc or zinc-magnesium.
 The hot dip coated material can be manufactured as hot strip as well as cold re-rolled hot strip or cold strip in the skin passed rolled (cold re-rolled) or non skin pass rolled state and/or stretch leveled or not stretch leveled state.
 Steel strips, in the present case as hot strips, cold re-rolled hot strip or cold strip made from the alloy composition according to the invention, are in addition characterized by a high resistance against edge-proximate crack formation during the further processing.
 The small differences in the characteristic values of the steel strip longitudinally and transversely to its rolling direction are advantageous in the subsequent material insertion, which as a result may be transversely, longitudinally and diagonally to the rolling direction.
 In order to ensure the cold rollability of a hot strip produced from the steel according to the invention, the hot strip is produced according to the invention with final rolling temperatures in the austenitic range above AC3 and coiling temperatures above the bainite start temperature.
 The following examples are representative for the industrial production for the hot dip coating according to process 7b (corresponds to method 2):
Cold Re-Rolled Hot Strip According to FIG. 7c
 A steel according to the invention with 0.101% C; 0.605% Si; 1.374% Mn; 0.327% Cr; 0.039% Al; 0.012% P; 0.0035% Nb; 0.003% Mo; 0.0063% N; 0.0009% S was melted in a converter steel plant, hot rolled at a final rolling temperature of 911 C and coiled at a coiling temperature of 484 C with a thickness of 2.37 mm. After pickling with hydrochloric acid the cold rolling occurred on a five-stand tandem path with a cold rolling degree of 16% from 2.37 mm to 1.99 mm.
 In a hot dip galvanizing system the steel was heated with about 6K/s to about 847 C according to FIG. 7c. A slow cooling out of the dual phase region occurred with about 2K/s to about 720° C., subsequently a fast cooling was performed with about 22K/s to abut 35=60° C. The strip was reheated prior to reaching the tuyere snout to about 450° C. After the hot dip galvanizing at about 450° C. the strip was cooled to room temperature with about 20K/s at ambient air. A cold re-rolling (skin passing) occurred inline with about 0.2%.
 The steel according to the invention has a microstructure after the heat treatment which consists of ferrite, martensite, bainite and residual austenite.
 This steel exhibits the following characteristic values (the values in brackets are transverse values according to EN 10346 and longitudinal values according to VDA239-100):
TABLE-US-00001 Proof strength transverse direction 419 MPa (340 Pa-420 MPa) (Rp0.2) longitudinal direction 408 Ma (330 MPa-430 MPa) Tensile transverse direction 646 MPa (min.-600 MPa) strength (Rm) longitudinal direction 636 MPa (590 MPa-700 MPa) Elongation at transverse direction 23.6% (min. 20%) fracture longitudinal direction 27.0% (min. 20%) n-value transverse direction 0.18 (min. 0.14) n-value longitudinal direction 0.19 (min. 0.14) Bake hardening transverse direction 77 MPa (min. 30 MPa) BH2 longitudinal direction 77 MPa (min. 30 MPa) Hole expansion according to ISO 16630 44% Planar Δr -0.14 (-0.15 to +0.15 = anisotropy quasi isotropic)
Cold Strip According to FIG. 7b
 A steel according to the invention with 0.101% C; 0.605% Si; 1.374% Mn; 0.327% Cr; 0.039% Al; 0.012% P; 0.0035% Nb; 0.003% Mo; 0.0063% N; 0.0009% S was melted in a converter steel plant, hot rolled in a hot rolling line at a final rolling temperature of 902 C and coiled at a coiling temperature of 676 C with a thickness of 2.02 mm. After pickling with hydrochloric acid the cold rolling occurred on a five-stand tandem path with a cold rolling degree of 50% from 2.02 mm to 0.99 mm.
 In a hot dip galvanizing system the steel was heated with about 54K/s to about 781 C according to FIG. 7b and subsequently further heated to about 890° C. with about 5K/s.
 A slow cooling out of the dual phase region occurred with about 1K/s to about 860° C., subsequently a fast cooling was performed with about 23.2K/s to abut 465° C. The strip reached the tuyere snout with about 465° C. After the hot dip galvanizing at about 450° C. zinc bath temperature the strip was cooled to room temperature with about 34.3K/s at ambient air. A cold re-rolling (skin passing) occurred inline with about 0.2%.
 The steel according to the invention has a microstructure after the heat treatment which consists of ferrite, martensite, bainite and residual austenite.
 This steel has the following characteristic values
TABLE-US-00002 Proof strength transverse direction 4382 MPa (340 Pa-420 MPa) (Rp0.2) longitudinal direction 380 Ma (330 MPa-430 MPa) Tensile transverse direction 657 MPa (min.-600 MPa) strength (Rm) longitudinal direction 654 MPa (590 MPa-700 MPa) Elongation at transverse direction 24.6% (min. 20%) fracture longitudinal direction 24.2% (min. 20%) n-value transverse direction 0.17 (min. 0.14) n-value longitudinal direction 0.17 (min. 0.14) Bake hardening transverse direction 56 MPa (min. 30 MPa) BH2 longitudinal direction 60 MPa (min. 30 MPa) Hole expansion according to ISO 16630 49% Planar Δr -0.06 (-0.15 to +0.15 = anisotropy quasi isotropic)
 The ultimate yield ratio Re/Rm is at 58% in longitudinal as well as in transverse direction (relative to the rolling direction)
Hot Strip According to FIG. 7b
 A steel according to the invention with 0.101% C; 0.605% Si; 1.374% Mn; 0.327% Cr; 0.039% Al; 0.012% P; 0.0035% Nb; 0.003% Mo; 0.0063% N; 0.0009% S was melted in a converter steel plant, hot rolled in a hot rolling line at a final rolling temperature of 916 C and coiled at a coiling temperature of 485° C. with a thickness of 2.02 mm.
 After pickling with hydrochloric acid no cold rolling occurred but a hot strip galvanizing. In a hot dip galvanizing system according to FIG. 7b the steel was heated to about 785° C. with about 25K/s and subsequently further heated to about 893 C with about 1 K/s.
 A slow cooling out of the dual phase region occurred with about 1 K/s to about 860° C., subsequently a fast cooling was performed with about 10K/s to abut 470° C. The strip reached the tuyere snout with about 465° C. After the hot dip galvanizing at about 450° C. zinc bath temperature the strip was cooled to room temperature with about 15K/s at ambient air. A cold re-rolling (skin passing) occurred inline with about 0.2%.
 The steel according to the invention has a microstructure after the heat treatment which consists of ferrite, martensite, bainite and residual austenite.
 This steel has the following characteristic values
TABLE-US-00003 Proof strength transverse direction 362 MPa (340 Pa-420 MPa) (Rp0.2) longitudinal direction 363 Ma (330 MPa-430MPa) Tensile transverse direction 620 MPa (min.-600 MPa) strength (Rm) longitudinal direction 622 MPa (590 MPa-700 MPa) Elongation at transverse direction 24.3% (min. 20%) fracture (A80) longitudinal direction 24.5% (min. 20%) n-value transverse direction 0.17 (min. 0.14) n-value longitudinal direction 0.17 (min. 0.14) Bake hardening transverse direction 58 MPa (min. 30 MPa) BH2 longitudinal direction 59 MPa (min. 30 MPa) Hole expansion according to ISO 16630 41% Planar Δr -0.08 (-0.15 to +0.15 = anisotropy quasi isotropic)
 The ultimate yield ratio Re/Rm is at 58.4% in longitudinal as well as in transverse direction (relative to the rolling direction)
 In tests the characteristic values shown in the following examples were achieved for the hot dip galvanizing according to FIG. 7c.
Hot Strip and Cold Re-Rolled Hot Strip According to FIG. 7c
 The steel according to the invention of example 2 (coiling temperature 676° C.) and of example 3 (coiling temperature 485° C.) was further processed after the pickling under approximately operating conditions. The cold rolling was performed in a test cold rolling system. The tested cold rolling degrees were 0% and 10%. The hot dip galvanizing cycle was simulated with an annealing simulator according to FIG. 7c.
 Depending on the strip thickness different heating and cooling rates were provide wherein the set temperatures were selected independent of the sample thickness. The steel was heated to 860° C., then first slowly cooled to 720° C. before the fast cooling with to 350° C. was initiated. Thereafter the samples were heated to 450° C. before they were cooled to room temperature with maximally 24K/s. The samples were not skin passed. The mechanical characteristic values were determined in longitudinal direction. The results of the tensile tests in longitudinal direction relative to the rolling direction are shown in FIG. 8.
 Further features, advantages and details of the invention will become apparent from the following description of exemplary embodiments shown in the drawing.
 It is shown in:
 FIG. 1: schematically the process chain (schematic) for the production of a strip from a steel according to the invention
 FIG. 2: results of a hole expansion test according to ISO 16630 (sheet thickness 1 mm) exemplary for the steel according to the invention relative to an hyper-peritectic, phosphorous alloyed and micro-alloy-free standard grade.
 FIG. 3: example for analytical differences of the steel according to the invention relative to an hyper-peritectic, phosphorous alloyed and microalloy-free standard grade
 FIG. 4: Example for mechanical characteristic values (transversely to the rolling direction) of the steel according to the invention compared to an hyper-peritectic, phosphorous alloyed and microalloy-free standard grade.
 FIG. 5: time temperature course (schematically) of the process steps hot rolling and continuous annealing, exemplary for the steel according to the invention.
 FIG. 6: ball impact test according to SEP1931.
 FIG. 7a, b, c: time temperature curve (annealing variants schematic).
 FIG. 8: mechanical properties (longitudinal probes) of the samples of a steel according to the invention annealed in the laboratory or cold rolled and annealed.
 FIG. 1 shows schematically the process chain for producing the steel according to the invention. Shown are the different process routes with regard to the invention. Up to position 5 (pickling) the process route is the same for all steels according to the invention, thereafter divergent process routes are followed depending on the desired results. For example the pickled hot strip can be galvanized or cold rolled and galvanized. Or it can be soft annealed, cold rolled and galvanized.
 FIG. 2 shows results of a hole expansion test (values relative to the standard grade). Shown are the results of the hole expansion test for a steel according to the invention compared to the standard grade, as reference serves standard grade process 2.
 The materials have a sheet thickness between 1 or 2 mm. The results apply to the test according to ISO 16630. It can be seen that the steels according to the invention achieve better expansion values in the case of punched holes than the standard grades with same processing. Process 2 corresponds hereby to an annealing for example a hot dip galvanization with combined directly fired furnace and radiant tube furnace, as described in FIG. 7b.
 The different temperature courses according to the invention within the mentioned range, result in different characteristic values or also in different hole expansion results which are both significantly improved compared to the standard grades. Principal differences are thus the temperature-time parameters in the heat treatment and the downstream cooling.
 FIG. 3: shows the relevant alloy elements of the steel according to the invention compared to the standard grade, which exemplifies the state of the art. In the comparison steel (standard grade) the main difference is in the carbon content, which lies in the hyper-peritectic range, but also in the elements silicone and chromium. In addition the standard grade is micro-alloyed with phosphorous. The steel according to the invention is clearly silicone alloyed.
 FIG. 4a: shows the mechanical characteristic values transversely and longitudinally to the rolling direction compared to the standard grade. All characteristic values, which were achieved by annealing in the dual phase region, correspond to the normative guidelines of a HCT600X or HDT580X.
 FIG. 5: schematically shows the time temperature course of the process steps hot rolling and continuous annealing of strips made of the alloy composition according to the invention. Shown is the time and temperature dependent transformation for the hot rolling process as well as for a heat treatment after the cold rolling.
 FIG. 6 shows the positive result of the performed ball impact test (according to SEP1931) at a hot dip galvanized material made from the steel according to the invention.
 The FIG. 7 show schematically three variants of the temperature time courses according to the invention during the annealing and cooling and respective different austenization conditions.
 The method 1 (FIG. 7a) shows the annealing and cooling of the produced cold or hot rolled steel strip in a continuous annealing system. First the strip is heated to a temperature in the range of about 700 to 950° C. The annealed steel strip is then cooled from the annealing temperature to an intermediate temperature of about 200 to 250° C. with a cooling rate between about 15 and 100° C./s. A second intermediate temperature (about 300 to 500° C.) is not shown in this schematic representation. Subsequently the steel strip is cooled to room temperature at air with a cooling rate between about 2 and 30° C./s or the cooling with a cooling rate between about 15 and 100° C./s is maintained until reaching room temperature.
 The method 2 (FIG. 7b) shows the process according to the method 1, however the cooling of the steel strip is shortly interrupted for the purpose of a hot dip galvanization during passage through the hot dip container in order to subsequently continue the cooling with a cooling rate between about 15 and 100° C./s until reaching an intermediate temperature of about 200 to 250° C. subsequently the steel strip is cooled at air with a cooling rate between about 2 and 30° C./s until reaching room temperature.
 The method 3 (FIG. 7c) also shows the process according to the method 1 in the case of a hot dip galvanization, wherein however the cooling of the steel strip is interrupted by a short break (about 1 to 20 s) at an intermediate temperature in the range of about 200 to 400° C. and reheated to the temperature which is required for the hot dip galvanization (about 420 to 470° C.). Subsequently the steel strip is again cooled to an intermediate temperature of about 200 to 250° C. The steel strip is then finally cooled with a cooling rate of about 2 and 30° C./s at air until reaching room temperature.
 FIG. 8 shows the mechanical characteristic values of steel according to the invention (example 4), which was annealed or cold rolled and annealed according to the method 3 (FIG. 7c)
 The values are average values of two longitudinal samples, which were determined in the tensile test. The low values for the yield strength can be attributed to the fact that the samples are not skin passed.
LIST OF REFERENCE SIGNS
 1. furnace process
 2. secondary metallurgy
 3. strip casting
 4. hot rolling
 5. pickling
 6. soft annealing hot strip (optional)
 7. cold rolling
 8. multi roller (optional)
 9. soft annealing cold strip (optional)
 10. hot dip galvanizing
 11. inline skin passing
 12. stretch bending alignment
Patent applications by Thomas Schulz, Salzgitter DE
Patent applications in class Titanium, zirconium or niobium containing
Patent applications in all subclasses Titanium, zirconium or niobium containing