EP2668302B1 - Method for manufacturing a strip from a high strength multi-phase steel having excellent forming properties - Google Patents

Description

The invention relates to a method for producing a cold-rolled or hot-rolled steel strip from a high-strength multiphase steel having excellent forming properties according to claim 1.

The hotly contested automotive market is forcing manufacturers to constantly seek solutions to reduce fleet consumption while maintaining maximum comfort and occupant safety. On the one hand, the weight saving of all vehicle components plays a decisive role, on the other hand, but also a most favorable behavior of the individual components with high static and dynamic stress during operation as well as in the event of a crash. The suppliers of raw materials are trying to meet this need by providing high-strength and ultra-high-strength steels that can reduce the thickness of the vehicle while at the same time improving forming and component behavior during production and operation. These steels must therefore meet comparatively high demands with regard to their strength and ductility, energy absorption and during their processing, for example during stamping, hot and cold forming, welding and / or surface finishing (eg metallically finished, organically coated).

In addition to the required weight reduction, newly developed steels must thus meet the high material requirements for yield strength, tensile strength and elongation at break with good formability, as well as the component requirements for high toughness, edge crack resistance, energy absorption and strength via the work hardening effect and the bake hardening effect.

In vehicle construction, therefore, dual-phase steels are increasingly being used, which consist of a ferritic basic structure in which a martensitic second phase and possibly a further phase with bainite and retained austenite are incorporated.

The steel grades determining processing characteristics of dual-phase steels, such as a very low yield ratio and at the same time very high tensile strength, a strong work hardening and a good cold workability, are well known.

Multiphase steels are also used with increasing tendency, such as complex-phase steels, ferritic-bainitic steels, bainitic steels and martensitic steels, which are characterized by different microstructural compositions.

Hot or cold rolled complex-phase steels are steels which contain small amounts of martensite, retained austenite and / or pearlite in a ferritic / bainitic matrix, whereby extreme grain refinement is caused by retarded recrystallization or precipitation of micro-alloying elements.

Hot-rolled ferritic-bainitic steels are steels that contain bainite or solidified bainite in a matrix of ferrite and / or solidified ferrite.

The solidification of the matrix is effected by a high dislocation density, by grain refining and the excretion of micro-alloying elements.

Hot rolled or cold rolled bainitic steels are steels 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 amounts of other phases, such as. As martensite and ferrite.

Hot-rolled martensitic steels are steels that contain small amounts of ferrite and / or bainite in a matrix of martensite due to thermomechanical rolling. The 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.

These steels are used, among other things, in structural, chassis and crash-relevant components, as well as flexibly cold-rolled strips. These T Ailor R olled B lank lightweight technology (TRB®) allows a significant weight reduction by the load customized choice of thickness greater than the length of the component.

However, the production of TRB®s with multi-phase structure is not without limitations with today's known alloys and available continuous annealing plants for widely varying sheet thicknesses such. B. for the heat treatment before cold rolling, possible. 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.

Usually, cold-rolled steel strips are annealed by recrystallization in a continuous annealing process to form a thin sheet which is easy to form. 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.

To set the dual-phase microstructure, the hot or cold strip is heated in a continuous annealing furnace to a temperature such that the required microstructure formation occurs during cooling. The same applies to the setting of a steel with a complex phase structure, martensitic, ferritic-bainitic and purely bainitic structure.

If, due to high corrosion protection requirements, the surface of the hot or cold strip is to be hot dip galvanized, the annealing treatment is usually carried out in a continuous annealing furnace upstream of the galvanizing bath.

Depending on the alloy concept, the required microstructure of the hot strip may also be adjusted in the continuous furnace during the annealing treatment, in order to achieve the required mechanical properties.

In the continuous annealing of hot or cold rolled steel strips with z. B. from the scriptures EP 0 152 665 B1 . EP 0691 415 B1 . EP 0510 718 B1 EP 1154028A1 or WO 2010 / 126161A1 known alloy concepts for dual-phase steels has the problem that only a narrow process window for the annealing parameters is present in order to ensure uniform mechanical properties over the strip length in cross-sectional jumps without adjustment of the process parameters.

In this case, a narrow process window means that, depending on the cross section of the strip to be annealed, the process parameters have to be adjusted in order to achieve the required microstructure and homogeneity through homogeneous temperature distribution in the strip and during cooling achieve mechanical and technological properties. With enlarged process windows, the required strip properties are possible with the same process parameters even with different cross sections of the strips to be annealed.

In addition to flexibly rolled strips with different sheet thicknesses over the tape length often bands with different thicknesses, z. B. with 1.5 and 2.0 mm and / or different width, such as. B. 900 and1400 mm, are annealed successively.

A homogeneous temperature distribution is difficult to achieve, especially at different thicknesses in the transition region from one belt to another. This can lead to alloy compositions with too small process windows in the continuous annealing that z. B. the thinner strip is either driven too slowly through the oven and thereby productivity is lowered, or that the thicker strip is driven too fast through the oven and the required annealing temperature and thus the required structure is not achieved. The consequences are increased rejects or even customer complaints.

The decisive process parameter is thus the setting of the speed in the continuous annealing, since the phase transformation is temperature- and time-controlled. The less sensitive the steel is to the uniformity of mechanical properties with changes in temperature and time in continuous annealing, the greater the process window.

Particularly serious is the problem of a too narrow process window in the annealing treatment when load-optimized components are to be made of hot or cold strip, which have over the tape length and bandwidth varying sheet thicknesses, eg. B. have been rolled flexibly.

A method for producing a steel strip with different thickness over the strip length is z. B. in the DE 100 37 867 A1 described.

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.

For flexibly rolled cold strips of steels of known compositions, because of the too small process window, the areas with lower sheet thickness due to the transformation processes during cooling either too high strengths due to excessive martensite or on areas with larger sheet thickness to low strengths due to low martensite. Homogeneous mechanical-technological properties over the strip length or width are virtually impossible to achieve with the known alloy concepts in continuous annealing.

The goal of achieving the final mechanical-technological properties in a narrow range over bandwidth and strip length by the controlled adjustment of the volume fractions of the structural phases has the highest priority and is therefore only possible through an enlarged process window. The known alloy concepts for multiphase steels are characterized by too narrow a process window and therefore unsuitable for solving the present problem, in particular in flexibly rolled strips. Currently, only steels of a strength class with defined cross-sectional areas can be represented with the known alloy concepts, so that altered alloy concepts are necessary for different strength classes and / or cross-sectional areas.

The invention is therefore based on the object to provide a different alloy concept for a high-strength multiphase steel with different microstructural compositions, with which the process window for the continuous annealing of hot or cold strips can be extended so that in addition to bands with different cross sections and steel bands with over tape length and, if necessary Bandwidth varying thickness can be produced with the most homogeneous mechanical and technological properties. Furthermore, an alloy concept is to be specified, with which different strength classes can be served.

According to the teachings of the invention, this object is achieved by a steel with the following contents in% by weight: C 0.060 to ≤ 0.115 al 0.020 to ≤ 0.060 Si 0.100 to ≤ 0.500 Mn 1,300 to ≤ 2,500 P ≤ 0.025 S ≤ 0.0100 Cr 0.280 to ≤ 0.480 Not a word ≤ 0.150 Ti ≥ 0.005 to ≤ 0.050 Nb ≥ 0.005 to ≤ 0.050 B ≥ 0.0005 to ≤ 0.0060 N ≦ 0.0100 The rest of the iron, including standard steel-accompanying elements not mentioned above.

The steel according to the invention offers the advantage of a significantly enlarged process window in comparison to the known steels. This results in an increased process reliability in the continuous annealing of cold and hot strip with multi-phase structure. Thus, homogenized mechanical-technological properties in the strip can be ensured for pass-annealed hot or cold strips even with different cross sections and otherwise identical process parameters.

This applies to the continuous annealing of successive belts with different belt cross-sections as well as belts with varying sheet thickness over belt length or belt width. For example, a processing in selected thickness ranges, such as smaller 1 mm strip thickness, 1 to 2 mm strip thickness and greater than 2 mm strip thickness is possible.

If, according to the invention, higher-strength hot or cold strips produced from multiphase steel with varying sheet thicknesses are produced in the continuous annealing process, stress-optimized components can advantageously be produced by deformation from this material.

The material produced can be produced both as a cold strip and as a hot 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 (intermediate annealing).

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 can be produced.

The steel strips produced with the alloy composition according to the invention are characterized in the production of a multi-phase or bainitic steel by a significantly wider process window in terms of temperature and flow rate at the intercritical annealing between A _c1 and A _c3 or austenitizing annealing over A _c3 with final controlled cooling compared to the known alloy concepts.

Annealing temperatures of 700 to 950 ° C. and cooling rates of 15 to 100 ° C./s have been found to be advantageous up to a temperature of 420 to 470 ° C. with holding at an intermediate temperature of 200 to 250 ° C. and also with optional reheating upstream, with which the required multiphase structures can be set uniformly over the strip length. This has a particularly advantageous effect in the annealing of flexibly rolled strips or in the successive annealing of strips of different cross sections, so that very uniform material properties are thereby achieved.

The basis for achieving a broad process window is the inventive combination of the micro-alloying elements titanium, niobium and boron with optional addition of molybdenum.

Fine titanium precipitates work in the same way as niobium carbides and together enhance the effect. Titanium binds off the nitrogen, which is therefore no longer available for the formation of boron nitride, whereby the boron alloy can act. In this case, the addition of boron, which is free, causes an increase in the hardenability.

Boron is one of the elements that is characterized not only by a high degree of hardening but also by a high hardening effect. The microstructure becomes more isotropic, because differences in the cooling rates caused by the process control or the geometry of the strip have less influence, which also leads to a larger process window

The free boron is capable of producing a relatively homogeneous microstructure (same microstructural proportions) over the sheet thickness. The same applies to the less pronounced influence of temperature gradients that occur over the length of the strip or in relation to its width.

In classical dual-phase steels, in addition to manganese, chromium and silicon, carbon is also responsible for the transformation of austenite to martensite. Boron therefore allows one to substitute part of the carbon. This also has a positive effect on the microstructure, as carbon is one of the most segregating elements in steel. As a result, segregations that lead to locally different thermodynamic driving forces are less pronounced, which in turn results in a higher degree of robustness with respect to process or geometry-related temperature fluctuations. Material characteristic is that the additional addition of titanium and boron in addition to niobium shifts the ferrite region very clearly at later times during cooling. This opens up the potential for complex-phase steels and bainitic steels.

Depending on the process parameters, the proportions of ferrite are more or less reduced by increased amounts of bainite. The combination of the three micro-alloying elements enables the material diversity described above.

Experiments have shown that the combination of niobium and boron alone is insufficient to achieve a broad process window and typically required tensile strength range of at least 750MPa for hot strip and at least 780MPa for cold rolled hot strip and cold strip. Only by the addition of titanium in the specified levels, this was possible.

By setting a low carbon content of ≤ 0.115%, the carbon equivalent can be reduced, thereby improving weldability and avoiding too much hardening during welding. In resistance spot welding, moreover, the electrode life can be significantly increased.

The carbide as well as the nitride formation of vanadium starts only at temperatures around 1000 ° C or even after the α / β-conversion, so much later than titanium and niobium. Vanadium thus has hardly any grain refining effect due to the small number of precipitates present in austenite. Even austenite grain growth is not inhibited by the late release of the vanadium carbides. Thus, the strength-enhancing effect is almost entirely due to precipitation hardening.

An advantage of vanadium is the high solubility in austenite and the large volume fraction of fine precipitates caused by the low precipitation temperature.

The effect of the elements in the alloy according to the invention is described in more detail below. The multiphase steels are typically chemically designed to combine alloying elements with and without micro-alloying elements. Accompanying elements complete the analysis concept.

Accompanying elements are elements that are already present in the iron ore or, due to their production, enter 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 that can diffuse through the iron lattice without creating lattice strains. As a result, the hydrogen in the iron grid is relatively mobile and can be absorbed relatively easily during production. 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 hydrogen-induced stress corrosion cracking. Even with the delayed crack, the so-called delayed-fracture, which occurs without external stresses, hydrogen is often mentioned as a reason.

Therefore, the hydrogen content in the steel should be as low as possible.

Oxygen (O) : In the molten state, the steel has a relatively high absorption capacity for gases, but at room temperature, 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 converted to safer conditions. So a setting of the oxygen over manganese, silicon and / or aluminum is usually common. However, the resulting oxides can cause negative properties as defects in the material. In a fine excretion, especially of aluminum oxides, but also a grain refinement can be done.

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

Nitrogen (N) is also a companion element of steelmaking. Steels with free nitrogen tend to have a strong aging effect. The nitrogen already diffuses at low temperatures at dislocations and blocks them. It causes an increase in strength combined with a rapid loss of toughness. Nitrogen bonding in the form of nitrides is possible by alloying aluminum or titanium.

For the abovementioned reasons, the nitrogen content is limited to ≦ 0.0100%, advantageously ≦ 0.0090% or optimally ≦ 0.0070% or unavoidable steel-accompanying amounts.

Like phosphorus, sulfur (S) is bound as a trace element in iron ore. It is undesirable in steel (except free-cutting steels), as it tends to segregate severely and has a strong embrittlement. It is therefore attempted to achieve the lowest possible amounts of sulfur in the melt (for example, by a deep vacuum treatment). 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. With diffusion-controlled transformation, this leads to a line-shaped structure and can lead to impaired mechanical properties in the case of pronounced bristleness (eg pronounced martensite parts instead of distributed martensite islands, no isotropic material behavior, reduced elongation at break).

For the above reasons, the sulfur content is limited to ≤ 0.0100% or unavoidable steel-accompanying quantities.

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 usually attempted to lower the phosphorus content as much as possible, since it is highly susceptible to segregation, among other things, by its low diffusion rate and greatly reduces the toughness. Due to the addition of phosphorus at the grain boundaries, grain boundary fractures usually 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.

In some steels, however, it is used as a micro-alloying element in small quantities (<0.1%) due to its low cost and high strength increase. Thus, in dual-phase steels, phosphorus is also used in part as a strength carrier.

For the aforementioned reasons, the phosphorus content is limited to ≤ 0.025% or unavoidable steel-accompanying amounts.

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.

For example, chromium in dissolved form can significantly increase the hardenability of steel even in small quantities. In the form of chromium carbides, it can bring about a direct increase in strength through particle hardening. By increasing the nucleation sites and by reducing the dissolved carbon content, however, the hardenability is reduced.

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. Its presence turns the iron into steel. Despite this fact, the carbon content is drastically lowered during steelmaking. For dual-phase steels for continuous hot-dip finishing, its proportion according to DIN EN 10346 is 0.23% depending on the quality, a minimum value is not specified.

Due to its small atomic radius, carbon is interstitially dissolved in the iron lattice. The solubility is a maximum of 0.02% in α-iron and a maximum of 2.06% in β-iron. Carbon in dissolved form considerably increases the hardenability of steel.

Due to the induced lattice strains in the dissolved state, diffusion processes are hindered and thus conversion processes are delayed. In addition, carbon favors the formation of austenite, thus expanding the austenite region to lower temperatures. With increasing positively dissolved carbon content, the lattice distortions and thus the strength values of the martensite increase.

Carbon is also required to form carbides. A representative occurring almost in every steel is the cementite (Fe3C). However, significantly harder special carbides with other metals such as chromium, titanium, niobium, 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. Therefore, to ensure sufficient strength on the one hand and good weldability on the other hand, the minimum C content is set to 0.060% and the maximum C content to 0.115%.

Silicon (Si) binds oxygen during casting, thus reducing segregation and impurities in the steel. In addition, silicon increases the strength and the yield strength ratio of the ferrite with solid solution hardening with only a slightly decreasing elongation at break. Another important effect is that silicon shifts the formation of ferrite to shorter times, thus allowing the formation of sufficient ferrite before quenching. The ferrite formation enriches the austenite with carbon and stabilizes it. In addition, silicon stabilizes austenite in the lower temperature range, especially in the area of bainite formation, by preventing carbide formation (no depletion of carbon).

During continuous galvanizing, silicon can diffuse to the surface during annealing and lead to silicon oxides there. During the immersion phase in the zinc bath, silicon oxides can interfere with the formation of a closed adhesion layer between steel and zinc (inhibiting layer). This manifests itself in a poor zinc adhesion and undigested places.

In addition, with high silicon contents during hot rolling, strongly adhering scale may form, which may impair the further processing.

For the above reasons, the minimum Si content is set at 0.100% and the maximum Si content at 0.500%.

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 / β conversion to lower temperatures.

One main reason for adding manganese into 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.

Manganese, similar to silicon, can lead to manganese oxides at high surface concentrations, which can adversely affect zinc adhesion and surface appearance.

The Mn content is therefore set at 1,300 to 2,500%.

Chromium (Cr) : In dual-phase steels, the addition of chromium mainly improves the hardenability. Chromium, when dissolved, shifts perlite and bainite transformation to longer times, while decreasing the martensite start temperature.

Another important effect is that chromium increases the tempering resistance considerably, so that almost no loss of strength occurs in the zinc bath.

Chromium is also a carbide former. If chromium is in carbide form, 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.

The Cr content is therefore set at values of 0.280 to 0.480%.

Molybdenum (Mo): The addition of molybdenum is similar to chromium to improve hardenability. The perlite and bainite transformation is pushed to longer times and the martensite start temperature is lowered.

In addition, molybdenum considerably increases the tempering resistance, so that no loss of strength is to be expected in the zinc bath and, as a result of solid solution hardening, increases the strength of the ferrite.

The Mo content is optionally added depending on the size, the equipment configuration and the microstructure setting, in which case the minimum addition should be 0.050% in order to have an effect. For cost reasons, the Mo content is limited to max. 0.150% set. Copper (Cu) : The addition of copper can increase the tensile strength and hardenability. In combination with nickel, chromium and phosphorus, copper can form a 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 in particular. The content of copper is therefore limited to unavoidable steel-accompanying quantities.

Other alloying elements such. As nickel (Ni) or tin (Sn) are limited in their contents to unavoidable, steel-accompanying amounts.

Microalloying elements are usually added only in very small amounts (<0.1%). They act in contrast to the alloying elements mainly by excretion formation but can also affect the properties in a dissolved state. Despite the small quantity additions, micro-alloying elements strongly influence the production conditions as well as the processing and final properties.

As micro-alloying elements, carbide and nitride formers which are generally soluble in the iron lattice are used. Formation of carbonitrides is also possible because of the complete solubility of nitrides and carbides in one another. The tendency to form oxides and sulfides is usually the most pronounced among the micro-alloying elements.

This property can be used positively by binding the generally harmful elements sulfur and oxygen. The setting may also have negative effects, if there are not enough micro-alloying elements for the formation of carbides available.

Typical micro-alloying elements are aluminum, vanadium, titanium, niobium and boron. These elements can be dissolved in the iron lattice and form carbides or nitrides with carbon and nitrogen because of a decrease in the free enthalpy.

Aluminum (Al) is usually added to the steel to bind the dissolved oxygen in the iron and nitrogen. The oxygen and nitrogen is thus converted into 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 therefore formed at higher temperatures.

When dissolved, aluminum, like silicon, shifts ferrite formation to shorter times, allowing the formation of sufficient ferrite in dual-phase steel. It also suppresses carbide formation and thus stabilizes austenite.

The Al content is therefore limited to 0.020 to a maximum of 0.060%.

Titanium (Ti) forms very stable nitrides (TiN) and sulfides (TiS2) even at high temperatures. These dissolve depending on the nitrogen content in part only in the melt. If the resulting precipitates are not removed with the slag, they form very coarse particles in the material due to the high formation temperature and are generally not conducive to the mechanical properties.

A positive effect on the toughness results from the setting of the free nitrogen and oxygen. For example, titanium protects other micro-alloying elements, such as niobium, from binding with nitrogen. These can then develop their effect optimally. Nitrides, which are formed by lowering the oxygen and nitrogen content only at lower temperatures, can also effectively inhibit austenite grain growth.

Unbonded titanium forms titanium carbides at temperatures above 1150 ° C and can thus effect grain refinement (inhibition of austenite grain growth, grain refinement by delayed recrystallization and / or increase in the number of nuclei in the case of α / β transformation) and precipitation hardening.

The Ti content therefore has values of more than 0.005 and less than 0.050%. Advantageously, Ti is limited to contents of ≦ 0.045 or ≦ 0.040%.

Niobium (Nb ) causes a strong grain refining because it most effectively retards recrystallization of all the micro-alloying elements and also inhibits austenite grain growth.

The strength-increasing effect is qualitatively higher than that of titanium due to the increased grain refining effect and the larger amount of strength-increasing particles (setting of the titanium to TiN at high temperatures).

Niobium carbides form from about 1200 ° C. In conjunction with titanium, which sets the nitrogen as already described, niobium can increase its strength-increasing effect by carbide formation in the lower temperature range (smaller carbide sizes).

Another effect of niobium is the retardation of the α / β conversion and the lowering of the martensite start temperature in the dissolved state. On the one hand this happens through the solute drag effect and on the other hand through the grain refining. This causes an increase in the strength of the structure and thus a higher resistance to expansion in the formation of martensite.

Limited is the use of niobium by the very low solubility limit. Although this limits the amount of excretions but mainly causes an early excretion with very coarse particles.

Precipitation hardening can thus become effective especially for steels with a low C content (greater supersaturation possible) and during hot forming processes (deformation-induced precipitation).

The Nb content is therefore limited to values between 0.005 and 0.050%, the maximum contents being advantageously restricted to ≦ 0.045 or ≦ 0.040%.

Vanadium (V) : The carbide as well as the nitride formation of vanadium starts only at temperatures of around 1000 ° C or even after the α / β transformation, which is much later than for titanium and niobium. Vanadium thus has hardly any grain refining effect due to the small number of precipitates present in austenite. Even austenite grain growth is not inhibited by the late release of vanadium carbides.

Thus, the strength-enhancing effect is almost entirely due to precipitation hardening. An advantage of vanadium is the high solubility in austenite and the large volume fraction of fine precipitates caused by the low precipitation temperature.

Since addition of vanadium is not necessary in the present alloy concept, the vanadium content is limited to unavoidable steel accompanying amounts.

Boron (B) forms nitrides or carbides with nitrogen as well as with carbon; however, this is usually not the goal. On the one hand, due to the low solubility, only a small amount of precipitates is formed and, on the other hand, these are usually concentrated on the Grain boundaries eliminated. An increase in hardness at the surface is not achieved (except boriding with formation of FeB (2) at the surface).

In order to prevent nitridation, it is usually attempted to bind the nitrogen by more affine elements. Nitrogen is in ascending order more affine to beryllium, aluminum, cerium, titanium and zirconium. Especially titanium can guarantee the setting of the entire nitrogen. Aluminum is less suitable.

Boron in the dissolved state in very small amounts leads to a significant improvement in hardenability. The mechanism of action of boron can be described as boron atoms preferentially attach to the grain boundaries and, by lowering the grain boundary energy, hinder diffusion and grain boundary slippage. In addition, the nucleation sites are reduced by reducing precipitation formation at the grain boundaries.

The effectiveness of boron is reduced with increasing grain size and increasing carbon content (> 0.8%). In addition, an amount exceeding 60 ppm causes a decrease in hardenability because boron carbides act as seeds on the grain boundaries.

Boron has a very high affinity for oxygen, which can lead to a lowering of the boron content in areas near the surface (up to 0.5 mm). In this connection, annealing at over 1000 ° C is not recommended. This is also recommended because boron can lead to a strong coarse grain formation at annealing temperatures above 1000 ° C.

For the above reasons, the B content is limited to values of 0.0005 to 0.0060%. However, these values are advantageously below 0.0050 or 0.0040%.

In addition, it was found in the tests that austenitizing annealing of a hot strip over A _{c3 can} achieve a complex phase steel with a minimum tensile strength of 750 MPa.

With an intercritical annealing between A _c1 and A _c3 and an austenitizing annealing over A _c3 with final controlled cooling, a multiphase steel strip with a thickness of 1 and 3mm was produced, which was characterized by a large tolerance to process variations and very uniform properties with the same process parameters had.

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

The annealing temperatures are between 700 and 950 ° C. for the steel according to the invention, so that a partially austenitic (two-phase area) or a fully austenitic structure (austenite area) is achieved, depending on the structure to be achieved (complex phase structure).

The experiments showed that the set microstructural fractions after the intercritical annealing between A _c1 and A _c3 and the austenitizing annealing via A _c3 with final controlled cooling, even after the process step "hot dip finishing" at temperatures between 420 to 470 ° C, for example at Z ( Zinc) and ZM (zinc-magnesium) were preserved.

The hot-dip coated material can be produced both as a hot strip, as a cold rolled hot strip or cold strip in the dressed (cold rolled) or stretch bent state (undressed).

Steel strips, in the present case as hot strip, cold-rolled hot strip or cold strip made of the alloy composition according to the invention, are further characterized by a high resistance to edge-near crack formation during further processing.

By a quasi-isotropy of the steel strip, it is also advantageously possible to use material 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 region above A _c3 and reeling temperatures above the recrystallization temperature.

Figur 1:

schematisch die Prozesskette für die Herstellung der erfindungsgemäßen Stähle,

Figur 2:

Ergebnisse des Lochaufweitungsversuches,

Figur 3:

Beispiele für analytische Unterschiede des erfindungsgemäßen Stahls gegenüber dem Stand der Technik,

Figur 4:

Beispiele für mechanische Kennwerte des erfindungsgemäßen Stahls im Vergleich mit dem Stand der Technik (Blechdicke t=2,0mm),

Figur 5:

schematisch den Zeit-Temperaturverlauf der Prozessschritte Warmwalzen und Durchlaufglühen,

Figur 6:

ZTU-Diagramm für einen erfindungsgemäßen Stahl,

Figur 7:

mechanische Kennwerte bei Variation der Abwalzgrade,

Figur 8:

Übersicht über die mit dem erfindungsgemäßen Legierungskonzept einstellbaren Festigkeitsklassen,

Figur 9:

Temperatur-Zeit-Kurve (schematisch).

Other features, advantages and details of the invention will become apparent from the following description of dargesteilten embodiments in a drawing. Show it:

FIG. 1:

schematically the process chain for the production of steels according to the invention,

FIG. 2:

Results of the hole expansion test,

FIG. 3:

Examples of analytical differences of the steel according to the invention over the prior art,

FIG. 4:

Examples of mechanical characteristics of the steel according to the invention in comparison with the prior art (sheet thickness t = 2.0 mm),

FIG. 5:

schematically the time-temperature curve of the process steps hot rolling and continuous annealing,

FIG. 6:

ZTU diagram for a steel according to the invention,

FIG. 7:

mechanical characteristics with varying degrees of rolling,

FIG. 8:

Overview of the strength classes that can be set using the alloy concept according to the invention,

FIG. 9:

Temperature-time curve (schematic).

FIG. 1 shows schematically the process chain for the production of the steels according to the invention. Shown are the different process routes relating to the invention. Up to position 5 (pickling), the process route is the same for all steels according to the invention, after which the corresponding processing takes place according to the desired results. For example, the pickled hot strip can be galvanized or cold rolled and galvanized. Or it can be annealed cold-rolled and galvanized.

FIG. 2 shows results of the hole expansion test (relative values in comparison). Shown are the results of the hole widening tests for a steel according to the invention compared to the standard grades. All materials have a sheet thickness of 2.00mm. The left panel shows the results for the test ISO TS 16630, right the results for the KWI test (Kaiser Wilhelm Institut). It can be seen that regardless of the type of processing, the steels according to the invention achieve the best expansion values for punched holes. Process 1 here corresponds to annealing, for example, on a hot-dip galvanizing with a combined directly fired furnace and radiant tube furnace. Process 2 corresponds, for example, to process control in a continuous annealing plant. In addition, a reheating of the steel can optionally be achieved directly in front of the zinc bath by means of an induction furnace. Due to the different sensed temperature guides within the specified range, different characteristic values or also different hole expansion results result, which are significantly improved for both processes compared to the standard grades. The basic difference is thus the temperature-time parameters during the heat treatment and the subsequent cooling.

FIG. 3 shows the relevant alloying elements of the steel according to the invention compared to steels of the same quality which correspond to the prior art. In the prior art steels, the main difference lies in the Carbon content that is in the overperature range. Occasionally there are steels that are individually micro-alloyed with Nb, Ti and B, but not in this combination.

FIG. 4 shows the mechanical characteristics of the steel according to the invention in comparison with those of the prior art. All characteristic values correspond to the normative specification.

FIG. 5 schematically shows the time-temperature curve of the process steps hot rolling and continuous annealing of strips of the alloy composition according to the invention. Shown is the time- and temperature-dependent conversion for the hot rolling process as well as for a heat treatment after cold rolling. Of particular interest here is the shift of the ferrite nose at later times. This opens up the potential for complex-phase steels and bainitic steels.

FIG. 6 shows a ZTU diagram for a steel according to the invention. Here, the determined ZTU diagram with the corresponding chemical composition and the A _c1 and A _c3 temperature is shown. By setting appropriate temperature-time profiles during cooling, it is advantageously possible to set a broad spectrum of microstructural compositions in the steel material.

FIG. 7 shows the mechanical characteristics with the same parameters of continuously annealed strips with varying degrees of rolling or different workpiece thickness. Shown are the characteristics of tensile strength, yield strength and elongation at break depending on selected degrees of reduction. Only the tensile strength increases slightly with increasing Abwalzgrad. All values are in the range of the standard for a HCT780XD and show that even with different sheet thicknesses after the continuous annealing practically identical mechanical properties are present.

FIG. 8 shows an overview of the adjustable with the alloy concept of strength classes. The alloy composition used corresponds to that in the FIG. 4 shown. Shown are the differently processed steel strips with their characteristic values and microstructural compositions. This makes clear the wide range of adjustable strength classes for hot and cold strip with the resulting structural components depending on the process steps performed and the set process parameters.

FIG. 9 shows schematically the temperature-time profiles in the annealing and cooling with 3 different variants.

Variant 1 ( FIG. 9a ) shows the annealing and cooling of the produced cold or hot rolled steel strip in a continuous annealing system. First, we heated the strip to a temperature in the range of 700 to 950 ° C. The annealed steel strip is then cooled from the annealing temperature at a cooling rate between 15 and 100 ° C / s to an intermediate temperature of 200 to 250 ° C. Subsequently, the steel strip is cooled at a cooling rate of 2 and 30 ° C / s until it reaches room temperature in air or the cooling at a cooling rate between 15 and 100 ° C / s is maintained up to room temperature, ie, the intermediate temperature corresponds to room temperature ,

Variant 2 ( FIG. 9b ) shows the process according to variant 1, however, the cooling of the steel strip for the purpose of hot dipping is briefly interrupted when passing through the hot dipping vessel, then the cooling at a cooling rate between 15 and 100 ° C / s up to an intermediate temperature of 200 to 250 ° C. continue. Subsequently, the steel strip is cooled at a cooling rate of 2 and 30 ° C / s until it reaches room temperature in air.

Variant 3 ( FIG. 9c ) also shows the process according to variant 1 in a hot dipping refinement, but the cooling of the steel strip is interrupted by a short pause (1 to 20 s) at an intermediate temperature of 200 to 250 ° C and up to the temperature necessary for hot dipping ( approx. 420-470 ° C). Subsequently, the steel strip is cooled to an intermediate temperature of 200 to 250 ° C. At a cooling rate of 2 and 30 ° C / s, the final cooling of the steel strip takes place until air reaches the room temperature.

Claims (17)

1. A method of producing a cold-rolled or hot-rolled steel strip from a higher-strength multiphase steel having excellent forming properties, in particular for lightweight vehicle construction,
   consisting of the elements (contents in wt.%): C 0.060 to ≤ 0.115 Al 0.020 to ≤ 0.060 Si 0.100 to ≤ 0.500 Mn 1.300 to ≤ 2.500 P ≤ 0.025 S ≤ 0.0100 Cr 0.280 to ≤ 0.480 Mo ≤ 0.150 Ti ≥ 0.005 to ≤ 0.050 Nb ≥ 0.005 to ≤ 0.050 B ≥ 0.0005 to ≤ 0.0060 N ≤ 0.0100 remainder iron, including common steel tramp elements not mentioned above, in which the required multi-phase structure is created during a continuous annealing, wherein
   the cold-rolled or hot-rolled steel strip is heated in the continuous annealing furnace to a temperature in the range of 700 to 950°C and the annealed steel strip is then cooled from the annealing temperature to an intermediate temperature of 200 to 250°C at a cooling rate of between 15 and 100°C/s, the steel strip is then cooled in air at a cooling rate of 2 to 30°C/s until ambient temperature is reached or the cooling is maintained at a cooling rate of between 15 and 100°C/s from the intermediate temperature to ambient temperature.
2. The method according to claim 1,
   characterised in that
   the Mo content is ≤ 0.100%.
3. The method according to claim 1 or 2,
   characterised in that
   the Mo content is ≥ 0.050%.
4. The method according to claims 1 - 3,
   characterised in that
   the Nb content is ≤ 0.045%.
5. The method according to claims 1 - 3,
   characterised in that
   the Nb content is ≤ 0.040%.
6. The method according to one of claims 1 - 5,
   characterised in that
   the Ti content is ≤ 0.045%.
7. The method according to one of claims 1 - 5,
   characterised in that
   the Ti content is ≤ 0.040%.
8. The method according to one of claims 1 - 7,
   characterised in that
   the B content is ≤ 0.0050%.
9. The method according to one of claims 1 - 7,
   characterised in that
   the B content is ≤ 0.0040%.
10. The method according to one of claims 1 - 9,
    characterised in that
    the N content is ≤ 0.0090%.
11. The method according to one of claims 1 - 9,
    characterised in that
    the N content is ≤ 0.0070%.
12. The method according to one of claims 1 - 11,
    characterised in that
    the total content of Ti, Nb and B is ≤ 0.106%.
13. The method according to one of claims 1 - 12,
    characterised in that
    the total content of Ti, Nb, B and Mo is ≤ 0.256%.
14. The method according to claim 1,
    characterised in that
    during a hot dip coating, after the heating and subsequent cooling, the cooling is halted before entry into the molten bath and, after the hot dip coating, the cooling is continued at a cooling rate of between 15 and 100°C/s to an intermediate temperature of 200 to 250°C and then the steel strip is cooled in air at a cooling rate of 2 and 30°C/s [sic] until ambient temperature is reached.
15. The method according to claim 1,
    characterised in that
    during a hot dip coating, after the heating and subsequent cooling to the intermediate temperature of 200 to 250°C, before entry into the molten bath the temperature is held for 1 to 20 seconds and then the steel strip is reheated to a temperature of 420 to 470°C and, after hot dip coating has taken place, a cooling takes place to the intermediate temperature of 200 to 250°C at a cooling rate of between 15 and 100°C/s and then cooling to ambient temperature takes place in air at a cooling rate of 2 and 30°C/s [sic].
16. The method according to one of claims 1, 14 and 15,
    characterised in that
    the steel strip is then dressed.
17. The method according to at least one of claims 1 and 14 to 16,
    characterised in that
    the steel strip is then straightened by stretch-bending.

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