KR102503990B1 - High-strength air-hardening multi-phase steel comprising outstanding processing properties and method for the production of a steel strip from said steel - Google Patents

Abstract

The present invention relates to a common-strength air-hardening multi-phase steel comprising the composition defined in claim 1, having a minimum tensile strength of 750 Mpa, and having excellent processing properties, from which the steel is as wide as possible in continuous annealing of hot or cold strip. Considering the process window, the total content of Mn+Si+Cr+Mo is adjusted according to the thickness of the strip produced as follows: below 1.00 mm, Mn+Si+Cr+Mo is 2.450 to 2.800% by weight; More than 1.00 and up to 2.00 mm, Mn+Si+Cr+Mo is 2.600 to 3.150% by weight; When it exceeds 2.00 mm, Mn+Si+Cr+Mo is 3.000 to 3.450% by weight.

Description

High-strength air-hardening multi-phase steel with excellent processing properties and method for producing steel strip from said steel

The present invention relates to a high-strength, air-hardenable, multi-phase steel with excellent processing properties according to claim 1. Advantageous improvements are the subject of dependent claims 2 to 26 .

The invention also relates to a method for producing hot-rolled strip and/or cold-rolled strip from such steel according to patent claims 27 to 34 and heat treatment thereof by air hardening and optionally subsequent tempering, and this method according to claims 35 to 41 It relates to a steel strip manufactured by

In particular, the present invention provides tensile strength in the range of 750 MPa or greater in the unannealed condition for the production of parts with improved deformability (e.g., increased hole expansion and increased bend angle) and improved welding properties. It is about a river with

Yield strength and tensile strength can be increased by heat treating such steels according to the present invention, for example by air hardening with optional post tempering.

The fiercely competitive automotive market forces manufacturers to constantly seek solutions to reduce consumption and CO ₂ emissions while maintaining the highest possible comfort and passenger protection. On the one hand, weight reduction of all vehicle components plays a decisive role, but on the other hand it also ensures optimum behavior of the individual components under conditions of high static and dynamic stress, both in operation and in a crash.

By providing high-strength to ultra-high-strength steel and reducing the thickness of the sheet metal, it is possible to reduce the weight of the vehicle while improving forming properties and part properties during manufacturing and operation.

Thus, high-strength to ultra-high-strength steels have strength and malleability, energy absorption and processing (e.g. punching, hot and cold forming, hot tempering (e.g. air hardening, press hardening), welding and/or surface treatment ( eg metal refining, organic coatings or varnishing), relatively high requirements have to be met.

Therefore, in addition to the required weight reduction through reduction of sheet thickness, the newly developed steel possesses excellent processing properties such as deformability and weldability while increasing the material properties such as yield strength, tensile strength, solidification behavior and elongation at break. must satisfy the necessary conditions.

Therefore, in the case of reducing the sheet thickness as mentioned above, in order to ensure sufficient strength of automobile parts, and in terms of toughness, edge crack resistance, improved bending angle and bending radius, energy absorption and hardening ability, and bake hardening effect, In order to meet the high requirements placed on the part in terms, high-strength to ultra-high-strength steels with single-phase or multi-phase microstructures must be used.

Also, when using resistance spot welding, improved conformity to the joint in the form of better general weldability, such as a larger usable weld area, and improved fracture behavior (failure pattern) of the weld line under conditions of mechanical stress, as well as delayed There is a growing demand for sufficient resistance to hydrogen embrittlement (ie no delayed fracture). This also applies to the suitability for welding of ultra-high strength steels, for example in the manufacture of pipes produced by the high-frequency induction welding method (HFI).

Hole expansion capability is a material property that describes a material's resistance to the risk of fracture and crack propagation during forming operations in areas close to the edge, such as, for example, collar formation.

Hole expansion testing is governed by, for example, the normative standard ISO 16630. For example, a prefabricated hole punched in a sheet is enlarged by a mandrel. The measured value is the change in hole diameter relative to the starting diameter at which the first crack occurs through the sheet at the edge of the hole.

Improved edge crack resistance means increased deformability of the sheet edge and can be explained by increased hole expansion capability. It is known by its synonyms "low edge crack (LEC)", "high hole expansion (HHE)" as well as xpand®.

The bending angle describes a material property from which conclusions can be drawn about material behavior during a forming operation with a dominant bending process (eg during folding), or even when subjected to a crash load. Thus, the increased flexion angle improves passenger compartment safety. Determination of the bending angle (α) is governed by the plate bending test described in the normative standard VDA 238-100.

The properties mentioned above are important in the case of parts that can be molded into very complex parts prior to heat treatment, eg air curing followed by tempering.

Improved weldability is achieved, as is known, in particular by reducing the carbon equivalent. Thus, synonyms are, for example, "underperitical (UP)" or the already known "low carbon equivalent (LCE)". Here, the carbon content is typically less than 0.120% by weight. In addition, the fracture behavior or fracture pattern of the weld line can be improved by alloying using microalloy elements.

High-strength components must have sufficient resistance to hydrogen-induced material embrittlement. The test for resistance to hydrogen-induced brittle fracture associated with the manufacture of high-strength steels (AHSS) used in automobile manufacturing is governed by SEP1970 and is tested through a bent beam test and a puncture tensile test. BACKGROUND OF THE INVENTION In vehicle manufacturing, dual-phase steels consisting of a ferritic basic microstructure incorporating a martensitic second phase are increasingly used. It has been found that in the case of low carbon microalloy steels the proportions of additional phases such as bainite and retained austenite have a beneficial effect on, for example, hole expansion behavior, flexural behavior and hydrogen induced brittle fracture behavior. Here, bainite may exist in various forms, for example, upper bainite and lower bainite.

The characteristic material properties of dual phase steels, such as low maximum yield ratio, predominant strain hardening and good cold formability associated with very high tensile strength, are well known, but are often no longer sufficient for more complex part geometries.

In general, multiphase steel groups are increasingly used. Multi-phase steels include, for example, multi-phase steels, ferritic-bainitic steels, TRIP-steels, as well as the above-described dual-phase steels characterized by different microstructural compositions.

Multiphase steels, according to EN 10346, are steels containing small proportions of martensite, retained austenite and/or pearlite within a ferritic/bainitic basic microstructure, in which intense grain refinement results in delayed recrystallization of the microalloy elements. or by precipitation.

Compared to biphasic steels, these composite phase steels have higher yield strength, higher yield-to-yield ratio, lower strain hardenability and higher hole expansion capability.

Ferritic-bainitic steels, according to EN 10346, are steels containing bainite or work-hardening bainite in a matrix of ferrite and/or work-hardening ferrite. The strength of this matrix is caused by high dislocation density and grain refinement and precipitation of microalloy elements.

Dual-phase steels, according to EN 10346, are steels with a ferritic basic microstructure, in which a martensitic second phase is included in the form of islands, optionally including some bainite as the second phase. . Dual phase steels exhibit a low maximum yield ratio and predominant strain hardening while having high tensile strength.

TRIP-steel, according to EN 10346, is a steel with a predominantly ferritic basic microstructure comprising bainite and retained austenite, which can transform to martensite during deformation (TRIP effect). Due to its predominant strain hardening, this steel obtains high values of uniform elongation and tensile strength. Together with the bake hardening effect, high part strength can be obtained. These steels are suitable for deep drawing as well as tension forming. However, higher sheet metal gripping and pressing forces are required during forming of the material. Relatively strong rebounding must also be taken into account.

High-strength steels with a single-phase microstructure include, for example, bainitic steels and martensitic steels.

Bainitic steels, according to EN 10346, are characterized by very high yield strength and tensile strength with a sufficiently high elongation for the cold forming process. Their chemical composition contributes to excellent weldability. The microstructure is typically composed of bainite. Small proportions of other phases, for example martensite and ferrite, may be included in this microstructure.

Martensitic steels, according to EN 10346, are steels containing a small proportion of ferrite and/or bainite in the basic microstructure of martensite as a result of thermo-mechanical rolling. These steel grades are characterized by very high yield strength and tensile strength with a sufficiently high elongation in the case of cold forming processes. Within the group of multiphase steels, martensitic steels have the highest tensile strength values. Suitability for deep drawing is limited. Martensitic steels are suitable primarily for bending processes such as roll forming.

A heat treatable steel is, according to EN 10083, a steel that obtains high tensile strength and durability by heat treatment (=quick hardening and tempering). When cooling during hardening in air produces bainite or martensite, the method is referred to as "air-hardening". Tempering after hardening allows the strength/toughness ratio to be affected in a targeted manner.

Applications and Manufacturing Process

High-strength and ultra-high-strength multiphase steels are used especially in structural parts, chassis parts and crash-related parts such as flexible cold-rolled strips, so-called TRB® or tailored strips, as well as sheet metal plates, tailored blanks.

Tailor Rolled Blank (TRB®) lightweight technology enables significant weight reduction by means of load-adaptive sheet thickness across part lengths and/or steel grades.

In continuous annealing plants, special heat treatments are carried out to adjust the desired microstructure, in which relatively soft constituents, for example ferrite or bainitic ferrite, lead to a low yield strength of the steel, martensitic or Hard constituents of the steel, such as carbon-rich bainite, contribute to the strength of the steel.

For economic reasons, cold rolled high strength to extra high strength steel strips are usually annealed in a continuous annealing process to form easily formable metal sheets. Depending on the alloy composition and the strip cross-section, process parameters such as processing speed, annealing temperature and cooling rate (cooling gradient) are adjusted depending on the mechano-technical properties required by the microstructure required for this purpose.

To tune the dual-phase microstructure, pickled hot strips of typical thicknesses of 1.50 mm to 4.00 mm or cold strips of typical thicknesses of 0.50 mm to 3.00 mm are subjected to recrystallization and continuous annealing to a temperature that forms the required microstructure during cooling. heated in the furnace.

In particular, it is particularly difficult to achieve a constant temperature when the thickness is different in the transition region from one strip to another. During continuous annealing of alloy compositions with too narrow a process window, this can lead to reduced productivity, for example, by passing thinner strips through the furnace too slowly, or by passing the thicker strips through the furnace too quickly, resulting in the desired microstructure. may fail to achieve the required annealing temperature to achieve As a result, defects increase and the cost of errors increases.

With the same process parameters, an enlarged process window is required so that the required strip characteristics can be obtained even when the cross-section of the strip to be annealed is larger.

A problem with very narrow process windows is particularly in annealing where load-optimized parts are made from hot-rolled or cold-rolled strip where the thickness of the strip varies over the strip length and width (eg by flexible rolling). stand out

However, if the sheet thickness varies significantly, producing TRB® with a multiphase microstructure using currently known alloys and available continuous annealing systems is costly, e.g. requires additional heat treatment prior to cold rolling. do it with In the region of different sheet thicknesses, i.e. when the degrees of rolling reduction are varied, a homogeneous multiphase microstructure in the cold-rolled steel strip and the hot-rolled steel strip results from the temperature difference in the conventional alloy-specific narrow process window. cannot be formed

A method for producing steel strips of different thickness over the entire length of the strip is described, for example, in DE 100 37 867 A1.

When the surface of a hot-rolled steel strip or cold-rolled steel strip has to be hot-dipped galvanized due to high requirements for corrosion protection, annealing is usually carried out in a continuous annealing furnace disposed upstream of the hot-dip galvanizing bath.

Also in the case of hot strip, according to the alloy concept, the microstructure required to realize the required mechanical properties is not formed prior to annealing in a continuous furnace.

Therefore, determining the process parameters is to control the cooling rate (cooling gradient) as well as the annealing temperature and rate in continuous annealing, since the phase transformation is temperature and time dependent. Thus, the lower the sensitivity of the steel with respect to uniformity of mechanical properties when temperature and time course change during continuous annealing, the larger the process window.

When continuously annealing hot-rolled or cold-rolled steel strips of different thicknesses using the alloy concept for dual phase steels known, for example, from published patent document EP 2 128 295 A1 or EP 2 154 028 A1, The problem is that even if the required mechanical properties can be satisfied by these alloy compositions, they are uniform over the entire length of the strip in the case of a change in cross section, for example in the case of a change in width or thickness, without applying process parameters. One is that only a narrow process window is available for the annealing parameters used to obtain one mechanical property.

When using known alloying concepts, this narrow process window makes it difficult to form uniform mechanical properties over the entire length and width of the strip during successive annealing of strips of different thicknesses.

In the case of flexible rolled cold strip made of multiphase steels of known composition, too narrow a process window results in regions of thinner sheet thickness having excessive strength due to excessive martensitic proportions due to the transformation process during cooling, or , areas of greater thickness will achieve insufficient strength due to insufficient martensite proportions. Practically homogeneous mechanical-technical properties over the entire strip length or width cannot be achieved with known alloy concepts in continuous annealing.

The goal of achieving mechano-technical properties obtained in narrow regions throughout the strip width and strip length through controlled adjustment of the volume fraction of the microstructure has the highest priority, and therefore this is only possible through an enlarged process window. . Known alloy concepts for multiphase steels are characterized by too narrow process windows and are therefore unsuitable for solving this problem, especially in the case of flexible rolled steel strips. With the alloy concepts known to date, only one strength class of steel with a finite cross-sectional area (sheet thickness and strip width) can be produced, thus requiring different alloy concepts for different strength classes or cross-sectional ranges.

Steel production has shown a tendency to reduce the carbon equivalent to achieve improved cold working (cold rolling, cold forming) and improved performance.

However, welding suitability characterized above all by carbon equivalent is an important evaluation factor.

For example, at the carbon equivalent of

- CEV(IIW) = C + Mn/6 + (Cu + Ni)/15 + (Cr + Mo + V)/5

- CET = C + (Mn + Mo)/ 10 + (Cr + Cu)/20 + Ni/40

- PCM = C + (Mn + Cu + Cr)/20 + Ni/60 + Mo/15 + V/10 + 5B

Characteristic standard elements such as carbon and manganese, as well as chromium or molybdenum and vanadium (content in % by weight) are contemplated.

Silicon plays only a secondary role in the calculation of carbon equivalent. This is of critical importance in relation to the present invention. The reduction in carbon equivalent through reduction of the content of manganese as well as carbon must be compensated by increasing the silicon content. Thus, edge crack resistance and weldability are improved while maintaining the same strength.

A low maximum yield ratio (Re/Rm) in the strength range above 750 MPa in the initial state is typical for dual phase steels and provides formability in drawing and deep drawing operations. This provides the fabricator with information about the distance between subsequent plastic deformation and failure of the material in a metastable load. Thus, a lower yield strength ratio provides a greater margin of safety against component failure.

The higher maximum yield ratio (Re/Rm) typical for composite phase steels is also characterized by high resistance to edge cracking. This can be attributed to the smaller differences in strength and hardness of the individual microstructured constituents and smaller microstructures, which have a beneficial effect on uniform deformation in the region of the edge being cut.

Regarding yield strength as well as maximum yield ratio (Re/Rm), there is an overlapping range in the standard that can be assigned to both composite and dual phase steels, which enhances the properties of the material.

The analytical situation for obtaining multiphase steels with a minimum tensile strength of 750 MPa is very diverse, not only for the strength-enhancing elements carbon, silicon, manganese, phosphorus, nitrogen, aluminum, but also for chromium and/or molybdenum. It shows a very wide alloy range for the addition of microalloys such as titanium, niobium, vanadium and boron.

The range of dimensions within this strength range is wide, ranging in thickness from about 0.50 to about 4.00 mm for strips for continuous annealing. The starting materials used may be hot rolled strip, cold rolled hot rolled strip and cold strip. Strips with a width of up to about 1600 mm are mainly used, but slit strip dimensions obtained by longitudinal splitting of the strip are also used. Sheet metal or plates are produced by cutting this strip crosswise.

EP 1 807 544 B1, WO 2011/000351 and EP 2 227 574 B1 with a minimum tensile strength of 800 (LH®800) and 900 MPa (LH®900) respectively in hot-rolled or cold-rolled version The air hardenable steel grades known from are characterized by very good formability (deep draw properties) in the soft state and high strength after heat treatment (tempering).

During hardening, the microstructure of the steel is transformed into the austenitic range by heating under a protective gas atmosphere to temperatures preferably exceeding 950°C. During subsequent cooling in air or protective gas, a martensitic microstructure is formed for high-strength components.

Residual stresses in the hardened part can be removed by subsequent tempering. At the same time, the hardness of the component is reduced to obtain the required toughness values.

Accordingly, an object of the present invention is a highly durable multi-phase air-, preferably having a dual-phase microstructure, having excellent processing properties in the longitudinal and transverse directions with respect to the rolling direction and a minimum tensile strength of 750 MPa in the non-heat treated state. It provides a new cost-effective alloying concept for conditioned steel, since the dual-phase microstructure enlarges the process window for continuous annealing of hot or cold strip, in addition to strips with different cross-sections, across the length of the strip and optionally can be produced such that a steel strip whose thickness varies over the strip width has the most uniform mechanical-technical properties.

In addition, the hot-dip treatment of the steel must be ensured, and the manufacturing process of the strip made of this steel must be initiated.

The present invention is also intended to ensure sufficient formability, HFI weldability, excellent general weldability, and resistance to hot dipping and tempering.

According to the teachings of the present invention, this object is achieved by a steel having the following chemical composition in weight percent:

C: 0.075 to 0.115

Si: 0.600 to 0.750

Mn: 1.000 to 1.950

Cr: 0.200 to 0.600

Al: 0.010 to 0.060

N: 0.0020 to 0.0120

S: 0.0030 or less

Mo: 0.0200 or more

Nb: 0.005 to 0.040

Ti: 0.005 to 0.030

B: 0.0005 to 0.0030

Ca: 0.0005 to 0.0060

Cu: 0.050 or less

Ni: 0.050 or less

Smelting-related impurities normally associated with the remainder of iron and steel. Here, in relation to the widest possible process window during the continuous annealing of hot and cold strips made of the steel, the total content of Mn + Si + Cr is adjusted as a function of the thickness of the strip produced as follows.

The sum of Mn + Si + Cr + Mo at a thickness of 1.00 mm or less is 2.450 to 2.800%,

The sum of Mn + Si + Cr + Mo in the thickness of more than 1.00 and not more than 2.00 mm is 2.600 to 3.150%,

The sum of Mn + Si + Cr + Mo at a thickness greater than 2.00 mm is 3.000 to 3.450%.

As a result of the possibilities described in claims 28 and 29 of hot-dip refining (eg hot-dip galvanizing) of a steel strip made of a steel according to the invention with a high silicon content of less than or equal to 0.500%, the tempering resistance The addition of vanadium may be omitted to ensure

According to the present invention, the microstructure consists of the main phases ferrite and martensite and the secondary phase bainite, which determines the improved mechanical properties of the steel.

The steel according to the invention is characterized by a low carbon equivalent, the carbon equivalent (CEV(IIW)) being limited to a maximum of 0.66% depending on the sheet thickness in order to achieve excellent weldability and further specific properties described below. . CEV(IIW) values of up to 0.50% for sheet thicknesses below 1.00 mm, up to 0.55% for sheet thicknesses below 2.00 mm and up to 0.60% above 2.00 mm have proven advantageous.

Due to its chemical composition, the steel according to the invention can be produced within a wide range of hot rolling parameters, eg with a coiling temperature exceeding the bainitic onset temperature (variant A). Furthermore, by controlling the process in a targeted manner, the microstructure can be adjusted such that the steel according to the present invention can be cold rolled without conventional softening annealing, wherein cold rolling reductions of 10 to 60% per cold rolling pass can be achieved. this can be used

The steel according to the invention is very suitable as a starting material for hot-dip refining and, due to the total amount of Mn, Si and Cr and Mo added according to the invention as a function of the strip thickness to be produced, a considerably wider process window compared to known steels have

Test results have shown that a wide process window for obtaining the required mechanical properties can be maintained if the total content of Mn + Si + Cr + Mo is adjusted according to the sheet thickness.

This increases process reliability during continuous annealing of cold and hot strips with dual-phase microstructures or multi-phase microstructures. Thus, more uniform mechanical-technical properties within the strip can be adjusted for continuously annealed hot or cold strips, even for different cross-sections and otherwise identical process parameters.

This applies to continuous annealing of strips having different strip cross-sections as well as having varying strip thicknesses throughout the strip length or width. For example, this allows processing at selected thickness ranges (eg, strip thickness less than 1.00 mm, strip thickness between 1.00 mm and 2.00 mm, and strip thickness greater than 2.00 mm).

If high-strength hot or cold strips with various strip thicknesses are produced from the multi-phase steel according to the invention in a continuous annealing process, components which are load-optimized components can be produced from these hot or cold strips.

The steel strip according to the present invention is applied in a hot dip-galvanizing line or a pure continuous annealing line in a skin-passed and non-skin-passed state, in a tensile-bend state and in a tensile-unbend state, and in a heat treatment (overaging) state. can be produced as cold strip and hot strip as well as cold re-rolled hot strip.

In the alloy composition according to the present invention, the steel strip is austenitized above A c3 by annealing in the transformation zone between A _c1 to A _c3 , _or by using a final controlled cooling to induce a dual-phase microstructure or multi-phase microstructure. It can be produced by annealing.

Annealing temperatures of about 700 to 950° C. have been found to be advantageous. Depending on the overall process (continuous annealing or further hot dip finishing) there are different approaches to heat treatment.

For continuous annealing plants without subsequent hot-dip refining, the strip is cooled from the annealing temperature to an intermediate temperature of about 160 to 250° C. at a cooling rate of about 15 to 100° C./sec. Optionally, it may be cooled beforehand to an intermediate temperature of 300 to 500° C. at a cooling rate of about 15 to 100° C./sec. Finally, cooling to room temperature is performed at a cooling rate of about 2 to 30° C./sec (see method 1, FIG. 6a).

For heat treatment in hot-dip tablets, two temperature profiles are possible. As described above, cooling is stopped before entering the hot dip bath, and cooling continues until an intermediate temperature of about 200 to 250 ° C is reached only after being discharged from the hot dip bath. As a result, depending on the temperature of the hot dip bath, the temperature of the hot dip bath is maintained at about 400 to 470 ° C. Cooling to room temperature is again carried out at a cooling rate of about 2 to 30° C./sec (modification 2, Fig. 6b).

A second variant of the temperature profile during hot-dip purification involves holding at an intermediate temperature of about 200-350°C for about 1-20 seconds, followed by reheating to the temperature required for hot-dip purification of about 400-470°C. . The strip is cooled again to about 200-250° C. after purification. Cooling to room temperature is again carried out at a cooling rate of about 2 to 30° C./sec (Variation 3, Fig. 6c).

In known dual phase steels, not only carbon, but also manganese, chromium and silicon participate in the transformation from austenite to martensite. The combination of the elements of carbon, silicon, manganese, nitrogen, molybdenum and chromium, as well as niobium, titanium and boron, added within certain limits, ensures the necessary mechanical properties, such as a minimum tensile strength of 750 MPa, It also significantly enlarges the process window during continuous annealing.

As a result of the addition of manganese in increasing weight percentages, the ferritic regions of the material are characterized by a displacement towards longer times and to lower temperatures during cooling. Hereby the proportion of ferrite is reduced to a greater or lesser extent by the increased amount of bainite, depending on the process parameters.

Adjusting the low carbon content to 0.115% by weight or less can reduce the carbon equivalent, thereby improving weldability and preventing excessive hardening during welding. In the case of resistance spot welding, electrode life can also be significantly increased.

Hereinafter, the effects of the elements in the alloy according to the present invention are described in more detail. Incidental elements are unavoidable and, if necessary, considered as analytical concepts in relation to their effects.

Incidental elements are elements already present in the iron ore or resulting from the manufacturing fixation of the steel. Incidental elements are generally undesirable because of their predominantly negative effects. These elements are sought to be removed to acceptable levels or transformed into more innocuous forms.

Hydrogen (H) is the only element that can diffuse through the iron lattice without causing lattice strain. As a result, hydrogen is relatively mobile within the iron lattice and can be absorbed relatively easily during steel processing. Hydrogen can only be absorbed in the iron lattice in atomic (ionic) form.

Hydrogen is highly embrittling and preferentially diffuses into energetically favorable sites (defects, grain boundaries, etc.), whereby bonds function as hydrogen traps and thus reduce the residence time of hydrogen in the material. can be significantly increased.

Recombination to molecular hydrogen can lead to cold cracking. This behavior occurs with hydrogen embrittlement or hydrogen-induced stress corrosion cracking. Hydrogen is often characterized as the cause of so-called delayed fractures that occur without external stress. Therefore, the hydrogen content in the steel should be as low as possible.

The more uniform structure achieved in the steel according to the invention, in particular by its wide process window, also reduces the susceptibility to hydrogen embrittlement.

Oxygen (O): In the molten state, steel has a relatively high absorption capacity for gases. However, at room temperature, oxygen is only soluble in very small amounts. Similar to hydrogen, oxygen can only diffuse into materials in atomic form. Due to the high embrittlement effect and negative influence on the aging resistance, attempts are made to reduce the oxygen content as much as possible during production.

To reduce oxygen, process-engineering approaches such as vacuum treatment and analytical approaches exist. By adding certain alloying elements, oxygen can be converted to a more harmless state. Thus, oxygen is typically bound by manganese, silicon and/or aluminum during the deoxidation of steel. However, the oxides obtained may cause negative properties as defects in the material.

Therefore, for this reason, the oxygen content in the steel should be as low as possible.

Phosphorus (P) is a trace element of iron ore and dissolves in the iron lattice as a substitution atom. Phosphorus increases hardness and improves hardenability by solid solution strengthening. However, due to the slow diffusion rate of phosphorus, phosphorus has a strong tendency to segregate and severely deteriorates toughness, so it is usually sought to lower the phosphorus content as much as possible. Phosphorus deposition at grain boundaries can lead to grain boundary cracking. Phosphorus also increases the transition temperature from toughness to brittle behavior by up to 300°C. During hot rolling, near-surface phosphorus oxide can cause segregation at the grain boundaries.

However, due to its low cost and high strength gain, phosphorus is used in small amounts as a microalloying element in some steels, for example in high-strength interstitial free (IF)-steel, bake-hardened steel or in some alloy concepts for dual-phase steels. (<0.1%). The steel according to the invention differs from the known analytical concept of using phosphorus as a solid solution former, in particular because phosphorus is not added but controlled as low as possible.

For the above reasons, the phosphorus content in the steel according to the present invention is limited to an amount unavoidable in steel production.

Sulfur (S), like phosphorus, is bound in iron ore as a trace element. Sulfur is highly brittle and highly brittle, making it undesirable in steels (exception, automation steel). Therefore, it is sought to lower the sulfur content in the melt as far as possible, for example by vacuum treatment. Additionally, by adding manganese, sulfur is converted to manganese sulfide, a relatively harmless compound. Manganese sulfide is rolled into a band during rolling and serves as a germination site for transformation. Particularly in the case of diffusion-controlled transformations, this leads to a microstructure organized in the form of bands, and in the case of very pronounced banding the mechanical properties are degraded (e.g., prominent martensite bands instead of islands of dispersed martensite, Anisotropic material behavior, reduced elongation at break).

For the above reasons, the sulfur content in the steel according to the present invention is limited to less than 0.0030% by weight, or optimally less than 0.0020% by weight, or to an amount unavoidable in steel production.

Alloying elements are typically added to affect specific properties in a targeted manner. Alloying elements can affect different properties of different steels. The effect generally depends on the amount and state of solid solution in the material.

Interactions can therefore be very diverse and complex. The effect of this alloying element is explained in more detail below.

Carbon (C) is the most important alloying element in steel. Targeted introduction of less than 2.06% by weight carbon is required to convert iron into steel. During steel production, the carbon content is often significantly reduced. In the case of dual-phase steel for continuous hot-dip coating, according to EN 10346 or VDA 239-100 the content is not more than 0.180% by weight, no minimum value specified.

Carbon dissolves interstitially in the iron lattice due to its relatively small atomic radius. The solubility in α-iron is up to 0.01% and in γ-iron up to 2.06%. Carbon in solubilized form significantly increases the hardenability of the steel and is therefore necessary to form a sufficient amount of martensite. However, excessive carbon content increases the hardness difference between ferrite and martensite and limits weldability.

For example, to satisfy the requirements regarding high hole expansion and bending angle, the steel according to the present invention contains less than 0.115% by weight of carbon.

Due to the different solubility of carbon in the phases, distinct diffusion processes are required during the phase transformation, which can lead to very different kinetic conditions. Carbon also increases the thermodynamic stability of austenite, which is shown in the phase diagram to expand the austenite region to lower temperatures. As the content of forcibly dissolved carbon in martensite increases, the lattice distortion increases, and with it the intensity of the non-diffusively produced phase increases.

Carbon also forms carbides. Cementite phase (Fe ₃ C) occurs in almost all steels. However, much harder specialty carbides can also be formed using other metals, such as chromium, titanium, niobium, vanadium, for example. The type as well as the distribution and size of the precipitates is critical to the strength increase obtained. In order to ensure on the one hand sufficient strength and on the other hand good weldability, improved hole expansion, improved bend angle and sufficient resistance to hydrogen-induced cracking (i.e. no delayed fracture), a minimum The C content is set to 0.075% by weight and the maximum C content to 0.115% by weight, advantageously the content is adjusted according to the cross section as in the following example.

Material thickness of less than 1.00 mm (C: 0.100% by weight or less)

Material thickness of 1.00 to 2.00 mm (C: 0.105% by weight or less)

Material thickness greater than 2.00 mm (C: 0.115% by weight or less).

Silicon (Si) is used for deoxidation of steel because it combines with oxygen during casting. It is important for subsequent steel properties that the segregation coefficient of silicon is significantly lower than that of manganese, for example (0.16 vs. 0.87). Segregation generally results in a band-like structure of the microstructured components, which deteriorates forming properties such as hole expansion and bending ability.

Characteristically, the addition of silicon induces strong solid solution hardening. The addition of 0.1% silicon causes an increase in tensile strength of about 10 MPa, where silicon below 2.2% only slightly lowers the expansion. This was investigated for different sheet thicknesses and annealing temperatures. Increasing silicon from 0.2% to 0.6% resulted in a strength increase of about 10 MPa in yield strength and about 25 MPa in tensile strength. This only reduces the elongation at break by about 1%. The latter is due in particular to the fact that silicon lowers the solubility of carbon in ferrite, whereby the ferrite becomes softer, resulting in improved moldability. Silicon is also a brittle phase and prevents the formation of carbides that reduce malleability. Within the scope of the steels according to the invention, the low strength-increasing effect of silicon forms the basis for a wide process window.

Another important effect is that silicon allows for sufficient ferrite formation prior to quenching as it shifts the formation of ferrite towards shorter times and temperatures. This provides improved cold rolling properties during hot rolling. In the hot dip coating process, austenite is enriched with carbon and stabilized by accelerated ferrite formation. Since silicon inhibits carbide formation, austenite is additionally stabilized. Thus, in accelerated cooling, the formation of bainite can be suppressed for martensite.

The addition of silicon to the scope according to the present invention has resulted in further surprising effects described below. The retardation of carbide formation described above can be caused, for example, by aluminum. However, since aluminum forms stable nitrides, sufficient nitrogen is not available for the formation of microalloy elements and carbonitrides. Because silicon does not form carbides or nitrides, this problem does not exist due to alloying with silicon. Thus, silicon has an indirect positive effect on the formation of precipitates by the microalloy, and thus a positive effect on the strength of the material. Since an increase in the transformation temperature by silicon tends to promote grain coarsening, microalloys with niobium, titanium, and boron are particularly suitable as targeted control of the nitrogen content in the steel according to the present invention.

As is known, strongly adhered red scale is formed in steels with high-silicon-alloy steels, and the risk of rolling scale occurring during hot rolling increases, which is the result of subsequent pickling and pickling can affect productivity. This effect is not detectable in steels according to the invention containing between 0.600% and 0.750% silicon, when pickling is advantageously carried out with hydrochloric acid instead of sulfuric acid.

Regarding the galvanizing ability of silicon-containing steels, DE 196 10 675 C1 states that due to the very poor wettability of the surface of the steel by liquid zinc, in particular steels containing up to 0.800% silicon or up to 2.000% silicon are hot-dipped. It states that it cannot be galvanized.

In addition to recrystallization of the hard strip, the atmospheric conditions in the continuous hot dip galvanizing plant during the annealing treatment lead to a reduction of iron oxides that may form on the surface, for example during cold rolling or due to storage at room temperature. However, in the case of oxygen-friendly alloying elements such as silicon, manganese, chromium, and boron, the overall atmosphere is oxidizing, which can lead to segregation and selective oxidation of these elements. Selective oxidation can take place externally, ie on the surface of the substrate as well as inside the metal matrix.

In particular, it is known that silicon can diffuse to the surface during annealing either alone or together with film-like oxides in the form of manganese. These oxides can prevent contact between the substrate and the melt, and can prevent or significantly reduce the wetting reaction. This can result in non-galvanized areas, so-called "bare spots" or even large surface areas without coating. Moreover, due to the reduced wetting reaction, the formation of the suppression layer may be insufficient, and thus the adhesion of the zinc layer or zinc alloy layer on the substrate may be reduced. The mechanism described above is also applied to pickled hot strip or cold rolled hot strip.

Contrary to this general knowledge, tests have unexpectedly revealed that good galvanizing properties and good zinc adhesion of the strip can be achieved only by properly operating the furnace during recrystallization annealing and passing through a zinc bath.

For this purpose, the surface of the strip must first be freed from residual scale, rolling oil or other dirt particles by chemical or thermo-water-mechanical pre-cleaning. In order to prevent silicon oxide from reaching the surface, measures must also be taken to promote internal oxidation of the alloying elements beneath the surface of the material. Depending on the configuration of the installation, different measures are used for this purpose.

In plant configurations where the annealing process step is performed entirely in a radiant tube furnace (RTF) (see method 3 in Fig. 6c), the internal oxidation of the alloying elements regulates the oxygen partial pressure in the furnace atmosphere (N ₂ -H ₂ protective gas atmosphere). By doing so, you can be affected in a targeted way. The oxygen partial pressure adjusted here must satisfy the following equation, wherein the temperature of the furnace is 700 to 950 ° C.

-12 > Log pO ₂ ≥ 5* Si ^-0.25 - 3* Mn ^-0.5 - 0.1* Cr ^-0.5 - 7*(-lnB) ^0.5

where Si, Mn, Cr, B are the corresponding alloying elements in the steel in weight percent, and pO ₂ is the oxygen partial pressure in mbar.

In the configuration of the plant in which the furnace zone consists of a combination of a direct fire furnace (DFF or non-oxidizing furnace (NOF)) followed by a radiant tube furnace (see method 2 in Fig. 6b), the selective oxidation is influenced by the gas atmosphere in the furnace zone may receive

Through the combustion reaction in the NOF, the oxygen partial pressure and with it the oxidation potential of the iron and alloy components can be controlled. The oxidation potential is controlled so that the oxidation of the alloying elements occurs inside, below the steel surface, and after passing through the NOF region, a thin iron oxide layer can be formed on the steel surface. This is achieved, for example, by reducing the CO value to less than 4%.

In the subsequent radiant tube furnace, the iron oxide layer that may be formed and also the alloying elements are further reduced under an N ₂ -H ₂ protective gas atmosphere. Here, the controlled oxygen partial pressure in this furnace region must satisfy the following equation, wherein the temperature of the furnace is 700 to 950 °C.

-18 > Log pO ₂ ≥ 5* Si ^-0.3 - 2.2* Mn ^-0.45 - 0.1* Cr ^-0.4 - 12.5*(-lnB) ^0.25

Here, Si, Mn, Cr, and B represent the corresponding alloy proportions in the steel in mass %, and pO ₂ represents the oxygen partial pressure in mbar.

In the transition region between the furnace → zinc pot (tuyere snout), the dew point of the gas atmosphere (N ₂ -H ₂ protective gas atmosphere) and together with the oxygen partial pressure are such that oxidation of the strip is prevented before immersion into the molten bath. It should be regulated. A dew point in the range of -30 to -40°C has proven advantageous.

The measures in the furnace area of the continuous hot dip galvanizing plant described above prevent the formation of oxides on the surface and achieve uniform and good wetting of the strip surface by the liquid melt.

If, instead of hot-dip galvanizing, the process of continuous annealing followed by electrolytic galvanizing is chosen (see method 1 in Fig. 6a), no special measures are required to ensure galvanizing properties. It is known that galvanizing of higher alloy steels can be realized considerably more easily by electrolytic galvanizing than by continuous hot-dip galvanizing. In electrolytic galvanizing, pure zinc is electrodeposited directly onto the strip surface. In order not to impair the flow of electrons between the strip and the zinc-ions and thus the galvanization, there should be no surface covering oxide layer present on the strip surface. These conditions are usually ensured by a standard reducing atmosphere during annealing and pre-cleaning prior to electrolysis.

To ensure a process window as wide as possible and sufficient galvanizing capability during annealing, the minimum Si-content is set to 0.600% and the maximum silicon content to 0.750%.

Manganese (Mn) is added to almost all steels for desulfurization to convert harmful sulfur to manganese sulphide. Also, as a result of solid solution strengthening, manganese increases the strength of ferrite and shifts the α/γ transformation towards lower temperatures.

The main reason for adding manganese to dual phase steel is a significant improvement in hardness penetration. Due to the diffusion hindrance, the pearlite and bainite transformations are shifted towards longer times and the martensitic initiation temperature is lowered.

At the same time, however, the addition of manganese increases the hardness ratio of martensite to ferrite. Also, the bending of the microstructure is increased. Due to the high hardness difference between the phases and the formation of martensitic bands, the hole expansion ability is lower, which adversely affects the edge crack resistance.

Manganese, like silicon, has a tendency to form oxides on the steel surface during annealing. Depending on the annealing parameters and the content of other alloying elements (particularly silicon and aluminum), manganese oxides (eg, MnO) and/or Mn mixed oxides (eg, Mn ₂ SiO ₄ ) may be formed. However, since manganese forms a spherical oxide instead of an oxide film at a low Si/Mn ratio or Al/Mn ratio, the importance is low. Nevertheless, high manganese content can negatively affect the zinc adhesion and appearance of the zinc layer.

For the reasons described above, the Mn-content is set to 1.000 to 1.900% by weight.

To achieve the required minimum strength, it is advantageous to vary the manganese content depending on the thickness.

For a strip thickness of less than 1.00 mm, the manganese content preferably ranges from 1.000 to 1.500% by weight, for a strip thickness of 1.00 to 2.00 mm, from 1.300 to 1.700% by weight, for a band thickness greater than 2.00 mm. , 1.600% to 1.900% by weight.

A further feature of the present invention is that a change in manganese content can be compensated for by a simultaneous change in silicon content. The increase in strength (here, yield strength, YS) due to manganese and silicon is generally well described by Pickering's equation.

YS(MPa) = 53.9 + 32.34 [wt%Mn] + 83.16 [wt%Si] + 354.2 [wt%Ni] + 17.402 d ^(-1/2)

However, this equation is mainly based on the effect of the solid solution transition, and according to this equation, the solid solution hardening is weaker in the case of manganese than in the case of silicon. At the same time, however, as mentioned above, manganese significantly increases the hardenability, which significantly increases the proportion of the second phase that increases the strength in multiphase steels. Therefore, as a first approximation in terms of strength increase, the addition of 0.1% silicon is set equal to the addition of 0.1% manganese. For steels of composition according to the invention and annealing with time-temperature parameters according to the invention, the following equations for yield strength (YS) and tensile strength (Ts) were experimentally determined.

YS(MPa) = 160.7 + 147.9 [wt.% Si] + 161.1 [wt.% Mn]

TS (MPa) = 324.8 + 189.4 [wt.% Si] + 174.1 [wt.% Mn]

Compared to the Pickering equation, the moduli of manganese and silicon are almost the same for yield strength as well as tensile strength, thus providing the possibility of replacing manganese with silicon.

Chromium (Cr) in solubilized form can, on the one hand, significantly increase the hardenability of an already small amount of steel. Chromium, on the other hand, causes precipitation hardening in the corresponding temperature profile in the form of chromium carbide. At the same time, an increase in the number of germination sites at a reduced carbon content leads to a decrease in hardenability.

In dual phase steels, the addition of chromium primarily improves the hardening depth. Chromium, in the dissolved state, shifts the pearlite and bainite transformations towards longer times while lowering the martensitic onset temperature.

Another important effect is that little strength loss occurs in the hot dip bath as chromium significantly increases the tempering resistance.

Chromium is also a carbide former. If chromium-iron mixed carbides are present, the austenitizing temperature before hardening should be chosen high enough to dissolve the chromium carbides. Otherwise, the increased number of nuclei may reduce the depth of hardening.

Chromium also tends to form oxides on the steel surface during the annealing process, which can reduce smelting quality. The formation of Cr oxide or Cr mixed oxides on the steel surface after annealing is reduced by the above-mentioned means for setting the furnace region in the case of continuous melt dip coating.

Therefore, the chromium content is set to a content of 0.200 to 0.600% by weight.

To achieve the required minimum strength, it is advantageous to maintain different chromium contents depending on the strip thickness.

For a strip thickness of less than 1.00 mm, the chromium content is in the range of 0.250 to 0.350 weight percent, for a strip thickness in the range of 1.00 mm to 2.00 mm, greater than 0.350 and up to 0.450 weight percent, for a strip thickness greater than 2.00 mm, It is preferable that it is more than 0.450 weight% and 0.550 weight% or less.

Molybdenum (Mo): Similar to chromium, molybdenum improves hardenability. The pearlite and bainite transformations are shifted towards longer times, and the martensitic onset temperature is lowered. Molybdenum is also a strong carbide former, which among other things also allows the formation of finely dispersed mixed carbides with titanium. Molybdenum also significantly increases the tempering resistance so no loss of steel is expected in the zinc bath and causes an increase in the strength of ferrite as a result of solid solution strengthening, but less effective than manganese and silicon.

Therefore, the content of molybdenum is adjusted to 0.200 or less. An advantageous range is 0.050 to 0.100% by weight.

As a compromise between the required mechanical properties and galvanizing ability, a combined Mo+Cr content of less than 0.800%, or optimally less than 0.700% by weight, has proven advantageous for the alloy concept of the present invention.

Copper (Cu): The addition of copper can increase tensile strength and depth of cure. In conjunction with nickel, chromium and phosphorus, copper can form a protective oxide layer on the surface, which significantly reduces the corrosion rate.

When combined with oxygen, copper can form oxides that are detrimental to the bonding interface, which can negatively affect the hot forming process. Therefore, the content of copper is set to 0.050% by weight or less, and thus limited to an amount unavoidable in steel production.

Nickel (Ni): In combination with oxygen, nickel can form harmful oxides at grain boundaries, which can negatively affect the hot forming process. Therefore, the content of nickel is set to 0.050% by weight or less, and is thus limited to an amount unavoidable in steel production.

Vanadium (V): Since the addition of vanadium is unnecessary in this alloying concept, the content of vanadium is limited to an unavoidable amount in the steel.

Aluminum (Al) is commonly added to steel to bind dissolved oxygen and nitrogen in the iron. Thus, oxygen and nitrogen are converted to aluminum oxide and aluminum nitride. These precipitates can cause grain refinement by increasing the number of nucleation points and thus increase toughness properties as well as strength values.

Aluminum nitride does not precipitate when titanium is present in a sufficient amount. Titanium nitride has a lower enthalpy of formation and forms at higher temperatures.

In the molten state, aluminum and silicon shift ferrite formation towards a shorter time and thus allow sufficient ferrite to form in the dual phase steel. It also inhibits carbide formation and therefore retards the transformation of austenite. For this reason, aluminum is also used as an alloying element in retained austenitic steels (TRIP steels) to replace part of silicon. The reason for this approach is that aluminum is slightly less important to the galvanizing reaction than silicon.

Therefore, the aluminum content is limited to 0.010 up to 0.060% by weight, or optimally 0.050% by weight, and is added to deoxidize the steel.

Niobium (Nb): Niobium has different effects on steel. This retards recrystallization by forming ultrafinely dispersed precipitates in finishing, which increases the density of germination points during hot rolling and produces finer grains after transformation. The proportion of dissolved niobium inhibits recrystallization. Precipitates in the final product increase strength. These precipitates may be carbides or carbonitrides. Often these precipitates are mixed carbides to which titanium may be incorporated. This effect begins to appear at 0.0050%, and is most prominent in niobium greater than 0.010% and less than or equal to 0.050% by weight. Precipitates also prevent grain growth during (partial) austenitization in hot dip galvanizing. Niobium exceeding 0.040% by weight is not expected to have an additive effect. It has proven advantageous with respect to the effect achieved by a niobium content of 0.005% to 0.040% by weight. A content of 0.015% to 0.035% by weight is optimal.

Titanium (Ti): Due to its high affinity for nitrogen, titanium precipitates mainly as TiN during solidification. It also occurs as a mixed carbide with niobium. TiN is very important for grain size stability in a kiln. Since these precipitates have high-temperature stability, unlike mixed carbides, they mostly exist as particles that hinder crystal grain growth at 1200°C. Titanium also has a retarding effect on recrystallization during hot rolling, but less so than niobium. Titanium works through precipitation hardening. However, larger TiN particles are less effective than more finely dispersed mixed carbides. The best efficacy is achieved in the range of 0.005 to 0.030% by weight of titanium, advantageously in the range of 0.005 to 0.025% by weight of titanium.

Boron (B): Boron is a very effective alloying agent for improving hardenability even in very small amounts (more than 5 ppm). The martensitic initiation temperature is not affected. To be effective, boron must be present in solid solution. Due to its high affinity for nitrogen, preferably nitrogen must first be combined with the stoichiometrically required amount of titanium. Due to its low solubility in iron, dissolved boron is preferentially present at austenite grain boundaries. Here, boron partially forms Fe-B carbide, which is coherent and lowers the grain boundary energy. Both effects retard the formation of ferrite and pearlite and thus increase the hardenability of the steel. However, excessive amounts of boron are detrimental as they can form iron borides which negatively affect the material's hardenability, formability and toughness. Boron also has a tendency to form oxides or mixed oxides during continuous hot dip coating that degrade the quality of hot dip galvanizing. The aforementioned measures for controlling the furnace area in continuous hot dip coating reduce the formation of oxides on the steel surface.

For the above reasons, the boron content for the alloy concept of the present invention is set to a value of 5 to 30 ppm, optimally 5 to 20 ppm.

Nitrogen (N) can be an alloying element and incidental element in steel production. An excessively high nitrogen content causes an increase in strength with an aging effect as well as a rapid loss of toughness. On the other hand, fine grain hardening through titanium nitride and niobium (carbon) nitride can be achieved by targeted addition of nitrogen together with the microalloy elements titanium and niobium. In addition, formation of coarse crystal grains is suppressed during reheating before hot rolling.

According to the present invention, therefore, the N-content is set to a value of 0.0020 to 0.0120% by weight.

In order to maintain the required properties of the steel, it has been found to be advantageous if the nitrogen content is adjusted as a function of the sum of Ti + Nb + B.

When the total content of Ti + Nb + B is 0.010 to 0.050% by weight, the content of nitrogen should be maintained at a value of 20 to 90 ppm. A nitrogen content of 40 to 120 ppm has proven advantageous if the total content of Ti + Nb + B exceeds 0.050% by weight.

For the combined content of niobium and titanium, a content of less than 0.065% by weight has proven advantageous, and a content of less than 0.055% by weight is particularly advantageous for cost reasons due to the fact that niobium and titanium are interchangeable up to a minimum niobium content of 10 ppm.

Regarding the interaction of the microalloy elements niobium and titanium with boron, a total content of 0.070% by weight or less is advantageous, and a total content of 0.060% by weight or less is particularly advantageous. A higher content has no further improving effect in the sense of the present invention.

In addition, a maximum total content of Ti+Nb+Mo+B of 0.175% by weight has proven advantageous for the reasons mentioned above.

Calcium (Ca): The addition of calcium in the form of a calcium-silicon mixed compound causes deoxidation and desulfurization of the molten phase during steel manufacture. The reaction products are thus transferred into the slag, and the steel is purified. Increased purity leads to better properties according to the invention of the final product.

For the reasons described above, a Ca content of 0.005 to 0.0060% by weight, advantageously 0.0030% by weight, is set.

In tests conducted on steels according to the present invention, a minimum tensile strength of 750 MPa with a thickness of 0.50 to 3.00 mm in the case of controlled cooling after annealing in the transformation zone between A _c1 and A _c3 or austenitizing annealing on top of A _c3 It has been found that a dual phase steel with strength (eg for cold strip) can be produced, which is characterized by sufficient tolerance for process variation.

The result is a significantly broader process window for the alloy composition according to the present invention compared to known alloy concepts.

The annealing temperature for the dual-phase structure to be achieved is about 700 to 950 ° C for the steel according to the invention, so that depending on the temperature range a partially austenitic (two-phase region) or fully austenitic structure (austenitic region) is achieved .

Test results The microstructure proportions formed after annealing in the transformation zone between A _c1 and A _c3 or after controlled cooling after austenitization annealing on top of A _c3 are obtained in an additional process step, for example, zinc at a temperature of 400 to 470 ° C. or even after "hot coating" with zinc-magnesium.

The continuously annealed, and optionally hot-dipped refined material may be in the skin-passed (cold re-rolled) condition or non-skin-passed rolled condition and/or in the tension-rectified condition or non-tension-rectified condition, as well as hot strip. condition, and also as cold re-rolled hot strip or cold strip in the heat-treated condition (overaged). Hereinafter, this state is referred to as an initial state.

The steel strip produced from the alloy composition according to the invention, in this case hot strip, cold rerolled hot strip or cold strip, is further characterized by a high resistance to near-edge crack formation during further processing.

A very small difference in the property values of the steel strip in the longitudinal and transverse directions relative to the rolling direction of the strip is advantageous in the subsequent use of the material. Thus, the plate can be cut from the strip regardless of the rolling direction (eg, transverse, longitudinal and diagonal directions, or at an angle to the rolling direction), and waste can be minimized.

In order to ensure the cold-rollability of the hot-rolled strip produced from the steel according to the present invention, the hot-rolled strip according to the present invention is subjected to a final rolling temperature in the austenite range exceeding _AC3 and coiling exceeding the bainite initiation temperature. temperature (variant A).

For example, in the case of hot strip or cold re-rolled hot strip with a cold reduction of about 10%, the hot-rolled strip according to the present invention has a final rolling temperature in the austenitic region and a bainite starting temperature exceeding A _c3 . (Variant B).

Additional features, advantages and details of the present invention will become apparent from the following description of exemplary embodiments illustrated in the drawings.

1 is a schematic process flow diagram for manufacturing a strip from a steel according to the present invention.
Figure 2 is a schematic time-temperature profile of process steps of exemplary hot rolling and cold rolling (optional) and continuous annealing, component manufacturing, heat treatment (air hardening) and tempering (optional) for a steel according to the present invention.
3 is the chemical composition of the studied steel.
FIG. 4A shows mechanical property values (along the rolling direction) as air-hardened, non-tempered target values.
Figure 4b is the mechanical properties (along the rolling direction) of the treated steel in the initial state.
Figure 4c shows the mechanical properties (along the rolling direction) of the treated steel in an air hardened, untempered state.
5 is a result of a hole expansion test according to ISO 16630 and a plate bending test according to VDA 238-100 for the steel according to the present invention.
6A is a temperature-time curve (annealing variant, schematic) of Method 1.
6B is a temperature-time curve (annealing variant, schematic) of Method 2.
6C is a temperature-time curve (annealing variant, schematic) of Method 3.

1 shows a schematic diagram of a process flow for producing a strip from steel according to the invention. Various process routes related to the present invention are illustrated. The process path up to hot rolling (final rolling temperature) is the same for all steels according to the invention, after which different process paths are implemented depending on the desired result. For example, pickled hot strip can be galvanized or cold rolled and galvanized with various reduction ratios. Also possible are cold rolled and galvanized hot annealed hot rolled strip or soft annealed cold strip.

Alternatively, it is also possible to treat the material only by continuous annealing without hot-dip refining, ie with or without subsequent electrolytic galvanization. Composite parts can now be made from selectively coated materials. Next, a curing process takes place, wherein cooling is carried out in air according to the invention. Optionally, the tempering step may complete the temperature treatment of the component.

Figure 2 schematically shows the time-temperature profile of the process steps of hot rolling and continuous annealing of a strip produced from an alloy composition according to the invention. Time and temperature dependent transformations are shown for the hot rolling process as well as heat treatment after cold rolling, part making, quenching, tempering and selective tempering.

Figure 3 shows the chemical composition of the steels studied in the upper half of the table. The LH®1100 alloy according to the present invention was compared to the reference grade LH®800/LH®900.

Compared to the reference grades, the alloy according to the invention has in particular a significantly increased content of Si and a lower content of Cr, and no V is added.

In the lower half of FIG. 3 , the total content of the various alloying elements is presented in weight percent, and the carbon equivalent weight (CEV(IIW)) determined for each is listed.

Figure 4 shows the mechanical characteristic values along the rolling direction of the studied steel, achieved in the air-hardened state (Fig. 4a), the initial state without air-hardening (Fig. 4b), and the air-hardened state (Fig. 4c). show The value to be achieved reaches the safe margin.

Figure 5 shows the results (absolute values) of the hole expansion test according to ISO 16630. The results of the hole expansion test of Variant A (coiling temperature above the bainite onset temperature) for Process 2 (FIG. 6B, 1.2 mm) and Process 3 (FIG. 6C, 2.0 mm) are shown.

The materials studied have sheet thicknesses of 1.2 and 2.0 mm, respectively. These results apply to tests according to ISO 16630.

Method 2 corresponds to annealing by hot-dip galvanizing using a combination of a direct fire furnace and a radiant tube furnace, for example, as shown in FIG. 6B.

Method 3 corresponds to process control in a continuous annealing system, for example as shown in FIG. 6C. Also in this case reheating of the steel can be achieved by means of an induction furnace just before the zinc bath.

Different characteristic values or different hole expansion results as well as bending angles are obtained by different temperature profiles according to the invention within the stated range. The main difference is therefore the temperature-time parameters during heat treatment and subsequent cooling.

Figure 6 schematically shows three variants of the temperature-time curve according to the invention upon annealing treatment and cooling, and in each case under various austenitizing conditions.

Method 1 (Fig. 6a) shows the annealing and cooling of a cold rolled or hot rolled or cold rerolled steel strip produced in a continuous annealing line. First, the strip is heated to a temperature in the range of about 700 to 950° C. (A _c1 to A _c3 ). The annealed steel strip is then cooled from this annealing temperature to an intermediate temperature (IT) of about 200 to 250° C. at a cooling rate of about 15 to 100° C./sec. The second intermediate temperature (about 300 to 500° C.) is not shown in this schematic.

The steel strip is then cooled in air at a cooling rate of about 2 to 30° C./sec until it reaches room temperature (RT), or cooling to room temperature is maintained at a cooling rate of about 15 to 100° C./sec.

Method 2 (FIG. 6B) shows the process according to Method 1, but for the purpose of hot-dip finishing, the cooling of the steel strip is intermittently stopped as it passes through the hot-dip vessel and cooled at a cooling rate of about 15 to 100° C./sec. Cool to an intermediate temperature of 250 °C. Next, the steel strip is cooled in air at a cooling rate of about 2 to 30° C./sec until it reaches room temperature.

Method 3 (Fig. 6c) also shows the process according to method 1 in the case of hot-dip refining, but the cooling of the steel strip is stopped briefly (about 1 to 20 seconds) at an intermediate temperature in the range of about 200 to 400 ° C, It is reheated to the temperature required for hot dip dipping (approximately 400 to 470° C.). Next, the steel strip is cooled again to an intermediate temperature of about 200 to 250°C. A final cooling of the steel strip is carried out in air at a cooling rate of about 2 to 30° C./sec until room temperature is reached.

The following examples are used for industrial production of hot dip galvanizing according to method 2 according to FIG. 6b and method 3 according to FIG. 6c involving a laboratory-borated coating process.

Example 1 (cold strip) (alloy composition, weight %)

Method 2 according to variant A/1.2 mm/ FIG. 6B

0.098% C, hot-dip purified according to Method 2 according to FIG. 6B; 0.712% Si; 1.623% Mn; 0.010% P; 0.001% S; 0.0046%N; 0.044 Al; 0.507% Cr; 0.081% Mo; 0.0181% Ti; 0.0306% Nb; 0.0008% B; The steel material according to the present invention with 0.0023% Ca was previously hot-rolled at a final rolling target temperature of 910 ° C, coiled to a thickness of 4.09 mm at a final rolling target temperature of 650 ° C, and additional heat treatment after pickling ( eg without batch annealing).

In the annealing simulator, the hot-dip refined air hardened steel strip was processed using the following parameters.

Annealing temperature 870℃

Hold time 120 seconds

Transfer time up to 5 seconds (no energy input)

subsequent air cooling

After tempering, the steel according to the invention has a microstructure composed of martensite, bainite and retained austenite.

This steel shows the following characteristic values along the rolling direction after air hardening (initial values in the unprocessed state in parentheses), and may correspond to, for example, Lh®1000.

- Yield strength (Rp0.2) 781 MPa (480 MPa)

- Tensile strength (Rm) 1085 MPa (781 MPa)

- Elongation at break (A80) 7.7% (21.1%)

- A5 Elongation 11.0% (-)

- Bake hardening index (BH ₂ ) 58 MPa

- Hole expansion ratio according to ISO 16630 - (42.1%)

- Bending angle (longitudinal, transverse) according to VDA 238-100 - (131°/127°)

The maximum yield ratio Re/Rm in the longitudinal direction was 62% of the initial state.

Example 2 (cold strip) (alloy composition, weight %)

Method 3 according to Variant B/2.0 mm/FIG. 6C

0.097% C, hot-dip purified according to Method 3 according to FIG. 6C; 0.665% Si; 1.894% Mn; 0.009% P; 0.001% S; 0.0052%N; 0.043 Al; 0.319% Cr; 0.005% Mo; 0.0159% Ti; 0.0296% Nb; 0.0017% B; The steel material according to the present invention with 0.0019% Ca was previously hot-rolled at a final rolling target temperature of 910 ° C, coiled at a core coiling temperature of 650 ° C to a thickness of 4.09 mm, and additional heat treatment after pickling ( eg, without batch annealing).

In the annealing simulator, the hot-dip refined steel was treated using the following parameters similar to the temperature treatment process (air hardening).

Annealing temperature 870℃

Hold time 120 seconds

Transfer time up to 5 seconds (no energy input)

subsequent air cooling

After tempering, the steel according to the invention has a microstructure composed of martensite, bainite and retained austenite.

This steel shows the following characteristic values along the rolling direction after air hardening (initial values in the unprocessed state in parentheses), and may correspond to, for example, Lh®1000.

- Yield strength (Rp0.2) 776 MPa (450 MPa)

- Tensile strength (Rm) 1074 MPa (781 MPa)

- Elongation at break (A80) 8.3% (22.3%)

- A5 Elongation 11.1% (-)

- Bake hardening index (BH ₂ ) 53 MPa

- Hole expansion ratio according to ISO 16630 - (35.7%)

- Bending angle (longitudinal, transverse) according to VDA 238-100 - (116°/116°)

The maximum yield ratio Re/Rm in the longitudinal direction was 62% of the initial state.

Reference numerals in FIG. 1
1. Secondary metallurgy
2. Steel plate casting
3. Hot rolling
4. Pickling
5. Softening annealing hot strip (optional)
6. Cold rolling (optional)
7. Double rollers (optional)
8. Softening annealing cold strip (optional)
9. Continuous annealing/hot dip galvanizing
10. Inline Skin Passing
11. Tensile calibration unit
12. Plate cutting
13. Parts Manufacturing
14. Heat treatment process
(air hardening with optional tempering)
Legend of Figure 6a:
ZT = medium temperature
RT = room temperature
Legend of Fig. 6B
ST = hot dip bath temperature
ZT = medium temperature
RT = room temperature

Claims (41)

An ultra-high-strength, air-hardenable, multiphase steel with excellent processing properties and a minimum tensile strength of 750 MPa in the non-air-hardened state for the production of cold-rolled or hot-rolled steel strip with a sheet thickness of 2.00 mm or less, in weight percent ,
C: 0.075 to 0.115
Si: 0.600 to 0.750
Mn: 1.000 to less than 1.700
Cr: 0.200 to 0.600
Al: 0.010 to 0.060
N: 0.0020 to 0.0120
S: 0.0030 or less
Mo: 0.0200 or more
Nb: 0.005 to 0.040
Ti: 0.005 to 0.030
B: 0.0005 to 0.0030
Ca: 0.0005 to 0.0060
Cu: 0.050 or less
Ni: 0.050 or less
the balance iron, and
Contains smelting-related impurities normally associated with steel;
The Mn content and the total content of Mn+Si+Cr+Mo in relation to the widest process window during continuous annealing of the cold-rolled or hot-rolled steel strip made of the above steel are as follows depending on the sheet thickness to be produced. regulated,
- in weight percent at a thickness of 1.00 mm or less, Mn is between 1.000 and 1.500% and the sum of Mn + Si + Cr + Mo is between 2.450 and 2.800%;
- in weight percent at a thickness greater than 1.00 and up to and including 2.00 mm, Mn is between 1.300 and less than 1.700%, and the sum of Mn + Si + Cr + Mo is between 2.600 and 3.150%;
Ultra high strength air hardenable multiphase steel. According to claim 1,
At a strip thickness of 1.00 mm or less, the C content is 0.100% or less, and the carbon equivalent (CEV(IIW)) is 0.50% or less;
Ultra high strength air hardenable multiphase steel. According to claim 1,
At a strip thickness greater than 1.00 and not greater than 2.00 mm, the C content is less than or equal to 0.105% and the carbon equivalent (CEV(IIW)) is less than or equal to 0.55%.
Ultra high strength air hardenable multiphase steel. delete delete delete delete According to claim 1 or 2,
At a strip thickness of 1.00 mm or less, the Cr content is 0.250 to 0.350%.
Ultra high strength air hardenable multiphase steel. According to claim 1 or 3,
At a strip thickness greater than 1.00 mm and less than or equal to 2.00 mm, the Cr content is greater than 0.350 and less than or equal to 0.450%;
Ultra high strength air hardenable multiphase steel. delete According to any one of claims 1 to 3,
When the sum of Ti + Nb + B is 0.010 to 0.050%, the N content is 0.0020 to 0.0090%,
Ultra high strength air hardenable multiphase steel. According to any one of claims 1 to 3,
When the sum of Ti + Nb + B exceeds 0.050%, the N content is 0.0040 to 0.0120%,
Ultra high strength air hardenable multiphase steel. According to any one of claims 1 to 3,
The S content is 0.0020% or less,
Ultra high strength air hardenable multiphase steel. According to any one of claims 1 to 3,
The Mo content is 0.050 to 0.100%,
Ultra high strength air hardenable multiphase steel. According to any one of claims 1 to 3,
The sum of Cr+Mo is 0.800% or less,
Ultra high strength air hardenable multiphase steel. According to any one of claims 1 to 3,
The sum of Cr+Mo is 0.700% or less,
Ultra high strength air hardenable multiphase steel. According to any one of claims 1 to 3,
The Nb content is 0.015 to 0.035%,
Ultra high strength air hardenable multiphase steel. According to any one of claims 1 to 3,
The Ti content is 0.005 to 0.025%,
Ultra high strength air hardenable multiphase steel. According to any one of claims 1 to 3,
The sum of Nb + Ti is 0.065% or less,
Ultra high strength air hardenable multiphase steel. According to any one of claims 1 to 3,
The sum of Nb + Ti is 0.055% or less,
Ultra high strength air hardenable multiphase steel. According to any one of claims 1 to 3,
The B content is 0.0005 to 0.0020%,
Ultra high strength air hardenable multiphase steel. According to any one of claims 1 to 3,
The sum of Ti + Nb + B is 0.070% or less,
Ultra high strength air hardenable multiphase steel. According to any one of claims 1 to 3,
The sum of Ti + Nb + B is 0.060% or less,
Ultra high strength air hardenable multiphase steel. According to any one of claims 1 to 3,
The sum of Ti + Nb + B + Mo is 0.175% or less,
Ultra high strength air hardenable multiphase steel. According to any one of claims 1 to 3,
The Ca content is 0.0030% or less,
Ultra high strength air hardenable multiphase steel. According to any one of claims 1 to 3,
The addition of silicon and manganese in relation to the strength properties to be achieved is
YS(MPa) = 160.7 + 147.9 [% Si] + 161.1 [% Mn]
TS(MPa) = 324.8 + 189.4 [% Si] + 174.1 [% Mn]
interchangeable according to the relationship of
Ultra high strength air hardenable polyphase steel. A method for producing a cold-rolled or hot-rolled steel strip from an air-filled multiphase steel as claimed in claim 1 in which the required microstructure is produced during continuous annealing, comprising:
The cold-rolled or hot-rolled steel strip is heated to a temperature in the range of 700 to 950°C during the continuous annealing, and then the annealed steel strip is heated from the annealing temperature to a temperature of 300 to 500°C at a cooling rate of 15 to 100°C/sec. Cooled to a first intermediate temperature, then cooled to a second intermediate temperature of 160 to 250 °C at a cooling rate of 15 to 100 °C/sec, then the steel strip reaches room temperature at a cooling rate of 2 to 30 °C/sec. or cooling from the first intermediate temperature to room temperature is maintained at a cooling rate of 15 to 100 ° C / sec.
Method for manufacturing steel strip. 28. The method of claim 27,
The cooling is stopped before entering the hot dip bath, and after hot dip refining, the cooling is continued at a cooling rate of 15 to 100° C./sec until an intermediate temperature of 200 to 250° C. is reached, and then The steel strip is cooled in air until room temperature is reached at a cooling rate of 2 to 30° C./sec.
Method for manufacturing steel strip. 29. The method of claim 28,
In the case of hot-dip refining, after heating and cooling to an intermediate temperature of 200 to 250 ° C, the temperature is maintained for 1 to 20 seconds before entering the hot dip bath, then the steel strip is reheated to a temperature of 400 to 470 ° C, , After the hot-dip purification, cooling is carried out at a cooling rate of 15 to 100 ° C / sec to an intermediate temperature of 200 to 250 ° C, followed by cooling in air to room temperature at a cooling rate of 2 to 30 ° C / sec,
Method for manufacturing steel strip. According to claim 28 or 29,
In continuous annealing using a plant configuration consisting of a directly fired furnace (NOF) and a jet tube (RTF), the oxidation potential is increased by controlling the CO content below 4% by volume in the NOF, and iron reduction in the RTF. The oxygen partial pressure in the furnace atmosphere is controlled according to the following equation,
-18 > Log pO ₂ ≥ 5* Si ^-0.3 - 2.2* Mn ^-0.45 - 0.1* Cr ^-0.4 - 12.5*(-lnB) ^0.25
where Si, Mn, Cr, B are the corresponding alloying parts in the steel expressed in weight percent, pO ₂ is the oxygen partial pressure expressed in mbar, and the dew point of the gaseous atmosphere prevents oxidation of the strip immediately before immersion in the hot dip bath. which is set below -30 ° C to prevent
Method for manufacturing steel strip. 31. The method of claim 30,
The oxygen partial pressure in the furnace atmosphere satisfies the following equation in the case of annealing using only one radiant tube furnace,
-12 > Log pO ₂ ≥ 5* Si ^-0.25 - 3* Mn ^-0.5 - 0.1* Cr ^-0.5 - 7*(-lnB) ^0.5
where Si, Mn, Cr, B are the corresponding alloying parts in the steel expressed in weight percent, pO ₂ is the oxygen partial pressure expressed in mbar, and the dew point of the gaseous atmosphere prevents oxidation of the strip immediately before immersion in the hot dip bath. which is set below -30 ° C to prevent
Method for manufacturing steel strip. According to claim 28 or 29,
For strips of different thicknesses during continuous annealing, the comparable microstructural state and mechanical property values of the strips are controlled by adjusting the system throughput during heat treatment.
Method for manufacturing steel strip. 33. The method of claim 32,
The steel strip is skin passed after the heat treatment or hot dip refining,
Method for manufacturing steel strip. 33. The method of claim 32,
The steel strip is stretch bent after the heat treatment or hot dip refining,
Method for manufacturing steel strip. A steel strip produced by the method according to any one of claims 27 to 29,
having a minimum hole-expansion value of 20% according to ISO 16630 in the non-air cured state,
river strip. 36. The method of claim 35,
having a minimum hole-expansion value of 30% according to ISO 16630 in the non-air cured state,
river strip. 36. The method of claim 35,
having a minimum bending angle of 60° according to VDA 238-100 in the longitudinal or transverse direction in the non-air-curing state,
river strip. 36. The method of claim 35,
having a minimum bending angle of 75° according to VDA 238-100 in the longitudinal or transverse direction in the non-air-curing state,
river strip. 36. The method of claim 35,
having a minimum multiplication value Rm x α (tensile strength x bending angle according to VDA 238-100) of 60,000 MPa ° in the non-air-curing state,
river strip. 39. The method of claim 38,
having a minimum multiplication value Rm x α (tensile strength x bending angle according to VDA 238-100) of 70,000 MPa ° in the non-air-curing state,
river strip. 36. The method of claim 35,
having no delayed fracture for at least 6 months meeting the requirements of SEP 1970 for perforated tensile and flexural beam testing;
river strip.

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