EP3221483B1 - High strength air hardenable multi phase steel with excellent workability and sheet manufacturing process thereof - Google Patents

Description

The invention relates to a method for producing a cold-rolled or hot-rolled steel strip from an air-hardenable multiphase steel according to claim 1, and advantageous developments according to claims 2 to 20.

In particular, the invention relates to steels with a tensile strength in the range of at least 950 MPa in the non-tempered state for the production of components which have improved formability (such as increased hole expansion and increased bending angle) and improved welding properties.

A tempering treatment of these steels according to the invention can increase the yield strength and tensile strength, for example by air hardening with optional subsequent tempering.

The hotly contested automotive market is constantly forcing manufacturers to find solutions to reduce fleet consumption and CO ₂ emissions while maintaining the greatest possible comfort and occupant protection. On the one hand, the weight reduction of all vehicle components plays a decisive role, on the other hand, the best possible behavior of the individual components with high static and dynamic loads both during use and in the event of a crash.

By providing high-strength to ultra-high-strength steels and reducing the sheet thickness, the weight of the vehicles can be reduced with improved forming and component behavior during production and operation.

High-strength to ultra-high-strength steels must therefore have comparatively high requirements with regard to their strength and ductility, energy absorption and processing, such as punching, hot and cold forming, thermal hardening (e.g. air hardening, press hardening), welding and / or surface treatment, e.g. metallic finishing, organic coating or painting are sufficient.

Newly developed steels must therefore face the increasing weight requirements due to reduced sheet thickness, the increasing material requirements for yield strength, tensile strength, strengthening behavior and elongation at break with good processing properties such as formability and weldability.

For such a reduction in sheet thickness, high-strength, high-strength steel with a single-phase or multi-phase structure must be used to ensure sufficient strength of the motor vehicle components and to meet the high component requirements with regard to toughness, edge crack resistance, improved bending angle and bending radius, energy absorption as well as strengthening and bake hardening Effect to suffice.

There is also an increasing demand for improved suitability for joining in the form of better general weldability, such as a larger usable welding area for resistance spot welding and an improved failure behavior of the weld seam (fracture pattern) under mechanical stress, as well as sufficient resistance to delayed hydrogen embrittlement (ie delayed fracture free). The same applies to high strength steels in the fabrication of pipes for the weldability, which are produced for example by means of H och f requenz- I nduktionsschweißverfahrens (HFI).

The hole expansion capacity is a material property that describes the resistance of the material to crack initiation and crack propagation during forming operations in areas close to edges, such as when pulling a collar.

The hole expansion test is regulated, for example, in ISO 16630. Then prefabricated holes punched into a sheet, for example, are expanded using a mandrel. The measured variable is the change in the hole diameter in relation to the initial diameter at which the first crack through the sheet occurs at the edge of the hole.

Improved edge crack resistance means an increased formability of the sheet edges and can be described by an increased hole expansion capacity. This situation is known under the synonyms "L ow E dge C rack" (LEC) or under "H igh E H ole xpansion" (HHE) and xpand®.

The bending angle describes a material property that gives conclusions about the material behavior during forming operations with dominant bending components (e.g. when folding) or also in the event of crash loads. Increased bending angles thus increase passenger compartment safety. The determination of the bending angle (α) is e.g. normatively regulated via the plate bending test in VDA 238-100.

The above-mentioned properties are important for components that e.g. can be formed into very complex components by air hardening with optional tempering.

As is known, improved weldability is achieved, among other things, by a reduced carbon equivalent. This is synonymous with synonyms such as " u erter p eritektisch" (UP) or the already well-known " L ow C arbon E equivalent" (LCE). The carbon content is usually less than 0.120% by weight. Furthermore, the failure behavior or the fracture pattern of the weld seam can be improved by alloying with micro-alloying elements.

High-strength components must be sufficiently resistant to embrittlement of hydrogen. Test of resistance of A dvanced H igh S trength S Teels (AHSS) for the automotive industry with respect to hydrogen-induced production-related brittle fractures is regulated in the SEP1970 and tested on the sample and the bracket Lochzugprobe. Dual-phase steels are increasingly being used in vehicle construction, which consist of a ferritic basic structure in which a martensitic second phase is embedded. It has been found that, in the case of low-carbon, micro-alloyed steels, portions of further phases such as bainite and residual austenite are advantageous, for example on the hole expansion behavior, the bending behavior and the affect hydrogen-induced brittle fracture behavior. The bainite can be present in different forms, such as upper and lower bainite.

The specific material properties of the dual-phase steels, e.g. Low yield strength ratios combined with very high tensile strength, strong work hardening and good cold formability are well known, but are often no longer sufficient with increasingly complex component geometries.

In general, the group of multi-phase steels is used more and more. Multi-phase steels include e.g. Complex phase steels, ferritic-bainitic steels, TRIP steels, as well as the previously described dual phase steels, which are characterized by different structural compositions.

Complex phase steels according to EN 10346 steels which small amounts of martensite, retained austenite and / or perlite contained in a ferritic / bainitic basic structure, being caused by a delayed recrystallization or by precipitation of micro-alloying elements, a strong grain refining.

Compared to dual-phase steels, these complex phase steels have higher yield strengths, a greater yield strength ratio, less strain hardening and a higher hole expansion capacity.

According to EN 10346, ferritic-bainitic steels are steels which contain bainite or solidified bainite in a matrix of ferrite and / or solidified ferrite. The strength of the matrix is brought about by a high dislocation density, grain refinement and the excretion of microalloying elements.

According to EN 10346, dual-phase steels are steels with a ferritic basic structure in which a martensitic second phase is embedded in the form of an island, sometimes also with parts of bainite as the second phase. With high tensile strength, dual-phase steels show a low yield ratio and strong strain hardening.

TRIP steels are, according to EN 10346, steels with a predominantly ferritic structure, in which bainite and residual austenite is embedded, which can convert to martensite during the forming (TRIP effect). Due to its strong strain hardening, the steel achieves high values of uniform elongation and tensile strength. In connection with the bake hardening effect high component strengths can be achieved. These steels are suitable for both stretch drawing and deep drawing. However, higher sheet metal holder forces and press forces are required for material forming. A comparatively strong springback must be taken into account.

The high-strength steels with a single-phase structure include e.g. bainitic and martensitic steels.

Bainitic steels are, according to EN 10346, steels which are characterized by a very high yield strength and tensile strength with sufficient elongation for cold forming processes. Due to the chemical composition, it is easy to weld. The structure typically consists of bainite. The structure may occasionally contain small amounts of other phases, such as martensite and ferrite.

According to EN 10346, martensitic steels are steels that contain small amounts of ferrite and / or bainite in a basic structure of martensite through thermomechanical rolling. This steel grade is characterized by a very high yield strength and tensile strength with sufficient elongation for cold forming processes. Within the group of multi-phase steels, the martensitic steels have the highest tensile strength values. The suitability for deep drawing is limited. The martensitic steels are primarily suitable for bending forming processes, such as roll forming.

Heat-treated steels are steels according to EN 10083, obtained by annealing (= hardening and tempering) a high tensile and fatigue strength. If cooling in air leads to bainite or martensite, the process is called "air hardening". A tempering after hardening can be used to influence the strength / toughness ratio.

Areas of application and manufacturing processes

High and high-strength multi-phase steels are used, among others. in structural, chassis and crash-relevant components, as sheet metal blanks, tailored blanks (welded blanks) and as flexible cold-rolled strips, so-called TRB®s or tailored strips.

The T ailor R olled B lank lightweight technology (TRB®) enables a significant weight reduction through a load-adjusted sheet thickness over the component length and / or steel grade.

In the continuous annealing plant, a special heat treatment takes place for the defined structure adjustment, where e.g. due to comparatively soft components such as ferrite or bainitic ferrite, the steel has its low yield strength and due to its hard components such as martensite or carbon-rich bainite, its strength.

For economic reasons, cold-rolled high-strength steel strips are usually recrystallized in a continuous annealing process to form sheet metal that 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 speed (cooling gradients), are set according to the required mechanical-technological properties with the necessary structure.

To set a dual-phase structure, the pickled hot strip in typical thicknesses between 1.50 to 4.00 mm or cold strip in typical thicknesses from 0.50 to 3.00 mm is heated to a temperature in the continuous annealing furnace that during recrystallization and cooling sets the required structure formation.

A constant temperature is difficult to achieve, especially with different thicknesses in the transition area from one belt to another belt. In the case of alloy compositions with process windows that are too small during continuous annealing, this can lead to e.g. the thinner strip is either passed through the furnace too slowly, which reduces productivity, or the thicker strip is passed through the furnace too quickly and the necessary annealing temperatures and cooling gradients are not achieved to achieve the desired structure. The consequences are increased waste and high costs of errors.

Broadened process windows are necessary so that, with the same process parameters, the required strip properties are also possible with larger changes in cross-section of the strips to be annealed.

The problem of a very narrow process window during annealing treatment becomes particularly serious if load-optimized components are to be produced from hot strip or cold strip, which have strip thicknesses that vary over the strip length and strip width (eg due to flexible rolling).

However, the production of TRB®s with a multi-phase structure is not without additional effort with today's known alloys and available continuous annealing systems for widely varying strip thicknesses, e.g. an additional heat treatment before cold rolling (hot strip soft annealing). In areas of different strip thickness, i.e. If there are different degrees of cold rolling, a homogeneous multi-phase structure cannot be set in cold as well as hot-rolled steel strips due to a temperature gradient occurring in the usual alloy-specific narrow process windows.

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

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

Depending on the alloy concept, even with hot strip, the required structure is only set during the annealing treatment in the continuous annealing furnace in order to achieve the required mechanical properties.

Decisive process parameters are therefore the setting of the annealing temperature and the speed, as well as the cooling rate (cooling gradient) in continuous annealing, since the phase change takes place depending on the temperature and time. The less sensitive the steel is to the uniformity of the mechanical properties when there are changes in temperature and time during continuous annealing, the larger the process window.

For continuous annealing of hot or cold rolled steel strips of different thicknesses, for example from the published documents EP 2 028 282 A1 , WO 2013/113304 A2 or EP 2 031 081 A1 Known alloy concepts for a dual-phase steel, there is the problem that with these alloy compositions the required mechanical properties are met, but only a narrow process window for the annealing parameters is available in order to achieve uniform mechanical properties in the event of cross-sectional changes, e.g. changes in width or thickness, without adapting the process parameters to be able to adjust the belt length.

When using the known alloy concepts, it is difficult to achieve uniform mechanical properties over the entire strip length and strip width, even during continuous annealing of strips of different thicknesses, due to the narrow process window.

In the case of flexibly rolled cold-rolled strips made of known steel alloys, the areas with a smaller strip thickness due to the conversion processes during cooling either have too high strengths due to excessively high martensite contents, or the areas with greater strip thickness achieve insufficient strengths due to insufficiently low martensite contents due to the process window being too small. Homogeneous mechanical-technological properties across the strip length or width can practically not be achieved with the known alloy concepts for continuous annealing.

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

In steel production, there is a trend towards reducing the carbon equivalent in order to achieve improved cold processing (cold rolling, cold forming) and better performance properties.

But the suitability for welding, which is also characterized by the carbon equivalent, is an important assessment parameter.

• CEV(IIW) = C + Mn/6 + (Cu + Ni)/15 + (Cr + Mo + V)/5
• CET = C + (Mn + Mo)/10 + (Cr + Cu)/20 + Ni/40
• PCM = C + (Mn + Cu + Cr)/20 + Ni/60 + Mo/15 + V/10 + 5 B

die charakteristischen Standardelemente, wie Kohlenstoff und Mangan, sowie Chrom bzw. Molybdän und Vanadium berücksichtigt (Gehalte in Gew.-%).For example, in the following carbon equivalents

• CEV (IIW) = C + Mn / 6 + (Cu + Ni) / 15 + (Cr + Mo + V) / 5
• CET = C + (Mn + Mo) / 10 + (Cr + Cu) / 20 + Ni / 40
• PCM = C + (Mn + Cu + Cr) / 20 + Ni / 60 + Mo / 15 + V / 10 + 5 B

the characteristic standard elements, such as carbon and manganese, as well as chromium or molybdenum and vanadium are taken into account (contents in% by weight).

Silicon only plays a subordinate role in the calculation of the carbon equivalent. This is crucial in relation to the invention. The lowering of the carbon equivalent due to lower carbon and manganese contents is to be compensated for by increasing the silicon content. Thus, the edge crack resistance and the weldability are improved with the same strength.

A low yield strength ratio (Re / Rm) in a strength range above 950 MPa in the initial state is typical for a dual-phase steel and is primarily used for the formability during stretching and deep-drawing processes. It gives the designer information about the distance between the onset of plastic deformation and failure of the material under quasi-static stress. Accordingly, lower yield strength ratios represent a greater safety margin from component failure.

A higher yield strength ratio (Re / Rm), as is typical for complex phase steels, is also characterized by a high resistance to edge cracks. This can be attributed to the smaller differences in the strength and hardness of the individual structural components and the finer structure, which has a favorable effect on a homogeneous deformation in the area of the cut edge.

Regarding the yield strength, there is an overlap area in the standards, as well as the yield strength ratio (Re / Rm), in which an assignment to both complex and dual-phase steels is possible and leads to improved material properties.

The analytical landscape for achieving multi-phase steels with a minimum tensile strength of 950 MPa is very diverse and shows very large alloy ranges for the strength-increasing elements carbon, silicon, manganese, phosphorus, nitrogen, aluminum as well as chromium and / or molybdenum as well as in the addition of micro alloys such as Titanium, niobium, vanadium and boron.

The range of dimensions in this strength range is wide and lies in the thickness range from approximately 0.50 to approximately 4.00 mm for strips which are intended for continuous annealing. Hot strip, cold-rolled hot strip and cold strip can be used as primary material. Tapes up to about 1600 mm wide are mainly used, but also Slit strip dimensions that result from slitting the strips lengthways. Sheets or sheets are made by cross-cutting the strips.

For example from the scriptures EP 1 807 544 B1 , WO 2011/000351 and EP 2 227 574 B1 Well-known air-hardenable steel grades with minimum tensile strengths of 800 (LH®800) or 900 MPa (LH®900) in hot or cold-rolled design are particularly characterized by their very good formability in the soft state (deep-drawing properties) and by their high strength after heat treatment ( Remuneration).

During hardening, the structure of the steel is transferred to the austenitic area by heating, preferably to temperatures above 950 ° C. in a protective gas atmosphere. Subsequent cooling in air or protective gas leads to the formation of a martensitic structure for a high-strength component.

The subsequent tempering enables the reduction of residual stresses in the hardened component. At the same time, the hardness of the component is reduced so that the required toughness values are achieved.

The invention is therefore based on the object of creating a new cost-effective alloy concept for a high-strength air-hardenable multiphase steel with excellent processing properties and with a minimum tensile strength of 950 MPa in the non-tempered state, lengthways and crosswise to the rolling direction, preferably with a dual-phase structure, with which the process window for continuous annealing of hot or cold rolled strips has been expanded so that in addition to strips with different cross-sections, steel strips with a thickness and strip width that varies over the strip length and the correspondingly varying degrees of cold rolling with the most homogeneous mechanical and technological properties can be produced.

In addition, the hot-dip coating of the steel is to be guaranteed and a method for producing a strip made from this steel is to be specified.

Adequate formability, HFI weldability, excellent general weldability and resistance to hot-dip and tempering should also be ensured.

• C ≥ 0,075 bis ≤ 0,115
• Si ≥ 0,400 bis ≤ 0,500
• Mn ≥ 1,900 bis ≤ 2,350
• Cr ≥ 0,200 bis ≤ 0,500
• Al ≥ 0,005 bis ≤ 0,060
• N ≥ 0,0020 bis ≤ 0,0120
• S ≤ 0,0030
• Nb ≥ 0,005 bis ≤ 0,060
• Ti ≥ 0,005 bis ≤ 0,060
• B ≥ 0,0005 bis ≤ 0,0030
• Mo ≥ 0,200 bis ≤ 0,300
• Ca ≥ 0,0005 bis ≤ 0,0060
• Cu ≤ 0,050
• Ni ≤ 0,050

According to the teaching of the invention, this object is achieved by a steel with the following chemical composition in% by weight:

• C ≥ 0.075 to ≤ 0.115
• Si ≥ 0.400 to ≤ 0.500
• Mn ≥ 1,900 to ≤ 2,350
• Cr ≥ 0.200 to ≤ 0.500
• Al ≥ 0.005 to ≤ 0.060
• N ≥ 0.0020 to ≤ 0.0120
• S ≤ 0.0030
• Nb ≥ 0.005 to ≤ 0.060
• Ti ≥ 0.005 to ≤ 0.060
• B ≥ 0.0005 to ≤ 0.0030
• Mo ≥ 0.200 to ≤ 0.300
• Ca ≥ 0.0005 to ≤ 0.0060
• Cu ≤ 0.050
• Ni ≤ 0.050

• bis 1,00 mm: Summe aus Mn+Si+Cr ≥ 2,800 und ≤ 3,000%
• über 1,00 bis 2,00 mm: Summe aus Mn+Si+Cr ≥ 2,850 und ≤ 3,100%
• über 2,00 mm: Summe aus Mn+Si+Cr ≥ 2,900 und ≤ 3,200%

Remainder iron, including common steel-related, melting-related impurities, in which the total content of Mn + Si + Cr is set as follows, with a view to the widest possible process window for continuous annealing of hot or cold strips made of this steel:

• up to 1.00 mm: sum of Mn + Si + Cr ≥ 2,800 and ≤ 3,000%
• over 1.00 to 2.00 mm: sum of Mn + Si + Cr ≥ 2.850 and ≤ 3.100%
• over 2.00 mm: sum of Mn + Si + Cr ≥ 2,900 and ≤ 3,200%

Due to the possibility of hot-dip coating (eg hot-dip galvanizing) of steel strips made from the steel according to the invention with high silicon contents of up to 0.500%, which is described in process claims 24 and 25, it is not necessary to add vanadium to ensure the tempering resistance.

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

The steel is characterized by low carbon equivalents and, with the carbon equivalent CEV (IIW), is dependent on the sheet thickness for the addition of max. 0.66% limited, so that excellent weldability and the further specific properties described below can be achieved. A CEV (IIW) value of max. 0.62%, for sheet thicknesses up to 2.00 mm a value of max. 0.64% and above 2.00 mm a value of max. Exposed to 0.66%.
Due to its chemical composition, the steel can be manufactured in a wide range of hot rolling parameters, for example with reel temperatures above the bainite start temperature (variant A). In addition, by means of a targeted process control, a microstructure can be set which then allows the steel according to the invention to be cold rolled without prior soft annealing, with cold rolling degrees of between 10 and 40% being used per cold rolling pass.

The steel is very well suited as a primary material for hot-dip coating and, due to the sum-related amount of Mn, Si and Cr added according to the invention depending on the strip thickness to be produced, has a significantly enlarged process window compared to the known steels.

Experiments have surprisingly shown that a wide process window with the required mechanical properties can be maintained if the total content of Mn + Si + Cr is set depending on the sheet thickness.
For a steel with a lower minimum tensile strength, a corresponding dependency in the DE 10 2012 013 113 A1 described.

This results in increased process reliability when continuous annealing cold and hot strip with dual or multi-phase structure. Therefore, for continuously annealed hot or cold strips, more homogeneous mechanical and technological properties can be set in the strip, even with different cross sections and otherwise the same process parameters. This applies to continuous annealing of successive strips with different strip cross sections, as well as for strips with varying strip thickness over strip length or strip width. For example, this means processing in selected thickness ranges possible (e.g. less than 1.00 mm tape thickness, 1.00 mm to 2.00 mm tape thickness and over 2.00 mm tape thickness).

If, according to the invention, high-strength hot or cold strips made of multi-phase steel with varying strip thicknesses are produced in the continuous annealing process, load-optimized components can advantageously be produced therefrom.

The steel strip according to the invention can be produced as cold and hot strip and as cold-rolled hot strip by means of a hot-dip galvanizing line or a pure continuous annealing system in the trained and undressed, in the stretch-bend-oriented and non-stretch-bend-oriented and also in the heat-treated (aged) state.

With the alloy composition according to the invention, steel strips can be produced by an intercritical annealing between A _c1 and A _c3 or in the case of an austenitizing annealing over A _c3 with a final controlled cooling, which leads to a dual or multi-phase structure.

Annealing temperatures of approximately 700 to 950 ° C. have proven to be advantageous. Depending on the overall process (only continuous annealing or additional hot-dip coating), there are different approaches to heat treatment.

In a continuous annealing plant without subsequent hot-dip coating, the strip is cooled from the annealing temperature with a cooling rate of approx. 15 to 100 ° C / s to an intermediate temperature of approx. 160 to 250 ° C. Optionally, you can cool down to a previous intermediate temperature of 300 to 500 ° C beforehand with a cooling rate of approx. 15 to 100 ° C / s. Cooling down to room temperature is finally carried out at a cooling rate of approx. 2 to 30 ° C / s (see also method 1, Figure 6a ).

With heat treatment as part of hot-dip coating, there are two options for temperature control. The cooling is stopped, as described above, before entering the molten bath and is continued only after the bath has exited until the intermediate temperature of about 200 to 250 ° C. has been reached. Depending on the molten bath temperature, the holding temperature in the molten bath is approximately 400 up to 470 ° C. Cooling down to room temperature takes place again at a cooling rate of approx. 2 to 30 ° C./s (see also method 2, Figure 6b ).

The second variant of the temperature control for hot-dip coating includes maintaining the temperature for approx. 1 to 20 s at the intermediate temperature of approx. 200 to 350 ° C and then reheating to the temperature required for hot-dip coating of approx. 400 to 470 ° C. After finishing, the strip is cooled again to approx. 200 to 250 ° C. The cooling to room temperature again takes place at a cooling rate of approx. 2 to 30 ° C./s (see also method 3, Figure 6c ).
In known dual-phase steels, manganese, chromium and silicon are responsible for the conversion of austenite to martensite in addition to carbon. Only the combination according to the invention of the elements carbon, silicon, manganese, nitrogen, molybdenum and chromium, as well as niobium, titanium and boron, which are added within the specified limits, on the one hand secures the required mechanical properties such as minimum tensile strengths of 950 MPa with a simultaneously broadened process window in continuous annealing.

It is also characteristic of the material that the addition of manganese with increasing percentages by weight shifts the ferrite area to longer times and lower temperatures during cooling. The proportions of ferrite are reduced to a greater or lesser extent by increasing the proportions of bainite, depending on the process parameter.

By setting a low carbon content of ≤ 0.115% by weight, the carbon equivalent can be reduced, which improves the weldability and prevents excessive hardening during welding. In the case of resistance spot welding, the electrode service life can also be significantly increased.

The effect of the elements in the alloy according to the invention is described in more detail below. Accompanying elements are inevitable and are taken into account in the analysis concept with regard to their effects, if necessary.

Instruction elements are elements that are already present in the iron ore or, due to the manufacturing process, get into the steel. Because of their predominantly negative influences, they are usually undesirable. An attempt is made to remove them to a tolerable level or to convert them into harmless forms.

Hydrogen (H) is the only element that can diffuse through the iron lattice without generating lattice strain. This means that the hydrogen in the iron lattice is relatively mobile and can be absorbed relatively easily during the processing of the steel. Hydrogen can only be absorbed into the iron lattice in an atomic (ionic) form.

Hydrogen has a strong embrittlement effect and diffuses preferentially to energetically favorable places (defects, grain boundaries etc.). Defects act as hydrogen traps and can significantly increase the length of time that hydrogen remains in the material.

Recycle to molecular hydrogen can result in cold cracks. This behavior occurs with hydrogen embrittlement or with hydrogen-induced stress corrosion cracking. Even in the case of a delayed crack, the so-called delayed fracture, which occurs without external stresses, hydrogen is often cited as the triggering cause. Therefore, the hydrogen content in the steel should be as low as possible.

A more uniform structure, which among other things in the steel according to the invention. achieved through its widened process window also reduces the susceptibility to hydrogen embrittlement.

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

On the one hand, there are procedural approaches to reducing oxygen, such as vacuum treatment, and on the other hand, analytical approaches. By adding certain alloying elements, the oxygen can be converted into less dangerous states. It is usually customary to set the oxygen in the course of deoxidizing the steel with manganese, silicon and / or aluminum. The resulting oxides, however, can cause negative properties as defects in the material.

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

Phosphorus (P) is a trace element from iron ore and is dissolved in the iron lattice as a substitute atom . Phosphorus increases hardness through solid-solution hardening and improves hardenability. However, attempts are generally made to lower the phosphorus content as much as possible, since, among other things, due to its low solubility in the solidifying medium, it tends to segregate and to a large extent reduces the toughness. Due to the accumulation of phosphorus at the grain boundaries, grain boundary breaks occur. In addition, phosphorus increases the transition temperature from tough to brittle behavior up to 300 ° C. During hot rolling, near-surface phosphorus oxides can cause tearing at the grain boundaries.

In some steels, however, phosphorus is used in small quantities (<0.1% by weight) as a microalloying element due to the low cost and the high increase in strength, for example in high-strength IF steels (interstitial free), bake hardening steels or in some alloy concepts for dual phase steels. The steel according to the invention differs from known analysis concepts which use phosphorus as a solid solution, inter alia in that phosphorus is not alloyed but is set as low as possible.

For the aforementioned reasons, the phosphorus content in the steel according to the invention is limited to amounts which are unavoidable in the production of steel.

Like phosphorus, sulfur (S) is bound as a trace element in iron ore. Sulfur is undesirable in steel (with the exception of free-cutting steels) because it tends to segregate and has a strong embrittlement effect. An attempt is therefore made to achieve the lowest possible sulfur content in the melt, for example by means of a vacuum treatment. Furthermore, the sulfur present is converted into the relatively harmless compound manganese sulfide (MnS) by adding manganese. The manganese sulfides are often rolled out in rows during the rolling process and act as germination points for the conversion. This leads to a stratified structure, especially in the case of diffusion-controlled conversion, and can lead to deteriorated mechanical properties in the case of pronounced stringency (e.g. pronounced marten seat lines instead of distributed martensite islands, anisotropic material behavior, reduced elongation at break).

For the reasons mentioned above, the sulfur content in the steel according to the invention is limited to ≤ 0.0030% by weight, advantageously to ≤ 0.0025% by weight or optimally to ≤ 0.0020% by weight or to quantities unavoidable in the production of steel .

Alloy elements are usually added to the steel in order to influence certain properties. An alloy element in different steels can influence different properties. The effect generally depends strongly on the amount and the state of the solution in the material.

The relationships can therefore be quite diverse and complex. The effect of the alloying elements will be discussed in more detail below.

Carbon (C) is the most important alloying element in steel. Due to its targeted introduction of up to 2.06% by weight, iron only becomes steel. The carbon content is often drastically reduced during steel production. In the case of dual-phase steels for continuous hot-dip coating, its proportion according to EN 10346 or VDA 239-100 is a maximum of 0.230% by weight, a minimum value is not specified.

Because of its comparatively small atomic radius, carbon is dissolved interstitially in the iron lattice. The solubility is a maximum of 0.02% in α-iron and a maximum of 2.06% in γ-iron. In dissolved form, carbon significantly increases the hardenability of steel and is therefore essential for the formation of a sufficient amount of martensite. Too high a carbon content, however, increases the difference in hardness between ferrite and martensite and limits weldability.

To meet the requirements e.g. To meet high hole expansion and bending angle, the steel according to the invention contains carbon contents of less than or equal to 0.115% by weight.

Due to the different solubility of carbon in the phases, pronounced diffusion processes are necessary during the phase transition, which can lead to very different kinetic conditions. In addition, carbon increases the thermodynamic stability of austenite, which is shown in the phase diagram in an expansion of the austenite area to lower temperatures. As the forcibly dissolved carbon content in the martensite increases, the lattice distortion increases and with it the strength of the diffusion-free phase.

• Materialdicke unter 1,00 mm (C von ≤ 0,100 Gew.-%)
• Materialdicken zwischen 1,00 bis 2,00 mm (C ≤ 0,105 Gew.-%)
• Materialdicken über 2,00 mm (C ≤ 0,115 Gew.-%).

Carbon also forms carbides. A structure phase that occurs in almost every steel is cementite (Fe ₃ C). However, much harder special carbides can form with other metals such as chromium, titanium, niobium, vanadium. Not only the type but also the distribution and size of the excretions is of crucial importance for the resulting increase in strength. In order to ensure sufficient strength on the one hand and good weldability, improved hole expansion, an improved bending angle and sufficient resistance to hydrogen-induced cracking (ie delayed fracture free), the minimum C content is therefore set to 0.075% by weight and the maximum C Content is set at 0.115% by weight, contents with a cross-section-dependent differentiation are advantageous, such as:

• Material thickness below 1.00 mm (C of ≤ 0.100% by weight)
• Material thicknesses between 1.00 and 2.00 mm (C ≤ 0.105% by weight)
• Material thicknesses over 2.00 mm (C ≤ 0.115% by weight).

Silicon (Si) binds oxygen during casting and is therefore used for calming during the deoxidation of the steel. It is important for the later steel properties that the segregation coefficient is significantly lower than, for example, that of manganese (0.16 compared to 0.87). Segregations generally lead to a line arrangement of the structural components, which deteriorate the forming properties, for example the widening of the holes and the ability to bend.

Characteristic of the material, the addition of silicon causes strong solid solution strengthening. Roughly, adding 0.1% silicon increases the tensile strength by approx. 10 MPa, with an addition of up to 2.2% silicon, the elongation deteriorates only slightly. This was investigated for different sheet thicknesses and annealing temperatures. The increase from 0.2% to 0.5% silicon caused an increase in strength of approx. 20 MPa in the yield strength and approx. 70 MPa in the tensile strength. The elongation at break decreases by about 2%. The latter is due, among other things, to the fact that silicon reduces the solubility of carbon in the ferrite and increases the activity of carbon in the ferrite, thus preventing the formation of carbides, which, as brittle phases, reduce ductility, which in turn improves the formability. The low strength-increasing effect of silicon within the range of the steel according to the invention creates the basis for a wide process window.

Another important effect is that silicon postpones the formation of ferrite at shorter times and temperatures, thus allowing sufficient ferrite to form before quenching. In hot rolling, this creates a basis for improved cold rolling. During hot dip refinement, the accelerated ferrite formation enriches the austenite with carbon and thus stabilizes it. Because silicon hinders carbide formation, the austenite is additionally stabilized. The accelerated cooling can thus suppress the formation of bainite in favor of martensite.

The addition of silicon in the range according to the invention has led to further surprising effects described below. The delay in carbide formation described above could e.g. can also be brought about by aluminum. However, aluminum forms stable nitrides, so that insufficient nitrogen is available for the formation of carbonitrides with microalloying elements. This problem does not exist due to the alloying with silicon, since silicon forms neither carbides nor nitrides. Silicon thus has an indirect positive effect on the formation of precipitates through microalloys, which in turn have a positive effect on the strength of the material. Since the increase in the transition temperatures due to silicon tends to favor grain coarsening, a microalloy with niobium, titanium and boron is particularly expedient, as is the targeted adjustment of the nitrogen content in the steel according to the invention.

As is known, hot rolling should result in the formation of strongly adhering red scale with higher silicon-alloyed steels and an increased risk of rolled scale, which can influence the subsequent pickling result and pickling productivity. This effect could not be determined in the steel according to the invention with 0.400 to 0.500% silicon if the pickling is advantageously carried out with hydrochloric acid instead of with sulfuric acid.

Regarding the galvanizability of silicon-containing steels, the DE 196 10 675 C1 stated that steels with up to 0.800 wt .-% silicon or up to 2.000 wt .-% silicon are not hot-dip galvanized due to the very poor wettability of the steel surface with the liquid zinc.

In addition to the recrystallization of the hard-rolled strip, the atmospheric conditions during the annealing treatment in a continuous hot-dip coating system result in a reduction in iron oxide, which is found, for example, in the Cold rolling or as a result of storage at room temperature on the surface. However, for alloy components with an affinity for oxygen, such as silicon, manganese, chromium, boron, the gas atmosphere is oxidizing, with the result that segregation and selective oxidation of these elements can occur. The selective oxidation can take place both externally, that is to say on the substrate surface, and internally within the metallic matrix.

It is known that silicon in particular diffuses to the surface during annealing and forms oxides on the steel surface alone or together with manganese. These oxides can prevent contact between the substrate and the melt and prevent or significantly worsen the wetting reaction. As a result, zinc-free spots, so-called "bare spots", or even large areas without coating can occur. Furthermore, the adhesion of the zinc or zinc alloy layer on the steel substrate can be reduced by a deteriorated wetting reaction with the consequence of an inadequate formation of an inhibitor layer. The above-mentioned mechanisms can also apply to pickled hot strip or cold-rolled hot strip.

Contrary to this general technical knowledge, it was surprisingly found in tests that a suitable hot-dip coating of the steel strip and good adhesion of the coating can be achieved solely by using a suitable furnace procedure during recrystallization annealing and when passing through the hot-dip bath.

To this end, it must first be ensured that the strip surface is free of scale residues, pickling or rolling oil or other dirt particles by means of a chemical-mechanical or thermal-hydromechanical pre-cleaning. In order to prevent silicon oxides from reaching the strip surface, methods must also be taken which promote the internal oxidation of the alloy elements below the material surface. Depending on the system configuration, different measures are used here.

In a system configuration in which the Glühprozessschritt exclusively in a radiant tube furnace: is conducted (see method 3 in (r adiant t ube f urnace RTF) Figure 6c ), the internal oxidation of the alloying elements can be influenced in a targeted manner by adjusting the oxygen partial pressure of the furnace atmosphere (N ₂ -H ₂ protective gas atmosphere). The set oxygen partial pressure must satisfy the following equation, the furnace temperature being between 700 and 950 ° C. $$-\mathrm{12th}>\log {\mathrm{pO}}_{\mathrm{2nd}}\ge -5*{\mathrm{Si}}^{-0.25}-\mathrm{3rd}*{\mathrm{Mn}}^{-0.25}-0.1{\mathrm{Cr}}^{-0.5}-7*{\left(-\ln \mathrm{B}\right)}^{0.5}$$

Si, Mn, Cr, B denote the corresponding alloy proportions in the steel in% by weight and pO ₂ the oxygen partial pressure in mbar.

In a system configuration in which the furnace section of a combination of a direct-fired furnace (d irect f ired f urnace: DFF and n on- o f xidizing urnace: NOF), and a subsequent radiant tube furnace (see procedure 2 in Figure 6b ), the selective oxidation of the alloy elements can also be influenced via the gas atmospheres of the furnace areas.

The oxygen partial pressure and thus the oxidation potential for iron and the alloying elements can be set via the combustion reaction in the NOF. This must be set so that the oxidation of the alloy elements takes place internally below the steel surface and, if necessary, a thin iron oxide layer forms on the steel surface after the passage through the NOF area. This is achieved e.g. by reducing the CO value below 4% by volume.

In the subsequent radiant tube furnace, the iron oxide layer that may be formed is reduced under an N2-H2 protective gas atmosphere and, likewise, the alloy elements are further oxidized internally. The oxygen partial pressure set in this furnace area must satisfy the following equation, the furnace temperature being between 700 and 950 ° C. $$-\mathrm{18th}>\log {\mathrm{pO}}_{\mathrm{2nd}}\ge -5*{\mathrm{Si}}^{-0.3}-2.2*{\mathrm{Mn}}^{-0.45}-0.1*{\mathrm{Cr}}^{-0.4}-12.5*{\left(-\ln \mathrm{B}\right)}^{0.25}$$

Si, Mn, Cr, B denote the corresponding alloy proportions in the steel in% by weight and pO ₂ the oxygen partial pressure in mbar.

In the transition area between the furnace → zinc pot (trunk), the dew point of the gas atmosphere (N ₂ -H ₂ protective gas atmosphere) and thus the oxygen partial pressure must be set so that oxidation of the strip before immersion in the molten bath is avoided. Dew points in the range of -30 to -40 ° C have proven to be advantageous.

The above-described measures in the furnace area of the continuous hot-dip coating system prevent the formation of oxides on the surface and achieve uniform, good wettability of the strip surface with the liquid melt.

If instead of hot-dip coating (here, for example, hot-dip galvanizing), the process route is selected via continuous annealing with subsequent electrolytic galvanizing (see process 1 in Figure 6a ), no special precautions are necessary to ensure the galvanizability. It is known that galvanizing higher-alloy steels is much easier to achieve by electrolytic deposition than by continuous hot-dip processes. With electrolytic galvanizing, pure zinc is deposited directly on the strip surface. In order not to hinder the flow of electrons between the steel strip and the zinc ions and thus the galvanizing, it must be ensured that there is no surface-covering oxide layer on the strip surface. This condition is usually guaranteed by a standard reducing atmosphere during annealing and pre-cleaning before electrolysis.

In order to ensure the widest possible process window during annealing and sufficient galvanizability, the minimum silicon content is set at 0.400% by weight and the maximum silicon content at 0.500% by weight.

Manganese (Mn) is added to almost all steels for desulfurization in order to convert the harmful sulfur into manganese sulfides. In addition, manganese increases the strength of the ferrite through solidification of the crystal and shifts the α- / γ-conversion to lower temperatures.

A main reason for alloying manganese in multi-phase steels, e.g. with dual-phase steels, is the significant improvement in hardenability. Due to the diffusion hindrance, the pearlite and bainite transformation is postponed for longer periods and the martensite start temperature is lowered.

At the same time, the addition of manganese increases the hardness ratio between martensite and ferrite. In addition, the structure of the structure is strengthened. A high difference in hardness between the phases and the formation of marten seat lines result in a lower hole expansion capacity, which is synonymous with increased sensitivity to edge cracking.

Like silicon, manganese tends to form oxides on the steel surface during the annealing treatment. Depending on the annealing parameters and the contents of other alloy elements (especially silicon and aluminum), manganese oxides (eg MnO) and / or Mn mixed oxides (eg Mn ₂ SiO ₄ ) can occur. However, with a low Si / Mn or Al / Mn ratio, manganese is to be regarded as less critical, since globular oxides form rather than oxide films. Nevertheless, high manganese levels can have a negative impact on the appearance of the zinc layer and the zinc adhesion. The above-mentioned measures for setting the furnace areas during continuous hot dip coating reduce the formation of Mn oxides or Mn mixed oxides on the steel surface after annealing.

The manganese content is set at 1,900 to 2,350% by weight for the reasons mentioned.

In order to achieve the required minimum strengths, it is advantageous to maintain a differentiation of the manganese content depending on the strip thickness.

With a strip thickness of less than 1.00 mm, the manganese content is preferably in a range between 1,9 1,900 and 2,2 2,200% by weight, with strip thicknesses of 1.00 to 2.00 mm between 2,0 2,050 and 50 2,250% by weight and for strip thicknesses over 2.00 mm between ≥ 2,100% by weight and ≤ 2,350% by weight.

Another special feature of the invention is that the variation in the manganese content can be compensated for by simultaneously changing the silicon content. The increase in strength (. Here, the yield strength, engl y ield s trength, YS) of manganese and silicon is described in generally well by the Pickering equation: $$\mathrm{YS}\left(\mathrm{MPa}\right)=53.9+32.34\left[\mathrm{Weight}.-\%\mathrm{Mn}\right]+83.16\left[\mathrm{Weight}.\%\mathrm{Si}\right]+354.2\left[\mathrm{Weight}.-\%\mathrm{N}\right]+\mathrm{17,402}{\mathrm{d}}^{\left(-1/\mathrm{2nd}\right)}$$

$$\mathrm{TS}\left(\mathrm{MPa}\right)=\mathrm{324,8}+\mathrm{189,4}\left[\mathrm{Gew}.-\%\mathrm{Si}\right]+\mathrm{174,1}\left[\mathrm{Gew}.-\%\mathrm{Mn}\right]$$

However, this is based primarily on the effect of mixed crystal hardening, which according to this equation is weaker for manganese than for silicon. At the same time, however, as mentioned above, manganese significantly increases the hardenability, which means that the proportion of strength-increasing second phase increases significantly in multi-phase steels. Therefore, the addition of 0.1% silicon is a first approximation with the addition of 0.1% manganese in the sense of increasing the strength equate. For a steel of the composition according to the invention and an annealing, which includes time-temperature parameter according to the invention, has the following relationship on an empirical basis for the yield strength and the tensile strength (engl t ensile s trength, TS.) Yield: $$\mathrm{YS}\left(\mathrm{MPa}\right)=160.7+147.9\left[\mathrm{Weight}.-\%\mathrm{Si}\right]+161.1\left[\mathrm{Weight}.-\%\mathrm{Mn}\right]$$

$$\mathrm{TS}\left(\mathrm{MPa}\right)=324.8+189.4\left[\mathrm{Weight}.-\%\mathrm{Si}\right]+174.1\left[\mathrm{Weight}.-\%\mathrm{Mn}\right]$$

Compared to the Pickering equation, the coefficients of manganese and silicon are approximately the same for both the yield strength and the tensile strength, which makes it possible to replace manganese with silicon.

Chromium (Cr) , on the one hand, can significantly increase the hardenability of steel in small quantities in dissolved form. On the other hand, with appropriate temperature control in the form of chromium carbides, chromium causes particle solidification. The associated increase in the number of germ sites with a simultaneously reduced carbon content leads to a reduction in the hardenability.

In dual-phase steels, the addition of chromium mainly improves hardenability. When dissolved, chromium shifts the pearlite and bainite transformation for longer times and at the same time lowers the martensite start temperature.

Another important effect is that chrome significantly increases the temper resistance, so that there is almost no loss of strength in the hot-dip bath.

Chromium is also a carbide former. If chromium-iron mixed carbides are present, the austenitizing temperature before hardening must be selected high enough to dissolve the chromium carbides. Otherwise, the increased number of bacteria can lead to a deterioration in the hardenability.

Chromium also tends to form oxides on the steel surface during the annealing treatment, which can degrade the hot dip quality. The above-mentioned measures for setting the furnace areas during continuous hot dip coating reduce the formation of Cr oxides or Cr mixed oxides on the steel surface after annealing.

The chromium content is therefore set at contents of 0.200 to 0.500% by weight.

Molybdenum (Mo): The addition of molybdenum leads to an improvement in hardenability, similar to that of chromium and manganese. The pearlite and bainite transformation is shifted to longer times and the martensite start temperature is lowered. At the same time, molybdenum is a strong chalk former, which produces finely divided mixed carbides, including with titanium. Molybdenum also significantly increases the tempering resistance, so that no loss of strength is to be expected in the hot-dip bath. Molybdenum also works through mixed crystal hardening, but is less effective than manganese and silicon.

The molybdenum content is therefore set between 0.200 to 0.300% by weight. Ranges between 0.200 and 0.250% by weight are advantageous.

As a compromise between the required mechanical properties and melt suitability, a total Mo + Cr content of von 0.725% by weight has proven to be advantageous for the alloy concept according to the invention.

Copper (Cu) : The addition of copper can increase 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 rate of corrosion.

In combination with oxygen, copper can form harmful oxides at the grain boundaries, which can have negative effects especially for hot forming processes. The copper content is therefore set at ≤ 0.050% by weight and is therefore limited to the amounts that are unavoidable in steel production.

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

Vanadium (V) : Since the addition of vanadium is not necessary in the present alloy concept, the vanadium content is limited to inevitable amounts accompanying the steel.

Aluminum (Al) is usually alloyed to the steel in order to bind the oxygen and nitrogen dissolved in the iron. Oxygen and nitrogen are thus converted into aluminum oxides and aluminum nitrides. These excretions can cause grain refinement by increasing the number of germs and thus increase the toughness properties and strength values.

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

In the dissolved state, aluminum and silicon shift the formation of ferrite at shorter times, thus enabling the formation of sufficient ferrite in dual-phase steel. It also suppresses carbide formation and thus leads to a delayed transformation of the austenite. For this reason, aluminum is also used as an alloying element in residual austenite steels (TRIP steels) to replace some of the silicon. The reason for this approach is that aluminum is somewhat less critical for the galvanizing reaction than silicon.

The aluminum content is therefore limited to 0.005 to a maximum of 0.060% by weight and is added to calm the steel.

Niobium (Nb ): Niobium acts in steel in different ways. When hot rolling in the finishing train, it delays recrystallization due to the formation of very finely divided precipitates, which increases the density of germination points and results in a finer grain after conversion. The proportion of dissolved niobium also inhibits recrystallization. The excretions increase strength in the final product. These can be carbides or carbonitrides. Often it is mixed carbides, in which titanium is also incorporated. This effect starts from 0.005% by weight and is most evident from 0.010% by weight of niobium. The precipitates also prevent grain growth during (partial) austenitization in hot-dip galvanizing. No additional effect is to be expected above 0.060% by weight of niobium. Contents of 0.025 to 0.045% by weight have proven to be advantageous.

Titanium (Ti) : Due to its high affinity for nitrogen, titanium is primarily excreted as TiN during solidification. It also occurs together with niobium as a mixed carbide. TiN is of great importance for grain size stability in the pusher furnace. The Excretions have a high temperature stability, so that, in contrast to the mixed carbides at 1200 ° C, they are mostly present as particles that hinder grain growth. Titanium also retards recrystallization during hot rolling, but is less effective than niobium. Titan works through precipitation hardening. The larger TiN particles are less effective than the more finely distributed mixed carbides. The best effectiveness is achieved in the range from 0.005 to 0.060% by weight of titanium, which is why this represents the alloy range according to the invention. For this, contents of 0.025 to 0.045% by weight have been found to be advantageous.

Boron (B) : Boron is an extremely effective alloying agent to increase hardenability, which is effective even in very small amounts (from 5 ppm). The martensite start temperature remains unaffected. To be effective, boron must be in solid solution. Since it has a high affinity for nitrogen, the nitrogen must first be set, preferably by the stoichiometrically necessary amount of titanium. Due to its low solubility in iron, the dissolved boron preferentially attaches to the austenite grain boundaries. There it partially forms Fe-B carbides, which are coherent and reduce the grain boundary energy. Both effects delay the formation of ferrite and pearlite and thus increase the hardenability of the steel. Excessive levels of boron are harmful, however, since iron boride can form, which has a negative impact on the hardenability, formability and toughness of the material. Boron also tends to form oxides or mixed oxides during annealing during the continuous hot dip coating, which deteriorate the galvanizing quality. The above-mentioned measures for setting the furnace areas during continuous hot dip coating reduce the formation of oxides on the steel surface.

For the aforementioned reasons, the boron content for the alloy concept according to the invention is set at values from 5 to 30 ppm, advantageously at ≤ 25 or optimally at ≤ 20 ppm.

Nitrogen (N) can be an alloying element as well as an accompanying element from steel production. Too high levels of nitrogen lead to an increase in strength combined with a rapid loss of toughness and aging effects. On the other hand, through a targeted addition of nitrogen in combination with the microalloying elements titanium and niobium, fine grain hardening can be achieved using titanium nitride and niobium (carbo) nitride. Coarse grain formation is also suppressed when reheating before hot rolling.

According to the invention, the N content is therefore set to values of 0,00 0.0020 to ≤ 0.0120% by weight.

It has turned out to be advantageous for maintaining the required properties of the steel if the nitrogen content is added as a function of the total of Ti + Nb + B.

With a total content of Ti + Nb + B of ≥ 0.010 to ≤ 0.070% by weight, the nitrogen content should be maintained at values of ≥ 20 to ≤ 90 ppm. For a total content of Ti + Nb + B of> 0.070% by weight, nitrogen contents of ≥ 40 to ≤ 120 ppm have proven to be advantageous.

For the total contents of niobium and titanium, contents of ≤ 0.100% by weight have proven to be advantageous and because of the principle interchangeability of niobium and titanium up to a minimum niobium content of 10 ppm and, for reasons of cost, particularly advantageous of ≤ 0.090% by weight.

In the interaction of the microalloy elements niobium and titanium with boron, total contents of 0 0.102% by weight have proven to be advantageous and particularly advantageous ≤ 0.092% by weight. Higher contents no longer have an improvement in the sense of the invention.

Furthermore, maximum contents of 0,3 0.365% by weight have proven to be total contents of Ti + Nb + V + Mo + B for the reasons mentioned above.

Calcium (Ca) : An addition of calcium in the form of calcium-silicon mixed compounds causes a deoxidation and desulfurization of the molten phase during the production of steel. In this way, reaction products are transferred to the slag and the steel is cleaned. The increased purity leads to better properties according to the invention in the end product.

For the reasons mentioned, a Ca content of ≥ 0.005 to ≤ 0.0060% by weight is established.

In tests carried out with the steel according to the invention, it was found that with an intercritical annealing between A _c1 and A _c3 or an austenitizing one Annealing over A _c3 with final controlled cooling can produce a dual-phase steel with a minimum tensile strength of 950 MPa in a thickness of 0.50 to 3.00 mm (e.g. for cold strip), which is characterized by a sufficient tolerance against process fluctuations.

This results in a significantly wider process window for the alloy composition according to the invention compared to known alloy concepts.

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

The tests also showed that the structural components set after the intercritical annealing between A _c1 and A _c3 or the austenitizing annealing over A _c3 with subsequent controlled cooling also after a further process step of hot-dip coating at temperatures between 400 to 470 ° C, for example with zinc or Zinc-magnesium are retained.

The continuously annealed and, in some cases, hot-dip coated material can be manufactured both as hot strip and as cold-rolled hot strip or cold strip in the trained (cold-rolled) or undressed state and / or in the stretch-oriented or non-stretch-bent state and also in the heat-treated state (aging). This state is referred to below as the initial state.

Steel strips, in the present case as hot strip, cold-rolled hot strip or cold strip, from the alloy composition according to the invention are also distinguished by a high resistance to edge cracking during further processing.

The very small differences in the characteristic values of the steel strip along and across its rolling direction are advantageous when materials are used later. This means that blanks can be cut from a strip regardless of the direction of rolling (for example, transversely, lengthways and diagonally or at an angle to the direction of rolling) and waste is minimized.

In order to ensure the cold rollability of a hot strip produced from the steel according to the invention, the hot strip is produced according to the invention with finish rolling temperatures in the austenitic region above A _r3 and at reel temperatures above the bainite start temperature (variant A).

In the case of hot strip or cold-rolled hot strip, for example with approximately 10% cold rolling degree, the hot strip is produced according to the invention with finish rolling temperatures in the austenitic region above A _r3 and coiling temperatures below the bainite start temperature (variant B).

Further features, advantages and details of the invention result from the following description of exemplary embodiments shown in a drawing.

Figur 1:

Prozesskette (schematisch) für die Herstellung eines Bandes aus dem erfindungsgemäßen Stahl

Figur 2:

Zeit-Temperatur-Verlauf (schematisch) der Prozessschritte Warmwalzen und Kaltwalzen (optional) sowie Durchlaufglühen, Bauteilfertigung, Vergüten (Lufthärten) und Anlassen (optional) beispielhaft für den erfindungsgemäßen Stahl

Figur 3:

Chemische Zusammensetzung der untersuchten Stähle

Figur 4a:

Mechanische Kennwerte (längs zur Walzrichtung) als Zielwerte, luftgehärtet und nicht angelassen

Figur 4b:

Mechanische Kennwerte (längs zur Walzrichtung) der untersuchten Stähle im Ausgangszustand

Figur 4c:

Mechanische Kennwerte (längs zur Walzrichtung) der untersuchten Stähle im luftgehärteten, nicht angelassenen Zustand

Figur 5:

Ergebnisse der Lochaufweitungsversuche nach ISO 16630 und des Plättchenbiegeversuchs nach VDA 238-100 an erfindungsgemäßen Stählen

Figur 6a:

Verfahren 1, Temperatur-Zeit-Kurven (Glühvarianten schematisch)

Figur 6b:

Verfahren 2, Temperatur-Zeit-Kurven (Glühvarianten schematisch)

Figur 6c:

Verfahren 3, Temperatur-Zeit-Kurven (Glühvarianten schematisch)

Show it:

Figure 1:

Process chain (schematic) for the production of a strip from the steel according to the invention

Figure 2:

Time-temperature curve (schematic) of the process steps hot rolling and cold rolling (optional) as well as continuous annealing, component production, tempering (air hardening) and tempering (optional) are examples of the steel according to the invention

Figure 3:

Chemical composition of the investigated steels

Figure 4a:

Mechanical parameters (along the rolling direction) as target values, air-hardened and not tempered

Figure 4b:

Mechanical parameters (longitudinal to the rolling direction) of the examined steels in the initial state

Figure 4c:

Mechanical parameters (longitudinal to the rolling direction) of the examined steels in the air-hardened, not tempered state

Figure 5:

Results of the hole expansion tests according to ISO 16630 and the plate bending test according to VDA 238-100 on steels according to the invention

Figure 6a:

Method 1, temperature-time curves (annealing variants schematically)

Figure 6b:

Method 2, temperature-time curves (annealing variants schematically)

Figure 6c:

Method 3, temperature-time curves (annealing variants schematically)

Figure 1 shows schematically the process chain for the production of a strip from the steel according to the invention. The different process routes relating to the invention are shown. The process route is the same for all steels according to the invention until hot rolling (final rolling temperature), after which process routes differ depending on the desired results. For example, the pickled hot strip can be galvanized or cold rolled and galvanized with different degrees of rolling. Soft-annealed hot strip or soft-annealed cold strip can also be cold-rolled and galvanized.

Optionally, material can also be processed without hot-dip coating, i.e. only in the context of continuous annealing with and without subsequent electrolytic galvanizing. A complex component can now be produced from the optionally coated material. This is followed by the hardening process, in which the air is cooled in accordance with the invention. Optionally, a tempering stage can complete the thermal treatment of the component.

Figure 2 shows schematically the time-temperature profile of the process steps hot rolling and continuous annealing of strips from the alloy composition according to the invention. It shows the time and temperature-dependent conversion for the hot rolling process as well as for heat treatment after cold rolling, component production, tempering and optional tempering.

Figure 3 shows the chemical composition of the investigated steels in the upper half of the table. Alloys LH®1100 according to the invention were compared with the reference grades LH®800 / LH®900.

Compared to the reference grades, the alloys according to the invention in particular have significantly higher Si contents and lower Cr contents and no V alloy.

In the lower half of the table Figure 3 the total contents of various alloy components in% by weight and the respectively determined carbon equivalent CEV (IIW) are listed.

Figure 4 shows the mechanical parameters along the rolling direction of the investigated steels, with target values to be achieved for the air-hardened state ( Figure 4a ), the determined Values in the non-air-hardened initial state ( Figure 4b ) and in air-hardened condition ( Figure 4c ). The specified values to be achieved are safely achieved.

Figure 5 shows results of hole expansion tests according to ISO 16630 (absolute values). The results of the hole expansion tests for variant A (reel temperature above bainite start temperature) are shown for method 2 ( Figure 6b , 1 , 2nd mm) and method 3 ( Figure 6c , 2.0 mm).

The investigated materials have a sheet thickness of 1.2 or 2.0 mm. The results apply to the test according to ISO 16630.

Method 2 corresponds to annealing, for example on hot-dip galvanizing with a combined direct-fired furnace and radiant tube furnace, as described in Figure 6b is described.

The method 3 corresponds, for example, to a process control in a continuous annealing system as shown in Figure 6c is described. In addition, the steel can optionally be reheated directly in front of the zinc bath using an induction furnace.

Due to the different temperature controls according to the invention within the specified range, different characteristic values or different hole expansion results as well as bending angles result from one another. The main difference is the temperature-time parameters during heat treatment and subsequent cooling.

The Figure 6 shows schematically three variants of the temperature-time profiles according to the invention in the annealing treatment and cooling and in each case different austenitizing conditions.

Procedure 1 ( Figure 6a ) shows the annealing and cooling of the cold or hot-rolled or cold-rolled steel strip produced in a continuous annealing plant. First, the tape is heated to a temperature in the range of about 700 to 950 ° C (Ac1 to Ac3). The annealed steel strip is then cooled from the annealing temperature with a cooling rate between approx. 15 and 100 ° C / s to an intermediate temperature (ZT) of approx. 200 to 250 ° C. This schematic representation does not show a second intermediate temperature (approx. 300 to 500 ° C).

Subsequently, the steel strip is cooled at a cooling rate of between about 2 and 30 ° C / s until reaching the R aum t emperature (RT) in air or cooling at a cooling rate between about 15 and 100 ° C / s up maintain at room temperature.

Procedure 2 ( Figure 6b ) shows the process according to method 1, however, the cooling of the steel strip is temporarily interrupted for the purpose of hot-dip coating when it passes through the hot-dip tank, in order to then cool at a cooling rate between approx. 15 and 100 ° C / s up to an intermediate temperature of approx. 200 continue up to 250 ° C. The steel strip is then cooled in air at a cooling rate between approx. 2 and 30 ° C / s until room temperature is reached.

Procedure 3 ( Figure 6c ) also shows the process according to method 1 for hot-dip coating, but the cooling of the steel strip is interrupted by a short pause (approx. 1 to 20 s) at an intermediate temperature in the range of approx. 200 to 400 ° C and up to the temperature ( ST), which is necessary for hot-dip coating (approx. 400 to 470 ° C), reheated. The steel strip is then cooled again to an intermediate temperature of approx. 200 to 250 ° C. The final cooling of the steel strip takes place at a cooling rate of approx. 2 and 30 ° C / s until the room temperature is reached in air.

For industrial production for hot dip galvanizing according to method 2 Figure 6b and according to method 3 Figure 6c with the laboratory-based remuneration process are the following examples:

Example 1 (cold strip) (alloy composition in% by weight) Variant A / 1.2 mm / method 2 according to FIG. 6b

A steel according to the invention with 0.099% C; 0.461% Si; 2.179% Mn; 0.009% P; 0.001% S; 0.0048% N; 0.040 AI; 0.312% Cr; 0.208% Mo; 0.0292% Ti; 0.0364% Nb; 0.0012% B; 0.0015% Ca according to method 2 accordingly Figure 6b hot-dip coated, the material was previously hot-rolled at a final rolling set temperature of 910 ° C and coiled at a reel set temperature of 650 ° C with a thickness of 2.30 mm and after Pickling without additional heat treatment (such as hood annealing) cold rolled twice with an intermediate thickness of 1.49 mm.

• Glühtemperatur 870°C
• Haltezeit 120 s
• Transportzeit max. 5 s (ohne Energiezufuhr)
• anschließende Abkühlung an Luft

A hot-dip coated, air-hardened steel strip was processed in a glow simulator with the following parameters.

• Annealing temperature 870 ° C
• Holding time 120 s
• Transport time max. 5 s (without energy supply)
• then cooling in air

After tempering, the steel according to the invention has a structure which consists of martensite, bainite and residual austenite.

This steel shows the following characteristic values after air hardening (initial values in brackets, unrefined condition): - yield strength (Rp0.2) 921 MPa (768 MPa) - tensile strength (Rm) 1198 MPa (984 MPa) - elongation at break (A80) 5.5% (10.7%) - A5 stretch 9.5% (-) - Bake hardening index (BH2) 52 MPa - Hole expansion ratio according to ISO 16630 - (49%) - Bending angle according to VDA 238-100 (lengthways, crossways) - (122 ° / 112 °) longitudinal to the rolling direction and would correspond to an LH®1100, for example.

The yield point ratio Re / Rm in the longitudinal direction was 78% in the initial state.

Example 2 (cold strip) (alloy composition in% by weight) Variant B / 2.0 mm / method 3 according to FIG. 6c

A steel according to the invention with 0.100% C; 0.456% Si; 2.139% Mn; 0.010% P; 0.001% S; 0.0050% N; 0.058 AI; 0.313% Cr; 0.202% Mo; 0.0289% Ti; 0.0337% Nb; 0.0009% B; 0.0021% Ca according to method 3 accordingly Figure 6c hot-dip coated, the material was previously hot-rolled at a final rolling temperature of 910 ° C and at a The nominal reel temperature of 650 ° C is coiled with a thickness of 2.30 mm and cold-rolled after pickling without additional heat treatment (such as hood annealing).

• Glühtemperatur 870°C
• Haltezeit 120 s
• Transportzeit: max. 5 s (ohne Energiezufuhr)
• Anschließende Abkühlung an Luft

Der erfindungsgemäße Stahl besitzt nach der Vergütung ein Gefüge, welches aus Martensit, Bainit und Restaustenit besteht.The hot-dip coated steel was processed in a glow simulator analogously to a tempering process (air hardening) with the following parameters.

• Annealing temperature 870 ° C
• Holding time 120 s
• Transport time: max. 5 s (without energy supply)
• Subsequent cooling in air

After tempering, the steel according to the invention has a structure which consists of martensite, bainite and residual austenite.

This steel shows the following characteristic values after air hardening (initial values in brackets, unrefined condition): - yield strength (Rp0.2) 903 MPa (708 MPa) - tensile strength (Rm) 1186 MPa (983 MPa) - elongation at break (A80) 7.1% (11.7%) - A5 stretch 9.1% (-) - Bake hardening index (BH2) 48 MPa - Hole expansion ratio according to ISO 16630 - (32%) - Bending angle according to VDA 238-100 (lengthways, crossways) - (104 ° / 88 °) longitudinal to the rolling direction and would correspond to an LH®1100, for example. The yield point ratio Re / Rm in the longitudinal direction was 72% in the initial state.

Claims (30)

1. Cold-rolled or hot-rolled steel strip made of an air-hardenable multiphase steel with a minimum tensile strength in the longitudinal and transverse directions to the rolling direction of 950 MPa before the air-hardening, with excellent processing properties consisting of the elements (contents in % by weight):

   C ≥ 0.075 to 0.115

   Si ≥ 0.400 to ≤ 0.500

   Mn ≥ 1.900 to ≤ 2.350

   Cr ≥ 0.200 to ≤ 0.500

   Al ≥ 0.005 to ≤ 0.060

   N ≥ 0.0020 to ≤ 0.0120

   S ≤ 0.0030

   Nb ≥ 0.005 to ≤ 0.060

   Ti ≥ 0.005 to ≤ 0.060

   B ≥ 0.0005 to ≤ 0.0030

   Mo ≥ 0.200 to ≤ 0.300

   Ca ≥ 0.0005 to ≤ 0.0060

   Cu ≤ 0.050

   Ni ≤ 0.050

   remainder iron, including usual steel-accompanying impurities arising from smelting, in which, with a view to as broad a process window as possible during continuous annealing of hot or cold bands from this steel, the total content of Mn+Si+Cr is set as a function of the strip thickness to be produced, as follows:

   up to 1.00 mm: Sum of Mn+Si+Cr ≥ 2.800 and ≤ 3.000% by weight

   over 1.00 to 2.00 mm: Sum of Mn+Si+Cr ≥ 2.850 and ≤ 3.100% by weight

   over 2.00 mm: Sum of Mn+Si+Cr ≥ 2.900 and ≤ 3.200% by weight
2. Steel strip according to claim 1,
   characterised in that with strip thicknesses of up to 1.00 mm, the C content is ≤ 0.100% and the carbon equivalent CEV (IIW) is ≤ 0.62%.
3. Steel strip according to claim 1,
   characterised in that
   for strip thicknesses of over 1.00 to 2.00 mm, the C content is ≤ 0.105% and the carbon equivalent CEV (IIW) is ≤ 0.64%.
4. Steel strip according to claim 1,
   characterised in that
   for strip thicknesses of over 2.00 mm, the C content is ≤ 0.115% and the carbon equivalent CEV (IIW) is ≤ 0.66%.
5. Steel strip according to claim 1 and 2,
   characterised in that
   with strip thicknesses of up to 1.00 mm, the Mn content is ≥ 1.900 to ≤ 2.200%.
6. Steel strip according to claim 1 and 3,
   characterised in that
   with strip thicknesses of over 1.00 to 2.00 mm, the Mn content is ≥ 2.050 to ≤ 2.250%.
7. Steel strip according to claim 1 and 4,
   characterised in that
   with strip thicknesses of over 2.00 mm, the Mn content is ≥2.100 to ≤ 2.350%.
8. Steel strip according to one of claims 1 to 7,
   characterised in that
   with a total of Ti+Nb+B of ≥ 0.010 to ≤ 0.070%, the N content is ≥ 0.0020 to ≤ 0.0090%.
9. Steel strip according to claim 8,
   characterised in that
   with a total of Ti+Nb+ B of ≥ 0.070% the N content is ≥ 0.0040 to ≤ 0.0120%.
10. Steel strip according to one of claims 1 to 9,
    characterised in that
    the S content is ≤ 0.0025%.
11. Steel strip according to claim 10,
    characterised in that
    the S content is ≤ 0.0020%.
12. Steel strip according to one of claims 1 to 11,
    characterised in that
    the Mo content is ≤ 0.250%.
13. Steel strip according to one of claims 1 to 12,
    characterised in that
    the Ti content is ≥ 0.025 to ≤ 0.045%.
14. Steel strip according to one of claims 1 to 13,
    characterised in that
    the Nb content is ≥ 0.025 to ≤ 0.045%.
15. Steel strip according to one of claims 1 to 14,
    characterised in that
    the sum of Nb+Ti is ≤ 0.100%.
16. Steel strip according to claim 15,
    characterised in that
    the sum of Nb+Ti is ≤ 0.090%.
17. Steel strip according to one of claims 1 to 16,
    characterised in that
    the sum of Cr+Mo is ≤ 0.725%.
18. Steel strip according to one of claims 1 to 17,
    characterised in that
    the sum of Ti+Nb+B is ≤ 0.102%.
19. Steel strip according to claim 18,
    characterised in that
    the sum of Ti+Nb+B is ≤ 0.092%.
20. Steel strip according to one of claims 1 to 19,
    characterised in that
    the Ca content is ≤ 0.0030%.
21. Heat treatment of a cold-rolled or hot-rolled steel strip made of an air-hardenable multiphase steel according to one of claims 1 to 20,
    characterised in that
    the cold- or hot-rolled steel strip is heated to a temperature in the range of approximately 700 to 950°C during the continuous annealing and that the annealed steel strip then cools down from the annealing temperature with a cooling rate between approximately 15 and 100°C/s to a first intermediate temperature of approximately 300 to 500°C, followed by a cooling rate between approximately 15 and 100°C/s to a second intermediate temperature of approximately 160 to 250°C, then the steel strip cools with a cooling speed of approximately 2 to 30°C/s in air until room temperature is reached or the cooling is maintained at a cooling rate between approximately 15 and 100°C/s from the first intermediate temperature to room temperature.
22. Heat treatment according to claim 21,
    characterised in that
    with hot-dip coating after heating and subsequent cooling, the cooling is stopped before entering the melt bath and after hot-dip coating cooling is continued with a cooling rate between approximately 15 and 100°C/s to an intermediate temperature of approximately 200 to 250°C and then the steel strip is cooled in air at a cooling rate of approximately 2 to 30°C/s until room temperature is reached.
23. Heat treatment according to claim 21,
    characterised in that
    in the case of hot-dip coating after heating and then cooling to the intermediate temperature of approximately 200 to 250°C, before entering the melt bath, the temperature is maintained for approximately 1 to 20 s and then the steel strip is reheated to a temperature of approximately 400 to 470°C and after the hot-dip coating has been carried out, cooling takes place at a cooling rate of between approximately 15 and 100°C/s to an intermediate temperature of approximately 200 to 250°C and then cooling takes place in air to room temperature at a cooling rate of approximately 2 to 30°C/s.
24. Heat treatment according to one of claims 22 to 23,
    characterised in that
    in continuous annealing, the oxidation potential of annealing with a system configuration consisting of a directly fired furnace area (NOF) and a radiant tube furnace (RTF) is increased by a CO content in the NOF of less than 4% by volume, wherein the oxygen partial pressure in the RTF of the iron-reducing furnace atmosphere is set according to the following equation, $$-18>\log {\mathrm{pO}}_2⩾-5*{\mathrm{Si}}^{-0.3}-2.2{\mathrm{Mn}}^{-0.45}-0.1*{\mathrm{Cr}}^{-0.4}-12.5*{\left(-\mathrm{lnB}\right)}^{0.25}$$

    

    wherein Si, Mn, Cr, B denote the corresponding alloy proportions in the steel in % by weight and p_O2denotes the oxygen partial pressure in mbar and, in order to avoid the oxidation of the strip, the dew point of the gas atmosphere is set at -30°C or below directly before immersion in the molten bath.
25. Heat treatment according to one of claims 22 to 23,
    characterised in that
    when annealing only with a radiant tube furnace, the oxygen partial pressure of the furnace atmosphere satisfies the following equation, $$-12>\log {\mathrm{pO}}_2\ge -5*{\mathrm{Si}}^{-0.25}-3*{\mathrm{Mn}}^{-0.5}-0.1*{\mathrm{Cr}}^{-0.5}-7*{\left(-\mathrm{lnB}\right)}^{0.5}$$

    

    wherein Si, Mn, Cr, B denote the corresponding alloy proportions in the steel in % by weight and pO₂ denotes the oxygen partial pressure in mbar and, in order to avoid oxidation of the strip, the dew point of the gas atmosphere is set at -30°C or below directly before immersion in the molten bath.
26. Method according to one of claims 21 to 25,
    characterised in that
    the steel strip is temper rolled after the heat treatment or hot-dip coating.
27. Method according to at least one of claims 21 to 26,
    characterised in that
    the steel strip is stretch-levelled after the heat treatment or hot-dip coating.
28. Steel strip produced by the method according to at least one of claims 21 to 27, having a minimum hole expansion value according to ISO 16630 of 25%.
29. Steel strip produced by the method according to at least one of claims 21 to 27, having a minimum bending angle according to VDA 238-100 of 65° in the longitudinal or transverse direction.
30. Steel strip produced by the method according to at least one of claims 21 to 27, having a minimum product value Rm x α (tensile strength x bending angle according to VDA 238-100) of 100000 MPa.

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