KR20200063167A - Ultra high strength multi-phase steel and method of manufacturing steel strips from the multi-phase steel - Google Patents

Abstract

The present invention relates to an ultra high strength multi-phase steel with a minimum tensile strength of 980 MPa, wherein the ultra high strength multi-phase steel (in weight percent) contains: 0.075 <C <0.115; 0.400 ≤ Si ≤ 0.500; 1.900 ≤ Mn ≤ 2.350; 0.250 ≤ Cr ≤ 0.400; 0.010 ≤ Al ≤ 0.060; 0.0020 ≤ N ≤ 0.0120; P ≤ 0.020; S ≤ 0.0020; 0.005 ≤ Ti ≤ 0.060; 0.005 ≤ Nb ≤ 0.060; 0.005 ≤ V ≤ 0.020; 0.0005 ≤ B ≤ 0.0010; 0.200 ≤ Mo ≤ 0.300; 0.0010 ≤ Ca ≤ 0.0060; Cu ≤ 0.050; Ni ≤ 0.050; Sn ≤ 0.040; H ≤ 0.0010; the rest is iron, generally containing steel-related smelting-related impurities, and the total content of Mn-Si+Cr is 1.750 weight in relation to the process window as wide as possible during the annealing of the cold strip of this steel, especially during continuous annealing. % To 2.250% by weight.

Description

Ultra high strength multi-phase steel and method of manufacturing steel strips from the multi-phase steel

The present invention has a dual percentage microstructure or a complex-phase microstructure and a small percentage of residual austenite with improved production and excellent material properties for subsequent processing, especially for lightweight automotive structures according to the preamble of claim 1. It relates to an ultra-high strength multi-phase steel having a. Advantageous development is described in the dependent claims 2 to 24.

The invention also relates to a method for producing a steel strip from such steel according to claim 25 and to a steel strip produced using it according to claim 38.

In particular, the present invention is in a non-quenched state to produce parts with improved strainability in relation to improved joining suitability and hole expansion, for example welding properties. It relates to a steel having a tensile strength in the region of 980MPa or more.

The competitive automotive market means producers must always find solutions to reduce fleet fuel consumption and CO2 emissions while maintaining the highest possible comfort and passenger protection. On the one hand, the weight savings of all vehicle parts play a decisive role as well as the most advantageous possible behavior of the individual components in case of high static and dynamic loads during operation and also in the event of a collision.

Steel suppliers take into account the aforementioned issues by providing ultra-high strength steel. In addition, by providing ultra-high strength steel with a thinner sheet thickness, the weight of vehicle parts can be reduced while the behavior of the parts remains the same or even improves.

This newly developed steel not only reduces the required weight, but also has high material requirements and toughness, edge crack insensitivity and improved bending angle with regard to elastic limit, tensile strength and elongation at break and bake hardening index. And parts requirements with limited solidification in relation to bending radius, energy absorption and work hardening effect and baking hardening effect.

In addition, it is necessary to ensure good processability. This is a process performed by automobile manufacturers such as stamping and deformation, selective heat quenching with subsequent optional tempering, welding and/or surface post-treatment such as phosphatizing and cathodic dip coating. And manufacturing processes performed by suppliers of semi-finished products, for example surface finishes with metal or organic coatings.

Improved bonding suitability, e.g., better weldability and delayed hydrogen embrittlement (for improved spot weld behavior and resistance spot welding), as well as in the form of better general welding capacity Sufficient resistance to delayed hydrogen embrittlement (ie, no delayed fracture) should also be increased. The same applies to welding suitability of ultra-high strength steel in the manufacture of pipes produced by, for example, high-frequency induction welding (HFI).

The ability to expand holes is a material property that describes the resistance of the material to crack initiation and crack propagation in deformation operations, for example, in areas near the edge during plunging.

The hole expansion test is specified, for example, in the ISO 16630 standard. According to this, for example, a prefabricated hole stamped with a metal sheet is expanded by a mandrel. The measurement parameter is the change in hole diameter relative to the initial diameter at which the first crack through the metal sheet occurs at the edge of the hole.

The improved edge crack sensitivity indicates the increased deformability of the sheet edge and can be explained by the increased hole expansion capability. This situation is also referred to synonymously as “low edge crack (LEC)” or “high hole expansion (HHE)” and xpand®.

The bending angle describes the material properties that can lead to conclusions regarding material behavior in the predominant bending ratio (eg during folding) or even in the case of impact loads in deformation motion. Thus, the passenger cabin safety increases as the bending angle increases.

The determination of the bending angle α is defined, for example, by the plate bending test of standard VDA 238-100.

The aforementioned characteristics are important for parts having a very complex shape.

It is known that improved welding ability is achieved in particular by reduced carbon equivalents.

Terms such as “underperitectic (UP)” or the already known “low carbon equivalent (LCE)” are used as synonyms for this. The carbon content is usually less than 0.120% by weight.

In addition, in the case of low carbon steels with reduced carbon equivalents, the fracture behavior or fracture pattern of the weld seam can be improved by significant addition by alloying with fine alloy elements.

High-strength parts must have sufficient resistance to material embrittlement against hydrogen.

Resistance testing of advanced high strength steel (AHSS) in automotive applications for production-related hydrogen-induced brittle fracture is specified in SEP1970 and yoke test piece and hole pull test piece It is tested using.

In vehicle structures, double phase steel is increasingly used, and the steel consists of a ferrite basic microstructure containing a second martensite phase. For low carbon microalloy steels, it has been found that the proportions of additional phases such as bare and residual austenite have beneficial effects on hole expansion behavior, bending behavior and hydrogen-induced brittle fracture behavior. In this case, bainite can appear with different manifestations, for example upper and lower bainite.

The characteristic processing properties of double phase steels such as strong cold solidification and good cold deformability, at the same time very low tensile strength ratio with very high tensile strength are well known.

The combination of properties required for steel ultimately represents a component-specific compromise of individual properties. However, for more complex part geometries, these properties are no longer suitable.

In addition, in the vehicle structure, the use of multi-phase steels such as composite phase steels, ferrite-vanite steels and martensite steels having different structural compositions is increased. According to EN 10346, the composite phase steel is a steel containing a small proportion of martensite, residual austenite and/or pearlite in the ferrite/bainite basic microstructure, and has strong grain refinement by recrystallization delay or precipitation of fine alloy elements. Is generated.

This composite phase steel has a high yield strength, has a higher yield strength ratio, has a low cold solidification and a high hole expansion ability.

Ferrite-bainite steel is a steel containing bainite or coagulated bainite in a matrix of ferrite and/or coagulated ferrite according to EN 10346. The strength of the matrix is created by high dislocation density by particle refinement and precipitation of fine alloy elements.

Double phase steel is a steel with a ferrite basic microstructure in which the martensite second phase is in the form of an island according to EN 10346 and is sometimes included in the proportion of bainite as the second phase. For high tensile strength, double phase steels exhibit a low yield strength ratio and strong cold solidification.

TRIP steel is a steel with predominantly ferrite basic microstructure containing bainite and residual austenite according to EN 10346, which can be converted to martensite during deformation (TRIP effect). Due to the cold solidification of the steel, the steel achieves high values for uniform elongation and tensile strength. A high part strength can be achieved in combination with the baking hardening effect. These steels are suitable for both stretch drawing and deep drawing. However, a higher seat holder force and press force are required during material deformation. Consider relatively strong resilience.

High strength steels comprising single phase microstructures include, for example, bainite and martensite steels.

Bainite steel is a steel characterized by very high yield strength and tensile strength with a sufficiently high elongation for cold forming processes according to EN 10346. Due to the chemical composition, effective welding capability is provided. The microstructure is usually composed of bainite. The microstructure can include small proportions of other isolated phases such as martensite and ferrite.

Martensitic steel is a steel containing a small proportion of ferrite and/or bainite in the basic microstructure of martensite due to thermomechanical rolling, according to EN 10346. This steel is characterized by a very high yield strength and tensile strength along with a sufficiently high elongation. Within the group of multi-phase steels, martensitic steels have the highest tensile strength values. The suitability for deep drawing is limited. Martensitic steel is mainly suitable for bending deformation methods such as roll forming.

Heat-treated steel is a steel that obtains high tensile and fatigue strength by quenching (= hardening and tempering) according to EN 10083. If cooling during curing in air produces bainite or martensite, this method is called “air curing”. The strength/tensile strength ratio can be influenced in a targeted manner by tempering affected after curing. The steel is used as a flexible cold rolled strip in structural parts, chassis parts and crash-related parts.

This Tailor Rolled Blank Lightweight Structural Technology (TRB®) enables significant weight reduction through selection of sheet thickness loads along the length of the part.

However, due to the currently known alloys and the continuous annealing equipment available for a wide variety of sheet thicknesses, the production of TRB® with multi-phase microstructures is not possible without limitations, for example heat treatment before cold rolling. In the region of different sheet thicknesses, a homogeneous multi-phase microstructure cannot be established in cold rolled and hot rolled steel strips due to temperature gradients that occur in established process windows.

When thin sheets are produced, it is common for cold rolled steel strips to be annealed in a continuous annealing method in a recrystallized manner to produce thin sheets that can be deformed in an effective manner for economic reasons.

Depending on the alloy composition and strip cross-section, process parameters such as treatment rate, annealing temperature and cooling rate are set corresponding to the required mechanical technical properties with the microstructure required for this purpose.

The aforementioned properties are, for example, during steel composition, process parameters during hot rolling, process parameters during acid cleaning (e.g. stretch-bend-straightening) and even during cold rolling before continuous annealing. It is greatly influenced by the process parameters of.

The steel composition is fixed by analytical rules defining the minimum and maximum ranges.

For example, standard slab thickness, slab lying time, slab output temperature, pass plan during preliminary strip rolling, standard preliminary strip thickness, entry temperature into hot strip line, pass plan during hot rolling, final rolling temperature, During hot rolling, such as hot strip cooling pattern, reeling temperature, process parameters are fixed according to the multi-phase steel produced.

During pickling, selective stretch-bending-correction (stretching) affects subsequent process steps.

During cold rolling, the hot strip thickness for producing the cold rolling thickness is already fixed by standard thinning by rolling upon order conversion to technical specifications (process parameters).

The thickness of the pre-strip during the hot rolling process represents the starting thickness before entering the multi-frame hot strip line, where the pre-strip is a reversing manner in multiple passes (runs) from one slab with a defined standard thickness. Was prepared.

Typical slab thickness is from 250 mm to 300 mm (standard 250 mm considered further here), and for multi-phase steel the preliminary strip thickness generally ranges from 40 mm to 60 mm.

In general, the preliminary strip thickness for subsequent hot rolling is relatively constant depending on the material composition, for example 45 mm (herein referred to as standard).

Values lower or higher affect subsequent deformation during cold rolling, such as cold solidification behavior, by creating modified technical hot strip property values such as tensile strength and yield strength.

According to the prior art, in order to achieve the final technical fine sheet property values required by the standard, the material-dependent pre-strip thickness is fixed in the case of a continuous annealing treatment to ensure normal recrystallization. In the case of classical steel, values above or below this value affect the final technical property value to the extent that significant batch variability (scattering range) can occur.

The degree of thinning by cold rolling describes the difference in percentage of the hot strip start thickness to the cold strip final thickness based on the hot strip start thickness.

In general, the degree of thinning by cold rolling is relatively constant, approximately up to 40% for thicker cold strips greater than 2 mm, and up to 60% for cold strips up to 1 mm thick.

According to the prior art, in order to achieve the technical characteristic values required by the standard, in the case of the continuous annealing treatment, on average, about 50% of cold rolling is required to ensure normal recrystallization. In the case of classical steel, the technical characteristic values fluctuate when the value is exceeded or lowered as described in the case of TRB®.

It is known that the minimum degree of cold rolling is set in accordance with the recrystallization temperature in order to set the corresponding dislocation density for recrystallization annealing, in order to achieve the microstructure of fine particles after the continuous annealing procedure.

If the degree of thinning by cold rolling is too low (even in the local region), a threshold for recrystallization cannot be overcome and a fine and relatively uniform microstructure cannot be achieved. After recrystallization, different particle sizes in the cold strip also cause different particle sizes in the final microstructure, resulting in fluctuations in the characteristic values. During cooling at the furnace temperature, different sized particles can be converted to different phase components and ensure additional non-uniformity.

To achieve each required microstructure, the cold strip is heated to a temperature during cooling to achieve the required microstructure formation (eg, double phase or composite phase microstructure) in a continuous annealing furnace.

If the surface of the cold strip needs to be hot-dip galvanized due to high corrosion protection requirements, annealing in a continuous hot-galvanising installation where heat treatment or annealing and downstream zinc plating takes place in a continuous process. Treatment is generally performed.

For example, continuous annealing of hot rolled or cold rolled steel strips using the alloy concept for ultra high strength double phase steel with a minimum tensile strength of approximately 980 MPa, known from documents EP 2 028 282 A1 and EP 2 031 081 A1 In the case, there is a problem that only a small process window is provided for the annealing parameters. Thus, adaptation of process parameters is required to achieve uniform mechanical properties even with minimal cross-sectional changes (thickness, width).

For extended process windows, if the process parameters are the same, the required strip properties are possible even with larger cross-sectional changes of the strip to be annealed.

This mainly concerns flexible rolled strips with different sheet thicknesses over the strip length, as well as strips with different thicknesses and/or different widths that must be continuously annealed.

Uniform temperature distribution can only be achieved when there is particular difficulty in the case of different thicknesses in the transition region from one strip to another. For alloy compositions where the process window is too small, during continuous annealing, this may result in, for example, thinner strips moving too slowly through the furnace, resulting in reduced productivity, or thicker strips moving too quickly through the furnace to achieve the desired microstructure. The required annealing temperature for this may result in not being achieved. As a result, the amount of scrap is increased.

Thus, a critical process parameter for materials with a relatively constant degree of thinning by cold rolling is the setting of the speed during continuous annealing because the phase transformation proceeds with temperature and time. Thus, the more insensitive the steel is with respect to uniformity of mechanical properties during changes in temperature and time profile during continuous annealing, the larger the process window.

The problem with a process window that is too narrow is the load-optimized component consisting of a cold strip and also composed of a cold strip, as well as the degree of thinning of wood low or too high by annealing or cold rolling of cold strips with too small or too large a spare strip thickness. It is especially serious in the annealing treatment of strips having sheet thicknesses varying over the length of the strip for the production of.

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

When the known alloy concept is applied to a group of multi-phase steels, a narrow process window during continuous annealing of strips of different thickness means that it is difficult to achieve uniform mechanical properties over the entire length of the strip. Composite phase steel also has a much narrower process window than double phase steel.

Using a known alloy concept during continuous annealing, it is not possible to substantially achieve the setting of relatively uniform mechanical technical properties of different cold strips with varying degrees of thinning by cold rolling or variable pre-strip thickness. The degree of thinning by cold rolling required for recrystallization annealing places very clear limits on the flexibility of material production within the entire process chain. The final cold strip thickness sets the thickness of the hot strip and thus the hot strip manufacturing parameters.

For flexible rolled cold strips composed of multi-phase steels of known composition, the very small process window leads to areas with smaller sheet thickness due to the conversion process during cooling resulting in very high levels of strength resulting in very large martensite proportions. This means that areas with or greater sheet thickness achieve very low levels of strength and consequently very small martensite ratios. During continuous annealing, it is virtually impossible to achieve uniform mechanical technical properties over the strip length or width.

The known alloy concept for multi-phase steel features a process window that is too narrow, and thus is not particularly suitable for cold strip manufacturing and flexible rolled strips with variable degree of thinning by cold rolling and variable pre-strip thickness. not.

Publication DE 10 2012 002 079 A1 discloses an ultra high strength multi-phase steel with a minimum tensile strength of 950 MPa, which is uniform even for this steel, even with a very wide process window for continuous annealing of hot or cold strips. It has been shown that it is not possible to achieve a variable degree of thinning or variable pre-strip thickness by cold rolling with a single hot strip thickness (master hot strip thickness) while creating material properties.

Publication DE 10 2015 111 177 A1 discloses an ultra-high strength multi-phase steel with a minimum tensile strength of 980 MPa, which already has a very wide process window for continuous annealing of hot or cold strips, and also a single hot strip thickness (master Having a hot strip thickness), a variable degree of thinning is achieved by cold rolling, and a continuous annealed cold strip with different thickness and uniform material properties can be achieved.

Published publication DE 10 2014 017 274 A1 discloses an ultra-high-strength air-curable multi-phase steel with a minimum tensile strength of 950 MPa in a non-air cured state, which already has a very wide process window for continuous annealing of hot or cold strips. , Also has a single hot strip thickness (master hot strip thickness), so that a variable degree of thinning is achieved by cold rolling, cold strips with different thicknesses are continuously annealed, uniform material properties are achieved and subsequent air Suitable for curing process.

The goal of achieving the resulting mechanical technical properties in a narrow area over the length and width of the strip by a controlled setting of the volume fraction of the microstructured parts is only possible with the highest priority and increased process window. Known alloy concepts feature too narrow process windows and are therefore unsuitable to solve this problem, especially for flexible rolled strips. With the known alloy concept, only one strength class steel with currently defined cross-sectional area (strip thickness and strip width) can be produced, so a modified alloy concept for different strength class and/or cross-sectional area is required.

Steel production is witnessing a tendency to reduce carbon equivalents to achieve improved cold working (cold rolling, cold forming) and better use properties.

However, welding suitability, particularly characterized by carbon equivalents, is an important evaluation variable.

For example, at the next carbon equivalent

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 are considered (% by weight).

It is a prior art that an increase in strength is achieved by increasing the amount of carbon and/or silicon and/or manganese, and an increase in strength is achieved by structural setting as well as mixed crystal hardening.

However, by increasing the amount of the elements described above, the degree to which the material processing properties increase, for example during welding, deformation and hot dip plating processing, is weakened.

However, steel production is witnessing a tendency to reduce carbon and/or manganese content to achieve improved cold working and better use properties.

One example is a hole expansion test to explain and quantify edge crack behavior. With the corresponding optimized steel grade adaptation, steel users expect higher values than standard materials. However, welding suitability, characterized by carbon equivalents, is also a focus.

The automotive industry is increasingly demanding grades of steel with requirements for the ratio of yield strength (Re) or elastic limit (Rp 0.2) to tensile strength, which varies considerably depending on the application. This leads to the development of steel with a relatively large yield strength spacing with standard tensile strength spacing.

The low yield strength ratio (Re/Rm) is a common double phase steel and is mainly used for deformability in stretching and deep drawing processes.

The typical high yield strength ratio (Re/Rm) of composite phase steel is also characterized by resistance to edge cracking. This may be due to a smaller difference in the strength of the individual microstructured parts, which has a favorable effect on uniform deformation in the region of the cut edge.

The analytical environment for achieving multi-phase steels with a minimum tensile strength of 980 MPa is highly variable and carbon as well as the addition of microalloys, either individually or in combination, in special properties that exhibit material properties such as hole expansion and reduced carbon equivalents. , Manganese, phosphorus, aluminum and chromium and/or molybdenum.

The dimensional spectrum is broad and ranges from 0.50 to 3.00 mm in thickness, and the range from 0.80 to 2.10 mm is quantitatively related.

Thickness ranges of less than 0.50 mm and greater than 3.00 mm are possible.

Overall, for known steel grades, for the minimum degree of thinning required by rolling during cold rolling for complete recrystallization after continuous annealing, it will be achieved at a given pre-strip strip thickness to produce the master hot strip thickness after hot rolling. The problem is that there is no longer manufacturing flexibility for different cold strip thicknesses (see FIG. 1, process steps 6, 8 and 9 are required). In particular, it is impossible to produce different cold strip thicknesses for constant master hot strip thicknesses with similar material properties on cold strips produced by process windows that are too small. In addition, the specification of a constant pre-strip strip thickness to produce a constant master hot strip thickness limits manufacturing flexibility.

Accordingly, it is an object of the present invention to provide a new alloy concept for a super high-strength multi-phase steel, a method for manufacturing a steel strip from this ultra-high-strength multi-phase steel, and to provide a steel strip produced according to this method, wherein The process window for continuous annealing is capable of producing different cold strip thicknesses from different pre-strip strip thicknesses, specific hot strip thicknesses (master hot strip thicknesses), and producing cold strip thicknesses (master cold strip thicknesses) from different hot strip thicknesses. It can be extended to make it possible. Also, a variable pre-strip thickness can be used before hot rolling instead of a constant pre-strip thickness.

In this case, the most uniform possible cold strip material properties should be achieved independently of the set degree of cold rolling and the set pre-strip thickness.

In addition, the process window for the annealing of cold rolled steel strips to the final thickness, in particular for continuous annealing, is in addition to strips with different cross-sections (cross-section jumps) over the strip length and optionally the strip width to enable the most uniform mechanical technical properties. Steel strips with varying thicknesses (TRB®) should also be expanded to allow production.

According to the teachings of the present invention, this object is achieved by an ultra high strength multi-phase steel with a minimum tensile strength of 980 MPa with the following content in weight percent:

C 0.075 or more and 0.115 or less

Si 0.400 or more and 0.500 or less

Mn 1.900 or more and 2.350 or less

Cr 0.250 or more and 0.400 or less

Al 0.010 or more and 0.060 or less

N 0.0020 or more and 0.0120 or less

P 0.020 or less

S 0.0020 or less

Ti 0.005 or more and 0.060 or less

Nb 0.005 or more and 0.060 or less

V 0.005 or more and 0.020 or less

B 0.0005 or more and 0.0010 or less

Mo 0.200 or more and 0.300 or less

Ca 0.0010 or more 0.0060 or less

Cu 0.050 or less

Ni 0.050 or less

Sn 0.040 or less

H 0.0010 or less

The remainder is iron, which generally contains steel-related smelting-related impurities, and the total content of Mn-Si+Cr is 1.750% to 2.250% by weight in relation to the process window as wide as possible during the annealing of the cold strip of this steel, especially during continuous annealing. %to be.

In mathematical terms, this means that the content indication of Mn and Cr is added, the content indication of Si is subtracted and the total (result) obtained accordingly is 1.750% by weight or more and 2.250% by weight or less. The same applies to the corresponding total content.

By means of the alloy concept according to the invention, mechanical technical properties are reliably achieved in a narrow range for cold strips with variable pre-strip thickness before hot rolling. Due to the variable pre-strip thickness, the cold rolling process can be cold rolled, double cold rolled, and soft annealed of the cold rolled strip before the next cold step, without adversely affecting the production of the master hot strip thickness or master cold strip thickness described above. It can be positively influenced by the fact that the steps are performed.

Important to this is a selected closely maintained alloy composition with a focus on chromium content that is limited and dependent on the cold strip thickness, which achieves uniform material properties with different pre-strip thicknesses as well as different degrees of thinning by cold rolling. It was proved very positive. In addition, the mechanical technical properties that can be produced are achieved in a narrow range over the strip width and strip length by setting the volume ratio on the microstructure in a controlled manner.

Also, the previous production philosophy of determining the hot strip thickness required for the final cold strip thickness (final thickness) and requiring a standard pre-strip thickness is that the selected pre-strip strip thickness and only one selected master hot strip thickness are different for different cold strip thicknesses. It can be neglected to the extent required. However, it is also advantageously possible to create a cold strip thickness to be achieved in a manner similar to the different hot strip thickness. This significantly increases manufacturing flexibility and reduces production costs.

Thus, a pre-strip can be produced from multi-phase steel in a slab state, which is then hot rolled to the hot strip thickness to be achieved.

Proceeding to a hot rolled hot strip having the same thickness as thinning by 72% to 87% rolling, for example to a final thickness achieved from a previously selected pre-strip with a pre-fixed slab thickness of 250 mm and a limited variable thickness. It is also possible to do.

In an advantageous manner, for steel strips of different thicknesses during continuous annealing, similar microstructure conditions and mechanical property values can be established by adjusting the plant treatment rate during heat treatment.

In addition, the steel according to the invention offers the advantage of a significantly increased process window compared to known steel. This increases the level of process reliability in continuous annealing of cold strips with multi-phase microstructures. Thus, for continuous annealed cold strips, uniform mechanical technical properties in the transition region of two strips with different cross-sections and otherwise with the same process parameters, or in strips and in strips with varying degrees of thinning by cold rolling. Can be guaranteed.

According to the present invention, a steel strip can be produced using the multi-phase steel of the present invention, where the multi-phase steel can be used to make a hot strip, and the steel strip from the hot strip is cold rolled to the final thickness to be achieved. And then the steel strip is annealed and in particular continuous annealing.

The properties of the multi-phase steel make it possible to proceed from various pre-strip strip thicknesses to selected master hot strips with specific thicknesses or to selected hot strips with different thicknesses in a wide range from 10% to 70% lamination by cold rolling. , The steel strip is cold rolled to the final thickness to be achieved.

In this case, according to the invention, the chemical composition of the multi-phase steel is selected according to the final thickness of the cold strip to be obtained. Thus, within the selectable thickness class of the cold strip to be obtained, it is possible to produce a corresponding cold strip having one or more final thicknesses from a master hot strip having a thickness, or a master cold strip having a constant thickness from different hot strip thicknesses. It is possible.

In order to achieve uniform mechanical properties, the steel strip is cold rolled to a final thickness of 0.50 to 3.00 mm and is selected as follows according to the final thickness to be achieved even when a pre-strip thickness with variable chemical composition of the multi-phase steel is used. It proved advantageous.

With regard to the possible use of variable pre-strip thickness, it has proven particularly advantageous that the total content of Mn-Si + Cr is selected as follows depending on the final thickness of the cold strip to be achieved:

Final thickness from 0.50 to 1.00 mm:

1.750 wt% ≤ Mn-sum of Si + Cr ≤ 2.030 wt%

Final thickness from 1.00 mm to 2.00 mm:

1.940 wt% ≤ Mn-sum of Si + Cr ≤ 2.110 wt%

Final thickness greater than 2.00 mm and less than 3.00 mm:

2.020 wt% ≤ Mn-sum of Si + Cr ≤ 2.220 wt%

In addition, the total content of Mn-Si + Cr + is advantageously chosen as follows depending on the final thickness of the cold strip to be achieved:

Final thickness from 0.50 to 1.00 mm:

1.950 wt% ≤ Mn-sum of Si + Cr + Mo ≤ 2.280 wt%

Final thickness from 1.00 mm to 2.00 mm:

2.140 wt% ≤ Mn-sum of Si + Cr + Mo ≤ 2.360 wt%

Final thickness greater than 2.00 mm and up to 3.00 mm:

2.220 wt% ≤ Mn-sum of Si + Cr + Mo ≤ 2.470 wt%

Thus, the final thickness of the steel strip to be achieved is related to the alloy composition of the pre-strip or hot strip made from multi-phase steel.

It has also proved advantageous that the carbon equivalent CEV (IIW) is selected as follows depending on the final thickness of the cold strip to be achieved.

Final thickness from 0.50 to 1.00 mm:

C content ≤ 0.100% by weight and carbon equivalent CEV(IIW) ≤ 0.62%

Final thickness from 1.00 mm to 2.00 mm:

C content ≤ 0.105% by weight and carbon equivalent CEV(IIW) ≤ 0.64%

Final thickness greater than 2.00 mm and up to 3.00 mm:

C content ≤ 0.115% by weight and carbon equivalent CEV(IIW) ≤ 0.66%

It has also proved advantageous that the Mn content is selected as follows depending on the final thickness of the cold strip to be achieved:

Final thickness from 0.50 to 1.00 mm:

1.900 wt% ≤ Mn content ≤ 2.200 wt%

Final thickness from 1.00 mm to 2.00 mm:

2.050% by weight ≤ Mn content ≤ 2.250% by weight

Final thickness greater than 2.00 mm and less than 3.00 mm:

2.100% by weight ≤ Mn content ≤ 2.350% by weight

With regard to the use of variable pre-strip thickness, it has proven particularly advantageous that the Cr content and the carbon equivalent CEV (IIW) are selected as follows depending on the final thickness of the cold strip to be achieved:

Final thickness from 0.50 to 1.00 mm:

0.260 wt% ≤ Cr content ≤ 0.330 wt and

Carbon equivalent CEV(IIW) ≤ 0.62%

Final thickness from 1.00 mm to 2.00 mm:

0.290 wt% ≤ Mn content ≤ 0.360 wt% and

Carbon equivalent CEV(IIW) ≤ 0.64%

Final thickness from 2.00 to 3.00 mm:

0.320 wt% ≤ Mn content ≤ 0.370 wt% and

Carbon equivalent CEV(IIW) ≤ 0.66%

This also applies to continuous annealing of continuous strips with different strip cross-sections and strips of varying sheet thickness over strip length or strip width. For example, it is possible to process cold strips with varying degrees of thinning by cold rolling.

According to the present invention, when the ultra-high strength cold strip produced by the continuous annealing method is made from multi-phase steel having various sheet thicknesses, it may be advantageous to manufacture a load-optimized part from this material by a deformation technique.

The material produced is a pure continuous annealing plant or hot dip galvanizing line in a temper-rolled and tambour-unrolled state and even in a heat-treated state (aging) and in a stretch and unstretched state (stretch-bending-correct). It can be prepared as a cold strip through.

At the same time, it is possible to set the microstructure fraction to produce steels of different strength grades, for example by yielding between 550 MPa and 950 MPa and tensile strength between 980 MPa and 1140 MPa by specifically changing the process parameters. .

The alloy composition according to the invention can be used to produce steel strips by inter-critical annealing between Ac1 and Ac3 or by austenitizing annealing over Ac3 by controlling cooling, which can be multi-phase or multi- This results in phase microstructure.

Annealing temperatures of about 700 to 950° C. have proven advantageous. According to the present invention, there are different approaches to heat treatment depending on the overall process (only continuous annealing or additional hot-dip finishing).

For a continuous annealing plant without subsequent hot dip treatment, the steel strip cold rolled to the final thickness is cooled from annealing temperature to an intermediate temperature of approximately 160 to 250°C at a cooling rate of approximately 15 to 100°C/s. Optionally, the cooling can be preformed to a previous intermediate temperature of 300 to 500°C at a cooling rate of approximately 15 to 100°C/s. Cooling to room temperature is finally performed at a cooling rate of approximately 2 to 30°C/s (method 1, see FIG. 8A). Alternatively, cooling may be performed at a cooling rate of approximately 15 to 100°C/s from a medium temperature of 300 to 500°C to room temperature.

For heat treatment as part of the hot dip treatment process, there are two temperature control options. Cooling is stopped as described above before entering the melting bath and continues only after exiting the bath until an intermediate temperature of approximately 200-250° C. is achieved. Depending on the melting bath temperature, a holding temperature of approximately 400 to 470°C is provided to the melting bath. Cooling to room temperature is performed at a cooling rate of approximately 2 to 30°C/s (method 2, see FIG. 8B).

The second variant for temperature control in the hot dip treatment procedure is to maintain the temperature for about 1 to 20 seconds at an intermediate temperature of about 200 to 350 °C and then reheat to a temperature of about 400 to 470 °C required for the hot dip treatment procedure. Includes. After the finishing procedure, the strip is cooled to approximately 200-250°C. Then, it is cooled to room temperature at a cooling rate of about 2 to 30°C/s (method 3, see FIG. 8C).

In the case of known double phase steels, what is required to convert austenite to martensite is manganese, chromium and silicon as well as carbon. Only a combination of the present invention of alloyed elements at the indicated limits of carbon, silicon, manganese, nitrogen, molybdenum and chromium as well as a narrow area of niobium, titanium and boron, a minimum tensile of 980 MPa with a significantly widened process window during the continuous annealing procedure It ensures the necessary mechanical properties such as strength.

It is also a feature of the material that the ferrite region is displaced at a lower temperature for a longer period during cooling by adding manganese as the weight percentage increases, and carbon, chromium, molybdenum and boron also behave in a similar manner. The proportion of ferrite decreases larger or smaller as the proportion of bainite increases depending on the process parameters.

By setting a low carbon content of 0.115% by weight or less, the carbon equivalent can be reduced, thereby improving welding suitability and preventing excessive hard spots during welding. In addition, in the case of resistance spot welding, the electrode life can be significantly increased.

The effect of the elements in the alloy according to the invention will be explained in more detail later. Related elements are inevitable and, if necessary, are considered in analytical concepts in terms of their effectiveness.

Related elements are those already present in iron ore or entering the river as a result of the production process. These are generally undesirable due to their negative effects. It is attempted to remove it to an acceptable content level or convert it to a less damaging form.

Hydrogen (H) can diffuse as a single element through the iron lattice without creating lattice tension. As a result, hydrogen in the iron lattice is relatively mobile and can be absorbed relatively easily during steel processing. Hydrogen can only be absorbed into the iron lattice in atomic (ionic) form.

Hydrogen exerts a significant embrittlement effect and preferably diffuses to an advantageous position in terms of energy (flaw, grain boundaries, etc.). Thus, the defect functions as a hydrogen trap and can significantly increase the dwell time of hydrogen in the material.

Cold cracks can be created by recombination with molecular hydrogen. This behavior occurs during hydrogen embrittlement or hydrogen-induced tension crack corrosion. It is often mentioned that hydrogen is also the cause of delayed cracks, so-called delayed fractures, which occur without external tension. Therefore, the hydrogen content of the steel should be kept as small as possible.

For the reasons mentioned above, the hydrogen content of the steel according to the invention is limited to 0.0010% by weight (10 ppm) or less or advantageously 0.0008% by weight or less, optimally 0.0005% by weight or less.

In the case of the steel according to the invention the more uniform microstructure achieved by a particularly widened process window also reduces the susceptibility to hydrogen embrittlement.

Oxygen (O): In the molten state, steel has a relatively large absorbency with respect to gas. However, only a very small amount of oxygen is dissolved at room temperature. In a similar way to hydrogen, oxygen can only diffuse into the material in atomic form. Due to the high embrittlement and negative effects, every attempt is made during production to reduce the oxygen content.

On the one hand, procedural approaches, such as vacuum processing, to reduce oxygen and on the other hand, analytical approaches are provided. Oxygen can be converted to a less dangerous state by adding certain alloying elements. It is generally customary to remove oxygen in the course of deoxidation of steel with manganese, silicon and/or aluminum, for example. However, the resulting oxide can produce negative properties as a defect in the material.

Therefore, for the reasons mentioned above, the oxygen content of the steel should be kept as small as possible.

Phosphorus (P) is a trace element, mainly derived from iron ore and dissolved in the iron lattice as a substitution atom. Phosphorus increases the hardness and improves the curability by curing the mixed crystals. However, attempts to lower the phosphorus content as much as possible are generally made because the low solubility in the solidification medium indicates a strong tendency to segregation and greatly reduces the level of toughness. The adhesion of phosphorus to the grain boundaries causes grain boundary destruction. In addition, phosphorus increases the transition temperature from toughness to brittle behavior to 300°C. During hot rolling, phosphorus oxide near the surface at the grain boundaries results in the formation of fracture.

However, in some steels, the cost is low and the strength is greatly increased, so phosphorus is a small amount (0.1) as a microalloy element in some alloy concepts for high strength IF steel (with no intrusive particles), sinter hardened steel or double phase steel. Less than weight percent) is used. The steel according to the invention differs from the known analytical concept of using phosphorus as a mixed crystal forming agent, in particular that phosphorus is not alloyed and instead is set as low as possible.

For the reasons mentioned above, the phosphorus content of the steel according to the invention is limited to an inevitable amount for the production of steel. Preferably, P should be 0.020% by weight or less.

Like phosphorus, sulfur (S) is bound by trace elements of iron ore. Sulfur is not generally required in steel (except for machining steel), since it exhibits a strong tendency to segregation and has a large embrittlement effect. Thus, an attempt is made to achieve the amount of sulfur in the melt as low as possible, for example by vacuum treatment. In addition, the sulfur present is converted by adding manganese to the relatively harmless compound manganese sulfide (MnS). Manganese sulfide is often rolled out to the line during the rolling process and serves as a nucleation site for conversion. Mainly in the case of diffusion control transformation, this creates a distinct line microstructure, and in the case of marked line formation, mechanical properties such as, for example, distinct martensite lines, anisotropic material behavior, reduced elongation at break instead of dispersed martensite islands. This can be damaged.

For the reasons mentioned above, the sulfur content in the steel of the present invention is limited to 0.0020% by weight or less or advantageously 0.0015% by weight or less, optimally 0.0010% by weight or less.

Alloying elements are usually added to the steel to influence certain properties in a targeted manner. Accordingly, the alloying elements can affect different properties of different steels. The effect generally depends largely on the amount of material and the state of the solution. Thus, relationships can be very diverse and complex.

The effect of the alloying elements will be discussed in more detail later.

Carbon (C) is considered the most important alloying element in steel. Its targeted introduction, up to 2.06% by weight, converts iron to steel first. The carbon fraction often decreases rapidly during steel manufacture. For double phase steels for continuous hot dip galvanizing, its fraction according to EN 10346 or VDA 239-100 is up to 0.230% by weight and the minimum value is not specified.

Carbon is interstitially dissolved due to its relatively small atomic radius. Solubility is up to 0.02% in α-iron and up to 2.06% in α-iron. The dissolved form of carbon significantly increases the hardenability of the steel and is therefore necessary to form a sufficient amount of martensite. However, an excessively high carbon content increases the difference in hardness between ferrite and martensite and limits weldability.

For example, in order to meet the requirements of high hole expansion and bending angles as well as improved weldability, the steel according to the invention comprises a carbon content of 0.115% by weight or less.

Different solubility of the carbon in the phase requires a distinct diffusion procedure during phase conversion, which can lead to very different kinetic conditions. In addition, carbon increases the thermodynamic stability of austenite, which is evidenced by a phase diagram of the expansion of the austenite region at low temperatures. As the carbon content that is forcibly dissolved in martensite increases, the lattice distortion and in this context the strength of the resulting phase without diffusion increases.

Carbon also forms carbides. The microstructural phase that occurs in almost all steels is cementite (Fe ₃ C). However, substantially harder special carbides can be formed from other metals, such as chromium, titanium, niobium, as well as vanadium, for example. Thus, it is not only the type of steel, but also the distribution and extent of precipitation, which is critical for the resulting increase in strength. Thus, on the one hand, to ensure sufficient strength and, on the other hand, effective weldability, improved hole glide, improved bending angle and sufficient resistance to hydrogen induced crack formation (without delayed fracture), the C content is 0.075 weight. Fixed at% and the maximum C content is fixed at 0.015% by weight, the content with cross-section-dependent differentiation is advantageous as follows:

Final thickness 0.50 mm to 1.00 mm (C ≤ 0.100 wt%)

Final thickness over 1.00 mm and under 2.00 mm (C ≤ 0.105 wt%)

Final thickness greater than 2.00 mm and less than 3.00 mm (C ≤ 0.115 wt%)

It is also advantageous to bond to the band thickness dependent differentiation of carbon content in combination with carbon equivalent CEV (IIW).

Final thickness 0.50 mm to 1.00 mm (C ≤ 0.100 wt%) with carbon equivalent CEV(IIW) of 0.62% or less

Final thickness greater than 1.00 mm and less than 2.00 mm (C ≤ 0.105 wt%) with carbon equivalent CEV(IIW) of 0.64% or less

Final thickness greater than 2.00 mm and less than 3.00 mm (C ≤ 0.115 wt%) with a carbon equivalent CEV (IIW) of 0.66% or less

During casting, silicon (Si) is used to kill oxygen during the deoxidation process of the steel in combination with oxygen. For subsequent steel properties, the segregation coefficient is considerably smaller, for example, than that of manganese (0.16 compared to 0.87). Segregation generally leads to an ordered arrangement of microstructured components that impair deformation properties such as, for example, hole expansion and bending capability.

In the anticorrosive properties of the material, the addition of silicon produces a strong mixed crystal hardening. Approximately, the addition of 0.1% silicon produces an increase in tensile strength by approximately 10 MPa, and the elongation is only slightly lowered for the addition of silicon below 2.2%. This was investigated for different sheet thicknesses and annealing temperatures. When the silicon increased from 0.2% to 0.5%, the yield strength increased by approximately 20 MPa and the tensile strength increased by approximately 70 MPa. The elongation at break is reduced by about 2%. The latter situation is due, in particular, to the fact that silicon reduces the solubility of carbon in ferrite and increases the activity of carbon in ferrite, thereby reducing the ductility as a brittle phase and consequently preventing the formation of carbides which improve the deformability. The effect of increasing the low strength of silicon within the span of the steel according to the invention provides the basis for a wide process window.

Another important effect is that silicon can transfer the formation of ferrite to a shorter time and temperature to produce enough ferrite before quenching hardening. During hot rolling, a basis for improving cold rolling properties is provided. During the hot dip treatment, the accelerated formation of ferrite is stabilized by enriching austenite with carbon. Since silicon interferes with the formation of carbide, austenite is further stabilized. Thus, the formation of bainite can be suppressed advantageously to martensite while cooling is accelerated.

The addition of silicon to the range according to the invention resulted in a further surprising effect which will be explained later. The delay described above in the formation of carbides can also be caused, for example, by aluminum. However, since aluminum forms a stable nitride, sufficient nitrogen cannot be used to form a carbonitride having a microalloy element. Alloying with silicon eliminates this problem because silicon does not form carbides or nitrides. Therefore, silicon has an indirectly positive effect on the formation of precipitates by the microalloy, which in turn has a positive effect on the strength of the material. Microalloys with niobium, titanium and boron are particularly convenient, such as targeting the nitrogen content in steels according to the invention, since the increase in conversion temperature caused by silicon tends to promote particle coarsening.

It is known that during hot rolling, higher silicon alloy steels form a red scale that adheres strongly and the risk of the rolling scale increases, which can affect subsequent acid cleaning results and acid-cleaning productivity. When acid washing is advantageously performed with hydrochloric acid instead of sulfuric acid, this effect cannot be established with 0.400 to 0.500% silicon in the steel according to the invention.

With regard to the zinc-plating ability of silicon-containing steels, DE 196 10 675 C1, in particular, steels containing up to 0.800% by weight of silicon or less than 2.000% by weight of silicon are hot dip galvanized for reasons of very poor wettability of the steel surface with liquid zinc. Note that it cannot be (hot-galvanised).

In addition to recrystallization of roll-hard cold strips, atmospheric conditions during annealing treatment in a continuous hot dip coating facility can be formed on the surface, for example during cold rolling or as a result of storage at room temperature. That produces a reduction in iron oxide. However, in the case of oxygen-affine alloy components such as silicon, manganese, chromium, and boron, for example, the gas atmosphere is oxidized, and as a result, selective oxidation and segregation of these elements may occur. Selective oxidation can occur both externally, ie on the substrate surface and internally within the metal matrix.

Silicon is known to diffuse to the surface, particularly during annealing, to form oxides on the steel surface by itself or with manganese. These oxides can inhibit the contact between the substrate and the melt and prevent or significantly reduce the wetting reaction. As a result, ungalvanized locations, so-called “bare spots” or large surface areas without coating may occur. In addition, the adhesion of the zinc or zinc alloy layer on the steel substrate may be reduced due to the reduced wetting reaction, resulting in insufficient formation of the suppression layer.

Contrary to this general knowledge in the art, surprisingly, effective hot dip treatment of the steel strip and effective adhesion of the coating can only be achieved by suitable furnace operations during recrystallization annealing and during passage through a hot-dip bath. It has been found in tests that can.

To this end, it is necessary to initially perform a chemical-mechanical or thermal-hydromechanical pre-cleaning procedure to ensure that there are no residues, acid cleaning oil or rolling oil or other contaminant particles on the strip surface. In addition, it is necessary to resort to a method that promotes the internal oxidation of the alloying element below the material surface in order to prevent the silicon oxide from reaching the strip surface. In this case, different measures are applied depending on the configuration of the equipment.

In the case of one configuration of an installation where the annealing process step is performed exclusively in a radiant tube furnace (RTF) (see method 3 in Figure 8c), the internal oxidation of the alloying elements is the oxygen partial pressure in the furnace atmosphere (N2-H2-protective gas atmosphere). By setting, you can be influenced in a targeted way. The set oxygen partial pressure must satisfy the following equation, and the furnace temperature is 700 to 950°C.

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

In this case, Si, Mn, Cr, and B define the corresponding alloy ratio of the steel by weight%, and pO2 defines the oxygen partial pressure in mbar.

In the case of a construction where the furnace zone consists of a direct combustion furnace (DFF) or a combination of a non-oxidizing furnace (NOF) and a downstream radiant tube furnace (see method 2 in Figure 8b), the selective oxidation of the alloying elements is The gas atmosphere in the zone can likewise be affected.

The oxygen partial pressure, and thus the oxidation potential for iron and alloying elements, can be set by the combustion reaction in NOF. This should be set in such a way that the oxidation of the alloying elements takes place internally below the steel surface, and possibly a thin iron oxide layer is formed on the steel surface after passing through the NOF region. This is achieved, for example, by reducing the CO value to less than 4% by volume.

Downstream copy N ₂ -H ₂ of the furnace tube - under protective gas atmosphere, possibly iron oxide layer is reduced and the similar alloy elements formed are further oxidized internally. The set partial pressure of oxygen in this furnace region must satisfy the following equation, and the furnace temperature 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 * (- ln B) 0.25

In this case, Si, Mn, Cr, and B define the corresponding alloy ratio of the steel by weight%, and pO ₂ defines the oxygen partial pressure in mbar.

The dew point of the gas atmosphere (N2-H2-protective gas atmosphere) in the transition region between the furnace and the zinc pot (muzzle) and the resulting oxygen partial pressure should be set in such a way as to avoid oxidation of the strip before plating in the melting bath. Dew points ranging from -30 to -40°C have proven advantageous.

The aforementioned measurements in the furnace area of a continuous hot dip coating facility prevent the formation of oxides on the surface and provide uniform and effective wetting of the strip surface with the liquid melt.

Instead of a hot dip treatment procedure (for example hot dip galvanizing in this case), a method comprising continuous annealing with a subsequent electrolytic galvanizing (see method 1 in Fig. 8a) is selected, to ensure a special galvanizing ability. It is not necessary to take action. It is known that galvanizing treatment of more highly alloyed steel can be achieved substantially more easily by electro-deposition than by a continuous hot dip plating method. In the case of electrolytic zinc plating, pure zinc is deposited directly on the strip surface. In order not to interfere with the electron flow between the steel strip and the zinc ions and thus not to the galvanization, it is necessary to ensure that no surface-coated oxide layer is present on the strip surface. This condition is usually ensured by a standard reducing atmosphere during annealing and pre-cleaning prior to electrolysis.

To ensure the widest possible process window and sufficient galvanizing capability during annealing, the minimum silicon content is fixed at 0.400% by weight and the maximum silicon content is fixed at 0.500% by weight.

Manganese (Mn) is added to almost all steels for the purpose of desulfurization in order to convert harmful sulfur into manganese sulfide. In addition, by means of mixed crystal hardening, manganese increases the strength of ferrite and shifts the α-/γ-conversion towards lower temperatures.

The main reason for adding manganese by alloying in multi-phase steels, for example in the case of double-phase steels, is a significant improvement in potential hardness increases. For reasons of suppression of diffusion, the pearlite and bainite conversion is shifted to a longer time and the martensite initiation temperature is reduced.

However, at the same time, the addition of manganese serves to increase the hardness ratio between martensite and ferrite. In addition, line formation of the microstructure is improved. The high hardness difference between the formation of the phase and the martensite line results in a low hole expansion capacity, which is equivalent to increased edge crack sensitivity.

Manganese, such as silicon, tends to form oxides on the steel surface during the annealing process. Manganese oxide (eg MnO) and/or Mn mixed oxide (eg Mn ₂ SiO ₄ ) may occur depending on the annealing parameters and the content of other alloying elements (particularly silicon and aluminum). However, manganese should be considered less important at small Si/Mn or Al/Mn ratios because spherical oxides are likely to form instead of oxide films. Nevertheless, high manganese content can negatively affect the appearance and zinc adhesion of the zinc layer. The above-mentioned measures for setting the furnace region during continuous hot dip coating serve to reduce the formation of Mn oxide or Mn mixed oxide on the steel surface after annealing.

For the reasons mentioned, the manganese content is fixed between 1.900% and 2.350% by weight.

In order to achieve the required minimum strength, it is advantageous to maintain strip thickness dependent differentiation of the manganese content.

When the final thickness is 0.50 mm to 1.00 mm, the manganese content is preferably in the range of 1.900 wt% to 2.200 wt%, and when the final thickness is 1.00 to 2.00 mm, the manganese content is 2.050 wt% to 2.250 wt% And in the case of a final thickness of 2.00 mm to 3.00 mm, in the range of 2.100% to 2.350% by weight.

Another feature of the invention is that the change in the manganese content can be compensated for by simultaneously changing the silicon content. The increase in strength due to manganese and silicon (in this case, yield strength, YS) is generally described in an effective manner by the Picking equation.

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

However, this is mainly based on the effect of mixed crystal hardening according to this equation, which is weaker for manganese than silicon. However, as mentioned above, manganese significantly increases the hardenability at the same time, and as a result, the proportion of the second phase that increases in strength increases significantly in the case of multi-phase steel. Therefore, the addition of 0.1% silicon should be equated to the first approximation for the addition of 0.1% manganese in relation to the increase in strength. For the annealing procedure comprising the steel of the composition according to the invention and the time-temperature parameter according to the invention, the following relationship was created on an empirical basis for yield strength and tensile strength (TS):

YS (MPa) = 185.7 + 147.9 [wt% Si] + 161.1 [wt% Mn]

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

Compared to the Pickering equation, the coefficients of manganese and silicon are almost the same for yield strength and tensile strength, providing a choice to replace manganese with silicon.

On the other hand, chromium (Cr), in dissolved form, can significantly increase the hardenability of the steel, even in small quantities. On the other hand, chromium in the form of chromium carbides, with corresponding temperature control, results in particle solidification. The increase associated with the number of nucleation sites with a simultaneously reduced carbon content leads to a decrease in sclerosis.

In double phase steel, the addition of chromium improves the potential increase in hardness. The dissolved chromium shifts the pearlite and bainite conversions to a longer time and at the same time lowers the martensite starting temperature.

Another important effect is that chromium significantly increases the tempering resistance, so there is little loss of strength in the hot dip bath.

In addition, chromium is a carbide former. If chromium-iron-mixed carbide is present, the austenitizing temperature must be selected to be high enough to dissolve the chromium carbide before curing. Otherwise, the increased number of nuclei can exacerbate the potential increase in hardness.

Chromium tends to form oxides on the steel surface during the annealing treatment, and as a result the hot dip quality may be impaired. The above-mentioned measures for setting the furnace region during the continuous hot dip coating serve to reduce the formation of Cr oxide or Cr mixed oxide on the steel surface after annealing.

Therefore, the chromium content is fixed at a content of 0.250% by weight to 0.400% by weight.

When the final thickness is 0.50 mm to 1.00 mm, the chromium content is preferably in the range of 0.260 wt% to 0.330 wt%, and when the final thickness is 1.00 to 2.00 mm, the chromium content is 0.290 wt% to 0.360 wt% And in the case of a final thickness of 2.00 mm to 3.00 mm, in the range of 0.320% to 0.370% by weight.

In order to achieve the required minimum strength, it is advantageous to comply with the band thickness dependent differentiation of the chromium content in combination with carbon equivalent CEV (IIW), in which case it is also particularly advantageous for processing with variable pre-strip thickness.

When the final thickness is 0.50 mm to 1.00 mm, the chromium content is preferably in the range of 0.260 wt% to 0.330 wt%, in which case the carbon equivalent is CEV(IIW) ≤ 0.62%, and the final thickness is 1.00 to 2.00 mm In the case, the chromium content is in the range of 0.290% by weight to 0.360% by weight, in this case the carbon equivalent is CEV(IIW) ≤ 0.66%, and in the case of a final thickness of 2.00 mm to 3.00 mm, 0.320% by weight to 0.370% by weight Range, in which case the carbon equivalent is CEV(IIW) ≤ 0.66%.

A chromium content of 0.250% by weight or more and less than 0.370% by weight can be used for a final thickness of less than 0.50 mm, and a chromium content of greater than 0.370% by weight and less than 0.400% by weight can be used for a final thickness greater than 3.00 mm.

Molybdenum (Mo): The addition of molybdenum improves the curability in a manner similar to the addition of chromium and manganese. The pearlite and bainite conversion is shifted to a longer time and the martensite initiation temperature is reduced. At the same time, molybdenum is a powerful carbide former that enables the production of finely distributed mixed carbides, especially with titanium. In addition, molybdenum significantly increases the tempering resistance, so no loss of strength is expected in the hot dip bath. Molybdenum also works by mixed crystal hardening, but is less effective than manganese and silicon.

Therefore, the content of molybdenum is set to more than 0.200% by weight and 0.300% by weight or less. For cost-related reasons, the Mo content is advantageously set in a range of greater than 0.200% by weight and up to 0.250% by weight.

As a compromise between the required mechanical properties and hot dip performance, the alloy concept according to the present invention has proven advantageous to have a total content of Mo + Cr of 0.650% by weight or less.

In order to achieve the required mechanical property values, mainly the minimum tensile strength, it is advantageous to comply with the total content of manganese, silicon and chromium via the total formula Mn-Si + Cr, where it is particularly treated with a variable pre-strip thickness. It should be limited to 1.750% by weight to 2.250% by weight.

In order to achieve the required mechanical property values, mainly the minimum tensile strength, it has proven advantageous to fix the total content of manganese, silicon, chromium and molybdenum through the total formula Mn-Si + Cr + Mo, where this is particularly variable It should be limited to 1.950% to 2.500% by weight to treat strips with pre-strip thickness.

Copper (Cu): Adding copper can increase tensile strength and increase potential hardness. Together with nickel, chromium and phosphorus, copper can form a protective oxide layer on the surface that can significantly reduce the rate of corrosion.

Together with oxygen, copper can form harmful oxides at the grain boundaries, which can negatively affect the thermal deformation process. Therefore, the content of copper is fixed at 0.050% by weight or less, which is limited to an inevitable amount in steel production.

Nickel (Ni): Together with oxygen, nickel can form hazardous oxides at the grain boundaries, which can negatively affect the thermal deformation process. Therefore, the nickel content is fixed at 0.050% by weight or less, which is limited to an inevitable amount in steel production.

Vanadium (V): In the case of the present alloy concept, the content of vanadium is fixed at 0.005% to 0.020% by weight, and optimally limited to 0.005% to 0.015% by weight.

Tin (Sn): In the case of this alloy concept, since the addition of tin is not required, the content of tin is fixed at 0.040% by weight or less, and thus is limited to the inevitable amount associated with steel.

Aluminum (Al) is usually added to the steel by alloying to combine oxygen and nitrogen dissolved in iron. Oxygen and nitrogen are thus converted to aluminum oxide and aluminum nitride. Such precipitates can perform particle refinement by increasing the nucleation sites, thus increasing toughness properties and strength values.

If titanium is present in sufficient quantity, aluminum nitride will not precipitate. Titanium nitride has a low enthalpy of formation and is formed at high temperatures.

In the dissolved state, aluminum, like silicon, shifts the formation of ferrite in a shorter time, allowing sufficient ferrite to be formed in the double phase steel. In addition, the formation of carbides is suppressed, and austenite conversion is delayed. For this reason, aluminum is used as an alloying element of residual austenitic steel (TRIP steel) to replace some of the silicon. The reason for this approach lies in aluminum, which is slightly less important for zinc plating reactions than silicon.

Therefore, the aluminum content is limited to 0.010% by weight to 0.060% by weight, and is added to calm the steel.

Niobium (Nb): Niobium acts differently in rivers. During hot rolling in the production line, this delays recrystallization by the formation of very finely distributed sediments, increasing nucleation site density and producing finer particles after conversion. The proportion of dissolved niobium also serves to inhibit recrystallization. The precipitate acts to increase the strength of the final product. These can be carbides or carbonitrides. These are often mixed carbides incorporating titanium. This effect starts at 0.005% by weight and is most pronounced at 0.010% by weight niobium. In addition, the precipitate prevents particle growth during (partial) austenitization in hot dip galvanizing. In the case of niobium of more than 0.060% by weight, additional effects cannot be expected. Contents of 0.025 to 0.045% by weight have proven advantageous.

Titanium (Ti): Due to its high affinity for nitrogen, titanium is mainly precipitated as TiN during solidification. Also, it appears together with niobium as a mixed carbide. TiN is very important for particle size stability in a pusher-type furnace. Since the precipitate has a high level of temperature stability, unlike mixed carbide, it exists mainly as a particle to inhibit particle growth at 1200°C. In addition, titanium has a retarding effect on recrystallization during hot rolling, but is less effective than niobium. Titanium works by precipitation hardening. Larger TiN particles are less effective than finely distributed mixed carbides. The best efficacy is achieved in the range of 0.005 to 0.060% by weight titanium, so this represents the alloy range according to the invention. For this purpose, a content of 0.025 to 0.045% by weight has proven advantageous.

Boron (B): Boron is a very effective alloying agent that achieves a variable degree of thinning by rolling. The tests surprisingly showed that the very narrow range of addition of boron according to the invention has a significant effect on the uniformity of the mechanical properties of the cold strip produced by the degree of variable lamination by cold rolling in subsequent treatments. This pronounced effect is also a process step (FIGS. 8A, 8B) with a material having a variable degree of lamination by cold rolling on a master hot strip thickness basis or a master cold strip thickness basis, instead of a relatively constant degree of thinning by cold rolling initially. Or 8c), resulting in the possibility of setting a limited range of characteristic values.

In addition, boron is an effective element for increasing effective hardenability even in a very small amount. The martensite starting temperature is thereby maintained unaffected. To be effective, boron must be present in the solid solution. Since it has a high affinity for nitrogen, nitrogen must first be removed by the stoichiometrically required amount of titanium. Since the solubility in iron is low, the dissolved boron is preferably attached to the austenite grain boundaries. In this position, it partially forms Fe-B carbides and aggregates to reduce grain boundary energy. Both effects work in a way that increases the hardenability of the steel by delaying the formation of ferrite and pearlite. However, an excessively high content of boron is dangerous because iron boron is formed and negatively affects the hardenability, deformability and toughness of the material. Boron also tends to form oxides or mixed oxides when annealing is performed during a continuous hot dip coating process, thereby inhibiting zinc plating quality. The above-described measures for setting the furnace region during the continuous hot dip coating serve to reduce the formation of oxides on the steel surface.

For the above reasons, the boron content for the alloy concept according to the invention is fixed to a value of more than 0.0005% by weight and less than 0.0010% by weight, advantageously less than 0.0009% by weight or optimally greater than 0.0006% by weight and less than 0.0009% by weight.

Nitrogen (N) can be an alloying element and related element in steel production. The excessively high nitrogen content produces an increase in strength with respect to the aging effect as well as a sharp loss of toughness. On the other hand, by targeted addition by alloying of nitrogen with the fine alloying elements titanium and niobium, fine particle curing can be achieved through titanium nitride and niobium (carbon) nitride. In addition, coarse particles are suppressed during reheating prior to hot rolling.

Therefore, according to the present invention, the N content is fixed at a value of 0.0020 to 0.0120% by weight.

It has proved advantageous for the content of hydrogen and nitrogen to have an optimal total sum of H+N of 0.0025% to 0.0130% by weight.

It has proven advantageous to maintain the required properties of the steel when nitrogen is added according to the total sum of Ti+Nb+B.

When the total content of Ti+Nb+B is 0.010 to 0.080% by weight, the nitrogen content should be maintained at a value of 0.0020 to 0.0090% by weight. When the total content of Ti+Nb+B is 0.070% by weight or more, a nitrogen content of 0.0040 to 0.0120% by weight has proved to be advantageous.

Contents of less than 0.100% by weight with respect to the total content of niobium and titanium have proven to be advantageous and due to the basic interchangeability of niobium and titanium up to a minimum niobium content of 0.0005% by weight and a content of less than 0.090% by weight in terms of cost It has proven particularly advantageous.

During the interaction of the microalloy element niobium with titanium and boron, a total content of 0.102% by weight or less has proven to be advantageous and a total content of 0.092% by weight or less has proved to be particularly advantageous. The high content no longer has the effect of providing an improvement in terms of the present invention.

In addition, for the above reasons, the maximum content of Ti+Nb+V+Mo+B of 0.365% by weight or less has been successfully demonstrated.

Calcium (Ca): The addition of calcium in the form of a calcium-silicon mixed compound causes deoxidation and desulfurization of the molten phase in steel production. For example, the reaction product is converted to slag and the steel is cleaned. According to the present invention, increased purity results in improved properties in the final product.

For the reasons mentioned, a Ca content of 0.0010 to 0.0060% by weight is set. Contents below 0.0030% by weight have proven advantageous.

According to the test for variable pre-strip thickness performed using the steel according to the invention, the master hot of 2.30 mm in a critical annealing between Ac1 and Ac3 or austenitizing annealing greater than Ac3 with controlled cooling. It has been found that starting from the strip, a biphasic steel with a minimum tensile strength of 980 MPa at a thickness of 1.50 mm was produced but could also be produced in a thickness range of 0.50 to 3.00 mm characterized by adequate resistance to process variations.

Thus, a clearly broadened process window is provided for the alloy composition according to the invention compared to known alloy concepts.

The annealing temperature of the dual phase microstructure to be obtained is about 700 to 950° C. in the steel according to the invention, and thus depending on the temperature range partially austenitic (double-phase region) or fully austenitic microstructure (austenite region) ) Is achieved.

The test can also be followed by a post-critical annealing between Ac1 and Ac3 or an austenitizing annealing above Ac3, even after further processing steps of hot dip galvanizing at a temperature between 400 and 470°C, for example with zinc or zinc-magnesium. It has been found that the set microstructure ratio is maintained with the cooled cooling.

The continuously annealed and sometimes melt-plated-treated material is in a rough-rolled (cold-post-rolled) or non-rough-rolled state and/or stretch-bending-corrected or non-stretch-bending-corrected state And also in a heat-treated state (aging).

In addition, steel strips composed of the alloy composition according to the invention are characterized during further processing by high edge crack sensitivity and high bending angle.

In an advantageous way, it is possible to produce steel strips having a minimum product value Rm×α (tensile strength×[bending angle according to VDA 238-100]) of 100000 MPa×°, especially 120000 MPa×°.

In addition, the steel strip according to the present invention has a delayed free state of fracture for at least 6 months and meets the requirements according to SEP 1970 for hole pull and hoop test pieces provided by steel producers.

The very small property value difference of the steel strip in the longitudinal and transverse directions with respect to the rolling direction of the steel strip is advantageous for the use of subsequent materials. For example, cutting a blank from a strip can be performed independently of the rolling direction (eg, horizontally, vertically and diagonally or at an angle to the rolling direction), thus minimizing waste.

To ensure the cold rolling properties of the hot strips made from the steel according to the invention, the hot strips are produced according to the invention at a final rolling temperature in the austenitic region above Ar3 and a reeling temperature higher than the bainite initiation temperature.

As part of the further processing of the steel strip according to the invention, it is thus possible, for example, to produce hardened parts for the automotive industry.

In this case, the blank is cut from the steel strip according to the invention, the blank is then heated to a temperature above Ac3, and the heated blank is formed of components and cured in a strain tool or air with optional subsequent tempering. .

In an advantageous way, the steel according to the present invention has the property that hardening occurs even during cooling in stationary air, so that separate cooling of the deformation tool can be omitted.

During hardening, the microstructure of the steel is converted to a temperature above 950° C. by heating in an austenite range, preferably a protective gas atmosphere. During subsequent cooling in air or a protective gas, martensite microstructures are formed for high strength components.

Subsequent tempering facilitates reduction of the intrinsic tension of the cured part. At the same time, the hardness of the part is reduced so that the required toughness values are achieved.

Other features, advantages and details of the present invention will become apparent from the following description of the exemplary embodiment shown in the drawings.
here:
1 shows (schematically) a process chain for the production of a strip made of steel according to the invention.
FIG. 2 is a schematic of the time-temperature curves of the process steps of continuous annealing (including optional hot dip plating treatment) and hot rolling as well as selective tempering, selective quenching (air hardening) and component manufacturing as an example of a steel according to the present invention. ) City.
Figure 3 shows the chemical composition (Examples 1 to 4) of the steel according to the present invention.
4a shows the mechanical property values (transverse to the rolling direction) of the steel according to the invention in the hot rolled state (HR).
4b shows the mechanical property values (transverse to the rolling direction) of the steel according to the invention in the cold rolled state (CR).
Figure 5a shows the characteristic values in the transverse direction to the rolling direction, the solidification behavior during cold rolling of the steel according to the invention.
5b shows the cold flow curve, the solidification behavior during cold rolling of the steel according to the invention.
Figure 6a shows the mechanical property values (transverse to the rolling direction) of the steel according to the invention in the fine sheet state (HDG).
6B shows the results of the plate bending test according to the steel VDA 238-100 according to the invention in the fine sheet state (HDG) and the hole expansion test according to ISO 16630.
FIG. 7A shows the mechanical properties (transverse to the rolling direction) of the steel according to the invention in Example 1 (preliminary strip thickness 40 mm), states HR, CR and HDG.
FIG. 7B shows the mechanical properties (transverse to the rolling direction) of the steel according to the invention in Example 2 (pre-strip thickness 45 mm), states HR, CR and HDG.
FIG. 7C shows the mechanical properties of the steel according to the invention (transverse to the rolling direction) in Example 3 (pre-strip thickness 50 mm), states HR, CR and HDG.
FIG. 7d shows the mechanical properties (transverse to the rolling direction) of the steel according to the invention in Example 4 (pre-strip thickness 55 mm), states HR, CR and HDG.
7E shows the mechanical properties of the steel according to the invention in states HR, CR and HDG as an overview
8A shows Method 1, temperature-time curve (annealing strain is schematically shown).
8B shows Method 2, temperature-time curve (annealing strain is schematically shown).
8C shows Method 3, temperature-time curve (annealing strain is schematically shown).

1 schematically shows a process chain for producing a strip from steel according to the invention. Possible process routes are shown in the present invention. Until pickling, the process route is the same for all steels according to the present invention, followed by a deviation process path depending on the desired result. For example, after pickling, the pickled hot strips can be cold rolled and hot-dipped to a different degree of thinning by rolling. The soft annealed hot strip or soft annealed cold strip can also be cold rolled and hot dip galvanized.

The material can also be selectively treated without a hot dip treatment procedure, ie only within the scope of continuous annealing with or without subsequent electrolytic zinc plating. Complex components can be produced from selectively coated materials. Subsequently, a quenching process such as, for example, air hardening, in which the heat-treated component is cooled in air, can be selectively performed. Optionally, the tempering step can end the heat treatment of the part.

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

FIG. 3 shows examples 1 to 4 from the melt to exclude analytical influences from the alloy composition of the steel according to the invention, depending on the preliminary strip thickness produced. Cold strips with a desired thickness of 1.50 mm cold strips are made from hot strips of the desired thickness of 2.30 mm. Depending on the thickness of the pre-strip to be produced, before hot rolling, Example 1 has a pre-strip thickness of 40 mm, Example 2 has a pre-strip thickness of 45 mm, Example 3 has a pre-strip thickness of 50 mm, Example 4 has a thickness of 55 mm. The alloy composition for the preliminary strips to have is shown.

Figure 4 shows the mechanical property values (transverse to the rolling direction) of the steel according to the invention in the hot rolled state (HR, hot rolling) of Figure 4a and the cold rolled state (CR, cold rolling) of Figure 4b.

5 shows the solidification behavior through mechanical properties in the transverse direction with respect to the rolling direction during cold rolling of the steel according to the invention, in the table of FIG. 5A and the graph as the cold flow curve of FIG. 5B.

Figure 6 shows the mechanical properties of the steel according to the invention in the fine sheet state (HDG, hot-dip galvanized) of Figure 6a (transverse to the rolling direction) and ISO 16630 longitudinally and transversely to the rolling direction in the fine sheet state (HDG) The results of the hole expansion test according to and the plate bending test according to VDA 238-100 as well as the corresponding products with tensile strength in FIG. 6B are shown.

FIG. 7 uses preliminary strip thicknesses of 40 mm in FIG. 7A, 45 mm in FIG. 7B, 50 mm in FIG. 7C, and 55 mm in FIG. 7D, and as seen in the state HR, CR and HDG in FIG. 7 as a summary graphic overview. The mechanical properties of the steel according to the invention (transverse to the rolling direction) are shown.

Figure 8 schematically shows three variations of the temperature-time curve according to the invention for annealing and cooling and in each case different austenitization conditions.

By the different temperature control according to the present invention within the above range, different characteristic values and/or different hole expansion results and bending angles are generated. Therefore, the main difference is the temperature-time parameter during heat treatment and next cooling.

Method 1 (FIG. 8A) shows the annealing and cooling of the manufactured steel strip cold rolled to the final thickness in a continuous annealing plant. First, the strip 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 to a median temperature (ZT) of approximately 200 to 250°C at a cooling rate of approximately 15 to 100°C/s. This schematic does not show the second intermediate temperature (approximately 300-500°C).

Subsequently, the steel strip is cooled in air at a cooling rate of approximately 2 to 30°C/s until it reaches room temperature (RT) or maintained at a cooling rate of 15 to 100°C/s until it reaches room temperature.

Method 2 (FIG. 8B) shows the process according to method 1 but the cooling of the steel strip is then hot dip galvanized to continue cooling to an intermediate temperature of approximately 200 to 250° C. at a cooling rate of approximately 15 to 100° C./s. For the purpose of the procedure, it is suspended for a while while passing through the hot dip plating vessel. The steel strip is cooled in air at a cooling rate of approximately 2 to 30°C/s until room temperature is reached. Method 2 corresponds to annealing as shown in FIG. 8B, for example hot dip galvanizing combined with a direct combustion furnace and a radiant tube furnace.

Method 3 (FIG. 8C) likewise shows the process according to method 1 in the hot dip treatment procedure, but cooling of the steel strip is stopped by a short stop (approximately 1 to 20 seconds) at an intermediate temperature in the range of approximately 200 to 400°C. The heating is performed at a temperature (ST) required for the hot-dip plating treatment procedure (approximately 400 to 470°C). The steel strip is then cooled to an intermediate temperature of approximately 200-250°C. Subsequent cooling in the air of the steel strip is carried out at a cooling rate of approximately 2 to 30°C/s until room temperature is reached.

Method 3 corresponds to a process performed in a continuous annealing plant, for example as shown in FIG. 8C. Further, in this case, by an induction furnace, reheating of the steel is optionally achieved just before the zinc bath.

The reduction from the slab for the preliminary strip varies from 78% to 84% for a subsequent hot rolling to a hot strip thickness of 2.30 mm to a corresponding reduction of 94% to 96% in subsequent examples. In a single cold rolling step, a cold strip of desired thickness of 1.50 mm is achieved with a degree of thinning by cold rolling of 35%. For very low pre-strip thicknesses and larger pre-strip thicknesses and ranges in between, it is impressive that relatively uniform values provided with conventional ranges of variation are achieved for tensile and yield strength in the transverse direction to the rolling direction. Is shown as The steel according to the invention similarly allows the use of a master hot strip thickness with varying degrees of thinning by cold rolling as well as the use of a master cold strip thickness without affecting the previous facts.

For example, in the case of industrial manufacturing for hot dip galvanizing (HDG) according to Method 3 shown in FIG. 8C, the following example is higher hot with a reduction in the pre-strip thickness resulting in significant variations in fine sheets (HDG). Forming part of a so-called feasibility test that demonstrates that variable pre-strip thickness, such as the required rolling force without strip strength (HR) and higher cold strip strength (CR), can significantly affect cold rollability. .

Example 1

(2.30 mm master hot strip and 1.50 mm cold strip from 40 mm pre-strip thickness)

Alloy composition in weight percent. The steel according to the invention is: 0.104% C according to method 3 corresponding to the hot dip treated FIG. 8c; 0.443% Si; 2.178% Mn; 0.012% P; 0.0004% S; 0.0045% N; 0.038 Al; 0.330% Cr; 0.208% Mo; 0.0372% Ti; 0.0332% Nb; 0.007% V; 0.0006% B; 0.0020% Ca; 0.027% Cu; 0.047% Ni; 0.008% Sn; 0.00038% H, and the 250 mm slab material is rolled into a 40 mm pre-strip in a reversing manner at a reduction rate of 84% before hot rolling in the preliminary line, followed by the desired final rolling at 910°C. Hot rolled in a hot wide strip line with a 94% reduction in temperature, reeled at a desired reeling temperature of 650° C. with a master hot strip thickness of 2.30 mm, and without acid heating followed by additional heating (eg batch annealing) It is cold rolled to 1.50 mm with a pass (thinning of about 35% by cold rolling).

Fine sheet condition (HDG)

The transverse yield strength ratio Re/Rm is 66%.

-Elastic limit (Rp0.2) 706 MPa

-Tensile strength (Rm) 1071 MPa

-Elongation at break (A80) 10.9%

-Annealed hardening index (BH2) 492 MPa

-Hole expansion ratio according to ISO 16630 39%

-Bending angle (vertical, horizontal) according to VDA 238-100 121˚/112˚

The material property value in the transverse direction with respect to the rolling direction corresponds to, for example, HC660XD.

Initial state (HR)

The transverse yield strength ratio Re/Rm is 77%.

-Yield strength (Re) 826 MPa

-Tensile strength (Rm) 1070 MPa

-Elongation at break (A80) 10.0%

Landscape in the middle (CR)

-Yield strength (Re) 1246 MPa

-Tensile strength (Rm) 1305 MPa

-Elongation at break (A80) 2.0%

Example 2

(2.30 mm master hot strip and 1.50 mm cold strip from 45 mm pre-strip thickness)

Alloy composition in weight percent.

The steel according to the invention is: 0.104% C according to method 3 corresponding to the hot dip treated FIG. 8c; 0.443% Si; 2.178% Mn; 0.012% P; 0.0004% S; 0.0045% N; 0.038 Al; 0.330% Cr; 0.208% Mo; 0.0344% Ti; 0.0372% Nb; 0.007% V; 0.0006% B; 0.0020% Ca; 0.027% Cu; 0.047% Ni; 0.008% Sn; 0.00038% H, and the 250 mm slab material is rolled into a 45 mm pre-strip in an inverted manner at a reduction rate of 82% before hot rolling in the preliminary line, and subsequently 95% at the desired final rolling temperature of 910°C. Hot rolled in a hot wide strip line with a reduction of 1.30 mm and a master hot strip thickness of 2.30 mm and reeled at the desired reeling temperature of 650° C., after acid cleaning 1.50 mm in one pass without further heating (eg batch annealing) It is cold-rolled to (thinning of about 35% by cold rolling).

Fine sheet condition (HDG)

The transverse yield strength ratio Re/Rm is 67%.

-Elastic limit (Rp0.2) 720 MPa

-Tensile strength (Rm) 1077 MPa

-Elongation at break (A80) 10.4%

-Annealed hardening index (BH2) 51 MPa

-Hole expansion ratio according to ISO 16630 35%

-Bending angle (vertical, horizontal) according to VDA 238-100 128˚/114˚

The material property value in the transverse direction with respect to the rolling direction corresponds to, for example, HC660XD.

Initial state (HR)

The transverse yield strength ratio Re/Rm is 70%.

-Yield strength (Re) 725 MPa

-Tensile strength (Rm) 1030 MPa

-Elongation at break (A80) 10.2%

Landscape in the middle (CR)

-Yield strength (Re) 1224 MPa

-Tensile strength (Rm) 1260 MPa

-Elongation at break (A80) 1.5%

Example 3

(2.30 mm master hot strip and 1.50 mm cold strip from 50 mm pre-strip thickness)

Alloy composition in weight percent.

The steel according to the invention is: 0.104% C according to method 3 corresponding to the hot dip treated FIG. 8c; 0.443% Si; 2.178% Mn; 0.012% P; 0.0004% S; 0.0045% N; 0.038 Al; 0.330% Cr; 0.208% Mo; 0.0344% Ti; 0.0372% Nb; 0.007% V; 0.0006% B; 0.0020% Ca; 0.027% Cu; 0.047% Ni; 0.008% Sn; 0.00038% H, and the 250 mm slab material is rolled into a 50 mm pre-strip in an inverted manner at a reduction rate of 80% before hot rolling in the spare line, and subsequently 96% at the desired final rolling temperature of 910°C. Hot rolled in a hot wide strip line with a reduction of 1.30 mm and a master hot strip thickness of 2.30 mm and reeled at a desired reeling temperature of 650° C., after acid cleaning 1.50 mm in one pass without further heating (eg batchwise annealing) It is cold-rolled to (thinning of about 35% by cold rolling).

Fine sheet condition (HDG)

The transverse yield strength ratio Re/Rm is 65%.

-Elastic limit (Rp0.2) 704 MPa

-Tensile strength (Rm) 1084 MPa

-Elongation at break (A80) 10.4%

-Annealed hardening index (BH2) 55 MPa

-Hole expansion ratio according to ISO 16630 38%

-Bending angle (vertical, horizontal) according to VDA 238-100 127˚/115˚

The material property value in the transverse direction with respect to the rolling direction corresponds to, for example, HC660XD.

Initial state (HR)

The transverse yield strength ratio Re/Rm is 69%.

-Yield strength (Re) 695 MPa

-Tensile strength (Rm) 1010 MPa

-Elongation at break (A80) 8.8%

Landscape in the middle (CR)

-Yield strength (Re) 1203 MPa

-Tensile strength (Rm) 1255 MPa

-Elongation at break (A80) 1.9%

Example 4

(2.30 mm master hot strip and 1.50 mm cold strip from 55 mm pre-strip thickness)

Alloy composition in weight percent.

The steel according to the invention is: 0.104% C according to method 3 corresponding to the hot dip treated FIG. 8c; 0.443% Si; 2.178% Mn; 0.012% P; 0.0004% S; 0.0045% N; 0.038 Al; 0.330% Cr; 0.208% Mo; 0.0344% Ti; 0.0372% Nb; 0.007% V; 0.0006% B; 0.0020% Ca; 0.027% Cu; 0.047% Ni; 0.008% Sn; 0.00038% H, and the 250 mm slab material is rolled into a 55 mm pre-strip in an inverted manner at a reduction rate of 78% before hot rolling in the preliminary line, and subsequently 96% at the desired final rolling temperature of 910°C. Hot rolled in a hot wide strip line with a reduction of 1.30 mm and a master hot strip thickness of 2.30 mm and reeled at the desired reeling temperature of 650° C., after acid cleaning 1.50 mm in one pass without further heating (eg batch annealing) It is cold-rolled to (thinning of about 35% by cold rolling).

Fine sheet condition (HDG)

The transverse yield strength ratio Re/Rm is 66%.

-Elastic limit (Rp0.2) 708 MPa

-Tensile strength (Rm) 1077 MPa

-Elongation at break (A80) 10.4%

-Annealed hardening index (BH2) 58 MPa

-Hole expansion ratio according to ISO 16630 40%

-Bending angle (vertical, horizontal) according to VDA 238-100 123˚/111˚

The material property value in the transverse direction with respect to the rolling direction corresponds to, for example, HC660XD.

Initial state (HR)

The transverse yield strength ratio Re/Rm is 70%.

-Yield strength (Re) 679 MPa

-Tensile strength (Rm) 967 MPa

-Elongation at break (A80) 9.6%

Landscape in the middle (CR)

-Yield strength (Re) 1158 MPa

-Tensile strength (Rm) 1230 MPa

-Elongation at break (A80) 2.5%

conclusion:

The large influence of the preliminary strip thickness on the mechanical properties on the fine sheet (HDG) cannot be seen.

This description applies to the degree of thinning by cold rolling of 35% used in the examples, but can be applied without limitation to the degree of thinning by variable cold rolling.

The invention has been described with the aid of a fine sheet steel sheet with a final thickness to be achieved of 1.50 mm in a thickness range of 0.50 to 3.00 mm. If desired, a final thickness in the range of 0.10 to 4.00 mm can also be produced.

Claims (41)

An ultra-high strength multi-phase steel with a minimum tensile strength of 980 MPa,
The ultra-high strength multi-phase steel (in weight percent) contains:
0.075 ≤ C ≤ 0.115
0.400 ≤ Si ≤ 0.500
1.900 ≤ Mn ≤ 2.350
0.250 ≤ Cr ≤ 0.400
0.010 ≤ Al ≤ 0.060
0.0020 ≤ N ≤ 0.0120
P ≤ 0.020
S ≤ 0.0020
0.005 ≤ Ti ≤ 0.060
0.005 ≤ Nb ≤ 0.060
0.005 ≤ V ≤ 0.020
0.0005 ≤ B ≤ 0.0010
0.200 ≤ Mo ≤ 0.300
0.0010 ≤ Ca ≤ 0.0060
Cu ≤ 0.050
Ni ≤ 0.050
Sn ≤ 0.040
H ≤ 0.0010
The remainder is iron, which generally contains steel-related smelting-related impurities, and the total content of Mn-Si+Cr is 1.750% to 2.250% by weight in relation to the process window as wide as possible during the annealing of the cold strip of this steel, especially during continuous annealing %sign,
Ultra high strength multi-phase steel. According to claim 1,
The total content of Mn-Si+Cr depends on the final thickness of the cold strip to be achieved ultra high strength multi-phase steel with the following content:
Final thickness from 0.50 to 1.00 mm:
1.750 wt% ≤ Mn-sum of Si + Cr ≤ 2.030 wt%
Final thickness from 1.00 mm to 2.00 mm:
1.940 wt% ≤ Mn-sum of Si + Cr ≤ 2.110 wt%
Final thickness greater than 2.00 mm and less than 3.00 mm:
2.020 wt% ≤ Mn-sum of Si + Cr ≤ 2.220 wt%. According to claim 1,
The total content of Mn-Si+Cr+Mo is ultra-high strength multi-phase steel with the following content depending on the final thickness of the cold strip to be achieved:
Final thickness from 50 to 1.00 mm:
1.950 wt% ≤ Mn-sum of Si + Cr + Mo ≤ 2.280 wt%
Final thickness from 1.00 mm to 2.00 mm:
2.140 wt% ≤ Mn-sum of Si + Cr + Mo ≤ 2.360 wt%
Maximum thickness of more than 2.00 mm to less than 3.00 mm... thickness:
2.220 wt% ≤ Mn-sum of Si + Cr + Mo ≤ 2.470 wt%. The method according to any one of claims 1 to 3,
If the final thickness of the cold strip is 0.50 to 1.00 mm:
C content ≤ 0.100% by weight,
If the final thickness of the cold strip is more than 1.00 mm and less than 2.00 mm:
C content ≤ 0.105% by weight, and
If the final thickness of the cold strip is greater than 2.00 mm to less than 3.00 mm:
C content ≤ 0.115% by weight,
Ultra high strength multi-phase steel. The method according to any one of claims 1 to 4,
Carbon equivalent CEV (IIW) is an ultra-high strength multi-phase steel with the following content depending on the final thickness of the cold strip to be achieved:
Final thickness from 0.50 to 1.00 mm:
C content ≤ 0.100% by weight and carbon equivalent CEV(IIW) ≤ 0.62%
Final thickness from 1.00 mm to 2.00 mm:
C content ≤ 0.105% by weight and carbon equivalent CEV(IIW) ≤ 0.64%
Final thickness greater than 2.00 mm and up to 3.00 mm:
C content ≤ 0.115% by weight and carbon equivalent CEV (IIW) ≤ 0.66%. The method according to any one of claims 1 to 5,
The Mn content is an ultra-high strength multi-phase steel with the following content depending on the final thickness of the cold strip to be achieved:
Final thickness from 0.50 to 1.00 mm:
1.900 wt% ≤ Mn content ≤ 2.200 wt%
Final thickness from 1.00 mm to 2.00 mm:
2.050% by weight ≤ Mn content ≤ 2.250% by weight
Final thickness greater than 2.00 mm and less than 3.00 mm:
2.100% by weight ≤ Mn content ≤ 2.350% by weight. The method according to any one of claims 1 to 6,
B content is 0.0009% by weight or less, especially 0.0006% by weight or more and 0.0009% by weight or less,
Ultra high strength multi-phase steel. The method according to any one of claims 1 to 7,
The sum of Nb+Ti is 0.100% by weight or less, especially the sum of Nb+Ti is 0.090% by weight or less,
Ultra high strength multi-phase steel. The method according to any one of claims 1 to 8,
The sum of Ti+Nb+B is 0.102% by weight or less, especially 0.092% by weight or less,
Ultra high strength multi-phase steel. The method of claim 9,
When the sum of Ti+Nb+B is 0.010 to 0.080% by weight, the N content is 0.0020 to 0.0090% by weight, or when the sum of Ti+Nb+B is 0.050% by weight or more, the N content is 0.0040 to 0.0120% by weight,
Ultra high strength multi-phase steel. The method according to any one of claims 1 to 10,
S content is 0.0015% by weight or less, especially S content is 0.0010% by weight or less,
Ultra high strength multi-phase steel. The method according to any one of claims 1 to 11,
The Mo content is greater than 0.200% by weight and less than 0.300% by weight, advantageously greater than 0.200% by weight and less than 0.250% by weight,
Ultra high strength multi-phase steel. The method according to any one of claims 1 to 12,
Ti content is 0.025 to 0.045% by weight,
Ultra high strength multi-phase steel. The method according to any one of claims 1 to 13,
Nb content is 0.025 to 0.045% by weight,
Ultra high strength multi-phase steel. The method according to any one of claims 1 to 14,
V content is 0.005 to 0.020% by weight, optimally 0.005 to 0.015% by weight,
Ultra high strength multi-phase steel. The method according to any one of claims 1 to 15,
The sum of Cr+Mo is less than 0.650% by volume,
Ultra high strength multi-phase steel. The method according to any one of claims 1 to 16,
The sum of Ti+Nb+B+Mo+V is 0.365% by weight or less,
Ultra high strength multi-phase steel. The method according to any one of claims 1 to 17,
Ca content is 0.0030% by weight or less,
Ultra high strength multi-phase steel. The method according to any one of claims 1 to 18,
H content is 0.00050% by weight or less,
Ultra high strength multi-phase steel. The method according to any one of claims 1 to 19,
The sum of H+N is 0.0025 to 0.0130% by weight,
Ultra high strength multi-phase steel. The method according to any one of claims 1 to 20,
When the Cr content is 0.260 to 0.330% by weight, the carbon equivalent CEV (IIW) is 0.62% or less, and when the Cr content is 0.290 to 0.360% by weight, the carbon equivalent CEV (IIW) is 0.64%, and the Cr content is 0.320 to 0.370% by weight. %, the carbon equivalent CEV(IIW) is 0.66%,
Ultra high strength multi-phase steel. The method of claim 21,
Cr content and maximum carbon equivalent CEV(IIW) are ultra high strength multi-phase steels selected as follows depending on the final thickness of the cold strip to be achieved:
Final thickness from 0.50 to 1.00 mm:
0.260 wt% ≤ Cr content ≤ 0.330 wt%, carbon equivalent CEV(IIW) ≤ 0.62%
Final thickness from 1.00 mm to 2.00 mm:
0.290 wt% ≤ Cr content ≤ 0.360 wt%, carbon equivalent CEV(IIW) ≤ 0.64%
Final thickness greater than 2.00 mm and up to 3.00 mm:
0.320% by weight ≤ Cr content ≤ 0.370% by weight, carbon equivalent CEV (IIW) ≤ 0.66%. The method according to any one of claims 1 to 22,
The addition of Si and Mn to the strength properties to be achieved is an ultra-high strength multi-phase steel that can be exchanged according to the following relationship:
YS (MPa) = 185.7 + 147.9 [% Si] + 161.1 [% Mn]
TS (MPa) = 574.8 + 189.4 [% Si] + 174.1 [% Mn] The method according to any one of claims 1 to 23,
The steel can be cured by air cooling,
Ultra high strength multi-phase steel. A method for manufacturing a steel strip from a multi-phase steel according to any one of claims 1 to 24,
The pre-strip is made from multi-phase steel in the slab state, after which the steel strip is hot-rolled to the hot strip thickness to be achieved,
How to make a steel strip. The method of claim 25,
Proceeding from a previously selected pre-strip with previously fixed slab thickness and limited but variable thickness, the hot strip is hot rolled to the same thickness as the degree of lamination by rolling between 72% and 87% to the final thickness to be achieved,
How to make a steel strip. The method of claim 25 and 26,
A hot strip is produced, from which the cold strip is cold rolled to the final thickness to be achieved and subsequently annealed the steel strip, in particular continuous annealing,
How to make a steel strip. The method of claim 27,
Proceeding from a selected master hot strip having a specific thickness or a selected hot strip having a different thickness, a cold strip having a degree of thinning by cold rolling of 10% to 70% is made to the final thickness to be achieved,
How to make a steel strip. The method according to any one of claims 25 to 28,
To produce the required multi-phase microstructure, the steel strip cold rolled to final thickness is heated to a temperature in the range of approximately 700 to 950° C. during continuous annealing,
The annealed steel strip is subsequently cooled from the annealing temperature to a first intermediate temperature of approximately 300 to 500°C at a cooling rate of approximately 15 to 100°C/s, and then approximately at a cooling rate of approximately 15 to 100°C/s. Cooled to a second intermediate temperature of 160 to 250°C, after which the steel strip is cooled in air until reaching room temperature at a cooling rate of approximately 2 to 30°C/s, or approximately 15 to 100°C/s Cooled from the first intermediate temperature to room temperature at a cooling rate,
How to make a steel strip. The method according to any one of claims 25 to 28,
To produce the required multi-phase microstructure, the steel strip cold rolled to the final thickness is heated to a temperature in the range of approximately 700 to 950°C during continuous annealing, and subsequently cooled to a temperature of approximately 400 to 470°C, and cooling is performed. It is stopped before entering the melting bath, after which it is subjected to hot dip treatment, and after the hot dip treatment procedure, cooling continues at an intermediate temperature of about 200 to 250 °C at a cooling rate of about 15 to 100 °C/s. After that, the steel strip is cooled in air until it reaches room temperature at a cooling rate of approximately 2 to 30°C/s,
How to make a steel strip. The method according to any one of claims 25 to 28,
To produce the required multi-phase microstructure, the steel strip cold rolled to the final thickness is heated to a temperature in the range of approximately 700 to 950°C during continuous annealing, subsequently cooled to an intermediate temperature of approximately 200 to 250°C, and melted. The temperature is maintained for approximately 1 to 20 seconds before entering the bath, after which the steel strip is heated to a temperature of approximately 400 to 470° C., followed by hot dip treatment, after the hot dip treatment procedure is performed, approximately 15 New cooling is performed at an intermediate temperature of approximately 200 to 250°C at a cooling rate of 100 to 100°C/s, followed by cooling in air at room temperature at a cooling rate of approximately 2 to 30°C/s. ,
How to make a steel strip. The method according to any one of claims 25 to 31,
During continuous annealing, the oxidation potential during annealing with an equipment configuration consisting of a direct combustion furnace region (NOF) and a radiant tube furnace (RTF) is increased by the CO content in the NOF of less than 4% by volume, and for iron in the RTF. The oxygen partial pressure of the reducing furnace atmosphere is set 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 * (- ln B) 0.25
Here, Si, Mn, Cr, and B represent the corresponding alloy ratio in steel by weight%, pO2 represents the oxygen partial pressure in mbar, and the dew point of the gas atmosphere is -30°C to avoid oxidation of the steel immediately before plating in the melting bath. Set below,
How to make a steel strip. The method of claim 32,
During annealing with a radiant tube furnace, the oxygen partial pressure in the furnace atmosphere satisfies the following equation,
-12> Log pO ₂ ≥ -5*Si ^-0.25-3 *Mn ^-0.5 -0.1*Cr ^-0.5 -7*(-ln B) ^0.5
Here, Si, Mn, Cr, and B represent the corresponding alloy ratio in steel by weight%, pO2 represents the oxygen partial pressure in mbar, and the dew point of the gas atmosphere is -30°C to avoid oxidation of the steel immediately before plating in the melting bath Set below,
How to make a steel strip. The method according to any one of claims 25 to 33,
For steel strips of different thickness during continuous annealing, similar microstructural conditions and mechanical property values of the strip are set by adjusting the plant treatment rate during heat treatment,
How to make a steel strip. The method of claim 27,
The steel strip is temper-rolled following a hot dip treatment procedure or annealing,
How to make a steel strip. The method of claim 27,
The steel strip is stretch-bend-straightened following a hot dip treatment procedure or annealing,
How to make a steel strip. The method according to any one of claims 25 to 36,
The blank is cut from the steel strip, then heated to a temperature above Ac3, the heated blank is deformed to make the part, and then cured in the tool or in air,
How to make a steel strip. A steel strip produced by the method according to any one of claims 25 to 37,
With a minimum hole expansion according to ISO 16630 of 20%, especially 25%,
River strip. The method of claim 38,
Comprising a minimum bending angle according to VDA 238-100 of 70° in the longitudinal or transverse direction, in particular 85°,
River strip. The method of claim 38 or 39,
Including the minimum production value Rm×α (tensile strength×[bending angle according to VDA 238-100]) of 100000 MPa˚, especially 120000 MPa˚,
River strip. The method according to any one of claims 38 to 40,
Delayed fracture-free condition for at least 6 months, thus meeting the requirements according to SEP 1970 for hole pull and hoop test pieces,
River strip.

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