KR20190142768A - High strength steel sheet with excellent ductility and elongation flangeability - Google Patents

Abstract

FIELD OF THE INVENTION The present invention relates to high strength steel sheets such as hot rolled and cold rolled products useful for vehicles and frame parts for automobiles, such as truck frames, and to methods of making such steel sheets as well. .

Description

High strength steel sheet with excellent ductility and elongation flangeability

FIELD OF THE INVENTION The present invention relates to high strength steel sheets such as hot rolled and cold rolled products useful in vehicles such as frames for trucks and frame parts for automobiles.

Recently, (advanced) high strength steel sheets, AHSS, are increasingly being used in automotive parts to reduce weight and fuel consumption. To meet growing requirements, HSLA, two phase (DP), ferrite-bainite including stretch-flangeable (SF), composite phase (CP), transformation-organic plasticity (TRIP), hot A series of (advanced) high strength steels have been developed, such as molding, twin-organic plasticity (TWIP).

However, AHSS sheet steels cannot be easily applied to a wide variety of automotive parts because of their relatively lack of formability. As steels become stronger and stronger, at the same time it becomes increasingly difficult to form steels into automotive parts. In fact, the practical application of AHSS steels (DP, CP, and TRIP) to automotive parts is still limited because of their formability. Therefore, improving moldability and manufacturability is an important issue in AHSS applications.

The relationship between elongation and strength of AHSS has been well defined through standard tensile tests, leading to the well-known strength-elongation banana curve. The microstructural parameters that determine the strength and ductility of the AHSS are understood qualitatively and only slightly quantitatively. However, elongation is not the only parameter that determines moldability in AHSS. AHSS grades include additional associated failure mechanisms when compared to mild steels. This is mainly due to local breakdowns that are more commonly observed in AHSS because of phase changes during polyphase structure and deformation. These local breaks are not necessarily associated with elongation and / or n-values. Thus, steels with higher (uniform and overall) elongation do not always have good formability. Microstructures that improve ductility differ from microstructures that improve formability. The position in the elongation-strength chart is not sufficient to select the appropriate materials for all parts. In most cases, for material selection, different relationships between formability and strength are needed. Research into the action of AHSS under all relevant molding conditions is essential. There are four basic tasks in automotive press molding with varying stress and strain conditions: deep drawing, stretching, stretch-flanging, and bending. . Each shaping scheme has certain dominant mechanical parameters such as r-value (ratio between in-plane plastic and thickness-through plastic deformation of the tensile test specimen), lambda (hole expansion ratio) value, and bending angle. For some mold-distressed parts, high punchability, stretch-flange and fatigue properties are required in application.

It is an object of the present invention to provide a steel grade that combines high yield and tensile strength with good elongation values and good hole expansion values.

It is also an object of the present invention to provide a steel grade having a yield strength of at least 570 MPa, a tensile strength of at least 760 MPa, and a hole expansion factor (λ) value of at least 70%.

This object is a cementless-free microstructure having a yield strength of at least 570 MPa, a tensile strength of at least 760 MPa, a total elongation (A50) of at least 10.3%, and a hole expansion rate (λ) value of at least 70%, including Made through high strength steel strip with:

* C, 0.005-0.08 wt.%;

* Mn, 1.30-2.30 wt.%;

* B, 2-35 ppm;

N, 5-65 ppm;

* Al_tot, 0.005-0.1 wt.%;

Ti, 0.03 to 0.20 wt.%;

Cu, 0-1.5 wt.%;

* Cr, 0-0.75 wt.%;

Mo, 0-0.05 wt.%;

Ni, 0-0.50 wt.%;

* V, 0-0.30 wt.%;

* Si, 0-0.6 wt.%;

* P, 0-0.03 wt.%;

* S, 0-0.01 wt.%; And

Residual iron and unavoidable impurities, where C / (Ti_sol + V) ≦ 0.25.

The unique and balanced combination of chemical elements ensures that the steel microstructure consists of only bainite and ferrite components, ideally including bainite and ferrite components. Preferably, the entire microstructure consists only of bainite components. Although sometimes the distinction between ferrite and bainite components is difficult, it is easy to distinguish between ferrite and bainite components on the one hand and structures such as martensite, residual austenite, cementite, pearlite and the like.

The absence of cementite (Fe ₃ C) in the microstructure is very important in the present invention. Through the addition of titanium (Ti) and vanadium (V), cementite formation is prevented, instead TiC and VC are formed. These latter carbides are much smaller (˜5-30 nm) and more finely dispersed than cementite (˜200 nm) that would normally be present. Cementite will also be plate-shaped, while VC and TiC are usually spherical or needle-shaped and located in ferrite rods, while located between ferrite laths in the bainite structure. This microstructure provides an improved combination of strength and fracture toughness. This microstructure draws high strength from ultrafine particle sizes of less than 1 μm and can be further strengthened by small carbide precipitates. In the present invention, the formation of Fe ₃ C or martensite / austenite microcomponents is suppressed using microalloying such as titanium and vanadium, in order to produce high moldability cementite free bainite steels. In addition, when microstructured tablets are used at the same time, microalloying precipitation strengthening is an effective way to increase strength without sacrificing toughness. At higher carbon contents, higher amounts of microalloying with V and Ti are needed to prevent (coarse) Fe ₃ C.

C is an element that forms cementite free bainite and contributes to the increase in strength. The low carbon content between 0.005 and 0.08 wt.% In the steel necessarily lowers the cooling rate dependence of the microstructure. In a preferred embodiment of less than 0.05 wt.%, Preferably less than 0.045 wt.%, More preferably less than 0.04 wt.%, Even more preferably less than 0.035 wt.%, When the carbon content is so low, No carbon distribution occurs during the transformation of austenite to ferrite. The minimum carbon content suitable to obtain the bainite structure and to ensure high strength with UTS of 760 MPa or more is 0.01 wt.%. By optimizing other alloying elements, it is possible to obtain uniform bainite microstructures which are formed very similarly over a very wide range of cooling rates. Since there are no carbide formations, some cementite will form. Bainite microstructures can be made cementite free by further alloying with Ti and / or V. In addition, low carbon concentrations can result in good low temperature impact toughness that balances proper weldability.

Titanium and vanadium play a very important role in the formation of high strength cementite free bainite. Ti combines with N, S and C to form nitrides, carbon sulfides and carbides, depending on the specific chemical composition of the steel. When the Ti content exceeds 0.20%, it is difficult to dissolve coarse Ti carbides during reheating of the slab before hot rolling. In this case, V is added to replace some amount of Ti, and V combines with the remaining carbon to form VC. The titanium content is at least 0.03%.

The formation of TiC and VC will completely prevent the formation of cementite. Fine carbides TiC and / or VC with a particle size of less than 10 nm in the ferrite phase during primary cooling (air cooling) following primary cooling after hot rolling thus contribute to the increase in strength. Ti also serves to fix N. All free nitrogen is detrimental to the hardenability to enhance the effect of B. Therefore, the nitrogen removal effect of titanium is required. Vanadium content is 0.30% or less.

In order to obtain hot-rolled steel sheets with high ductility and high hole expansion rates, cementite formation should be prevented. That is, C / (Ti_sol + V) must satisfy the following equation:

Where Ti_sol represents the amount of Ti that can form Ti carbides:

The appropriate minimum value of C / (Ti_sol + V) is 0.15. In order to promote the curing effect of boron, Ti_sol> 0, preferably> 0.01 wt.%, More preferably> 0.02 wt.%.

The term cementite free is intended to mean that no cementite (Fe ₃ C) is present in the microstructure at all. A composition that satisfies the equation given above will ensure this. However, due to local compositional variations in the steel strip, it can happen accidentally and unintentionally that trace amounts of cementite are identified in the microstructure that do not affect the properties and performance of the entire steel strip.

Manganese (Mn) is an essential element for promoting low carbon bainite microstructures and improving the balance between strength and low temperature toughness. The Mn content is at least 1.30 wt.% And up to 2.30 wt.%. Mn stabilizes austenite, delays bainite transformation at a certain temperature and ensures good hardenability. The austenite region extends to lower temperatures, providing a wide temperature gap for proper controlled rolling. In addition, Mn promotes the formation of fine needle-like ferrites and bottom-bainite. Disadvantages of very high Mn contents are poor surface quality after hot rolling due to lower HAZ toughness, increased centerline segregation of continuous cast steel slabs, and increased internal oxidation. The Mn content for the steels of the invention is preferably at least 1.5 wt%, and preferably at most 2.0 wt.%. More preferably, Mn is at least 1.65 wt.% And at most 1.95 wt.%.

Boron (B) is a strong hardening enhancer in low carbon, low alloyed steels. A small amount of B is added to the low carbon steels to ensure that bainite microstructures can be produced at lower cooling rates without the formation of cornerstone ferrite. The B content is at least 2 ppm and at most 35 ppm. B is the most effective alloying element for increasing the yield strength. The B content should preferably be 25 ppm or less so as not to impair low temperature toughness. In order for boron to play this role, it is essential that no free nitrogen is present to prevent BN formation. This is where the nitrogen removal effect of titanium is introduced.

Nitrogen (N) is inevitably present in the BOF steel manufacturing process. N has a strong affinity for Ti, forming Ti nitrides that act as a dispersoid for controlling austenite particle size during reheating. The N content is at least 5 ppm and at most 65 ppm. If the N content exceeds 50 ppm (= 0.005 wt.%), A relatively large amount of Ti is required to secure the ability to form Ti carbides that contribute to strengthening and protect free boron, which leads to an increase in cost. . Therefore, the N content is 50 ppm or less. The N content is preferably reduced as much as possible. A suitable and substantial minimum N content is 10 ppm. The titanium content is at least 0.03 wt.% And at most 0.20 wt.%, Preferably at least 0.06 wt.% And preferably at most 0.18 wt.%, More preferably at most 0.16 wt.%.

The steels are preferably additionally alloyed with one or both of copper (Cu) or chromium (Cr), with a maximum content of 1.50 wt.% Cu and 0.75 wt.% Cr. Suitable maximum content is 1.25 wt.% Cu. Cu can promote low carbon bainite structures and provide solid solution hardening. The strength of the steel is increased by precipitation hardening of the nano-sized Cu precipitates. Through the hot-process precipitation control process (TPCP), Cu precipitation can be obtained during coil cooling after hot rolling, so no additional thermal process is necessary. Cr increases strength mainly due to transformational transformation.

Nickel (Ni) improves hardenability as well as toughness. Nickel imparts good toughness to the steel at high strength levels. Ni not only increases the strength and toughness of the steel, but also copes with the hot shortness due to all Cu alloys. Ni is preferably present only as impurities, mainly in terms of cost. Ni may be added up to 0.5 wt.% To prevent thermal brittleness when the Cu content exceeds 0.5%. In one embodiment, nickel is not added to the steel.

Silicon (Si) is added to improve strength in spite of solution hardening and transformation hardening. However, when Si is excessive, HAZ toughness, weldability and coating property are impaired. Silicon is maximized up to 0.6 wt.%, Preferably up to 0.5 wt.%.

Aluminum (Al) is used as a deoxidation element and is an effective element to improve steel cleanliness. In order to achieve this effect, the total Al content in the steel must be set at 0.005 wt.% Or more. The Al content is maximized up to 0.1 wt.% And preferably up to 0.05 wt.%. The higher the content, the higher the probability of surface defects and the higher the alloy cost.

Molybdenum (Mo) has been found to promote low carbon bainite structure at low concentrations of about 0.1 wt.%, But at higher concentrations it may degrade the toughness of high strength cementite free bainite steels. Since Mo is not an economically preferred alloying element, it is not recommended to use it as an alloying element in these steels.

Sulfur (S) is present as an impurity in the steel. Initial MnS particles will be formed in the Mn containing steels during casting. These coarse MnS particles are very harmful because they elongate in the rolling direction. When Ti is added, Ti ₄ S ₂ C ₂ and / or MnS is formed during casting depending on the concentration of Ti, C and S. Ti ₄ S ₂ C ₂ will be present as initial coarse particles and should be prevented as much as possible. The S content should be at most 0.012 wt.%, Preferably at most 0.01 wt.% And more preferably less than 0.005 wt.%.

Phosphorus (P) is present as an impurity in the steel. If the P content exceeds 0.03%, segregation becomes noticeable at grain boundaries, leading to deterioration of toughness and weldability. P content is preferably reduced as much as possible. The P content should be 0.012 wt.% Or less, preferably 0.01 wt.% Or less, and most preferably less than 0.005 wt.%.

The steel according to the invention can be hot-rolled steel and can be used as such, or can subsequently be cold-rolled and annealed steel. Hot-rolled or cold-rolled steel may have a metal coating. The coating may be provided through hot dip plating, preferably the metal coating is a zinc-based or aluminum-based coating. Selective coating of steel is carried out by conventional means, including but not limited to hot dip coating, electrocoating, PVD or CVD.

It should be noted that the steel according to the invention does not contain niobium (Nb) as an alloying element. The values in Table 1 indicate that they are present only as impurities and niobium is not added to the steel.

According to a second aspect, the present invention also provides a cementite free structure, a yield strength of at least 570 MPa, a tensile strength of at least 760 MPa, a total elongation (A50) of at least 10.3%, and a hole expansion ratio (λ) of at least 70%. It is embodied in a method of making a high strength steel strip having:

Casting the melt into slabs or strips having the following composition;

* C, 0.005-0.08 wt.%;

* Mn, 1.30-2.30 wt.%;

* B, 2-30 ppm;

N, 5-65 ppm;

* Al_tot, 0.005-0.1 wt.%;

Ti, 0.03 to 0.20 wt.%;

Cu, 0-1.5 wt.%;

* Cr, 0-0.75 wt.%;

Mo, 0-0.05 wt.%;

Ni, 0-0.50 wt.%;

* V, 0-0.30 wt.%;

* Si, 0-0.6 wt.%

* P, 0-0.03 wt.%;

* 0-0.01 wt.% S;

* Residual iron and unavoidable impurities. Where C / (Ti_sol + V) ≦ 0.25;

Reheating the slab to a slab reheating temperature of at least 1,200 ° C. and hot rolling the slab or strip into a hot-rolled strip with a hot-rolling finishing temperature above Ar 3; And

Cooling the hot-rolled strip to a winding temperature of less than 500 ° C. at a run-out table at an average cooling rate of 15 ° C./s to 100 ° C./s, followed by coil cooling to room temperature. .

In the process according to the invention, the steel melt is cast in the form of thick slabs, thin slabs or strips in a conventional manner. After casting, the (re) heating and / or homogenization results in hot-rolling temperatures and hot-rolling. The last hot-rolling step is carried out for steels which are still completely austenite, ie when the finish rolling temperature exceeds Ar3. After finishing rolling, the steel is cooled to a winding temperature of 500 ° C. or less at an average cooling rate of 15 ° C./s to 100 ° C./s on the run-out table of the hot rolling mill, followed by natural coil cooling of the coil to ambient temperature. The slab reheating temperature of the steel should be high enough to dissolve the coarse Ti and V carbides deposited in the slab during casting. The inventors found that an SRT of 1,200 ° C. or higher is required. A suitable maximum SRT is 1,300 ° C. The hot rolling finish temperature should be in the austenite range, preferably between 850 ° C and 950 ° C. This range of 850 ° C. to 950 ° C. is applied to produce fine austenite particle sizes in the strip after the last rolling step and to keep Ti and V in solid solution. Immediately after the last hot rolling step (maximum, within a time of 2 seconds between finish rolling and start of cooling), the strips are cooled by accelerated cooling to a winding temperature of 500 ° C. or less at a speed in the range of 15 ° C./s to 100 ° C./s. After winding, the coil is cooled to ambient temperature without further accelerated cooling. To avoid misunderstandings, mention that ambient temperature has the same meaning as room temperature.

In order to avoid the formation of pearlite, and to avoid the formation of ferrite and coarse Ti and V carbides, an average cooling rate after hot rolling of 15 ° C./s to 100 ° C./s is necessary.

After cooling, the hot-rolled sheet is wound at winding temperatures up to 500 ° C. The wound strip cools slowly, which allows for bainite phase transformation to occur. The bainite phase formed in the steels according to the invention while winding in this winding temperature range is cementite free, which is desirable for the steel sheet to have good elongation flangeability. Instead of the formation of Fe ₃ C, precipitation of fine TiC and / or VC carbides can occur within this winding temperature range, which allows for further curing.

If the winding temperature exceeds 500 ° C., cementite may be formed of pearlite or degraded pearlite, and the resulting stretch flangeability will be significantly lower than in the process according to the invention.

Preferably, the winding temperature is at least 420 ° C. If the winding temperature value is lower, there is a risk of too much martensite and residual austenite being formed, which will reduce the formability of the resulting steel. This risk is more important if you have a higher carbon content. For steels with a C content in the range of 0.03 wt.% To 0.08 wt.%, The winding temperature should be in the range of 420 ° C. to 500 ° C., because the microstructure and properties, especially elongation flangeability, are very sensitive to the process paths. do. The CCT plot for the VS72 alloy shown in FIG. 1 shows that different microstructures grow with cooling rate.

For steels with a carbon content of less than 0.03 wt.%, The cooling rate after hot rolling should be in the range 15 ° C./s to 100 ° C./s and the winding temperature should be less than 500 ° C. Since the risk of martensite formation or the presence of residual austenite is very low, the minimum winding temperature is not specified. However, for caution, the winding temperature is preferably not lower than 420 ° C. The inventors found that the mechanical properties are relatively insensitive to cooling rate and cooling temperature. As the CCT diagram (FIG. 2) and the microstructures (FIG. 3) for the VS74 alloy already suggest, the final structure is relatively insensitive to the cooling trajectory. This is especially true for steel grades containing Cu and / or Cr.

The cooling rate range can be obtained by spraying water or air / water mixture at the exit of the finishing mill, depending on the thickness of the sheet.

In view of ensuring stable properties, the winding temperature is preferably at least 440 ° C and / or at most 480 ° C. Preferably, after rolling, the cooling rate before winding is at least 25 ° C./s.

If the hot-rolled strip is subsequently cold rolled, cold rolling reduction is applied to obtain the required thickness. The total cold-roll reduction rate of the hot-rolled strip is preferably 50% to 90%. The cold-rolled fully-cured strip is reheated to a dissolution temperature above Ac 3, preferably in the temperature range of 850 ° C.-1,000 ° C., and maintained at the dissolution temperature for 2 to 8 minutes and then at 440 ° C. to 480 ° C. Cooled to a holding temperature at a cooling rate in the range of 15 ° C./s to 50 ° C./s and held for up to 30 minutes, and preferably 0.5 to 30 minutes, resulting in bainite transformation and adversely affecting formability. To avoid the risk of martensite formation. The steel must then be cooled to room temperature at a cooling rate of 0.5 ° C./s to 100 ° C./s. Cooling rates for alloys containing less than 0.03 wt.% C should preferably be between 10 ° C./s and 100 ° C./s. Higher reheat temperatures are desirable to allow more TiC to dissolve into austenite.

According to a third aspect, the invention is also embodied in an automobile or truck part, such as an automobile chassis part, a part of a body in white, a frame or an auxiliary frame part, which part is made of a steel sheet according to the invention. Has been.

1 is a CCT diagram for a VS72 alloy,
2 is a CCT diagram for a VS74 alloy,
3 is a microstructure of the VS74 alloy (taken from the samples of FIG. 2).

The present invention is now described with reference to the following non-limiting examples.

Steels with the compositions shown in Table 1 were cast into 30 kg ingots with dimensions of 200 mm × 110 mm × 110 mm. Steel VS71 is a comparative example because the C / (T_sol + V) ratio is outside the composition range of the present invention. The ingots were reheated to 1,250 ° C. and immersed for one hour, and then hot hot rolled to 35 mm thick. The contraction and segregation areas at both ends were cut. The cut mass was reheated at 1,200 ° C. for 30 minutes and then hot rolled to 3 mm thickness in five passes. The finishing rolling temperature was about 900 ° C.

Immediately after hot rolling, the strips were cooled at a run-out table at a rate of 30-60 ° C./s to 500 ° C./s, then transferred to a furnace preheated to 440 ° C. or 480 ° C. and held for one hour to mimic the winding process. . Materials were removed from the furnace and cooled to room temperature in air. The hot rolled strips were then pickled with 85 ° C. HCl to remove oxide layers. Specimens for microstructure observation, tensile test, and hole expandability test were processed for hot rolled strips.

The hot rolled strips were then cold-rolled at a cold rolling reduction of 67%. The cold rolled strips of 1 mm were then heat treated at 900 ° C. for 2 minutes and then cooled to three different conditions, as specified in Table 3.

Tensile Test —From the resulting hot-rolled, cold-rolled, and annealed sheets, European standard test pieces (dimension length = 50 mm; width = 12.5 mm) were processed so that the tensile direction was parallel to the rolling direction. Room temperature tensile tests were performed with a Sxhenk TREBEL test instrument conforming to the NEN10002 standard to confirm the tensile properties (yield strength YS (MPa), final tensile strength UTS (MPa), total elongation TE (%)). For each condition, three tensile tests were performed and the average value of the mechanical properties reported.

Hole Expandability Test (Elongation Flange Evaluation Test) —Samples of test pieces (size: 90 × 90 mm) were tested from the obtained rolled sheet to test hole expandability. According to the Japanese Iron and Steel Association Standard JFS T 1001, a 10 mm diameter hole was drilled in the center of the test piece and a 60 ° conical punch was pushed into the hole. When the crack penetrated, the sheet thickness and hole diameter d (mm) were measured. The hole expansion ratio λ (%) was calculated from the following equation: λ (%) = {(dd ₀ ) / d _o } × 100, where d ₀ is 10 mm.

Flexural Test —A 3-point “inductive flexural test” was performed on specimens with a size of 40 mm × 30 mm. The longitudinal direction of the specimens was parallel to the rolling direction of the steel sheets. Parallel flexural tests were performed in which the flexural axis was perpendicular to the rolling direction of the sheets. In this method, a molding machine and two supporting cylinders were used to bend the steel sheet. Cylinders and punches were mounted in a tensile test instrument. A load cell was used to measure the force of the punch and moved the punch through the movement of the crosshead. The experiments were discontinued at different deflection angles, and the curved surface of the specimen was examined for burst identification to determine the deflection angle.

CCT-graphs (FIGS. 1 and 2) of steels VS72 and VS74 were established using a standard dilatometer (heating rate 5 ° C./s up to 900 ° C., held for 5 minutes). Microstructure evaluation (FIG. 3) revealed that, for different cooling rates, the microstructure was consistently ferrite-bainite at VS74.

Claims (15)

A high strength steel strip having a cementite free microstructure,
* C, 0.005-0.08 wt.%;
* Mn, 1.30-2.30 wt.%;
* B, 2-35 ppm;
N, 5-65 ppm;
* Al_tot, 0.005-0.1 wt.%;
Ti, 0.03 to 0.20 wt.%;
Cu, 0-1.5 wt.%;
* Cr, 0-0.75 wt.%;
Mo, 0-0.05 wt.%;
Ni, 0-0.50 wt.%;
* V, 0-0.30 wt.%;
* Si, 0-0.6 wt.%;
* P, 0-0.03 wt.%;
* S, 0-0.01 wt.%;
* Residual iron and unavoidable impurities
Wherein C / (Ti_sol + V) ≦ 0.25 (Ti_sol = Ti− (48/14) · N),
A steel having a yield strength of at least 570 MPa, a tensile strength of at least 760 MPa, a total elongation (A50) of at least 10.3%, and a hole expansion rate (λ) value of at least 70%. The method of claim 1,
* 0-1.5 wt.% Cu; And
* Steel containing at least one of 0-0.75 wt.% Cr. The method according to any one of claims 1 to 2,
Steel, wherein the microstructure comprises bainite and ferrite particles. The method according to any one of claims 1 to 3,
Steel with a C of 0.045 wt.% Or less. A method of making a high strength steel strip having a cementite free structure, a yield strength of at least 570 MPa, a tensile strength of at least 760 MPa, a total elongation (A50) of at least 10.3%, and a hole expansion rate (λ) value of at least 70%.
Casting the melt into slabs or strips having the following composition;
* C, 0.005-0.08 wt.%;
* Mn, 1.30-2.30 wt.%;
* B, 2-35 ppm;
N, 5-65 ppm;
* Al_tot, 0.005-0.1 wt.%;
Ti, 0.03 to 0.20 wt.%;
Cu, 0-1.5 wt.%;
* Cr, 0-0.75 wt.%;
Mo, 0-0.05 wt.%;
Ni, 0-0.50 wt.%;
* V, 0-0.30 wt.%;
* Si, 0-0.6 wt.%
* P, 0-0.03 wt.%;
* 0-0.01 wt.% S;
Residual iron and unavoidable impurities, where C / (Ti_sol + V) ≦ 0.25) (Ti_sol = Ti- (48/14) .N);
Reheating the slab to a slab reheating temperature of at least 1,200 ° C. and hot rolling the slab or strip into a hot-rolled strip, wherein the hot-rolling finishing temperature exceeds Ar 3; And
After cooling the hot-rolled strip in a run-out table at an average cooling rate of 15 ° C./s to 100 ° C./s, within a 2 second time between the finish rolling and the start of cooling to a winding temperature of less than 500 ° C. And coil cooling to ambient temperature. The method of claim 5,
The slab or strip
* 0-1.5 wt.% Cu; And
* Containing at least one of 0-0.75 wt.% Cr. The method according to claim 5 or 6,
Wherein the carbon content of the steel is at least 0.03% and the winding temperature is at least 420 ° C. The method according to any one of claims 5 to 7,
C is 0.045 wt.% Or less. The method according to any one of claims 5 to 8,
The winding temperature is at least 440 ° C and / or at most 480 ° C. The method according to any one of claims 5 to 9,
Wherein the hot-rolled strip is subsequently cold-rolled to obtain a cold-rolled strip. The method of claim 10,
Wherein the cold-rolled strip is annealed by reheating the strip to a temperature above Ar ₃ , maintaining the strip, and subsequently cooling the strip to ambient temperature. The method of claim 10,
The total cold rolling reduction is between 50% and 90%. The method according to any one of claims 10 to 12,
The cold-rolled fully-cured strip is reheated to a dissolution temperature above Ac3 within the range of 850 ° C.-1,000 ° C., maintained at the dissolution temperature for 2 to 8 minutes, and cooled in the range of 15 ° C./s to 50 ° C./s And cooled to a holding temperature between 440 ° C. and 480 ° C. at a rate, held at holding temperature for 0-30 minutes to cause bainite transformation, and then cooled to room temperature. An automobile or truck part, such as an automobile chassis part, a part of a white car body, a part of a frame or subframe, which is made of a steel sheet according to any one of claims 1 to 4. An automobile or truck part, such as an automobile chassis part, a part of a white car body, a part of a frame or subframe, which is manufactured by the method according to any one of claims 5 to 13.

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