EP3612650A1

EP3612650A1 - High strength steel sheet having excellent ductility and stretch flangeability

Info

Publication number: EP3612650A1
Application number: EP18717386.9A
Authority: EP
Inventors: Shangping Chen; Richard MOSTERT
Original assignee: Tata Steel Nederland Technology BV
Current assignee: Tata Steel Nederland Technology BV
Priority date: 2017-04-20
Filing date: 2018-04-19
Publication date: 2020-02-26
Anticipated expiration: 2038-04-19
Also published as: WO2018193032A1; US20200071789A1; EP3612650B1; CN110621794A; CN110621794B; KR20190142768A

Abstract

This invention relates to a high strength steel sheet as hot rolled and cold rolled products useful for frame components for vehicles and automobiles such as frames for trucks, and to a method of producing the steel sheet, as well as a use thereof.

Description

HIGH STRENGTH STEEL SHEET HAVING EXCELLENT DUCTILITY AND STRETCH FLANGEABILITY

This invention relates to a high strength steel sheet as hot rolled and cold rolled products useful for frame components for vehicles and automobiles such as frames for trucks.

In recent years, (advanced) high strength steel sheets, AHSS, are increasingly used in car components to reduce weight and fuel consumption. A series of (advanced) high strength steels, such as HSLA, Dual phase (DP), Ferritic-bainitic (FB) including stretch- flangeable (SF), Complex phase (CP), Transformation-induced plasticity (TRIP), Hot- formed, Twinning-induced plasticity (TWIP) has been developed to meet the growing requirements.

However, AHSS sheet steels cannot be applied easily to a wide variety of car components because their formability is relatively poor. As steels became increasingly stronger, they simultaneously became increasingly difficult to form into automotive parts. Actually, the real application of AHSS steels (DP, CP and TRIP) to car components is still limited by their formability. Therefore, improving formability and manufacturability become an important issue for AHSS application.

The relationship between elongation and strength of AHSS has been well established from the standard tensile tests and leads to the well-known strength-elongation banana curve. The microstructural parameters that govern the strength and the ductility of AHSS are understood qualitatively and to a lesser degree quantitatively. However, elongation is not the only parameter governing formability in AHSS. AHSS grades have additional relevant failure mechanisms compared to mild steels. This is mainly caused by local failure which is observed more commonly in AHSS due to multi-phase structure and phase changes during deformation. These local failures do not necessarily correlate with elongation and/or n-value. Therefore, steels having higher (uniform and total) elongations do not always have a good formability. The microstructures improving ductility are different from those improving formability. The position in the diagram of the elongation-strength is not sufficient to select the proper materials for all parts. In most cases, another relationship between formability and strength is needed for material selection. It is essential to study the behaviour of AHSS under all relevant forming conditions. There are four basic operations in automotive press forming with various stress and strain states: deep drawing, stretching, stretch-flanging and bending. Each forming mode has a specific governing mechanical parameter such as r-value (the ratio between plastic strain in-plane and the plastic strain through-the-thickness of a tensile test sample), λ (hole expansion ratio) value, and bending angle. For some difficult-to-form parts, high punchability, stretch-flangeability and fatigue properties are demanded in the application. It is an object of the present invention to provide a steel grade which combines high yield and tensile strength with a good elongation and excellent hole expansion ratio values.

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

This object is reached by a high strength steel strip having a cementite-free microstructure comprising :

0.005 - 0.08 wt.% C;

1.30 - 2.30 wt.% Mn;

· 2 - 35 ppm B;

• 5 - 65 ppm N;

• 0.005 - 0.1 wt.% AI_tot;

• 0.03 to 0.20 wt.% Ti;

0 - 1.5 wt.% Cu;

0 - 0.75 wt.% Cr;

• 0 - 0.05 wt.% Mo;

• 0 - 0.50 wt.% Ni;

0 - 0.30 wt.% V;

0 - 0.6 wt.% Si;

0 - 0.01 wt.% P;

0 - 0.01 wt.% S;

C/(Ti_sol+V)<0.25

remainder iron and inevitable impurities, the steel strip 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 ratio (λ) value of at least 70%.

The unique and balanced combination of chemical elements ensures that the microstructure of the steel comprises bainitic and ferritic components, and ideally consists only of bainitic and ferritic components. Preferably the entire microstructure consists of bainitic components only. Even though it is sometimes hard to distinguish between ferritic and bainitic components, it is easy to distinguish between ferritic and bainitic components on the one hand, and structures like martensite, retained austenite, cementite, pearlite, etc. on the other hand.

Crucial in the invention is the absence of cementite (Fe3C) in the microstructure. By the addition of titanium (Ti) and vanadium (V) cementite formation is prevented and TiC and VC are formed instead. These latter carbides are much smaller (~ 5-30 nm) and more finely dispersed than the cementite (~ 200 nm) that would normally have been present. The cementite would also be plate shaped and located between the ferrite laths in the bainite structure, whereas the VC and TiC are usually spherical or needle like and located inside the ferrite lath. This microstructure offers an improved combination of strength and fracture toughness. The microstructure derives its high strength from the ultra-fine grain size, less than Ιμιτι, which can additionally be strengthened by the small carbide precipitates. In the present invention, to produce cementite-free bainitic steels with high formability, the formation of Fe3C or martensitic/austenitic microconstituents is suppressed by using microalloying such as titanium and vanadium. Moreover, microalloying precipitation strengthening is an effective way to increase strength without sacrificing toughness, when microstructural refinement is used simultaneously. At higher carbon contents, higher amount of micro alloying with V and Ti is necessary to avoid (coarse) Fe₃C.

C is an element that forms cementite-free bainite, thus contributing to an increase in strength. The low carbon content of between 0.005 and 0.08 wt.% in the steel ensures that the cooling rate dependence of the microstructure is low. At these low carbon contents, in a preferable embodiment of less than 0.05 wt.%, preferably less than 0.045 wt.%, more preferably less than 0.04%, even more preferably less than 0.035 wt.% carbon partitioning does not occur during the austenite to ferrite transformation. A suitable minimum carbon content is 0.01 wt.% to obtain the bainitic structure and ensure the high strength of a UTS of 760 MPa or more. By optimizing the other alloying elements, it is possible to obtain a uniform bainitic microstructure that will form across a very large range of cooling rates in a very similar manner. In the absence of carbide formers, some cementite will form. The bainitic microstructures can be made cementite free by further alloying with Ti and/or V. Also, the low carbon concentrations may bring about good low temperature impact toughness balanced with adequate 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, carbosulphides 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 prior to hot rolling. In this case, V is added to replace some amount of Ti and V will combined with rest of 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 grain size of less than 10 nm in the ferrite phase during natural cooling (air cooling) subsequent to primary cooling after hot rolling, thus contributing to an increase in strength. Ti also plays the role of fixing N. Any free nitrogen is detrimental to the hardenability improving effect of B, and therefore the nitrogen scavenging effect of the titanium is desired. The vanadium content is at most 0.30%.

To obtain a hot-rolled steel sheet having a high ductility and a high hole expansion ratio, it is necessary to prevent the formation of cementite. That is, C/(Ti_sol+V) must satisfy the following expression : where Ti_sol represents the amount of Ti that can form Ti carbides:

48\

Ti_sol = Ti - -) - N

A suitable minimum value for C/(Ti_sol+V) is 0.15. To promote the hardening effect of boron, the Ti_sol > 0, preferably > 0.01 wt.%, more preferably > 0.02 wt.%.

The term cementite-free is intended to mean that the aim is that no cementite (Fe3C) whatsoever is present in the microstructure. A composition satisfying the equation presented above should ensure that this is the case. However, due to local compositional fluctuations in the steel strip it may inadvertently and unintentionally occur that a minute amount of cementite is discernible in the microstructure that does not affect the properties and performance of the steel strip as a whole.

Manganese (Mn) is an essential element for promoting low carbon bainitic microstructures and in improving the balance between strength and low temperature toughness. The Mn content is at least 1.30 and at most 2.30 wt.%. Mn stabilizes austenite and delays the bainite transformation at a given temperature and ensures a good hardenability. The austenite field is extended to lower temperatures, which offers a wide temperature gap for proper controlled rolling. Furthermore, Mn promotes the formation of fine acicular ferrite and lower-bainite. Disadvantages of very high Mn contents are, deterioration of HAZ toughness, increased centreline segregation of the continuous cast steel slabs, and poor surface quality after hot rolling due to increased interior oxidation. For the steel of this invention, the Mn content is preferably at least 1.5, and preferably at most 2.0 wt.%. More preferably Mn is at least 1.65 and at most 1.95 wt.%.

Boron (B) is a potent hardenability enhancer in low C, low alloy steels. A small amount of B is added to low carbon steels to ensure that bainitic microstructures can be produced at lower cooling rates without formation of proeutectoid ferrite. The B content is at least 2 and at most 35 ppm. B is the most effective alloying element in increasing the yield strength. The B content should preferably be at most 25 ppm so as not to impair low temperature toughness. For the boron to be able to perform this role, it is essential that no free nitrogen is present so that the formation of BN is avoided. This is where the nitrogen scavenging effect of the titanium comes in.

Nitrogen (N) is inevitably present in the BOF steel making process. N is an element that has a strong affinity for Ti, and forms Ti nitrides which acts as dispersoids for austenite grain size control during reheating. The N content is at least 5 and at most 65 ppm. When 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 which contribute to strengthening and protect the free boron, which results in a rise in costs. Therefore, the N content is 50 ppm or less. Desirably, the N content is decreased as much as possible. A suitable and practical minimum N content is 10 ppm. The titanium content is at least 0.03 and at most 0.20 wt.% Ti, and preferably at least 0.06 and at preferably at most 0.18, more preferably at most 0.16 wt.%.

The steels are preferably additionally alloyed with one or both of copper (Cu) or chromium (Cr) to a maximum of 1.50 wt.% Cu and 0.75 wt.% Cr. A suitable maximum is 1.25 wt.% Cu. Cu can promote low carbon bainitic structures, and provide solid solution hardening. The strength of the steel is increased by precipitation hardening of nano-sized Cu precipitates. Through a thermo-mechanical precipitation control process (TPCP) it is possible to obtain Cu precipitation during coil cooling after hot rolling, and therefore no extra heat treatment is necessary. Cr increases the strength mainly due to transformation strengthening.

Nickel (Ni) improves toughness as well as hardenability. Nickel imparts good toughness to the steel material at a high level of strength. In addition to increasing the strength and toughness of the steel, Ni counters the hot shortness caused by any Cu alloying. Ni is preferably present only as an impurity, mainly from a cost perspective. Ni can be added up to 0.5 wt% to prevent hot shortness when the Cu content exceeds 0.5%. In an embodiment no nickel is added to the steel.

Silicon (Si) is added to improve the strength though solution hardening and transformation hardening. However, with excess of Si, the HAZ toughness, weldability and coatability are impaired. The silicon is maximised at 0.6 wt.%, preferably maximised at 0.5%.

Aluminium (Al) is utilized as a deoxidizing element and is an element effective for improving the steel cleanliness. It is necessary to set the total Al content in the steel at 0.005 wt.% or more to obtain such an effect. The Al content is maximised at 0.1 wt.% and preferably at 0.05 wt.% because the higher the content the higher the likelihood of cause surface defects and the higher the costs of the alloy.

Molybdenum (Mo) has been found to promote low carbon bainitic structure at small concentrations of about 0.1 wt%, but at higher concentration it can deteriorate toughness of high strength cementite free bainitic steels. Mo is not an economically preferred alloying element and it is not recommended to use as an alloying element in these steels.

Sulphur (S) is present in steel as an impurity. Primary MnS particles will be formed in Mn containing steels during casting. These coarse MnS particles are very detrimental because they are elongated in the rolling direction. When Ti is added, T14S2C2 and/or MnS are formed during casting depending on the concentrations of Ti, C and S. T14S2C2 will be present as primary coarse particles and needs to be avoided as much as possible. The S content should be at most 0.012, preferably at most 0.01 wt.% and most preferably below 0.005 wt.%.

Phosphorus (P) is present in steel as impurity. When the P content exceeds 0.03%, segregation in the grain boundaries becomes marked, resulting in degradation in toughness and weldability. Desirably, the P content is decreased as much as possible. The P content should be at most 0.012, preferably at most 0.01 wt.% and most preferably below 0.005 wt.%.

The steel according to the invention may be a hot-rolled steel and used as such, or a subsequently cold-rolled and annealed steel. The hot-rolled or cold-rolled steel may be provided with a metallic coating. The coating may be provided by means of hot-dipping, and preferably the metallic coating is a zinc or aluminium based coating. The optional coating of the steel is performed by conventional means and includes, but is not limited to, hot dip coating, electrocoating, PVD or CVD.

It is noted that the steel according to the invention does not contain niobium (Nb) as alloying element. The values in Table 1 indicate a presence as impurity only, and no niobium is added to the steel.

According to a second aspect the invention is also embodied in a process for producing a high strength steel strip having a cementite-free microstructure, 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) value of at least 70%, said process comprising :

casting a melt into a slab or strip having the following composition;

0.005 - 0.08 wt.% C;

1.30 - 2.30 wt.% Mn;

· 2 - 30 ppm B;

5 - 65 ppm N;

0.005 - 0.1 wt.% AI_tot;

0.03 to 0.20 wt.% Ti;

0 - 1.5 wt.% Cu;

0 - 0.75 wt.% Cr;

0 - 0.05 wt.% Mo;

0 - 0.50 wt.% Ni;

0 - 0.30 wt.% V;

0 - 0.6 wt.% Si;

0 - 0.01 wt.% P;

0 - 0.01 wt.% S;

C/(Ti_sol+V)<0.25,

remainder iron and inevitable impurities,

reheating the slab to a slab reheating temperature of at least 1200 °C, hot rolling the slab or strip to a hot-rolled strip wherein the hot-rolling finishing temperature is above Ar3;

cooling the hot-rolled strip at an average cooling rate of 15 to 100 °C/s on the runout table to a coiling temperature below 500 °C followed by coil cooling to room temperature. In the process according to the invention a steel melt is conventionally cast in the form of a thick slab, a thin slab or a strip. After casting it is brought to hot-rolling temperatures by (re-)heating and/or homogenising and hot-rolled. The last hot-rolling pass is performed on the steel while still fully austenitic, i.e. the finish rolling temperature is above Ar3. After finish rolling the steel is cooled on the run-out table of the hot strip mill at an average cooling rate of between 15 and 100 °C/s to a coiling temperature of at most 500 °C followed by cooling of the coil by natural cooling down to ambient temperature. The slab reheating temperature for the steel has to be sufficiently high to dissolve coarse Ti and V carbides precipitated in the slab during casting. The inventors found that a SRT of at least 1200 °C was required. A suitable maximum SRT is 1300 °C. The hot rolling finishing temperature has to be in the austenite range, and is preferably between 850 °C to 950 °C. This range of 850 °C to 950 °C is applied to produce a fine austenite grain size in the strip after the last rolling pass, and to keep Ti and V in solid solution. Immediately after the last hot rolling pass (at most within a time period of 2 seconds between finish rolling and start of cooling) the strips are cooled by accelerated cooling at rates in the range of 15 to 100 °C/s to a coiling temperature of at most 500°C. After coiling the coil is allowed to cool to ambient temperature without further accelerated cooling. For the sake of avoiding misunderstanding it is noted that ambient temperature has the same meaning as room temperature.

The average cooling rate after hot rolling of 15 to 100 °C/s is needed to avoid the formation of pearlite and to avoid the formation of ferrite and coarse Ti and V carbides.

After cooling, the hot-rolled sheet is coiled at a coiling temperature up to 500°C. The coiled strip cools slowly, which allows for the bainitic phase transformation to occur. The bainite phase formed in the steels according to the invention during coiling in this coiling temperature range is cementite-free, which is preferable for the steel sheet to exhibit excellent stretch flangeability. Instead of formation of Fe3C, the precipitation of fine TiC and/or VC carbides may occur within this coiling temperature range, enabling additional hardening to be obtained.

If the coiling temperature exceeds 500 °C, cementite might be formed as pearlite or degenerated pearlite and the resulting stretch flangeability is markedly lower than in the process according to the invention.

Preferably the coiling temperature is at least 420 °C, because for lower values of the coiling temperature there is a risk of the formation of too much martensite and retained austenite and the formability of the resulting steel will be reduced as a consequence thereof. This risk is more prominent for higher carbon contents. For steels with a C content in the range of 0.03 to 0.08 wt%, the coiling temperature should be in the range of 420 to 500 °C , as the microstructure and properties, especially stretch flangeability are rather sensitive to the process routes. The CCT diagram for the VS72 alloy shown in Figure 1 indicates that depending on the cooling rate different microstructures develop. For steels with the C contents being less than 0.03 wt%, the cooling rate after hot rolling has to be in the range of 15 to 100 °C/s and the coiling temperature has to be below 500 °C. No minimum coiling temperature is specified, because the risk of forming martensite or the presence of retained austenite is very low. However, to be on the safe side, the coiling temperature is preferably not below 420 °C. The inventors found that the mechanical properties were relatively insensitive to cooling rate and the cooling temperature. As the CCT diagram for VS74 alloy (Figure 2) and the microstructures (Figure 3) already suggest, the final structure is relatively insensitive to the cooling trajectory. This is particularly the case for the steel grades containing Cu and/or Cr.

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

From the viewpoint of securing stable properties, the coiling temperature is preferably at least 440 °C and/or at most 480 °C. Preferably the cooling rate after rolling and prior to coiling is at least 25 °C/s.

If the hot-rolled strip is subsequently cold rolled, a cold rolling reduction is applied to obtain the required thickness. The total cold-rolling reduction of the hot-rolled strip is preferably between 50 and 90 %. The cold-rolled full-hard strip is reheated to a solution temperature above Ac3, preferably in the temperature range of 850-1000 °C, and held at the solution temperature for 2 to 8 minutes, and then cooled with a cooling rate in the range of 15 to 50 °C/s to a holding temperature between 440 and 480 °C, and held for up to 30 min, and preferably for 0.5 to 30 min, to allow the bainitic transformation to take place and avoid the risk for formation of martensite in amounts that adversely affect the formability. Thereafter, the sheet should be cooled with a cooling rate of 0.5 to 100 °C/s to room temperature. For alloys containing less than 0.03 wt% C, the cooling rate should preferably be between 10 and 100 °C/s. A higher reheating temperature is preferable to allow more TiC dissolve in to austenite.

According to a third aspect the invention is also embodied in a car or truck component, such as an automotive chassis component, a component of the body in white, a component of the frame or the subframe, said component having been produced from the steel sheet according to the invention.

The invention will now be described with reference to the following non-limiting examples.

Steels having compositions shown in Table 1 were cast into 30 kg ingots of 200 mm x 110 mm x 110 mm in dimensions. Steel VS71 is a comparable example because the C/(T_sol+V) ratio is out of the invented composition range. The ingots were reheated to 1250 °C and soaked for 1 hour and then rough hot rolled to 35 mm thickness. The shrinkage and segregation zone from both ends were cut off. The cut blocks were reheated at 1200 °C for 30 min and then hot rolled to 3 mm thickness in 5 passes. The finish rolling temperature was about 900 °C. Immediately after hot rolling, the strips were cooled at 30-60 °C/s to 500 °C in the run-out table and were then transferred to a preheated furnace at 440 °C or 480 °C and held for 1 hour to simulate the coiling process. The materials were then taken out the furnace and cooled in the air to the room temperature. The hot rolled strips were then pickled in HCI at 85 °C to remove the oxide layers. Samples for microstructure observations, tensile tests and hole expansion tests were machined for hot rolled strips.

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

Tensile tests - Euronorm test pieces (gauge length = 50 mm; width = 12.5 mm) were machined from the obtained hot-rolled and cold rolled and annealed sheets such that the tensile direction was parallel to the rolling direction. Room temperature tensile tests were performed in a Schenk TREBEL testing machine following NEN 10002 standard to determine tensile properties (yield strength YS (MPa), ultimate tensile strength UTS (MPa), total elongation TE (%)). For each condition, three tensile tests were performed and the average values of mechanical properties are reported.

Hole Expansion Test (Stretch Flanqeability Evaluation Test) - Test pieces for testing hole expandability (size: 90x90 mm) were sampled from the obtained rolled sheet. In accordance with The Japan Iron and Steel Federation Standards JFS T 1001, a 10 mm diameter punch hole was punched in the centre of the test piece and a 60° conical punch was pushed up and inserted into the hole. When a crack penetrated the sheet thickness, the hole diameter d (mm) was measured . The hole expansion ratio λ (%) was calculated by the following equation : A(%)={(d-do)/d₀} l00, with do being 10 mm.

Bending test - The 3-point "guided bending tests" were conducted on samples with dimensions 40 mm x 30 mm. The length direction of the samples was parallel to the rolling direction of steel sheets. Parallel bending tests where the bending axis is perpendicular to the rolling direction of the sheets were carried out. For this method, a former and two supporting cylinders were used in order to bend the steel sheets. The cylinders and the punch were mounted in a tensile testing machine. The load cell is used to measure the punch force and the displacement of the crosshead gives the punch displacement. The experiments were stopped at different bending angles and the bent surface of the specimen was inspected for identification of failure in order to determine the bending angle.

The CCT-diagrams (Figure 1 and Figure 2) of steels VS72 and VS74 were determined by means of standard dilatometry (heating rate 5 °C/s to 900 °C, holding for 5 minutes). Microstructural evaluation (Figure 3) revealed that the microstructure was consistently Ferritic-Bainitic for VS74 alloy after different cooling rate. Table 1. Chemical compositions of the cast steels in wt.%, except B and N in ppm.

x = impurity level.

Table 2. Mechanical properties of the hot rolled 3 mm strips.

Table 3. Properties of the cold-rolled and annealed steels (thickness 1.0 mm)

Figure 1. The CCT diagram for VS72 alloy.

Figure 2. The CCT diagram for VS74 alloy.

Figure 3. The microstructure of the VS74 alloy (taken from the samples of figure 2).

Claims

High strength steel strip having a cementite-free microstructure comprising :

• 0.005 - 0.08 wt.% C;

• 1.30 - 2.30 wt.% Mn;

• 2 - 35 ppm B;

• 5 - 65 ppm N;

• 0.005 - 0.1 wt.% AI_tot;

• 0.03 to 0.20 wt.% Ti;

• 0 - 1.5 wt.% Cu;

. 0 - 0.75 wt.% Cr;

• 0 - 0.05 wt.% Mo;

• 0 - 0.50 wt.% Ni;

• 0 - 0.30 wt.% V;

• 0 - 0.6 wt.% Si;

• 0 - 0.01 wt.% P;

. 0 - 0.01 wt.% S;

• ——≤ 0.25 wherein Ti_sol=Ti-((48/14)-N)

Steel according to claim 1 containing at least one of

0 - 1.5 wt.% Cu;

• 0 - 0.75 wt.% Cr;

Steel according to any one of the preceding claims wherein the microstructure comprises bainitic and ferritic grains.

Steel according to any one of the preceding claims wherein C is at most 0.045 wt.%.

Process for producing a high strength steel strip having a cementite-free microstructure, 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) value of at least 70%, said process comprising :

casting a melt into a slab or strip having the following composition;

0.005 - 0.08 wt.% C;

1.30 - 2.30 wt.% Mn;

• 2 - 35 ppm B;

• 5 - 65 ppm N;

0.005 - 0.1 wt.% AI_tot;

0.03 to 0.20 wt.% Ti;

0 - 1.5 wt.% Cu;

0 - 0.75 wt.% Cr;

0 - 0.05 wt.% Mo;

0 - 0.50 wt.% Ni;

0 - 0.30 wt.% V;

0 - 0.6 wt.% Si;

0 - 0.01 wt.% P;

0 - 0.01 wt.% S;

——≤ 0.25 wherein Ti_sol=Ti-((48/14)-N);

remainder iron and inevitable impurities,

hot rolling the slab or strip to a hot-rolled strip;

cooling the hot-rolled strip at an average cooling rate of 15 to 100 °C/s on the run-out table within a time period of 2 seconds between finish rolling and start of cooling to a coiling temperature below 500 °C and then coil cooling by natural cooling to ambient temperature.

Process according to claim 5 wherein the slab or strip contains at least one of

0 - 1.5 wt.% Cu;

• 0 - 0.75 wt.% Cr;

Process according to claim 5 or claim 6 wherein the carbon content of the steel is at least 0.03% and wherein the coiling temperature is at least 420 °C.

Process according to any one of claims 5 to 7 wherein C is at most 0.045 wt.%.

Process according to any one of claims 5 to 8 wherein the coiling temperature is at least 440 °C and/or at most 480 °C.

10. Process according to any one of claims 5 to 9 wherein hot-rolled strip is subsequently cold-rolled to obtain a cold-rolled strip.

11. Process according to claim 10 wherein the cold-rolled strip is annealed by reheating the strip to a temperature above Ar₃, holding it and subsequently cooling it to ambient temperatures.

12. Process according to claim 10 wherein the total cold rolling reduction is between 50 and 90%.

13. Process according to any one of claim 10 to 12 , wherein the cold-rolled full-hard strip is reheated to a solution temperature above Ac3 in the range of 850-1000 °C, held at the solution temperature for 2 to 8 minutes, cooled with a cooling rate in the range of 15 to 50 °C/s to a holding temperature between 440 and 480 °C, held at the holding temperature for 0 to 30 min to allow the bainitic transformation to take place, and then cooled to room temperature.

14. Car or truck component, such as an automotive chassis component, a component of the body in white, a component of the frame or the subframe, said component having been produced from the steel sheet according to any one of claim 1 to 4.

15. Car or truck component, such as an automotive chassis component, a component of the body in white, a component of the frame or the subframe, said component having been produced by means of the process according to any one of claim 5 to 13.