This invention relates to a method for producing a high strength strip steel with a good deep drawability and a high strength steel produced thereby.

Car manufacturers focus the main design criteria of a modern car on reduced weight and high safety in order to fulfil the customer's expectations and the legal requirements and standards. In this context, new steel grades have been developed not only for weight saving but also to improve crash safety of vehicles. These steels combine an increased formability with a high strength level at a wide strain rate spectrum. Advanced high strength steels, among them especially dual phase and TRIP steels, feature promising results in this field, while their extraordinary mechanical properties can be tailored and adjusted by alloying and processing. Different phases with a wide variety of strength levels together with an extraordinary fine microstructure contribute to the excellent cold formability and crash behaviour of AHSS.

AHSS are multiphase steels which contain phases like martensite, bainite and retained austenite in quantities sufficient to produce unique mechanical properties. Compared to conventional high strength steels, AHSS exhibit higher strength values or a superior combination of high strength with good formability (Bleck, W. ; Phiu-On, K., Grain Refinement and Mechanical Properties in Advanced High Strength Sheet Steels, HSLA Steels 2005, Fifth International Conference on HSLA Steels, 08.-10.11.2005, Sanya, Hainan, China). In principle one can distinguish four types of AHSS:

· Dual phase: microstructure comprising ferrite and 5-30% martensite

• Transformation Induced Plasticity steels with a microstructure of ferrite, bainite and retained austenite

• PM : partly or fully martensitic steels

• Complex Phase steels with a mixture of strengthened ferrite, bainite and martensite.

Although these advanced high strength steel sheets have excellent strength and formability in terms of the strength-ductility relationship, it is known that these AHSS suffer from a low r-value and therefore a poor deep drawability. For complex automotive stamped parts, the deep drawability can be critical. It is therefore an object of the invention to provide a method to produce AHSS with an improved r-value.

It is also an object of the invention to provide an AHSS with an improved r- value.

It is another object of the invention to provide an AHSS with an improved r- value while keeping the excellent strength-ductility relationship.

One or more of these objects may be reached by a method for producing a high- strength steel sheet having excellent deep drawability, the method comprising producing a hot-rolling precursor by melting and casting a precursor comprising (in wt.%):

• C: 0.01 to 0.2%;

• Si : 2.0% or less;

• Mn : 1.0 to 3.0%;

• P: 0.005 to 0.1%;

· S : 0.03% or less;

• ALsol : 0.01 to 0.1%;

• Cr: 2.0% or less

. N : 0.002 to 0.01%;

• Remainder Fe and inevitable impurities

wherein Ti and B as impurities should be kept as low as possible and wherein N-14/48*Ti-14/l l*B > 0.0015%

the precursor optionally also comprising

• Ca : 0.0005 to 0.01%;

• Mo : 0.8% or less

· Cu : 0.8% or less

• Ni : 0.8% or less

- reheating the precursor material to a temperature of at least 1100°C to dissolve all, or at least most, AIN-precipitates or in case of a thin slab cast precursor or strip-cast precursor to be rolled immediately after casting, maintaining the temperature at a temperature of at least

1100°C to prevent, or at least substantially reduce, AIN from forming;

- hot rolling the precursor to a hot-rolled strip by hot rolling at a finishing temperature above Ar3;

- cooling the hot-rolled strip to a temperature so as to substantially prevent precipitation of AIN in the hot-rolled strip followed by:

a. coiling the hot rolled strip at a temperature of at most 650°C, or by b. coiling the hot rolled strip at a temperature of at most 700°C and subsequently cooled by full immersion of the coil in a water bath and cooled to a temperature of 600°C or lower

- pickling the cooled hot-rolled strip;

- cold-rolling the hot-rolled strip with a reduction in thickness of between

50 and 90%;

- batch recrystallisation annealing the cold-rolled strip at a temperature of between 550°C and Acl to achieve a favourable crystallographic texture for deep-drawing;

- cooling the recrystallised strip;

- continuous intercritical annealing of the recrystallised strip by reheating the strip to a temperature between Acl and Ac3, holding the strip between Acl and Ac3 for at most 600 seconds to obtain a microstructure comprising at most 50% austenite, preferably at most 30%, followed by cooling from the holding temperature to an intermediate temperature of

350-500°C at a cooling rate of between 1 and 100°C/s and followed by a. cooling without delay from the intermediate temperature to a temperature between ambient temperature and 200°C at a cooling rate of between 1 and 100°C/s, or

b. holding the strip at the intermediate temperature of between 350 and 500°C for a period of time which is not more than 600s;

- wherein the final microstructure of the strip after batch annealing and continuous annealing is a microstructure comprising ferrite and one or more second phases such as martensite, residual austenite and/or bainite.

Preferred embodiments are provided in the dependent claims.

This method achieves a high r-value of the cold-rolled and annealed product by first batch annealing a cold rolled steel to obtain a recrystallised strip followed by a continuous annealing to produce the mixed microstructure of the AHSS as described above. This method deviates from the normal (i.e. conventional) processing route of reheating → hot rolling → cold rolling → continuous annealing. Due to their composition (high C, Mn etc) and annealing in the intercritical (i.e. two-phase α-γ) temperature region, AHSS produced by this (conventional) route have a low r-value. The process according to the invention is as follows: reheating → hot rolling → cold rolling → batch recrystallisation annealing → continuous annealing. The batch annealing is to produce a favourable crystallographic texture (high fraction of (111) oriented crystals) and the continuous annealing is to achieve a mixed microstructure of ferrite and the desired second phase (martensite, retained austenite, bainite). Figure 6 gives a schematic representation of subsequent processing steps, wherein HR represents hot rolling, followed by cooling (C), coiling, uncoiling, pickling (P), cold rolling (CR) either in line with pickling or in a separate CR-line, followed by batch recrystallisation annealing (BA) and continuous annealing (CA) to produce the desired final microstructure, whereas figure 7 shows the conventional (prior art) route.

Alternative ways to achieve higher r-value is by cold rolling→ annealing→ cold rolling → annealing. The process according to the invention eliminates one cold rolling step and is therefore cost effective. Still another processing is as follows: ferritical hot-rolling with good lubrication → annealing → cold rolling → continuous annealing. This alternative process is complicated and needs substantially modifications of existing hot rolling facility. Moreover, the anisotropy of the ferritically hot-rolled products is quite restrictive.

It is important that the composition of the steel produced using the method according to the invention contains no nitride forming elements other than aluminium, i.e. no Ti, Nb, V, Zr or B-additions. These nitride forming elements Ti, B, Nb, V, and Zr may only be present as impurities. It is necessary that N- (14/48)*Ti-(14/l l)*B > 0.0015%. This guarantees that free nitrogen is available to bind to aluminium and influence the texture favourably during the annealing process. Preferably N-(14/48)*Ti-(14/l l)*B-(14/93)*Nb-(14/51)*V-(14/91)*Zr > 0.0015%.

The aluminium content in the steel is expressed as AI_sol (acid soluble), meaning that the aluminium is not bound to oxygen. The killing of the steel during steelmaking requires the addition of a deoxidant, usually aluminium and sometimes silicon, to bind the oxygen into alumina. The majority of that alumina is removed from the melt, but any alumina remaining behind in the steel is not available any longer for binding to nitrogen. The process from preparing the melt and producing a hot-rolling precursor by melting and casting a precursor, such as a thick slab, a thin slab or a cast strip, is fairly straightforward. Austenitic finish hot-rolling is important to create a fine equiaxed structure of the hot rolled strip, and in the cooling stage of the hot rolled strip it is essential that the precipitation of aluminium-nitride is prevented. AIN that precipitates in the hot strip is no longer available for precipitation during the batch-annealing after cold rolling and is therefore no longer beneficial in achieving the right crystallographic texture. The aluminium and nitrogen should be in solid solution in the cooled hot-rolled strip. Also any cementite that precipitates in the hot strip after hot-rolling should be as fine as possible. After pickling and cold rolling of the hot rolled strip the cold-rolled full-hard strip is subjected to recrystallisation annealing by batch annealing. Precipitation of aluminium nitride (AIN) inhibits nucleation so that the most favoured (l l l)-grains are promoted, (l l l)-grains represent 'cube on corner' grains and this orientation of the crystal planes in the microstructure is favourable for deep-drawing, i.e. for the r-value. This requires optimal slow heating to allow precipitation to occur at an early stage of recrystallisation, as in batch annealing. The effect of the AIN precipitation is to reduce the number of nuclei so that the recrystallised structure has larger (pancake-shaped) grains resulting in a stronger (l l l)-texture which is favourable for good r-values. In an embodiment the batch annealed and subsequently continuous annealed strip has a gamma-fibre texture (i.e. {l l l}-fibre) of at least 5 times random. Preferably the ratio of gamma-fibre texture over 'cube-on-face' texture ({100}) is at least 2, preferably at least 3, more preferably at least 4.

When continuously annealing the recrystallised strip the ferritic or cementite, or ferritic- pearlitic microstructure starts to transform into austenite, and depending on the maximum temperature of the annealing, the ratio of newly formed austenite and untransformed ferrite differs. By keeping the annealing temperature down, the fraction transforming to austenite is limited. If present, the pearlite fraction and cementite will transform to austenite first. The amount of ferrite, with the already favourable texture, transforming to austenite is to be as small as possible, and therefore the annealing temperature should not be too high. Too high an annealing temperature will also reduce the average carbon content in the austenite, and therefore the critical cooling rate for the austenite transforming to the desired microstructure needs to be higher as well. As a result of the fast cooling a part of the austenite, enriched in carbon, will transform into the desired second phase (martensite, bainite) or stay untransformed as retained austenite. The latter depends on the enrichment of the austenite and the presence of any austenite stabilising alloying elements such as manganese and nickel. The cooling after intercritical annealing has to be fast enough to induce the desired second phase as martensite or bainite to form. The fast cooling is continued to an intermediate temperature of between 350 and 500°C. At that stage of the process the choice is to continue cooling to a temperature between 200°C and ambient temperature (for instance on a continuous annealing line without an overageing section), to hold the strip at this intermediate temperature for a short period of time (for instance on a continuous annealing line with an overageing section), and/or to subject the strip to hot dip galvanizing by dipping the strip in a bath of molten metal, usually a zinc bath or zinc based alloy. The zinc based alloy may be alloyed with aluminium and or magnesium such as Magizinc ^® . The strip may also be coated by means of electrocoating such as electrogalvanizing. In that case the cooling does not need to halt at the intermediate temperature but the cooling can be continued without delay to a temperature between 200°C and ambient temperature. Although the cooling rate is prescribed by the fact that it must be sufficiently high to achieve the desired microstructure, the inventors found that a cooling rate between l-100°C/s, and preferably of at most 50°C/s is adequate to achieve this aim.

For the sake of avoiding misunderstanding : upon reaching the intermediate temperature of between 350 and 500°C three options are available:

1. The strip is held at a temperature between 350 and 500°C for a period of time (overageing), and this period of time is not more than 600 seconds, and preferably not more than 300 seconds (see figure 1) or

2. The strip is not held at a temperature between 350 and 500°C for a period of time and the cooling is continued without delay to a temperature to between ambient temperature and 200°C.

3. Hot dip galvanising the strip by dipping the strip in a molten zinc or molten zinc-alloy bath. This can be combined with either option 1 or

2.

In an embodiment of the invention the hot rolled strip is coiled at a temperature of at most 600°C or even 550°C.

After finishing the hot rolling, the hot rolled strip is cooled to the coiling temperature. The cooling rate on the run-out table (i.e. after finish hot-rolling and before coiling) is at least l°C/s and preferably at least 10°C/s.

The coiling temperature should not be chosen too low. The formation of lower bainite structures (forming between around 250 to 400°C) in the hot-rolled coil should be prevented (upper bainite structures forming around between 550 and 400°C are considered allowable) and the formation of pearlite promoted (provided the carbon content is chosen high enough to allow pearlite to be formed). This low coiling temperature ensures that no aluminium nitride precipitates in the coil upon cooling of the hot-rolled strip to ambient temperature. A low coiling temperature is also beneficial to the homogeneous distribution of the aluminium and nitrogen in solid solution along the length of hot rolled coils. In an embodiment of the invention the hot rolled strip is coiled at a temperature of at most 700°C and subsequently cooled by full immersion of the coil in a water bath and allow it to cool to a temperature of not higher than 600°C, i.e. to a temperature below which the precipitation of AIN does not occur within the timeframe of the process, more preferably to a temperature of not higher than 500°C. The immersion in water should take care of sufficiently fast cooling of the strip to prevent AIN precipitation. As the centre of the coil cools down very slowly, even when immersed in water, the combination of coiling temperature, size of the coil (width of the strip and weight of the strip) should be carefully chosen to avoid any precipitation of AIN in the hot-rolled condition also in the centre of the coil. From a practical point of view, low temperature coiling is preferable over this high temperature option followed by immersion in water. The cooling rate on the run-out table (i.e. after finish hot-rolling and before coiling) to the coiling temperature of at most 700°C is at least l°C/s and preferably at least 10°C/s. In an embodiment the coiling temperature prior to immersion is at most 650°C.

In an embodiment of the invention the strip during continuous intercritical annealing is held between Acl and Ac3 to obtain a microstructure comprising at most 40% austenite, preferably at most 30%. By keeping the amount of austenite down, the amount of ferrite grains transforming is kept down, and the enrichment in carbon of the austenite is increased, thereby maximising the chances of obtaining the desired second phase(s) and retaining as much as possible of the favourable crystallographic texture.

In an embodiment of the invention the reheating temperature of the precursor material is at least 1100°C, preferably at least 1150°C and more preferably at least 1200°C. This high temperature ensures that all, or at least most, of the aluminium nitride formed during the cooling down of the cast slab is dissolved. When hot-rolling in a thin slab-casting-direct rolling facility such as a CSP or a strip-casting - direct rolling facility it is important that the temperature of the cast thin slab or strip does not drop to a level that AIN-precipitation starts. Therefore the homogenisation temperature between casting and rolling must be such that the aluminium and nitrogen stay in solid solution. For this purpose it is preferable that the homogenisation temperature of the precursor material is at least 1100°C, preferably at least 1125°C. A suitable maximum homogenisation temperature would be 1150°C or even 1200°C. The higher the reheating or homogenisation temperature, the higher the likelihood that all Al and N go into or remain in solution. The temperature where all AIN is dissolved depends on the chemistry of the steel, but the kinetics of the precipitation play a role as well.

In an embodiment of the invention the heating rate to the holding temperature during intercritical annealing is between 1 to 40°C/s.

In an embodiment of the invention the strip is provided with a metallic coating by hot dip coating the continuously annealed strip or electrogalvanising.

In an embodiment the cold rolling reduction is at least 60%.

In an embodiment of the invention the heating rate of the batch annealing is at least 5°C/hr and/or at most 50°C/hr. A preferable minimum heating rate is 10 °C/hr. A preferably maximum heating rate is 40°C/hr or even 30°C/hr. In the batch annealing process, 2-5 coils separated by a convector plate are stacked on a furnace base. A protective cover encloses the coil where an inert or reducing gas is circulated. H ₂ or HNX is commonly used for this purpose. This enclosure is externally heated by a gas or oil fired furnace. The outer and inner surfaces of the coils are heated by convection from the circulating gas and by radiation between cover and coil. The inner portion of the coil is heated through conduction, which is retarded due to the air gaps between the sheets. As a result, during the annealing operation, different locations in the coil undergo different thermal cycles. In an industrial operation, the two locations of special interest are hot spot (near the coil surface with the highest temperature during the heating cycle) and cold spot (in the coil core with the lowest temperature), where by monitoring these two extreme locations, the product quality is controlled. The heating rate of at most 50°C/hr, preferably at most 40°C/hr relates to the heating of the hot spot.

Precipitation of aluminium nitride (AIN) inhibits nucleation by pinning the deformed grain boundaries so that the most favoured (111) grains are promoted. This requires optimal slow heating to allow precipitation to occur at an early stage of the annealing to allow the pan-cake grains to form. Too slow a heating rate will result in too long an annealing process, which is not economic. Moreover, the risk of recovery of the full hard microstructure is avoided. Too high a heating rate means that the recrystallisation starts well before the AIN precipitation is able to pin the grain boundaries. To avoid excessive grain growth it is important to limit the soaking time during batch annealing. Soaking time is defined as the time during which the hot spot is at the annealing temperature. During that time the cold spot also reaches the annealing temperature. After the cold spot reaches the annealing temperature the cooling of the coil can start. After the required temperature is reached at the cold spot, the furnace is removed from over the bell and the charge is cooled, still under protective atmosphere, by forced circulation of the atmosphere gas, which is cooled by heat exchangers. According to The Book of Steel, Lavoisier Publishing, ©1996, p. 1272' for HNX furnaces the duration of cooling is about 28 h, and for a hydrogen furnace 22.5h, resulting in an average cooling of less than l°C/min. This cooling rate can be accelerated by using water cooled bells, or by removing the bell prematurely, but oxidation of the coil should be prevented.

In an embodiment of the invention the hot-rolling finishing temperature is at least Ar3+20°C. This ensures that the strip is finished while still being austenitic and austenitic finishing hot-rolling is important to create a fine equiaxed structure of the hot rolled strip.

The invention is now further explained by means of the following, non- limitative example.

A steel composition as given in table 1 was used. This is a composition that can be used for producing cold-rolled and continuously annealed DP500.

This chemical composition results in 0.0023% N (i.e. 23 ppm) available for precipitation with aluminium as AIN if Ti, Nb, V, Zr and B is taken into account and 0.0026 (i.e. 26 ppm) if Ti and B are taken into account.

From this composition hot-rolled strip was produced at a thickness of 4 mm, finished at about 870°C, and coiled at a temperature of 580°C. These hot-rolled strips were subsequently pickled and cold rolled to 1 mm (i.e. 75% cold reduction) or to 1.32 mm (67% CR). The microstructure after hot rolling consisted of ferrite and pearlite.

The cold-rolled samples were subsequently batch annealed by two step heating : 200°C/h to 500°C, holding at 500°C for 5 hours, then heating to 680°C at a heating rate of 20°C/h, soaking temperature is 700°C and soaking time is 15 hours. These values relate to the hot spot and the holding at 500°C is intended to limit or prevent any differences between the microstructure of the hot spot and that of the cold spot. For the cold spot the corresponding heating rate is 23°C/h. This is to simulate H ₂ -annealing. For HNX annealing : the corresponding cold spot heating rate is about 15°C/h. The cooling rate after BA is known to be below l°C/min in the conventional process. The continuous intercritical annealing was performed as follows: heating to 720°C at 29°C/s, then slow heating to 790°C at l . l°C/s, soaking for 44 seconds, then slow cooling to 740°C at 1.9°C/s, then fast cooling to 470°C at 61°C/s, holding at the intermediate temperature of 470°C for 48 seconds and followed by hot dip galvanising, after galvanising cooling to 25°C at 33°C/s. The microstructure of the continuously annealed steel is presented in figure 2 (703C without BA) and 3 (703C-BA with BA), and the crystallographic texture in Figure 4A (Orientation Distribution Function (ODF) section at Φ ₂ =45° where φΐ is on the horizontal axis and φ on the vertical axis). Figure 4 reveals a well developed γ- fibre (at φ1 =0°, and φ=55°) for the samples B and C which is indicated as 703C- BA in comparison to 703C A. The intensities on the continuous γ-fibre for 703C-BA are at least 5 (times random) with a maximum above 7 (figure 4B) or 9 (figure 4C), whereas for 703C (figure 4A) the continuous value is at least 3 with a maximum of 5. The disadvantageous (lOO)-intensity, representing the 'cube on face' orientation is smaller for 703C-BA (between 1 and 2) than for 703C (between 2 and 3). The latter sample was not batch-annealed in between cold rolling and continuous annealing and is therefore representative of the known (prior art) process. The difference in texture in Figure 4 is reflected in the r- values of both specimens (see table 2 and figure 5). Figure 5 shows a number of specific crystallographic orientations where E, F and T* are favourable textures for deep-drawing, and e.g. H and G are unfavourable. The BA-samples have significantly larger amounts of E, F and T* than the 703C-sample, whereas the opposite is the case for H and G. The effect of the AIN precipitation is retardation of the recrystallisation reaction while the nucleation process is more selective towards { 111} grains, so that the recrystallised structure has larger (pancake- shaped) grains resulting in a stronger { l l l}-texture which is favourable for good r-values. The subsequent continuous annealing retains this favourable texture by annealing in the intercritical range and strengthens the material by rapid cooling after annealing.

Figure 4 shows the ODF with isolines of 1-2-3-4-etc. This reveals that the 703C-BA has a more favourable texture than the 703C, particularly visible in the top half of the . Also in Figure 4C s 703C- BA (87% CR). This shows a very nice texture with strong intensities on the γ- fibre, and very low intensities on the rotated cube and cube component (at φ1 =0°, and Φ=0° and at φ2=45°, and φ1 =45°, Φ=0° respectively). Although the r-value in the rolling direction (r-value in figure 5 (right) at alpha(RD)=0°) as such is lower (see table 2), the Ar-value, which is a measure for planar anisotropy, is better because of the higher r45 (r-value in figure 5 (right) at alpha(RD)=45°) for 703C-BA (87% CR).

Table 2: Mechanical properties of 703C-BA in comparison to 703C.

The tensile properties and the r-value were determined in the rolling direction of the sheet in accordance with EN10002 and ISO10113 on a tensile specimen of 50mm gauge length.