EP3790999A1 - Variably rolled steel strip, sheet or blank and production method therefor - Google Patents

Abstract

The invention relates to a variably rolled steel strip, sheet or blank, having at least one portion having a high thickness and at least one portion having a low thickness, wherein the variation in thickness between the high thickness and the low thickness is at least 35%, and wherein in the portion with high thickness the yield strength is 350 MPa or higher and the tensile elongation A50 (in %) and sheet thickness t (in mm) satisfy the equation of A50 / t^0.2 ≥ 14, and in the portion with low thickness the yield strength is 450 MPa or higher and the tensile elongation A50 (in %) and sheet thickness t (in mm) satisfy the equation A50 / t^0.2 ≥ 17.

Description

VARIABLY ROLLED STEEL STRIP, SHEET OR BLANK

AND PRODUCTION METHOD THEREFOR

The invention relates to a variably rolled steel strip, sheet or blank and production method therefor.

A variably rolled steel strip, sheet or blank is used for the production of automotive chassis parts having a variable thickness. The batch-annealed variably rolled steel should have an excellent balance between recrystallization on the one hand and strength and formability on the other hand at low cold-rolling reductions in the range of 30% to 60%. A method of manufacturing said steel strip, sheet or blank is provided.

Ever increasing pressure on the automotive industry to reduce greenhouse gas emissions and to improve fuel economy of their vehicle fleet has led to the use of higher strength steels in order to reduce vehicle weight via down-gauging. Frequently, this is done by using steel with increased strength and reduced, albeit constant, thickness gauge. However, a step further to save component weight via down-gauging at minimum penalty in terms of functionality is to use steel strip or sheet with variable thickness to produce tailor- rolled blanks (TRBs). This in turn, enables a variation in thickness of the final formed component produced out of the TRB that in the end has an optimized balance between weight reduction and functional optimization. The TRB manufacturing route is particularly suited to produce automotive chassis parts with reduced component weight.

A common industrial practice to produce TRBs for automotive chassis parts is to make use of flexible rolling of hot-rolled steel strip or sheet in which the local thickness of the final steel strip or sheet can be controlled via adaption of the roll gap during cold rolling. A common degree of cold-rolling reduction for the incoming hot-rolled and pickled steel strip or sheet for the flexible rolling process is within a range of 20% to 60%. Subsequently, the tailor-rolled steel strip or sheet is batch annealed to promote recrystallization of the work-hardened microstructure and to recover formability after cold rolling, preferably at minimum penalty in terms of loss in strength due to exposure to the thermal cycle of the batch annealing process. The top temperature for the batch annealing in the aforementioned industrial approach is commonly in the ferrite-phase temperature region, i.e., the top temperature during batch annealing is below the Ac1 transformation point.

From the batch-annealed steel strip or sheet, TRBs will be obtained via a blanking process. Typically, a TRB suitable for automotive chassis components will have a minimum thickness of 1 mm (to ensure a certain degree of stiffness) and a variation in sheet thickness of 35% or higher. Current TRB manufacturing practice is to use hot-rolled HSLA steels as input for the production of TRBs for automotive chassis components and to use HSLA grades with increased strength in hot-rolled condition to achieve increased strength in the final TRBs. The increase in strength of hot-rolled HSLA steels is commonly achieved by a combination of increased grain refinement and increased precipitation strengthening, which comes from the use of increased levels of one or more micro-alloying elements, including Niobium (Nb), Titanium (Ti), and Vanadium (V). Since the solubility products of Nb-based and Ti-based precipitates are considerably lower than those of V-based precipitates, Nb and Ti are considerably more effective as strengthening agents than V. For this reason hot-rolled HSLA steels are commonly alloyed with Nb and/or Ti, and increased strength for these grades is achieved by using increased levels of Nb and/or Ti.

The increased levels of micro-alloying elements in hot-rolled HSLA steel will lead to a higher density of precipitates in the microstructure. This is excellent to achieve increased (precipitation) strength in a hot-rolled product, but the increased precipitate density can seriously hinder recrystallization during batch annealing in the TRB manufacturing process to get sufficient formability.

Reason is that precipitates hinder the migration of grain boundaries during the growth stage of recrystallization. This delay in recrystallization can provide increased opportunity for recovery, leading to a loss of dislocation density and a decrease in stored energy to drive forward recrystallization. As a result, a higher top temperature and/or extended holding time at top temperature may be required during batch annealing to achieve the desired degree of recrystallization and formability. However, this increased top temperature and/or holding time can be at the expense of precipitation strengthening due to increased precipitate coarsening during batch annealing.

For all three aforementioned micro-alloying elements, precipitation strengthening comes from the formation of carbide, nitride, and/or carbo-nitride precipitates. Each type of precipitate will have its own solubility product, meaning that some precipitates will form, coarsen, or dissolve earlier than others. The solubility of V-based precipitates in ferrite is known to be considerably higher than that of Nb- and Ti-based precipitates, making V precipitates considerably more prone to coarsening during - for instance - batch annealing to achieve recrystallization (³650 °C). This increased coarsening behaviour of V precipitates promotes increased loss of precipitation strengthening with batch annealing. For that reason, V is considered as an inappropriate micro-alloying element for steels that are subjected to batch annealing to achieve full or substantial recrystallization and high strength. In particular, when increased top temperature and/or extended holding time at top temperature are required to obtain sufficient recrystallization and formability for the production of TRB sheets with increased micro-alloy content and high strength. Instead, Nb and Ti are generally considered to be the most suitable micro-alloying elements in this case. In this context, it is common practice to use hot-rolled Nb- or (Nb,Ti)- based HSLA steels for the production of TRB sheets via flexible cold rolling and high- temperature batch annealing (³650 °C).

A challenge to produce high-strength TRB sheets via flexible cold rolling and subsequent batch annealing with hot-rolled (Nb,Ti)-HSLA steels with high precipitate density already in hot-rolled condition, is to balance sufficient recrystallization of the work-hardened ferrite microstructure and hence sufficient formability on the one hand and to minimise the loss in precipitation strengthening of the ferrite due to precipitate coarsening on the other. These two features mutually can contradict as recrystallization may require a high top temperature during batch annealing to promote nucleation and growth of new ferrite grains, whereas precipitation strengthening is likely to benefit from a relatively low top temperature during batch annealing to prevent precipitate coarsening.

An optimum balance between sufficient recrystallization and formability on the one hand and (retaining) sufficient strength on the other with batch annealing is needed to produce a final high-strength TRB sheet with specific targets on performance and component weight.

To produce high-strength TRB sheets, commonly a more heavily alloyed (Nb,Ti)-based HSLA steel is used as input for the flexible cold rolling and batch annealing process. However, the problem with this approach is that the high precipitate density inherited from the hot-rolling process interferes with recrystallization and lead to insufficient recrystallization and too low formability at low rolling reductions, e.g., 30% to 40%, after batch annealing. Furthermore, insufficient recrystallization can lead to regions with elongated grains and locally hardened microstructural regions, which can promote delamination or splitting when the steel is subjected to shearing operations during the production of TRB blanks and/or the manufacturing of final automotive chassis components.

To achieve sufficient recrystallization and formability across the full range flexible cold- rolling reduction for high-strength TRB sheets based on a heavily alloyed (Nb,Ti)-based HSLA, high batch annealing top temperature and potentially longer dwell times at top temperature are required. However, such changes may either be in conflict with maximum heating capability of the furnace and/or by the fact that this reduces strength too much as precipitates will increasingly coarsen with higher top temperatures. This can lead to an unacceptable loss in strength and not meeting TRB product targets set for minimum yield strength and maximum component weight.

It is an object of the invention to provide a variably rolled steel strip, sheet or blank suitable for manufacturing TRBs, which after variable cold-rolling with reductions of 30% or higher, followed by batch annealing, has a high yield strength and a high tensile elongation (A50 in %) relative to the sheet thickness. It is a further object of the invention to provide a method for producing such a steel strip having a variable thickness.

According to a first aspect of the invention one or more of these objects is reached by a variably rolled steel strip, sheet or blank, consisting of the following elements (in wt%):

0.05 - 0.20 C

0.10 - 0.70 Si

0.80 - 2.50 Mn

0.01 - 0.10 Al

0.07 - 0.25 Ti

0.10 - 0.35 V

0.05 - 0.40 Mo

optionally 0.02 - 0.10 Nb

optionally 0.01 - 0.80 Cr

at most 0.06 P

at most 0.01 S

at most 0.01 N

at most 0.005 Ca

the balance consisting of inevitable impurities and Fe,

wherein the strip, sheet or blank has at least one portion having a high thickness and at least one portion having a low thickness, wherein the variation in thickness between the high thickness and the low thickness is at least 35%, and wherein in the portion with high thickness the yield strength is 350 MPa or higher and the tensile elongation A50 (in %) and sheet thickness t (in mm) satisfy the equation of A50 / t^0.2 ³ 14, and in the portion with low thickness the yield strength is 450 MPa or higher and the tensile elongation A50 (in %) and sheet thickness t (in mm) satisfy the equation A50 / t^0.2 ³ 17.

The invention provides a batch-annealed variably rolled steel strip or sheet with an excellent balance between recrystallization behaviour, strength, and ductility and solves the problem of too little recrystallization in a cold-rolling reduction range of 30% to 60% and too high a loss in strength in these reduction ranges. The excellent balance between recrystallization and strength is obtained with an essential and substantial addition of V of at least 0.1 wt.% to a steel composition that further contains as precipitating elements Ti, Mo, and optionally Nb.

The V in solid solution during the early stages of batch annealing will form - predominantly - V-based carbide precipitates in addition to V-based nitride and/or carbo-nitride precipitates. The V precipitation during batch annealing is accelerated by the presence of dislocations induced by the cold-rolling operation as the increased dislocation density will increase the diffusional rate of elements. At the same time, the dislocations will act as preferential nucleation sites for precipitation during the early stages of batch annealing. This in turn will suppress annihilation of dislocations and hence reduce the degree of recovery as the newly formed V- based precipitates will pin dislocations and hinder their movement. The reduced degree of recovery will increase the driving force for the nucleation of recrystallized ferrite grains and increase the density of ferrite nuclei, leading to more impingement during recrystallization and ultimately stimulate grain refinement of the final microstructure. Subsequent growth of the newly formed ferrite nuclei / grains is promoted by ensuring a sufficiently high batch-annealing top temperature. This top temperature should be equal to or above the dissolution temperature of V-based precipitates in order to dissolve the V-based precipitates and to lift their pinning force, allowing migration of grain boundaries. Inventors found that in the present invention a top temperature of 700 °C or higher is sufficient to promote substantial recrystallization.

Such a high top temperature during batch annealing will impair precipitation strengthening as precipitates in general will be prone to substantial coarsening at those temperatures. The loss in precipitation strengthening at these temperatures can be suppressed by using Ti and Mo, optionally in combination with Nb. The element Mo is known to combine with Ti and Nb to form composite carbide and/or carbo-nitride precipitates that have increased thermal stability and hence an increased resistance to coarsening. In the present invention, the V brought in solid solution in the ferrite matrix during batch annealing after dissolution may again partially precipitate in the ferrite matrix upon cooling down, contributing to some extent to precipitate strengthening.

However, more important is that the V addition provides grain refinement by stimulating recrystallization. This ensures that loss in precipitation strengthening due to the use of an elevated batch annealing temperature (³700 °C) is mitigated by grain refinement and hence increased grain-boundary or so-called Hall-Petch strengthening.

Evidence of recrystallization can be determined either via analysis of the microstructure by means of light-optical microscopy (LOM) or electron backscatter diffraction (EBSD). These techniques have been employed to determine the fraction recrystallized ferrite of batch- annealed steels and to determine the average grain size of the recrystallized ferrite in the microstructure of batch-annealed steel sheets. The followed procedures are disclosed in Example 1 .

An alternative method to assess if substantial recrystallization of the microstructure has been achieved after batch annealing is to record the evolution of the yield strength after batch annealing for a range of cold rolling reductions. At low cold-rolling reductions, from 0% upwards, the Rp0.2 after batch annealing will increase due to work hardening that is not (significantly) compromised as the amount of dislocations does not provide sufficient driving energy for recrystallization. However, depending on the batch annealing parameters, at some point the Rp0.2 after batch annealing will again start to decrease with increasing cold-rolling reduction as locally recrystallization will start to occur, leading to a loss in dislocation hardening. Depending on the batch annealing parameters again, at some point the Rp0.2 may start to remain stable or increase again after a region in which it decreased with increasing cold-rolling reduction. This is the region of interest for the present invention and is the region in which recrystallization becomes dominant. The increase in Rp0.2 with increasing cold-rolling reduction in this region is the result of increasing grain refinement, which results from recrystallization and the increasing amount of dislocations and hence increasing amount of potential nuclei that are present to form new, recrystallized ferrite grains. Hence, increased impingement and grain refinement will lead to an increase in Rp0.2 due to grain refinement. In this context, inventors found that if the yield strength (Rp0.2) is constant or - preferably - increases with an increase in cold-rolling reduction (CR%) in the range of 30 to 60% or higher sufficient recrystallization is achieved to have sufficient formability for forming operations during manufacturing and to avoid or substantially suppress the issue of delamination or splitting as a result of shearing operations, including cutting or punching, when the steel is used to manufacture (automotive) components.

In case of the present invention, a preferred objective is that the yield strength (Rp0.2) of the inventive steel after cold rolling with a reduction of 30% or higher and subsequent batch annealing is constant or preferably increases with increasing cold-rolling reduction prior to batch annealing. A suitable measure according to the inventors to assess this is to fit the evolution of Rp0.2 (in longitudinal direction) of the batch-annealed steels as a function of the cold-rolling reduction (CR%) from 30 to 60% reduction according to

with a and b as fitting parameters. Preferred objective of the present invention is that fitting parameter a is 0 or higher.

The role of the alloying elements for the present invention is as follows.

Carbon (C) is added to form carbide and/or carbo-nitride precipitates together with V, Ti, Mo and - in the present invention - optionally with Nb. The amount of C depends on the amount of V, Ti, Nb and/or Mo used and should be at least 0.05 wt.%. However, the maximum content is 0.20 wt.% to prevent excessive segregation and to prevent a too high fraction of cementite and/or pearlite. For the present invention, the fraction of pearlite and/or cementite in the microstructure of the batch-annealed steel is preferably at most 10%, or more preferably at most 5%, or most preferably at most 3%. Segregation and an excessive amount of pearlite and/or cementite are considered to be deleterious for hole-expansion capacity. A more preferable C content range for the present invention is between 0.06 and 0.17 wt.%, or most preferably between 0.07 and 0.14 wt.%.

Silicon (Si) provides significant solid-solution strengthening, which is desirable as its contribution to strength is not compromised by the thermal cycle of the batch annealing process. Furthermore, it retards the formation of cementite and pearlite, thus suppressing the formation of coarse carbides, which can impair hole-expansion capacity. However, too high Si will lead to an undesired increase in rolling loads and may lead to surface issues and reduced fatigue properties. For these reasons, the Si content is at least 0.10 wt.% may not exceed 0.70 wt.%. A more preferable Si content range for the present invention is between 0.20 and 0.60 wt.%, or most preferably between 0.30 and 0.60 wt.%.

Manganese (Mn) provides solid-solution strengthening, which is desirable as its contribution is not compromised by the thermal cycle of the batch-annealing process. Therefore, Mn content should be at least 0.8 wt.%. However, a too high Mn content may lead to excessive segregation, which can impair hole-expansion capacity and promote delamination or splitting during shearing operations. Furthermore, a too high Mn content will suppress the ferritic transformation temperature and promote hardenability, leading to hard carbon-rich phase constituents in the intermediate hot-rolled feedstock (e.g., martensite and retained-austenite) which in turn can lead to unacceptable high strength and too high rolling loads for the cold mill. Hence, a suitable maximum Mn content for the present invention is 2.5 wt.%. A more preferable Mn content range for the present invention is between 0.9 and 2.30 wt.%, or most preferably between 1 .20 and 2.00 wt.%.

Phosphorus (P) provides solid-solution strengthening, However, at high levels, P segregation will promote delamination or splitting during shearing operations and impair hole- expansion capacity. Therefore, the P content should be at most 0.06 wt.%, or preferably at most 0.04 wt.%, and more preferably at most 0.02 wt.%.

Sulphur (S) is known to be detrimental for formability, in particular for hole-expansion capacity. Therefore, the S content should be at most 0.01 wt.%, or preferably at most 0.005 wt.%, or more preferably at most 0.003 wt.%.

Aluminium (Al) is added as a deoxidizer. A suitable minimum Al content is 0.01 wt.%. However, too high Al can be deleterious as it forms AIN particles during solidification of the molten steel, which can provoke surface issues during casting. Furthermore, a too high Al content can impair hole-expansion capacity as it may lead to a too high fraction of Al_xO_y inclusions in the steel matrix, which can promote internal fractures upon shearing the steel. Hence, the Al content should be at most 0.10 wt.%. A suitable Al content range for the present invention is between 0.01 and 0.10 wt.%, or more preferably between 0.02 and 0.09 wt.%, and most preferably between 0.04 and 0.08 wt.%.

The Nitrogen (N) content should be low, i.e., at most 0.01 wt.%. Too high N content, in particular when too much N is free and in solid solution in the ferrite matrix, is deleterious for formability in general. Furthermore, too high N content in the presence of Ti can lead to an excessive amount of large cuboid TiN particles, which impair formability in general and hole- expansion capacity in particular. On the other hand, N can be beneficial to promote nitride and/or carbo-nitride precipitates, which in general are more thermally stable than carbide precipitates. In this context, N can be beneficial to suppress coarsening during the thermal cycle of the batch annealing process. A more preferable range for N content for the present invention is at most 0.008 wt.%, or most preferably between 0.002 and 0.007 wt.%.

Titanium (Ti) is used in the present invention to realise precipitation strengthening and to some degree grain refinement. As such, Ti is an essential element in the alloy composition of the present invention to achieve a desired strength level for the steel strip or sheet after batch annealing. A suitable minimum Ti content is 0.07 wt.% or more preferably 0.08 wt.% or even 0.10 wt.%. A too high Ti content can lead to undesired segregation-related phenomena, to too high rolling loads during hot rolling and subsequent cold rolling, and to too lowformability due to insufficient recrystallization achieved after batch annealing. This insufficient recrystallization of the steel after batch annealing may lead to issues with splitting or delamination resulting from shearing the steel during manufacturing operations. Hence, a suitable maximum Ti content is 0.25 wt.%. A more preferable Ti maximum content for the present invention is 0.22 wt.%, or most preferably 0.20 wt.%.

Niobium (Nb) is used in the present invention to realise a certain degree of precipitation strengthening as well as to achieve grain refinement and hence strength via the Hall-Petch effect. As the degree of precipitation hardening is relatively limited compared to that of Ti, the use of Nb is considered as optional for the present invention. However, when used, a suitable minimum Nb content is 0.02 wt.%, or more preferably 0.03 wt.%, and a suitable maximum Nb content is 0.10 wt.%, more preferably 0.09 wt.%, and most preferably 0.08 wt.%.

Molybdenum (Mo) is known to be a carbide-forming element and can form together with Ti, V and/or Nb composite carbide and/or carbo-nitride precipitates. These composite precipitates comprising Mo, are reported to be more thermally stable than their counterparts without Mo and hence more resistant to coarsening during exposure to a thermal cycle at temperatures above 600 °C. Hence, Mo is beneficial to suppress precipitate coarsening during batch annealing at top temperatures above 600 °C and to reduce the loss in precipitation strengthening due to batch annealing above 600 °C. The desired strength level of the final batch annealed steel in the end will determine to what extent Mo, which is an expensive alloy element, is required. For the present invention, a suitable Mo content is at least 0.05 wt.% and at most 0.40 wt.%. A more preferable Mo content range for the present invention is between 0.08 and 0.35 wt.%, or most preferably between 0.10 and 0.30 wt.%.

Vanadium (V) is an essential element for the present invention as it acts as an agent to stimulate recrystallization during batch annealing, providing grain refinement, and provides precipitation strengthening. The former - i.e., the aspect of recrystallization - is achieved by the formation of V-based carbide precipitates during the initial stages of batch annealing which nucleate on dislocations and hence pin dislocations, reducing their mobility and suppressing recovery. As a consequence, the driving force for the onset of recrystallization is increased as the pool of surviving dislocations at the start of recrystallization increases. By using a top temperature during batch annealing that ensures that sufficient V-based precipitates again start to dissolve later on during the batch annealing, the increased driving force for recrystallization is released and the growth of new, recrystallized ferrite grains is stimulated.

As mentioned, a substantial V addition stimulates recrystallization already at low cold- rolling reductions as it suppresses annihilation of dislocations and thus maintains an increased level of stored energy as driving force for recrystallization. At the same time, this leads to an increased density of nuclei for recrystallization and hence an increased degree of impingement, promoting grain refinement of the final microstructure. This grain refining effect will bring increased strength. This will mitigate to some extent loss in precipitation strengthening with batch annealing at a top temperature of 700 °C or higher.

V in solid solution in the ferrite matrix at elevated temperature during the batch annealing cycle may again precipitate later on during the final stages of the batch annealing cycle, contributing again to precipitation strengthening of the ferrite microstructure of the final steel strip or sheet after batch annealing.

The V in the present invention is not only believed to be beneficial to achieve strength via grain refinement and direct precipitation strengthening via the formation of freshly V-based precipitates during batch annealing as mentioned above, but also indirectly by suppressing the coarsening kinetics of Ti-based precipitates in the steel matrix during batch annealing. The latter is believed to be the result of a relatively high V content in solid solution in the ferrite matrix, which will reduce Ti solubility and hence suppress coarsening kinetics of Ti-based precipitates. Furthermore, part of the V that will precipitate during batch annealing will correspond with co-precipitation, i.e., V precipitating on existing Ti-based precipitates. This can promote a V-rich shell surrounding the Ti-based precipitate, which will acts as a barrier, suppressing coarsening of Ti-based precipitates covered with a V-rich shell.

The amount of V should be sufficiently high enough to promote a sufficient degree of recrystallization. Inventors found that a suitable minimum V content is 0.10 wt.%, or preferably 0.12 wt.%, and more preferably 0.13 wt.%. At the same time, the amount of V and the corresponding amount of VC precipitates should correspond with a dissolution temperature for VC precipitates in ferrite that is within the industrial capacity of the batch annealing furnace used. As a rule of thumb, inventors used the equation below - based on the Arrhenius relationship - for an estimation of the dissolution temperature (T_dis in °C) for VC precipitates in ferrite with the assumption that all V ties up with C to form VC precipitates with a 1 :1 atomic ratio in which A and B are constants with values of 5500 and 3.39 K^-1, respectively, and with [V] in wt.%. The value for T_dis should be in line with the heating capacity of the batch annealing furnace in order to ensure that VC can be sufficiently dissolved during the batch annealing cycle to promote substantial recrystallization. A suitable maximum V content is 0.35 wt.%, or more preferably 0.30 wt.%, or most preferably 0.25 wt.%.

Chromium (Cr) is an optional element for the present invention and can be used to promote the formation of ferrite, in particular when elevated levels of Mo and/or Mn are used that can suppress the formation of ferrite. If used, a suitable Cr content is 0.01 - 0.80 wt.%, or preferably 0.01 - 0.60 wt.%, or more preferably 0.01 - 0.40 wt.%.

Calcium is an optional element for the present invention and may be used to modify MnS-type of inclusions to improve formability and/or to modify Al_xO_y-type of inclusions to reduce the risk of clogging and to improve cast ability of the steel during steel making. However, a too high Ca content can lead to excessive wear of the refractory lining in the installations of the steel-making plant In case a Calcium treatment is used during steel making for inclusion control, a suitable maximum Ca content is 50 ppm, or more preferable maximum 35 ppm. In case of a Calcium treatment, a suitable minimum Ca content in the steel is 20 ppm. In the absence of a Calcium treatment during the steel making process, the Ca content in the steel is at most 20 ppm, or preferably at most 10 ppm, or most preferably at most 5 ppm.

According to a preferred embodiment the variably rolled steel strip, sheet or blank contains Ti, and Mo and optionally Nb represented by weight percentage (wt.%) satisfying the equation of

or preferably

With the amounts of Nb, Ti and Mo satisfying these equations, a suitable balance between these elements is provided.

An even more preferred embodiment is provided when the steel strip, sheet or blank contains C, N, Ti, Mo, V and optionally Nb represented by weight percentage (wt.%) satisfying the equation of

with and

Using such a balance between these elements provides an optimal balance between the elements that are essential for the invention.

Preferably the steel has a precipitation strengthened ferrite microstructure, consisting of at least 50% recrystallized ferrite at ¼ depth of the portion with high thickness, and which optionally further contains cementite and/or pearlite, and wherein the precipitates in said microstructure consist of Ti, V, Mo and optionally Nb. According to a preferred embodiment, the steel has a microstructure that consists of at least 60% recrystallized ferrite at ¼ depth in the portion with high thickness, preferably at least 70% recrystallized ferrite at ¼ depth in the portion with high thickness, more preferably at least 80% recrystallized ferrite at ¼ depth in the portion with high thickness, most preferably 90% recrystallized ferrite at ¼ depth in the portion with high thickness. The amount of recrystallized ferrite is an indication of the stability of the microstructure after batch annealing.

Preferably, the steel in the portion with high thickness has a yield strength of 400 MPa or higher, preferably of 450 MPa or higher, or more preferably of 500 MPa or higher, or most preferably 550 MPa or higher, and/or wherein the steel in the portion with low thickness has a yield strength equal to or higher than that of the portion with high thickness, and preferably has a yield strength of 500 MPa or higher, more preferably of 550 MPa or higher. The higher the yield strength, the better the performance of the TRBs in automotive structures.

Furthermore it is preferable that the steel in the portion with high thickness satisfies the equation A50 / t^0.2 ³ 16, preferably A50 / t^0.2 ³ 18, most preferably A50 / t^0.2 ³ 20, and wherein the steel in the portion with low thickness satisfies the equation: A50 / t^0.2 ³ 18, preferably A50 / t^0.2 ³ 20, most preferably A50 / t^0.2 ³ 22, with A50 / t^0.2 in the portion with low thickness equal to or higher than A50 / t^0.2 in the portion with high thickness. The value of equation A50 / t^0.2 is an indication of the tensile elongation A50 (in%) relative to the sheet thickness t (in mm) in the portion with high thickness and low thickness, respectively. The higher this value, the better the crashworthiness of the TRB in a car.

According to a second aspect of the invention one or more objects of the invention are reached using a method for producing a steel strip having a variable thickness, comprising the steps of:

• casting a slab having the composition according to the first aspect of the invention,

• reheating the solidified slab to a temperature between 1 150 and 1300° C,

• finishing the hot rolling at a finish hot rolling temperature of Ar3 transformation point or higher, • cooling the hot-rolled steel strip to the coiling temperature at an average cooling rate of 10 to 150° C/s,

• coiling the hot-rolled steel strip in the temperature range between 450 and 580° C,

• cold-rolling the strip as variable rolling, such that a cold-rolling reduction between 30% and 60% is performed and a variation of strip thickness of at least 35% is obtained,

• batch annealing the steel strip.

The role of the processing steps for the present invention is as follows.

The reheating temperature of the slabs in the furnaces of the hot-strip mill prior to rolling should be high enough h to ensure that practically all carbide and carbo-nitride precipitates containing Ti and V, and optionally Nb, have dissolved in the steel matrix. This is required to maximise the amount of Ti and V, and optionally Nb, in solid solution prior to hot rolling and further down-stream processing. The optimum reheating temperature depends on the amount of Ti and V, and optionally Nb. However, inventors found that a suitable range of the reheating temperature is between 1 150 and 1300 °C.

Finish rolling in the hot-strip mill should be done at Ar3 transformation point or higher in order to finish the hot rolling sequence in the austenite phase region prior to actively cooling down the steel strip or sheet to enforce austenite-to-ferrite phase transformation.

The average cooling rate on the run-out-table of the hot-strip mill to cool the steel strip or sheet just after finish rolling should be in the range of 10 to 150 °C/s.

The temperature to coil the steel strip or sheet in the hot-strip mill should be low enough to suppress precipitation in general, but in particular that of V. At the same time, the coiling temperature should not be too low as this leads to too much transformation hardening. The microstructure of the intermediate hot-rolled feedstock in the present invention is preferably ferrite and/or bainitic in nature, preferably without the presence of a substantial amount of martensite. Inventors found that a suitable coiling temperature of the steel strip or sheet in the hot-strip mill is between 450 and 580 °C.

When these process steps are performed, cold-rolling of the strip as variable rolling can be performed with a cold-rolling reduction between 30% and 60% such that a variation of strip thickness of at least 35% is obtained. After annealing of the variably rolled strip, a strip according to the first aspect of the invention is obtained.

Preferably the hot-rolled steel strip is hot-rolled with a finish rolling temperature of 870 °C or higher, preferably with a finish rolling temperature of 900 °C or higher, more preferably with a finish rolling temperature of 940 °C or higher, and most preferably with a finish rolling temperature of 980 °C or higher. To reduce the rolling loads and suppress strain-induced precipitation of micro-alloy elements during the final rolling passes, the temperature set for finish rolling may be chosen higher. Another benefit of a higher finish rolling temperature is its beneficial influence on texture development and hence mechanical properties and isotropy. Hence for the present invention, the finish rolling temperature should preferably be 900 °C or higher, or more preferably 940 °C or higher, or most preferably 980 °C or higher.

According to a preferred embodiment the hot-rolled steel strip after finish rolling is cooled to the coiling temperature with an average cooling rate of 40 to 100° C/s.

Preferably the hot-rolled steel strip is coiled in the temperature range between 480 and 560° C, or more preferably between 500 and 540 °C to provide the preferred microstructure of the intermediate feedstock.

Preferably, the hot-rolled steel strip is cold-rolled and batch annealed

• for at least 8 hours at a top temperature of 740 °C or higher, or

• for at least 10 hours at a top temperature of 720 °C or higher, or

• for at least 14 hours at a top temperature of 700 °C or higher.

This batch annealing process results in a minimal 50% recrystallization of the cold rolled strip, resulting in a precipitation strengthened ferrite microstructure with an adequate balance between strength and formability.

The invention will be elucidated by means of the following, non-limitative examples, referring to the attached figure.

Figure 1 shows the time-temperature curve of the batch annealing cycle used in the examples.

EXAMPLE 1 :

(1) Alloys, process conditions, testing and microstructural analyses procedures

Examples are performed using laboratory cast ingots.

Steels 1A to 1 H having chemical compositions shown in Table 1.1 were hot rolled after reheating the ingots to 1250 °C for 45 minutes to ensure optimum dissolution of carbide and carbo-nitride precipitates, which, depending on alloy composition, comprise Mo, Nb, Ti, and V. The hot-rolled steel sheets were rolled in 5 passes from a thickness of 35 to 3.5±0.5 mm with an exit temperature for the final rolling pass in the range of circa 900 to 1000 °C. After the final rolling pass, the steel sheet was transferred to a run-out-table (ROT) and cooled down from a start ROT temperature in between 850 to 900 °C with an average cooling rate of circa 40 to 50 °C/s to an exit ROT temperature around 600 or 540 °C. Next, the hot-rolled steel sheet was transferred to a furnace to replicate slow coil cooling from a start temperature of 600 or 540 °C to ambient temperature. The hot-rolled steel sheets were pickled prior to tensile testing of the hot-rolled steel sheet or further processing in terms of cold rolling and subsequent batch annealing followed by tensile testing of cold-rolled and batch-annealed steel sheets. Batch annealing was done on hot-rolled steel sheets with no cold-rolling reduction (CR% equals 0%) and on cold-rolled steel sheets after a 10, 20, 30, 40, 50, or 60% cold-rolling reduction post hot rolling. To suppress decarburisation during batch annealing, plates were wrapped in stainless steel foil and a protective H₂ atmosphere was used in the batch anneal furnace. A number of different settings were used to carry out batch annealing (BA) simulations. These included for all steels 1 A to 1 H (T/t = top temperature in °C / holding time in hours at top temperature): BA-675/3, BA-700/3, BA-740/3, BA-700/10, and BA-740/10. In addition, for steel 1 G also the following BA cycles were carried out; BA-740/7, BA-740/8, and BA-740/9. Details about a number of batch annealing curves used are shown in Table 1.2, showing the time-Temperature (t-T) profiles for the following batch annealing simulations: BA- 675/3, BA-700/3, BA-740/3, BA-700/10, and BA-740/10. Although the top temperature (Ttop) and holding time at top temperature (thoid) for the batch annealing cycle are variable, the gradients for the heating and cooling stages in the batch annealing curves were held fixed in all simulations. The time-temperature curve of the batch annealing cycle used for the steels in this example is shown in Figure 1.

The tensile properties were in all cases, i.e., for hot-rolled as well as batch-annealed steel sheets, measured parallel to rolling direction by means of taking out A50 test pieces and applying a tensile load to the test pieces according to EN 1002-1/ISO 6892-1 (Rp0.2 is the 0.2% offset proof or yield strength; Rm is the ultimate tensile strength; Ag is the uniform elongation; A50 is the total tensile elongation).

To determine the fraction re-crystallized ferrite and the grain size of this re-crystallized fraction (after deformation and annealing) the microstructures were characterised with Electron Back Scatter Diffraction (EBSD). To this purpose the following procedures were followed with respect to sample preparation, EBSD data collection and EBSD data evaluation.

The EBSD measurements were conducted on cross sections parallel to the rolling direction (RD-ND plane) mounted in a conductive resin and mechanically polished to 1 mm. To obtain a fully deformation free surface, the final polishing step was conducted with colloidal silica (OPS).

The Scanning Electron Microscope (SEM) used for the EBSD measurements is a Zeiss Ultra 55 machine equipped with a Field Emission Gun (FEG-SEM) and an EDAX PEGASUS XM 4 HIKARI EBSD system. EBSD scans were collected on the RD-ND plane of the sheets. The samples were placed under a 70° angle in the SEM. The acceleration voltage was 15kV with the high current option switched on. A 120 mm aperture was used and the working distance was 15 mm during scanning. To compensate for the high tilt angle of the sample, the dynamic focus correction was used during scanning. The EBSD scans were captured using the TexSEM Laboratories (TSL) software OIM (Orientation Imaging Microscopy) Data Collection version 7.2 Typically, the following data collection settings were used: Hikari camera at 6 x 6 binning combined with standard background subtraction. The scan area was in all cases located at a position of ¼ the sample thickness and care was taken to avoid as much as possible to include non-metallic inclusions in the scan area.

The EBSD scan size was in all cases 100 x 100 mm, with a step size of 0.1 mm, and a scan rate of approximately 100 frames per second. Fe(a) was used to index the Kikuchi patterns. The Hough settings used during data collections were: Binned pattern size of circa 96; theta set size of 1 ; rho fraction of circa 90; maximum peak count of 10; minimum peak count of 5; Hough type set to classic; Hough resolution set to low; butterfly convolution mask of 9 x 9; peak symmetry of 0.5; minimum peak magnitude of 10; maximum peak distance of 20.

The EBSD scans were evaluated with TSL OIM Analysis software version“8.0 x 64 [12-14- 16]”. Typically, the data sets were 90° rotated over the RD axis to get the scans in the proper orientation with respect to the measurement orientation. A standard grain dilation clean-up was performed (Grain Tolerance Angle (GTA) of 5°, a minimum grain size of 5 pixels, criterion used that a grain must contain multiple rows for a single dilation iteration clean up). Next to this a pseudo-symmetry clean-up (GTA 5, axis ang 30@1 1 1 ) was applied.

Partitions of the re-crystallized fractions were created by evaluation of the grain average misorientation maps and average IQ maps. From these created partitions, the re-crystallized fraction was determined and the grain size (Grain tolerance angle = 15°, minimum number of pixels 10, grains must contain multiple rows).

(2) Tensile properties hot-rolled and batch-annealed steel sheets

Hot-rolled steel sheets : Table 1 .3 gives the A50 tensile properties of the hot-rolled steel sheets coiled at 600 °C of steels 1A to 1 D (corresponding hot-rolled steel sheets labelled as 1 A-HR600, 1 B-HR600, 1 C-HR600, 1 D-HR600). Table 1 .4 gives the A50 tensile properties of the hot-rolled steel sheets coiled at 540 °C of steels 1A to 1 H with labelling of the corresponding hot-rolled steel sheets in a similar fashion as done in Table 1 .3.

Batch-annealed steel sheets : Tables 1 .5 and 1 .6 give the tensile properties of batch- annealed steel sheets without any intermediate cold rolling (CR% equals 0%) and corresponding with hot-rolled steel sheets coiled at 600 and 540 °C, respectively. Tables 1 .7 to 1 .13 provide the tensile properties of all batch-annealed steel sheets with intermediate cold- rolling reductions of 0 to 60% for all aforementioned batch-annealing conditions.

(3) Microstructures batch-annealed steel sheets Tables 1 .7 to 1.13 provide the fraction recrystallized ferrite (in %) and the average grain size (in mm) of the recrystallized ferrite based on EBSD measurements.

(4) Interpretation of results: control over precipitation strengthening

Table 1 .4 shows data for steels 1A to 1 D corresponding with the difference in Rp0.2 and Rm between coiling at 600 and 540 °C. The data demonstrates that lowering the coiling temperature from 600 to 540 °C leads to a reduction in strength, in particular for Rp0.2. This decrease in Rp0.2 with a decrease in coiling temperature is most pronounced for steels 1 B and 1 D with - compared with steel 1A - increased content in Ti and Mo as well as for steel 1 C with increased Ti, Mo, and V. This reduction in strength is largely attributed to a loss in precipitation strengthening as the reduction in coiling temperature reduces the kinetics needed to nucleate and form precipitates. In turn, this implies that for steels 1 B, 1 C, and 1 D a certain amount of Ti and - in particular - a substantial amount of V is not precipitated in the ferrite in the final microstructure of the hot-rolled steel sheet, but instead remains in solid solution. This Ti and V in solid solution can be allowed to precipitate in a subsequent thermal cycle, such as a batch annealing cycle, given that the top temperature (Ttop) is sufficiently high to provide the necessary kinetics to allow carbide and/or carbo-nitride precipitates to nucleate and grow.

Tables 1 .5 and 1 .6 show the tensile data of hot-rolled steel sheets coiled at 600 and 540 °C, respectively, of steels 1 A to 1 H when subjected to a batch annealing with different values for Ttop and thoid and without intermediate cold rolling.

Coiling at 600 °C and having most of the micro-alloying elements precipitated in the ferrite of the final microstructure of the hot-rolled steel sheet, will lead to a subsequent loss in strength when the hot-rolled steel sheet is batch annealed for 3 hours at a top temperature of 675 °C. The measured loss in Rp0.2 and Rm (Table 1 .5) upon batch annealing is roughly the same. The loss in strength after batch annealing can be explained by a loss in precipitation strengthening. This latter will be the result of coarsening of precipitates originating from the hot-rolling stage and the fact that no significant fraction of new precipitates could be formed during batch annealing as most micro-alloying content was consumed in precipitation during the hot-rolling stage.

In contrast, when the hot-rolled steel sheets corresponding with steels 1 A to 1 H and coiled at 540 °C are subjected to a batch annealing with a Ttop of 675 or 700 °C for 3 or 10 hours thoid, a substantial increase in Rp0.2 is measured (see Table 1 .6). This increase in Rp0.2 will be largely linked to precipitation of micro-alloying elements that remained in solid solution in the hot-rolled steel sheet due to low-temperature coiling, but have precipitated during subsequent batch annealing with a Ttop of 675 or 700 °C for 3 or 10 hours thoid. If T_top is raised above 700 °C, i.e., 740 °C in the examples shown in Table 1.6, the Rp0.2 decreases after batch annealing with thoid of 3 or 10 hours. This decrease in Rp0.2 is seen for all steels, i.e., steels 1A to 1 H, and is believed to be related to a loss in precipitation strengthening due to substantial precipitate coarsening above 700 °C.

The observations above imply that it is possible to control precipitation during batch annealing by stimulating nucleation and growth of freshly formed precipitates during batch annealing on the one hand and by promoting coarsening of precipitates on the other hand. Depending on Ttop and thoid as critical input parameters for the batch annealing cycle, the Rp0.2 of the steel sheet after batch annealing can in this way be increased or decreased compared with that of the Rp0.2 of the corresponding hot-rolled steel sheet. This control over the degree of precipitation strengthening during batch annealing may be used to control the strength of the final batch-annealed steel sheet after hot-rolling without any intermediate cold- rolling step, but can also be used to control and improve recrystallization behaviour of cold- rolled steel sheets during batch annealing, i.e., promote substantial/partial (i.e., ³50%) or - preferably - full recrystallization already at a relatively low cold-rolling reduction (e.g., CR% ³ 30%). This onset of substantial recrystallization is indicated by an increase in yield strength and tensile elongation as a function of cold-rolling reduction, e.g., CR% ³ 30%. Since thickness varies in a TRB sheet, the tensile elongation is normalised to 1 mm thickness by using the equation A50 / t^0.2 in which A50 is the tensile elongation (in %) and t is the sheet thickness (in mm). This is done to assess properly if tensile elongation increases truly as a result of an increase in the fraction recrystallized ferrite due to an increase in cold-rolling reduction.

(5) Interpretation of results: control over recrystallization

Tables 1.7 to 1.13 show - apart from the tensile data - the fraction recrystallized ferrite and average ferrite grain size of the recrystallized ferrite of all batch-annealed steel sheets. The former microstructural parameter is a clear and direct indication of the degree of recrystallization realized with batch-annealing. Another indicator of recrystallization is the evolution of Rp0.2 after batch annealing as a function of cold-rolling reduction. Tables 1 .7 to 1.13 show the fitting parameters a and b corresponding with a logarithmic fit through Rp0.2 as a function of cold-rolling reduction (CR%) from 30 to 60% (considered here as a typical TRB range) according to

If the value of fitting parameter a is 0 or higher, the evolution of Rp0.2 in the cold-rolling reduction range from 30 to 60% is regarded as an indication of substantial or full recrystallization with an increase in Rp0.2 with increasing cold-rolling reduction coming from additional grain refinement. Hence, it is preferred for the present invention that fitting parameter a is at least 0. Another indicator of increased recrystallization with increased cold- rolling reduction, is an increase in tensile elongation according to A50 / t^0.2.

Also shown in Tables 1 .7 to 1 .13 is an indication if a batch-annealed steel in combination with a particular cold-rolling reduction is suitable as an element for TRB application (i.e. TRB element), for which it is required that the yield strength (Rp0.2) is at least 350 MPa, A50 / t^0.2 ³ 14, and the fraction recrystallized ferrite at ¼ depth is at least 50%. Batch-annealed steels in Tables 1.7 to 1 .13 for which these conditions are fulfilled are marked with an“O” in the TRB column. Batch-annealed steels in Tables 1.7 to 1.13 for which these conditions are not fulfilled are marked with an“X” in the TRB column.

In case of a TRB application relevant for the present invention, an inventive example corresponds with a set of adjacent and suitable TRB elements, wherein each element has a thickness of at least 1 mm and wherein the thickness of the elements obtained with the lowest and highest cold-rolling reduction varies at least 35%, and for which the element with the highest cold-rolling reduction has a yield strength (Rp0.2) of at least 450 MPa, and A50 / t^0.2 ³ 17. Table 1.12 shows an inventive example for a TRB application corresponding with NbTiMo- V alloyed steel 1 G as claimed in the present invention and processed according the invention, which in the present case is coiling at 540 °C, cold-rolling with reductions ranging from 30 to 60%, and batch-annealing for 10 hours at 740 °C (steels 1 G-158BA to 1 G-161 BA).

Most comparative examples in Tables 1 .7 to 1.13 correspond with batch-annealed steels that either have A50 / t^0.2 < 14 and/or a fraction recrystallized ferrite at ¼ depth below 50%. However, in addition to these comparative examples, there are comparative examples in Tables 1 .7 to 1.13 that do have A50 / t^0.2 ³ 14 and a fraction recrystallized ferrite of at least 50%, but which do not form with adjacent suitable TRB elements with a thickness of at least 1 mm a collective of TRB elements wherein the thickness of the elements obtained with the lowest and highest cold-rolling reduction varies at least 35%. All these comparative examples have in common that: (1 ) either the alloy does not contain a substantial amount of V and/or (2) that hot-rolling processing conditions were not adequate, and/or that the batch-annealing conditions were not adequate to provide a solution for a TRB application in which aforementioned elements are met, including a variation in thickness of at least 35%. Table 1.1 : Composition of steels (in wt.%)

Table 1.2: Batch-annealing (BA) cycles for a number of annealing cycles. Shown in this Table, examples of batch annealing cycles from room temperature (RT) to 675, 700, or 740 °C top temperature with 3 or 10 hours holding time at top temperature.

Table 1.3: Tensile properties (longitudinal direction - A50 test piece geometry) of hot-rolled steels coiled at 600 °C

Table 1.4: Tensile properties (longitudinal direction - A50 test piece geometry) of hot-rolled steels coiled at 540 °C and difference (D) in Rp0.2 and Rm compared with steels with identical composition, but coiled at 600 °C (see Table 1 .2).

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Table 1.5: Tensile properties (longitudinal direction - A50 test piece geometry) of batch-annealed (BA) steel sheets coiled at 600 °C after hot rolling (no intermediate cold-rolling - CR% is zero) and difference (D) in Rp0.2 and Rm compared with hot-rolled steels with identical composition but not subjected to batch annealing (see Table 1 .2)

Table 1.6: Tensile properties (longitudinal direction - A50 test piece geometry) of batch-annealed (BA) steel sheets coiled at 540 °C after hot rolling (no intermediate cold-rolling - CR% is zero) and difference (D) in Rp0.2 and Rm compared with hot-rolled steels with identical composition but not subjected to batch annealing (see Table 1.3)

Table 1.7: Tensile properties (longitudinal direction) of batch-annealed (BA with 675 °C top temperature and 3 hours holding at top temperature) steel sheets coiled at 600 °C

Table 1.8: Tensile properties (longitudinal direction) of batch-annealed (BA with 675 °C top temperature and 3 hours holding at top temperature) steel sheets coiled at 540 °C

Table 1.9: Tensile properties (longitudinal direction) of cold-rolled and batch-annealed (BA with 700 °C top temperature and 3 hours holding at top temperature) steel sheets coiled at 540 °C

Table 1.10: Tensile properties (longitudinal direction) of cold-rolled and batch-annealed (BA with 740 °C top temperature and 3 hours holding at top temperature) steel sheets coiled at 540 °C

Table 1.11 : Tensile properties (longitudinal direction) of cold-rolled and batch-annealed (BA with 700 °C top temperature and 10 hours holding at top temperature) steel sheets coiled at 540 °C

Table 1.12: Tensile properties (longitudinal direction) of cold-rolled and batch-annealed (BA with 740 °C top temperature and 10 hours holding at top temperature) steel sheets coiled at 540 °C

Table 1.13: Tensile properties (longitudinal direction) of cold-rolled and batch-annealed (BA with 740 °C top temperature and 7,8, and 9 hours holding at top temperature) steel sheets coiled at 540 °C

EXAMPLE 2:

(1) Alloys, process conditions, testing and microstructural analyses procedures

Steels 2A to 2G having chemical compositions shown in Table 2.1 were hot rolled and further processed in a similar fashion as reported in Example 1 . The tensile properties were measured in an identical way as reported in Example 1. The procedures followed to determine fraction recrystallized ferrite and the average grain size of the recrystallized ferrite were identical to those reported in Example 1 .

(2) Tensile properties hot-rolled and batch-annealed steel sheets

(2A) Hot-rolled steel sheets : Table 2.2 gives the A50 tensile properties of the hot-rolled steel sheets coiled at 600 or 540 °C of steels 2A to 2G. Labelling of the corresponding hot- rolled steel sheets is done in a similar fashion as previously in Example 1 .

(2B) Batch-annealed steel sheets Tables 2.3 gives the tensile properties of batch-annealed steel sheets without any intermediate cold rolling (CR% equals 0%) and corresponding with hot-rolled steel sheets coiled at 600 and 540 °C, respectively. Tables 2.4 and 2.5 provide the tensile properties of the batch-annealed steel sheets with intermediate cold-rolling reductions of 0 to 60% for associated with hot-rolled feedstock coiled at 600 and 540 °C, respectively.

(3) Microstructures batch-annealed steel sheets

Tables 2.4 and 2.5 provide the fraction recrystallized ferrite (in %) and the average grain size (in mm) of the recrystallized ferrite based on EBSD measurements.

(4) Interpretation of results: control over recrystallization

Tables 2.4 and 2.5 show - apart from the tensile data - the fraction recrystallized ferrite and average ferrite grain size of the recrystallized ferrite of all batch-annealed steel sheets. The former microstructural parameter is a clear and direct indication of the degree of recrystallization realized with batch-annealing. Another indicator of recrystallization is the evolution of Rp0.2 after batch annealing as a function of cold-rolling reduction. Tables 2.4 and 2.5 show the fitting parameters a and b corresponding with a logarithmic fit through Rp0.2 as a function of cold-rolling reduction (CR%) from 30 to 60% (considered here as a typical TRB range) according to

If the value of fitting parameter a is 0 or higher, the evolution of Rp0.2 in the cold-rolling reduction range from 30 to 60% is regarded as an indication of substantial or full recrystallization with an increase in Rp0.2 with increasing cold-rolling reduction coming from additional grain refinement. Hence, it is preferred for the present invention that fitting parameter a is at least 0. Another indicator of increased recrystallization with increased cold- rolling reduction, is an increase in tensile elongation according to A50 / t^0.2.

Also shown in Tables 2.4 and 2.5 is an indication if a batch-annealed steel in combination with a particular cold-rolling reduction is suitable as an element for TRB application (i.e. TRB element), for which it is required that the yield strength (Rp0.2) is at least 350 MPa, A50 / t^0.2 ³ 14, and the fraction recrystallized ferrite at ¼ depth is at least 50%. Batch-annealed steels in Tables 1.7 to 1 .13 for which these conditions are fulfilled are marked with an“O” in the TRB column. Batch-annealed steels in Tables 1.7 to 1.13 for which these conditions are not fulfilled are marked with an“X” in the TRB column.

In case of a TRB application relevant for the present invention, inventive examples in Tables 2.4 and 2.5 correspond with a set of adjacent and suitable TRB elements, wherein each element has a thickness of at least 1 mm and wherein the thickness of the elements obtained with the lowest and highest cold-rolling reduction varies at least 35%, and for which the element with the highest cold-rolling reduction has a yield strength (Rp0.2) of at least 450 MPa, and A50 / t^0.2 ³ 17. Table 2.5 shows a large number of inventive examples for a TRB application corresponding all with NbTiMo-V alloyed steels (steels 2C to 2G) as claimed in the present invention and processed according the invention, which in the present case is coiling at 540 °C, cold-rolling with reductions ranging from 20 to 60% or 30 to 60%, and batch- annealing for at least 10 hours at a top temperature of at least 700 °C.

Most comparative examples in Tables 2.4 and 2.5 correspond with batch-annealed steels that either have A50 / t^0.2 < 14 and/or a fraction recrystallized ferrite below 50%. However, in addition to these comparative examples, there are a number of comparative examples in Tables 2.4 and 2.5 that do constitute TRB elements with A50 / t^0.2 ³ 14 and a fraction recrystallized ferrite at ¼ depth of at least 50%, but which do not form with adjacent suitable TRB elements with a thickness of at least 1 mm a collective of TRB elements wherein the thickness of the elements obtained with the lowest and highest cold-rolling reduction varies at least 35%. All these comparative examples have in common that: (1 ) either the alloy does not contain a substantial amount of V (i.e., steels 2A and 2B) and/or (2) that hot-rolling processing conditions were not adequate, and/or that the batch-annealing conditions were not adequate to provide a solution for a TRB application in which aforementioned elements are met, including a variation in thickness of at least 35%. Table 2.1 : Composition of steels (in wt.%)

Table 2.2: Tensile properties (longitudinal direction - A50 geometry) of hot-rolled steels coiled at 600 or 540 °C

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Table 2.3: Tensile properties (longitudinal direction) of batch annealed steels coiled at 540 °C (no intermediate cold rolling) and difference (D) in

Rp0.2 and Rm compared with identical steels but not subjected to batch (see Table 2.2)

Table 2.4: Tensile properties (longitudinal direction) of cold-rolled and batch-annealed (BA with 675 and 720 °C top temperature and 16 hours holding at top temperature) steel sheets coiled at 600 °C

Table 2.5: Tensile properties (longitudinal direction) of cold-rolled and batch-annealed steel sheets coiled at 540 °C

Claims

1. Variably rolled steel strip, sheet or blank, consisting of the following elements (in wt%):

0.05 - 0.20 C

0.10 - 0.70 Si

0.80 - 2.50 Mn

0.01 - 0.10 Al

0.07 - 0.25 Ti

0.10 - 0.35 V

0.05 - 0.40 Mo

optionally 0.02 - 0.10 Nb

optionally 0.01 - 0.80 Cr

at most 0.06 P

at most 0.01 S

at most 0.01 N

at most 0.005 Ca

the balance consisting of inevitable impurities and Fe,

wherein the strip, sheet or blank has at least one portion having a high thickness and at least one portion having a low thickness, wherein the variation in thickness between the high thickness and the low thickness is at least 35%, and wherein in the portion with high thickness the yield strength is 350 MPa or higher and the tensile elongation A50 (in %) and sheet thickness t (in mm) satisfy the equation of A50 / t^0.2 ³ 14, and in the portion with low thickness the yield strength is 450 MPa or higher and the tensile elongation A50 (in %) and sheet thickness t (in mm) satisfy the equation A50 / t^0.2³ 17.

2. Variably rolled steel strip, sheet or blank according to claim 1 , wherein for one or more elements the following range is valid:

0.06 - 0.17 C

0.20 - 0.60 Si

0.90 - 2.30 Mn

0.02 - 0.09 Al

0.08 - 0.22 Ti

0.12 - 0.30 V

0.08 - 0.35 Mo

optionally 0.03 - 0.09 Nb

optionally 0.01 - 0.60 Cr at most 0.04 P

at most 0.005 S

at most 0.008 N

at most 0.003 Ca.

3. Variably rolled steel strip, sheet or blank according to claim 1 or 2, wherein for one or more elements the following range is valid:

0.07 - 0.14 C

0.30 - 0.60 Si

1 .20 - 2.00 Mn

0.04 - 0.08 Al

0.10 - 0.20 Ti

0.13 - 0.25 V

0.10 - 0.30 Mo

optionally 0.03 - 0.08 Nb

optionally 0.01 - 0.40 Cr

at most 0.02 P

at most 0.003 S

0.002 - 0.007 N

at most 0.001 Ca.

4. Variably rolled steel strip, sheet or blank according to any one of the preceding claims, wherein Nb is present in an amount between 0.02 and 0.10%, preferably between 0.03 and 0.09%, more preferably between 0.03 and 0.08%.

5. Variably rolled steel strip, sheet or blank according to any one of the preceding claims, wherein the contents of Ti and Mo and optionally Nb represented by weight percentage (wt.%) satisfy the equation of

or preferably

6. Variably rolled steel strip, sheet or blank according to any one of the preceding claims, wherein the contents of C, N, Ti, Mo, V and optionally Nb represented by weight percentage (wt.%) satisfy the equation of

with

and

7. Variably rolled steel strip, sheet or blank according to any one of the preceding claims, wherein the steel has a precipitation strengthened ferrite microstructure, consisting of at least 50% recrystallized ferrite at ¼ depth of the portion with high thickness, and which optionally further contains cementite and/or pearlite, and wherein the precipitates in said microstructure consist of Ti, V, Mo and optionally Nb.

8. Variably rolled steel strip, sheet or blank according to any one of the preceding claims, wherein the steel has a microstructure that consists of at least 60% recrystallized ferrite at ¼ depth in the portion with high thickness, preferably at least 70% recrystallized ferrite at ¼ depth in the portion with high thickness, more preferably at least 80% recrystallized ferrite at ¼ depth in the portion with high thickness, most preferably 90% recrystallized ferrite at ¼ depth in the portion with high thickness.

9. Variably rolled steel strip, sheet or blank according to any one of the preceding claims, wherein the steel in the portion with high thickness has a yield strength of 400 MPa or higher, preferably of 450 MPa or higher, or more preferably of 500 MPa or higher, or most preferably 550 MPa or higher, and/or wherein the steel in the portion with low thickness has a yield strength equal or higher than that of the portion with high thickness, and preferably has a yield strength of 500 MPa or higher, more preferably of 550 MPa or higher.

10. Variably rolled steel strip, sheet or blank according to any one of any one of the preceding claims, wherein the steel in the portion with high thickness satisfies the equation A50 / t^0.2 ³ 16, preferably A50 / t^0.2 ³ 18, most preferably A50 / t^0.2 ³ 20, and wherein the steel in the portion with low thickness satisfies the equation: A50 / t^0.2³ 18, preferably A50 / t^0.2 ³ 20, most preferably A50 / t^0.2 ³ 22, with A50 / t^0.2 in the portion with low thickness equal to or higher than A50 / t^0.2 in the portion with high thickness.

1 1. Method for producing a steel strip having a variable thickness, comprising the steps of:

• casting a slab having the composition according to any one of the preceding claims,

• reheating the solidified slab to a temperature between 1 150 and 1300° C,

• finishing the hot rolling at a finish hot rolling temperature of Ar3 transformation point or higher,

• cooling the hot-rolled steel strip to the coiling temperature at an average cooling rate of 10 to 150° C/s,

• coiling the hot-rolled steel strip in the temperature range between 450 and 580° C,

• cold-rolling the strip as variable rolling, such that a cold-rolling reduction between 30% and 60% is performed and a variation of strip thickness of at least 35% is obtained,

• batch annealing the steel strip.

12. Method according to claim 12, wherein the hot-rolled steel strip or sheet is hot-rolled with a finish rolling temperature of 870° C or higher, preferably with a finish rolling temperature of 900° C or higher, more preferably with a finish rolling temperature of 940° C or higher, and most preferably with a finish rolling temperature of 980° C or higher.

13. Method according to claim 12 or 13, wherein the hot-rolled steel strip or sheet after finish rolling is cooled to the coiling temperature with an average cooling rate of 40 to 100° C/s.

14. Method according to claim 12,13 or 14, wherein the hot-rolled steel strip or sheet is coiled in the temperature range between 480 and 560° C.

15. Method according to any one of claims 12 to 15, wherein the cold-rolled steel strip or sheet is batch annealed

• for at least 8 hours at a top temperature of 740 °C or higher, or

• for at least 10 hours at a top temperature of 720 °C or higher, or

• for at least 14 hours at a top temperature of 700 °C or higher.