AND METHOD TO PRODUCE SUCH STRIP

The invention relates to a steel strip, sheet or blank having improved formability, and to a method for producing such a steel strip, sheet or blank.

In the automotive industry there exists a constant wish for improved steel types having a high strength and a good formability. For such steel types, special compositions with increased levels of one or more micro-alloying elements may be needed. The strength of such steels can for a large part be achieved by precipitation strengthening. However, with conventional hot rolling and coiling, the time span within the suitable temperature window for precipitation in ferrite can be too short when austenite-to-ferrite transformation is sluggish and a substantial amount of transformation occurs during coiling and coil cooling at temperatures which no longer provide sufficient kinetics for precipitation. This can lead to too low strength due to insufficient precipitation since precipitation is not fully developed.

An object of the invention is to provide a formable steel strip, sheet or blank suitable for automotive chassis components.

It is another object of the invention to provide a formable steel strip, sheet or blank showing sufficient recrystallization, preferably a fully recrystallized microstructure, at a low rolling reduction, such as 30 - 40% cold rolling reduction.

A further object of the invention is to provide a formable steel strip, sheet or blank suitable for a tailor-rolled blank for automotive chassis components, which after flexible cold-rolling with reductions of 30% and higher, followed by batch annealing, has a variable thickness with a high yield strength and suitable tensile elongation (A50 in %).

A further object of the invention is to provide a batch-annealed formable steel strip or sheet according to the objects described hereinabove wherein the steel has a microstructure that consists of at least 50% recrystallized ferrite.

A further object of the invention is to provide a method of manufacturing a formable steel strip or sheet according to the objects described hereinabove.

According to a first aspect of the invention, one or more of these objects can be reached by a 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 steel has a microstructure of at least 90% ferrite, the remainder being cementite and/or pearlite.

The invention provides a formable steel strip or sheet with an excellent balance between recrystallization behaviour, strength, and ductility after batch-annealing and cold rolling. 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. Aim of the present invention is that after hot rolling, the steel is cooled down and coiled at a relatively low temperature to suppress the formation of precipitates in general, and that of V-based precipitates in particular. The intention is to have as much V in solid solution in the ferrite matrix of the hot-rolled intermediate feedstock prior to cold rolling (either to produce a uniform thickness via conventional cold rolling or a variable thickness via flexible rolling) and subsequent batch annealing within the ferrite phase region (below Ac1 transformation point).

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 is mitigated by an increase in strength from grain refinement.

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.

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.%. The microstructure thus contains at least 90% ferrite, which is the sum of recrystallised ferrite and non-recrystallised ferrite.

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 _x O _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 low formability 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 (Tdis 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 273.15

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 Tdis 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 _x O _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 steel strip, sheet or blank has a yield strength of 350 MPa or higher, preferably of 400 MPa or higher, more preferably of 450 MPa or higher, still more preferably of 500 MPa or higher, most preferably 550 MPa or higher after batch annealing. Such a yield strength is suitable for automotive applications.

Preferably the steel strip, sheet or blank has a precipitation strengthened ferrite microstructure containing recrystallized ferrite, cementite and/or pearlite, and wherein the precipitates in said microstructure consist of Ti, V, Mo and optionally Nb after batch annealing. The invention provides a steel strip or sheet with a predominantly ferrite microstructure that is strengthened with precipitates consisting of Ti, V, Mo, and optionally Nb after batch annealing. The indication“predominantly” in this case means that the ferrite microstructure comprises at least 90%, or preferably at least 95%, or more preferably at least 97% ferrite, or most preferably 100% ferrite. The remainder of the microstructure after batch annealing can be cementite and/or pearlite. Hence, for the present invention, a predominantly ferrite microstructure may consist of at most 10%, or preferably at most 5%, or more preferably of at most 3% of cementite and/or pearlite.

According to a preferred embodiment the steel strip, sheet or blank contains Nb, Ti, and Mo 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

48 X N

Ti ^* = Ti -

14

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

Preferably, the steel after hot rolling and annealing has a yield ratio of at least 0.9, and/or a tensile strength of 900 MPa or higher, preferably 950 MPa or higher, and/or an elongation A50/t ^{0 2} that is 9% or higher. Such mechanical properties are often required by the automotive industry for high strength steel.

According to a preferred embodiment, the steel after cold rolling with a cold rolling reduction of at least 30% and annealing has a yield strength of 350 MPa or higher, preferably of 400 MPa or higher, more preferably of 450 MPa or higher, still more preferably of 500 MPa or higher, most preferably 550 MPa or higher, and an elongation A50 / 1 ^{0 2} that is 14% or higher. These mechanical properties are advantageous for such high strength steel types.

Preferably, the steel has a precipitation strengthened ferrite microstructure containing recrystallized ferrite, cementite and/or pearlite, and wherein the precipitates in said microstructure consist of Ti, V, Mo and optionally Nb, and wherein fraction recrystallized ferrite at ¼ depth is at least 50%, preferably at least 60%, more preferably at least 70%, most preferably at least 80%. Such a microstructure can provide the mechanical properties that are looked for.

According to a second aspect of the invention a method is provided for producing the steel according to the first aspect of the invention, 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.

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.

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.

According to a preferred embodiment the hot-rolled steel strip is batch annealed after hot rolling for at least 1 hour at a top temperature between 550 and 700 °C, preferably at a top temperature between 600 and 700 °C, more preferably at a top temperature between 650 and 700 °C. The heating of the hot rolled strip promotes the formation of new precipitates and thus increases the yield strength of the hot rolled strip to provide an improved balance between strength and formability.

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 at least 50% recrystallized cold rolled strip, resulting in a precipitation strengthened ferrite microstructure with an adequate balance between strength and formability.

The cold-rolled steel strip or sheet is thus batch annealed, preferably in an inert and protective atmosphere, consisting of hydrogen and/or nitrogen, to prevent excessive decarburisation and/or Fe-based oxide scale formation, and/or to promote more efficient heat transfer during batch annealing.

The batch annealing cycle preferably uses an overall slow heating rate to top temperature to provide sufficient time and thermal energy to promote precipitation in general and V precipitation in particular, in order to pin dislocations to suppress recovery.

The preferred top temperature of the batch annealing cycle depends partially on the amount of V and C in solid solution in the ferrite matrix available for the formation of VC precipitates and consequently is linked to the minimum temperature needed to dissolve the VC precipitates formed during the initial stages of the batch annealing. Hence the top temperature (Ttop in °C) during batch annealing should be at least equal to the dissolution temperature (Tdis in °C), or

5500

T ° > ° 273.15

with the V content [V] expressed in wt.%.

The inventors found that for the present invention a top temperature of 700 °C or higher will provide an adequate degree of recrystallization in order to ensure that the object of the present invention is realised; the yield strength (Rp0.2) of the steel strip or sheet after batch annealing is constant or preferably increases with an increase in cold-rolling reduction (CR%) of 30% and higher, or preferably 30% to 60%. Furthermore, inventors found that a dwell time of at least 14 hours at a top temperature of 700 °C or higher during the batch annealing cycle is preferred to reach aforementioned objects of the present invention.

Furthermore, the batch annealing cycle requires an overall slow cooling rate from top temperature to circa 200 °C to provide sufficient time and thermal energy to promote again precipitation in general and V precipitation in particular.

After batch annealing, the steel strip or sheet can optionally be provided with a zinc- based coating to the surface of the steel strip or sheet to provide corrosion protection. This can be done via a coating process, such as heat-to-coat or electro-galvanizing, wherein a zinc or zinc alloy coating is applied to the surface of the steel strip or sheet. In case of the use of a zinc alloy coating for corrosion protection, the alloy preferably contains Aluminium and/or Magnesium as its main alloying elements. The invention will now be explained 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 use 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, Ti, Nb 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.

In view of their composition, it will be clear that steel 1A, 1 B, 1 D, 1 E, 1 F and 1 H are comparative examples, since they contain less than 0.10 wt% V.

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 Fh 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. Figure 1 shows an example for the BA-740/10 curve.

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).

(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.

(3) 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 1 A to 1 H coiled at 600 and 540 °C, respectively, 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-alloy content was consumed in precipitation during the hot-rolling stage.

In contrast, when the hot-rolled steel sheets corresponding with steels 1A 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 Ttop 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., >50%) or - preferably - full recrystallization already at a relatively low cold-rolling reduction (e.g., CR% > 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% > 30%. Table 1.1 : Composition of steels (in wt.%)

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Table 1.2: Batch-annealing (BA) cycles for a number of annealing cycles. Shown in this Table, examples of batch annealing cycles from room 5 temperature (RT) to 675, 700, or 740 °C top temperature with 3 or 10 hours holding time at top temperature.

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

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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

o rolling (no intermediate cold-rolling - CR% is zero) and difference (D) in Rp0.2 and Rm compared with hot-rolled steels with identical composition

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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)

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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

In view of their composition, it will be clear that steel 2A and 2B are comparative examples, since they contain less than 0.10 wt% V.

. The procedures followed to determine fraction recrystallized ferrite and the average grain size of the recrystallized ferrite are as follows.

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 pm. 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 pm 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 pm, with a step size of 0.1 pm, 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 x64 [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

(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 pm) 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. Table 2.1 : Composition of steels (in wt.%)

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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 annealing (see Table 2.2)

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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

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

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