The invention relates to a process for producing batch annealed tailor rolled steel. The invention also relates to a tailor rolled strip produced thereby, and a tailor rolled blank cut from the tailor rolled strip.

To produce tailor rolled blanks usually a tailor rolled strip is cut into parts having at least one thicker and one thinner portion, as is known in the art. The tailor rolled strip is produced from cold rolled steel having a standard thickness, which strip is differentially rolled to provide alternating different thicknesses over the length of the strip. After the differential rolling the strip is usually batch annealed, because it is difficult to continuous anneal a strip having different thicknesses. The batch annealing normally does not lead to high yield strength.

Conventional high strength steel types, such as dual phase steel, make use of continuous annealing lines. Such conventional high strength steels do not contain sufficient manganese to get martensite in the microstructure during the batch annealing, resulting in an yield strength that is too low.

It is an object of the invention to provide a process for producing batch annealed tailor rolled strip having high strength.

It is another object of the invention to provide a process for producing batch annealed tailor rolled strip having a good elongation.

It is a further object of the invention to provide a process for producing batch annealed tailor rolled strip having a composition that is suitable for providing tailor rolled strip having a dual phase microstructure.

It is furthermore an object of the invention to provide the tailor rolled strip produced according to the process, and the tailor rolled blanks that are cut from the tailor rolled blanks.

According to a first aspect of the invention a process for producing batch annealed tailor rolled steel strip is provided, comprising the following steps:

a) casting a slab having a composition of (in wt%):

C = 0 - 0.20; Mn = 2.0 - 4.0; Si £ 1.0; P = 0.001 - 0.05: S = 0.001 - 0.01 ; N £ 0.02; Al = 0.01 - 0.10; optionally: Ti £ 0.04; Nb £ 0.04; V £ 0.20; Mo £ 0.10; B £ 0.005; Cr £ 1.0; Ca £ 0.005, wherein Ti + Nb £ 0.05, wherein V + Mo £ 0.15, and wherein 2.0 £ Cr + Mn £ 4.0, the remainder being Fe and unavoidable impurities

b) hot rolling the slab to a finish rolling temperature between 910 and 1 100 °C c) cooling the hot rolled strip to a coiling temperature between 200 and 750 °C d) cold rolling the hot rolled strip with a reduction of at least 30% using differential rolling to produce different thicknesses within the strip, so as to produce a tailor rolled steel strip, to be cut into tailor rolled blanks

e) batch annealing the differential cold rolled strip at a temperature between 650 °C and the Ac3 temperature of the strip, to obtain a batch annealed tailor rolled steel strip having a Rm of at least 570 MPa.

The invention shows that control of the hot-rolling process is vital to the microstructural development during batch-annealing. In the first instance a high finishing rolling temperature is required to ensure that the structure remains in the austenite state during rolling to avoid unfavourable development in texture leading to poor formability [Bodin et. Al, Metal Mater Trans. A. (2000) 33. A. 1589-1603] and to retain precipitating elements such as Nb, Ti, Mo and V from forming dynamic precipitates during rolling. It is essential that these elements remain in solid solution so that they can form precipitates of the small size required for strengthening during the batch- annealing process. For this reason, the cooling rate after hot-rolling should be sufficient to avoid the formation of precipitates during cooling (³1 °C/s) where these elements are present. Thus a finish rolling temperature between 910 and 1100 °C is needed.

Cooling at a moderate rate (~1 to ~20°C/s) to a fairly high coiling temperature (600 - 750°C) results in a ferrite-pearlite hot-rolled microstructure. The ferrite is polygonal and free from carbides and there are pearlite colonies distributed throughout, such that all of the carbon available is contained within the randomly distributed pearlite colonies. Manganese has a stabilising effect on pearlite, so these colonies will be predominately associated with areas where the local manganese content is increased, due to segregation and/or normal variability. Both manganese and carbon serve to reduce the ferrite to austenite transformation temperature so that during batch annealing, the first austenite to be formed will nucleate within the pearlite colonies. This austenite will become further stabilised due to manganese and carbon partitioning and will eventually transform to martensite during cooling after annealing. The resultant microstructure will be a dual phase mixture of polygonal ferrite with randomly distributed martensite islands. Due to the long annealing times, there is significant ferrite grain growth which leads to a reduction in yield strength with respect to that in the hot-rolled condition.

By increasing the cooling rate to above 20°C/s and in combination with coiling temperature below the eutectoid temperature (~650°C or higher for steels with Mn>1.0 wt.%) but above the martensite start temperature, it is possible to achieve not only the retention of precipitating elements in solution, but more importantly to suppress the formation of pearlite and promote the formation of acicular ferrite and/or bainite. This has the advantage that the carbon is evenly distributed throughout the microstructure.

The cold-rolling reduction of the hot rolled strip is at least 30% to different thicknesses with in the same coil.

Upon reheating during batch annealing after cold-rolling, carbon trapped in ferrite/bainite diffuses towards the grain boundaries which become nucleation sites for austenite transformation. As the temperature increases, the austenite grains grow epitaxially, thereby inhibiting ferrite grain growth and results in an effective reduction of ferrite grain size due to the impingement of the austenite growth at the grain boundaries. Upon cooling, the austenite transforms to martensite leaving all of the ferrite grains ringed with martensite. This has a dual strengthening effect: (a) the yield strength is increased as a result of the reduced ferrite grain size according to a Hall-Petch relationship and (b) the increase in martensite-ferrite grain boundaries results in an increase in load partitioning between martensite and ferrite leading to increased work hardening capacity of the ferrite phase prior to yielding which is characteristic for multiphase type steels. To reach the right grain size, the batch annealing temperature must be between 650 °C and the Ac3 temperature of the strip, preferably during a time of at least 2 hours and at most 24 hours.

The composition of the steel that is used in the process is chosen for the following reasons.

Carbon (C) is added to form carbide and/or carbo-nitride precipitates, optionally together with Ti, Nb, V and Mo. The amount of C depends on the amount of Ti, Nb, V and/or Mo used. However, the maximum content is 0.20 wt.% to prevent excessive segregation and to prevent a too high fraction of cementite and/or pearlite. A more preferable C content range for the present invention is between 0.001 and 0.15 wt.%, or most preferably between 0.002 and 0.12 wt.%.

Silicon (Si) provides significant solid-solution strengthening, which can be used 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. 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 may not exceed 1.0 wt.%. A more preferable Si content range for the present invention is at most 0.7 wt.%, or more preferably at most 0.60 or 0.50 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 2.0 wt.%. However, a too high Mn content may lead to excessive segregation, which can 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 4.0 wt.%. A more preferable Mn content range for the present invention is between 2.1 and 3.8 wt.%, or most preferably between 2.2 and 3.7 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 0.05 wt.%, or preferably at most 0.04 wt.%, and more preferably at most 0.02 wt.%. Due to the production process of steel there will always be some P present, at least 0.001 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.%. Due to the production process of steel there will always be some S present, at least 0.001 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. 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 effect of Nitrogen (N) on mechanical properties is similar to that of carbon being increased strength as a result of (i) interstitial solid solution strengthening by the free nitrogen (ii) precipitation strengthening by aluminium and other nitrides (V, Nb, Ti) and (iii) grain refinement due to the presence of nitride precipitates. As well as nitrides, nitrogen can also be present in more complex precipitate species in combination with carbon and/or boron e.g. carbo-nitrides, boro-nitrides etc. Small amount of nitrogen present as precipitates have a beneficial effect on impact properties. Nitrides of aluminium, vanadium, niobium and titanium result in the formation of fine grained ferrite. Finer grain size lowers the transition temperature and improves the toughness. However nitrogen content of more than 0.02 wt.% (200 ppm) may result in inconsistent mechanical properties in hot rolled steels, embrittlement of the heat affected zone (HAZ) of welded steels, and poor cold formability. In particular, nitrogen can result in strain ageing and reduced ductility of cold rolled and annealed low carbon aluminium killed steels. A more preferable range for N content for the present invention is at most 0.012 wt.%, or most preferably between 0.002 and 0.010 wt.%.

Titanium (Ti) can be used in the present invention to realise precipitation strengthening and to some degree grain refinement. As such, Ti is an optional 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 maximum Ti content is 0.04 wt.%. A more preferable Ti maximum content for the present invention is 0.02 wt. However, for strength that is relatively low, no Ti needs to be added.

Niobium (Nb) can be 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. The use of Nb is considered as optional for the present invention. However, when used, a suitable maximum Nb content is 0.04 wt.%, more preferably 0.02 wt.%.

Since Ti and Nb result in the same effects in the steel, the total of Ti + Nb is set at max 0.05 wt.%.

Vanadium (V) 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. However, since there is no need to add V to the steel in the process according to the present invention, a suitable maximum of V is 0.20 wt.%. Preferably, the maximum V content is 0.10 wt.%.

Molybdenum (Mo) is known to be a carbide-forming element and can form together with Ti, Nb and/or V 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 most 0.10 wt.%. However, there is no need to add Mo.

Since V and Mo result in the same effects in the steel, the total of V + Mo is set at at most 0.15 wt.%.

Chromium (Cr) is an element that can be added as a replacement of Mn to provide solid-solution strengthening. Since Cr can have detrimental effects for the end product, Cr should be added in an amount of at most 1.0 wt.%. Since it is a replacement of Mn, the total of Cr + Mn should be between 2.0 and 4.0 wt.% and follow the preferred ranges for Mn.

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 castability 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 first preferred process, the slab has a composition of: C = 0 - 0.10; Mn = 2.2 - 3.2; Si £ 0.4; Al = 0.01 - 0.10; P = 0.001 - 0.05; S = 0.001 - 0.01 ; N £ 0.02 and optionally: B £ 0.005; Cr £ 1.0; Ca £ 0.005 wherein 2.0 £ Cr + Mn £ 4.0 wherein the finish rolling temperature is between 910 and 1100 °C, wherein the coiling temperature of the hot rolled strip is between 200 and 700 °C, and wherein the batch annealing temperature is between 650 and 760 °C, to obtain a batch annealed tailor rolled steel strip having a Rm between 570 and 750 MPa.

Since according to this preferred embodiment the strip that is produced has a strength between 570 and 750 MPa, lower amounts of C, Mn and Si can be used, and there is no need to add Ti, Nb, V or Mo. To produce the hot rolled steel also the annealing temperature can be lower, that is between 650 and 760 °C.

According to a second preferred embodiment the slab has a composition of: C = 0.02 - 0.11 ; Mn = 2.2 - 3.2; Si £ 0.4; Al = 0.01 - 0.10; P = 0.001 - 0.05; S = 0.001 - 0.01 ; N £ 0.02, and optionally: Ti £ 0.04; Nb £ 0.04; V £ 0.20; Mo £ 0.10; B £ 0.005; Cr £ 1.0; Ca £ 0.005; wherein Ti + Nb £ 0.05, and wherein 2.0 £ Cr + Mn £ 4.0, the remainder being Fe and unavoidable impurities, wherein the finish rolling temperature is between 910 and 1100 °C, wherein the coiling temperature of the hot rolled strip is between 200 and 580 °C, and wherein the batch annealing temperature is between 670 and 790 °C, to obtain a batch annealed tailor rolled steel strip having a Rm between 700 and 900 MPa.

In this preferred embodiment the strip that is produced has a strength between 700 and 900 MPa, so the amounts of C, Mn and Si can be somewhat higher than according to the previous embodiment. Moreover, it is an option to add Ti, Nb, V and/or Mo. Furthermore, the annealing temperature can be slightly higher, between 670 and 790 °C.

According to a third preferred embodiment the slab has a composition of: C = 0.06 - 0.14; Mn = 3.0 - 4.0; Si = 0.2 - 0.7; Al = 0.01 - 0.10; P = 0.001 - 0.05; S = 0.001 - 0.01 ; N £ 0.02 and optionally: Ti £ 0.04; Nb £ 0.04; V £ 0.20; Mo £ 0.10; B £ 0.005; Cr £ 1.0; Ca £ 0.005, wherein Ti + Nb £ 0.05, and wherein 2.0 £ Cr + Mn £ 4.0, wherein the finish rolling temperature is between 910 and 1100 °C, wherein the coiling temperature of the hot rolled strip is between 200 and 580 °C, wherein the batch annealing temperature is between 670 and 790 °C, to obtain a batch annealed tailor rolled steel strip having a Rm between 900 and 1200 MPa.

The strip produced according to this third embodiment has a high strength between 900 and 1200 MPa, for which higher amounts of C, Mn and Si can be added. It is also possible to add Ti, Nb, V and/or Mo. Here as well the annealing temperature can be slightly higher then for the first embodiment.

Preferably the coiling temperature of the hot rolled strip is at least 300 °C, preferably at least 400 °C, more preferably at least 500 °C. Since the higher limit for the coiling temperature is most relevant, it is advantageous to coil the strip at a higher temperature.

According to a preferred embodiment, the composition contains B £ 0.003 wt%. There is no specific need to add more boron.

Since high strength steels are mostly used for automotive purposes, it is preferred when the batch annealed tailor rolled dual phase steel strip is coated with an organic coating or a metallic coating, preferably by heat-to-coat, electro-galvanising, PVD or CVD.

When using a metallic coating, this coating is preferably a zinc or zinc alloy coating or an aluminium or aluminium alloy coating, such as a Gl coating, a GA coating, a AlSi coating or a ZnAIMg coating. These are the coatings mainly used for automotive purposes.

According to a second aspect of the invention there is provided a tailor rolled strip obtained by using the process in accordance with the process described above, wherein the strip has alternating portions having a high thickness and portions having a low thickness, wherein the variation in thickness between the high thickness and the low thickness is at least 15%.

According to a first embodiment this strip has an ultimate tensile strength Rm between 570 and 750 MPa, a 0.2% proof strength Rp0.2 between 270 and 400 MPa and a total elongation A80 of at least 18%. Such a strip can be used as a DP600 dual phase tailor rolled strip.

According to a second embodiment this strip has an ultimate tensile strength Rm between 700 and 900 MPa, a 0.2% proof strength Rp0.2 between 350 and 550 MPa and a total elongation A80 of at least 14%. Such a strip can be used as a DP800 dual phase tailor rolled strip.

According to a third embodiment this strip has an ultimate tensile strength Rm between 900 and 1200 MPa, a 0.2% proof strength Rp0.2 between 500 and 700 MPa and a total elongation A80 of at least 8%, preferably a total elongation of at least 10%, more preferably a total elongation of at least 12%. Such a strip can be used as a DP1000 dual phase tailor rolled strip. DP1000 has a total elongation of at least 12%.

The strip according to the second aspect of the invention can also be a multiphase steel. Usually the total elongation of multiphase steel is somewhat lower than the total elongation of dual phase steel.

Preferably the steel of the tailor rolled strip according to the first embodiment of the second aspect of the invention has a microstructure consisting of at most 10% bainite and at most 15% martensite, the remainder being ferrite. The ferrite fraction is acicular ferrite. With this microstructure, the strength and elongation of the batch annealed strip according to the first embodiment are sufficient. In dual phase steel usually no bainite is present.

Preferably the steel of the tailor rolled strip according to the second or third embodiment of the second aspect of the invention has the microstructure consisting of 10 - 20% martensite, 0 - 20% bainite, the remainder being ferrite. The ferrite fraction is acicular ferrite. With this microstructure, the strength and elongation of the batch annealed strip according to the second and third embodiment are sufficient. In dual phase steel usually no bainite is present.

According to a third aspect of the invention a tailor rolled blank (TRB) cut from a tailor rolled strip according to the second aspect of the invention is provided, wherein the TRB has at least one portion having a high thickness and at least one portion having a low thickness. This TRB has a strength and elongation in accordance with that of the tailor rolled strip from which it is cut. This combination of strength and elongation makes the TRB very suitable for use in automotive products. Preferably the tailor rolled strip or TRB according to the third aspect of the invention has a zinc or zinc alloy coating or an aluminium or aluminium alloy coating, such as a Gl coating, a GA coating, a AlSi coating or a ZnAIMg coating. The TRB with a metallic coating is especially suitable for automotive purposes.

The invention will be elucidated in view of a series of steel types as provided in the tables hereinbelow.

Table 1 shows the composition of 21 different steels types in accordance with the present invention. Table 2 shows the process conditions for producing these steel types, and the resulting mechanical properties thereof.

The composition of and process conditions for the steel types 1 to 9 result in steel types according to the first embodiment as described above. All these nine steel types a DP600 steels.

The composition of and process conditions for the steel types 10 to 18 result in steel types according to the second embodiment as described above. These nine steel types are DP800 steels.

The composition of and process conditions for the steel types 19, 20 and 20 result in steel types according to the third embodiment as described above. These nine steel types are DP1000 or multiphase steels having a strength level of a DP1000 steel.

In Table 1 the main alloying elements C, Mn, Si and B are always shown. The other elements are not always measured, because in the steel types of these examples Mo, Cr, Cu, Ni, Nb, P, S and N are impurities. However, the amount of N can be higher than usual in high strength steels. Ti and V are added in steel types 17, 18 and 21. The amount of Al is low because in laboratory experiments no Al needs to be added as deoxidising element.

Though it is not exemplified, Nb can be added instead of or in combination with Ti, since it results in the same effects for the steel. This is common knowledge for the skilled person. In the same way, Mo can be added instead of or in combination with V.

Furthermore, it is known to the skilled person that Chromium can be added instead of Mn, though this is not exemplified.

Table 2 shows that with a finish rolling temperature (FRT) of at least 910 °C and a coiling temperature (CT) of at most 650 °C, these steel types result in the required mechanical properties, when the reduction (CR red., cold rolling reduction) is at least 30% and the batch annealing temperature (Soak) is more than 650 °C. For information purposes, the calculated austenisation temperatures, Ae3 and Ae1 , the bainite start temperature Bs and the martensite start temperature Ms are given for all 21 steel types.

Though not indicated in the tables, the cooling rate for the strip on the run-out table is at least 20 °C/s, the heating rate before batch annealing was approximately 40 °C/s and the cooling rate after batch annealing was approximately 30 °C/s. However, these heating and cooling rates are not decisive, as long as the cooling rate on the run out table is more than 20 °C.

Table 2 shows that the required Rm, Rp and total elongation T.EI are met in all cases. Example 19 provides a T.EI of 12,5 %, so this example shows that a DP1000 with the required strength and elongation can be provided by the invention.

Examples 20 and 21 show multiphase steel types having the same strength as a DP1000, but having a slightly lower total elongation.

Table 1 : Composition of 21 steel types in weight%, only N and B in ppm (parts per million).

Table 2: Process conditions and mechanical properties of the 21 steel types of Table 1.