EP3790999B1 - Variably rolled steel strip, sheet or blank and production method therefor - Google Patents
Variably rolled steel strip, sheet or blank and production method therefor Download PDFInfo
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- EP3790999B1 EP3790999B1 EP19721631.0A EP19721631A EP3790999B1 EP 3790999 B1 EP3790999 B1 EP 3790999B1 EP 19721631 A EP19721631 A EP 19721631A EP 3790999 B1 EP3790999 B1 EP 3790999B1
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Images
Classifications
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- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
- C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
- C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
- C21D8/0236—Cold rolling
-
- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
- C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
- C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
- C21D8/0226—Hot rolling
-
- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
- C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
- C21D8/0247—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
- C21D8/0273—Final recrystallisation annealing
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/001—Ferrous alloys, e.g. steel alloys containing N
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/12—Ferrous alloys, e.g. steel alloys containing tungsten, tantalum, molybdenum, vanadium, or niobium
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/14—Ferrous alloys, e.g. steel alloys containing titanium or zirconium
-
- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D2211/00—Microstructure comprising significant phases
- C21D2211/004—Dispersions; Precipitations
-
- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D2211/00—Microstructure comprising significant phases
- C21D2211/005—Ferrite
Definitions
- 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.
- TRBs tailor-rolled blanks
- 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%.
- 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.
- TRBs will be obtained via a blanking process.
- 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.
- Hot-rolled HSLA steels are used 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.
- 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.
- 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 ( ⁇ 650 °C). This increased coarsening behaviour of V precipitates promotes increased loss of precipitation strengthening with batch annealing.
- 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.
- 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.
- Nb and Ti are generally considered to be the most suitable micro-alloying elements in this case.
- 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.
- 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.
- DE 10 2005 031461 A1 discloses a variable cold rolled steel sheet and the corresponding method for producing a micro-alloyed strip with sections of varying thickness, starting from a hot rolled sheet, and batch annealing the cold rolled sheet.
- WO 2017/050790 relates to hot rolled high strength formable steel sheets for automotive applications, and discloses a steel sheet with similar composition.
- a variably rolled steel strip, sheet or blank consisting of the following elements (in wt%):
- 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.
- 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.
- 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.
- 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.
- 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 ( ⁇ 700 °C) is mitigated by grain refinement and hence increased grain-boundary or so-called Hall-Petch strengthening.
- 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.
- 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.
- 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.
- 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.
- 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.
- 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.
- 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.
- 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%.
- 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 between 0.30 and 0.60 wt.%.
- Mn Manganese
- Mn content should be at least 0.8 wt.%.
- a too high Mn content may lead to excessive segregation, which can impair hole-expansion capacity and promote delamination or splitting during shearing operations.
- 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.
- 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.%.
- S Sulphur
- 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.%.
- Al is added as a deoxidizer.
- a suitable minimum Al content is 0.01 wt.%.
- too high Al can be deleterious as it forms AIN particles during solidification of the molten steel, which can provoke surface issues during casting.
- 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.
- 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.
- 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.
- 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.
- 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.
- the degree of precipitation hardening is relatively limited compared to that of Ti.
- 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 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.
- 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.
- the driving force for the onset of recrystallization is increased as the pool of surviving dislocations at the start of recrystallization increases.
- 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.
- 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.
- T dis in °C the dissolution temperature
- a B 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
- T dis ° C A B ⁇ 10 log V 2 ⁇ 12 51 ⁇ 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 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 x O y -type of inclusions to reduce the risk of clogging and to improve cast ability of the steel during steel making.
- a too high Ca content can lead to excessive wear of the refractory lining in the installations of the steel-making plant
- a suitable maximum Ca content is 50 ppm, or more preferable maximum 35 ppm.
- a suitable minimum Ca content in the steel is 20 ppm.
- the Ca content in the steel is at most 20 ppm, or preferably at most 10 ppm, or most preferably at most 5 ppm.
- the variably rolled steel strip, sheet or blank contains Ti, and Mo and optionally Nb represented by weight percentage (wt.%) satisfying the equation of 0.8 ⁇ Nb 93 + Ti 48 Mo 96 ⁇ 2.2 or preferably 0.8 ⁇ Nb 93 + Ti 48 Mo 96 ⁇ 1.3
- the steel has a microstructure that consists of at least 60% recrystallized ferrite at 1 ⁇ 4 depth in the portion with high thickness, preferably at least 70% recrystallized ferrite at 1 ⁇ 4 depth in the portion with high thickness, more preferably at least 80% recrystallized ferrite at 1 ⁇ 4 depth in the portion with high thickness, most preferably 90% recrystallized ferrite at 1 ⁇ 4 depth in the portion with high thickness.
- the amount of recrystallized ferrite is an indication of the stability of the microstructure after batch annealing.
- 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.
- the steel in the portion with high thickness satisfies the equation A50 / t 0.2 ⁇ 16, preferably A50 / t 0.2 ⁇ 18, most preferably A50 / t 0.2 ⁇ 20, and wherein the steel in the portion with low thickness satisfies the equation: A50 / t 0.2 ⁇ 18, preferably A50 / t 0.2 ⁇ 20, most preferably A50 / t 0.2 ⁇ 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.
- 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.
- 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:
- 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 1150 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.
- 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.
- a strip according to the first aspect of the invention is obtained.
- 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 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.
- 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.
- 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.
- 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.
- 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.
- Figure 1 shows the time-temperature curve of the batch annealing cycle used in the examples.
- Examples are performed using laboratory cast ingots.
- Steels 1A to 1H 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.
- 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.
- ROT run-out-table
- 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.
- CR% cold-rolling reduction
- plates were wrapped in stainless steel foil and a protective H 2 atmosphere was used in the batch anneal furnace.
- 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).
- 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.
- EBSD Electron Back Scatter Diffraction
- 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 ⁇ m. To obtain a fully deformation free surface, the final polishing step was conducted with colloidal silica (OPS).
- OPS colloidal silica
- 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 ⁇ m 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 ⁇ 6 binning combined with standard background subtraction. The scan area was in all cases located at a position of 1 ⁇ 4 the sample thickness and care was taken to avoid as much as possible to include non-metallic inclusions in the scan area.
- TSL TexSEM Laboratories
- OIM Orientation Imaging Microscopy
- the EBSD scan size was in all cases 100 ⁇ 100 ⁇ m, with a step size of 0.1 ⁇ m, and a scan rate of approximately 100 frames per second. Fe( ⁇ ) 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 ⁇ 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 ⁇ 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).
- GTA Gram Tolerance Angle
- a pseudo-symmetry clean-up (GTA 5, axis ang 30@111) was applied.
- Hot-rolled steel sheets Table 1.3 gives the A50 tensile properties of the hot-rolled steel sheets coiled at 600 °C of steels 1Ato 1D (corresponding hot-rolled steel sheets labelled as 1A-HR600, 1B-HR600, 1C-HR600, 1D-HR600).
- Table 1.4 gives the A50 tensile properties of the hot-rolled steel sheets coiled at 540 °C of steels 1A to 1H 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.
- Tables 1.7 to 1.13 provide the fraction recrystallized ferrite (in %) and the average grain size (in ⁇ m) of the recrystallized ferrite based on EBSD measurements.
- Table 1.4 shows data for steels 1A to 1D 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 1B and 1D with - compared with steel 1A - increased content in Ti and Mo as well as for steel 1C 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.
- Tables 1.5 and 1.6 show the tensile data of hot-rolled steel sheets coiled at 600 and 540 °C, respectively, of steels 1A to 1H when subjected to a batch annealing with different values for T top and t hold 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.
- the Rp0.2 decreases after batch annealing with t hold of 3 or 10 hours. This decrease in Rp0.2 is seen for all steels, i.e., steels 1A to 1H, 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.
- 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%.
- 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.
- 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 are 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 ⁇ 14, and the fraction recrystallized ferrite at 1 ⁇ 4 depth is at least 50%.
- TRB element an element for TRB application
- 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 ⁇ 17.
- Table 1.12 shows an inventive example for a TRB application corresponding with NbTiMo-V alloyed steel 1G 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 1G-158BA to 1G-161BA).
- Table 1.1 Composition of steels (in wt.%) Steel C Mn Si P S Al Nb V Ti Mo N Ratio A* Ratio B* 1A 0.079 1.820 0.430 0.001 0.0015 0.065 0.037 ⁇ 0.001 0.096 0.150 0.0040 1.53 0.00 1B 0.105 1.820 0.430 ⁇ 0.001 0.0013 0.060 0.036 ⁇ 0.001 0.180 0.240 0.0060 1.65 0.00 1C 0.097 1.820 0.430 ⁇ 0.001 0.0013 0.060 0.036 0.150 0.120 0.240 0.0050 1.15 0.96 1D 0.097 1.810 0.430 ⁇ 0.001 0.0013 0.062 ⁇ 0.001 0.002 0.180 0.240 0.0060 1.50 0.02 1E 0.079 1.855 0.456 0.002 0.0020 0.074 0.041 0.005 0.118 0.154 0.0057 1.81 0.04 1F 0.097 1.861 0.462 0.002 0.0017 0.081 0.042 0.008 0.204
- Example 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.
- 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.
- 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.
- Tables 2.4 and 2.5 provide the fraction recrystallized ferrite (in %) and the average grain size (in ⁇ m) of the recrystallized ferrite based on EBSD measurements.
- 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.
- 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 ⁇ 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.
- Table 2.1 Composition of steels (in wt.%) Steel C Mn Si P S Al Nb V Ti Mo N Ratio A Ratio B 2A 0.082 1.849 0.454 0.011 0.0013 0.071 0.044 0.006 0.125 0.156 0.0042 1.89 0.05 2B 0.081 1.873 0.456 0.011 0.0014 0.068 0.046 0.005 0.129 0.147 0.0054 2.08 0.04 2C 0.105 1.858 0.462 0.013 0.0010 0.072 0.046 0.149 0.133 0.245 0.0053 1.28 0.88 2D 0.102 1.850 0.452 0.012 0.0005 0.069 0.046 0.147 0.133 0.150 0.0049 2.09 0.72 2E 0.094 1.877 0.466 0.011 0.0010 0.071 0.046 0.101 0.134 0.248 0.0050 1.27 0.85 2F 0.080 1.422 0.450 0.012 0.0011 0.061 0.045 0.100 0.136 0.151 0.0046 2.11 0.93 2
Description
- 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 (≥650 °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 (≥650 °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.
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DE 10 2005 031461 A1 discloses a variable cold rolled steel sheet and the corresponding method for producing a micro-alloyed strip with sections of varying thickness, starting from a hot rolled sheet, and batch annealing the cold rolled sheet. -
WO 2017/050790 relates to hot rolled high strength formable steel sheets for automotive applications, and discloses a steel sheet with similar composition. - 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.
- The invention is defined in the appended claims.
- 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
- 0.30 - 0.60 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
- 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 / t0.2 ≥ 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 / t0.2 ≥ 17,
- and 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 further contains cementite and/or pearlite, and wherein the precipitates in said microstructure consist of Ti, V, Mo and Nb.
- 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 (≥700 °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
- 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 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 AlxOy 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. The degree of precipitation hardening is relatively limited compared to that of Ti. 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
- 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 AlxOy-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.
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- With the amounts of Nb, Ti and Mo satisfying these equations, a suitable balance between these elements is provided.
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- Using such a balance between these elements provides an optimal balance between the elements that are essential for the invention.
- 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 / t0.2 ≥ 16, preferably A50 / t0.2 ≥ 18, most preferably A50 / t0.2 ≥ 20, and wherein the steel in the portion with low thickness satisfies the equation: A50 / t0.2 ≥ 18, preferably A50 / t0.2 ≥ 20, most preferably A50 / t0.2 ≥ 22, with A50 / t0.2 in the portion with low thickness equal to or higher than A50 / t0.2 in the portion with high thickness. The value of equation A50 / t0.2 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 1150 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, 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.
- 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 1150 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.
- 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.
- 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.
- 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. - Examples are performed using laboratory cast ingots.
- Steels 1A to 1H 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 1A to 1H (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 1G 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 (thold) 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 µm. 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 µm 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 × 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 × 100 µm, with a step size of 0.1 µm, and a scan rate of approximately 100 frames per second. Fe(α) 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 × 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 × 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@111) 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). - Hot-rolled steel sheets: Table 1.3 gives the A50 tensile properties of the hot-rolled steel sheets coiled at 600 °C of steels 1Ato 1D (corresponding hot-rolled steel sheets labelled as 1A-HR600, 1B-HR600, 1C-HR600, 1D-HR600). Table 1.4 gives the A50 tensile properties of the hot-rolled steel sheets coiled at 540 °C of steels 1A to 1H 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.
- Tables 1.7 to 1.13 provide the fraction recrystallized ferrite (in %) and the average grain size (in µm) of the recrystallized ferrite based on EBSD measurements.
- Table 1.4 shows data for steels 1A to 1D 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 1B and 1D with - compared with steel 1A - increased content in Ti and Mo as well as for steel 1C 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 1B, 1C, and 1D 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 1A to 1H when subjected to a batch annealing with different values for Ttop and thold 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 1A to 1H and coiled at 540 °C are subjected to a batch annealing with a Ttop of 675 or 700 °C for 3 or 10 hours thold, 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 thold.
- 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 thold of 3 or 10 hours. This decrease in Rp0.2 is seen for all steels, i.e., steels 1A to 1H, 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 thold 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%. Since thickness varies in a TRB sheet, the tensile elongation is normalised to 1 mm thickness by using the equation A50 / t0.2 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.
- 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
- 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 / t0.2 ≥ 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 / t0.2 ≥ 17. Table 1.12 shows an inventive example for a TRB application corresponding with NbTiMo-V alloyed steel 1G 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 1G-158BA to 1G-161BA).
- Most comparative examples in Tables 1.7 to 1.13 correspond with batch-annealed steels that either have A50 / t0.2 < 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 / t0.2 ≥ 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.%) Steel C Mn Si P S Al Nb V Ti Mo N Ratio A* Ratio B* 1A 0.079 1.820 0.430 0.001 0.0015 0.065 0.037 <0.001 0.096 0.150 0.0040 1.53 0.00 1B 0.105 1.820 0.430 <0.001 0.0013 0.060 0.036 <0.001 0.180 0.240 0.0060 1.65 0.00 1C 0.097 1.820 0.430 <0.001 0.0013 0.060 0.036 0.150 0.120 0.240 0.0050 1.15 0.96 1D 0.097 1.810 0.430 <0.001 0.0013 0.062 <0.001 0.002 0.180 0.240 0.0060 1.50 0.02 1E 0.079 1.855 0.456 0.002 0.0020 0.074 0.041 0.005 0.118 0.154 0.0057 1.81 0.04 1F 0.097 1.861 0.462 0.002 0.0017 0.081 0.042 0.008 0.204 0.241 0.0042 1.87 0.13 1G 0.103 1.863 0.459 0.001 0.0024 0.068 0.042 0.148 0.122 0.246 0.0062 1.17 0.84 1H 0.101 1.858 0.454 0.002 0.0023 0.073 0.003 0.009 0.202 0.245 0.0052 1.66 0.09 * with and and and 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. BA-675/3 BA-700/3 BA-740/3 BA-700/10 BA-740/10 Section ΔT/Δt (°C/min) T (°C) t (hrs) T (°C) t (hrs) T (°C) t (hrs) T (°C) t (hrs) T (°C) t (hrs) RT 0 RT 0 RT 0 RT 0 RT 0 1 1.19 500 6.8 500 6.8 500 6.8 500 6.8 500 6.8 2 0.50 620 10.8 620 10.8 620 10.8 620 10.8 620 10.8 3 0.27 675 14.2 700 15.7 740 18.2 700 15.7 740 18.2 4 0.00 675 17.2 700 18.7 740 21.2 700 25.7 740 28.2 5 -0.25 645 19.2 670 20.7 710 23.2 670 27.7 710 30.2 6 -0.92 450 22.7 450 24.7 450 27.9 450 31.7 450 34.9 7 -0.63 300 26.7 300 28.7 300 31.9 300 35.7 300 38.9 8 -0.42 200 30.7 200 32.7 200 35.8 200 39.7 200 42.8 Table 1.3: Tensile properties (longitudinal direction - A50 test piece geometry) of hot-rolled steels coiled at 600 °C Sheet (HR) Steel Gauge (mm) Rp0.2 (MPa) Rm (MPa) Ag (%) A50 (%) 1A-HR600 1A 3.65 800 867 10.4 18.1 1B-HR600 1B 3.72 870 929 9.6 16.9 1C-HR600 1C 3.84 967 1018 8.4 13.2 1D-HR600 1D 3.83 921 983 8.9 14.0 Table 1.4: Tensile properties (longitudinal direction - A50 test piece geometry) of hot-rolled steels coiled at 540 °C and difference (Δ) in Rp0.2 and Rm compared with steels with identical composition, but coiled at 600 °C (see Table 1.2). Sheet (HR) Steel Gauge (mm) Rp0.2 (MPa) Rm (MPa) Ag (%) A50 (%) ΔRp0.2* (MPa) ΔRm* (MPa) 1A-HR540 1A 3.75 749 847 8.6 15.3 -51 -20 1B-HR540 1B 3.92 679 889 8.6 14.8 -191 -40 1C-HR540 1C 3.80 799 953 6.9 13.0 -168 -65 1D-HR540 1D 3.90 784 932 7.5 13.9 -137 -51 1E-HR540 1E 3.66 705 844 7.9 14.2 - - 1F-HR540 1F 3.59 828 917 6.9 12.1 - - 1G-HR540 1G 3.72 777 916 7.3 12.9 - - 1H-HR540 1H 3.63 795 923 6.9 12.5 - - * with ΔRp0.2 and ΔRm defined as ΔRp0.2 = Rp0.2(CT 540 °C) - Rp0.2(CT 600 °C) and ΔRm = Rm(CT 540 °C) - Rm(CT 600 °C) 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 (Δ) in Rp0.2 and Rm compared with hot-rolled steels with identical composition but not subjected to batch annealing (see Table 1.2) Sheet (BA) Sheet (HR) T (°C) t (hrs) Rp0.2 (MPa) Rm (MPa) Ag (%) A50 (%) ΔRp0.2* (MPa) ΔRm* (MPa) 1A-1 BA 1A-HR600 675 3 748 817 10.7 19.0 -52 -50 1B-8BA 1B-HR600 675 3 784 848 9.0 15.8 -86 -81 1C-15BA 1C-HR600 675 3 896 957 8.4 14.3 -71 -61 1D-22BA 1D-HR600 675 3 866 929 9.6 16.8 -55 -54 * with ΔRp0.2 and ΔRm defined as ΔRp0.2 = Rp0.2(BA) - Rp0.2(HR) and ΔRm = Rm(BA) - Rm(HR) 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 (Δ) in Rp0.2 and Rm compared with hot-rolled steels with identical composition but not subjected to batch annealing (see Table 1.3) Sheet (BA) Steel T (°C) t (hrs) Rp0.2 (MPa) Rm (MPa) Ag (%) A50 (%) ΔRp0.2* (MPa) ΔRm* (MPa) 1A-29BA 1A 675 3 782 829 7.3 14.2 +33 -20 1A-57BA 700 3 775 835 7.0 14.0 +26 -12 1A-85BA 740 3 655 774 8.2 15.6 -94 -73 1E-113BA 1E 700 10 799 845 7.5 13.9 +94 +1 1E-141BA 740 10 650 755 7.5 13.5 -55 -89 1B-36BA 1B 675 3 827 878 7.0 15.6 +148 -11 1B-64BA 700 3 854 907 7.4 13.8 +175 +18 1B-92BA 740 3 676 788 8.2 13.3 -3 -101 1F-120BA 1F 700 10 877 911 6.7 12.7 +49 -6 1F-148BA 740 10 727 807 7.0 13.1 -101 -110 1C-43BA 1C 675 3 948 986 6.4 14.7 +149 +33 1C-71BA 700 3 925 967 6.9 13.0 +126 +14 1C-99BA 740 3 756 840 7.0 14.1 -43 -113 1G-127BA 1G 700 10 832 897 6.7 12.1 +55 -19 1G-169BA 740 7 764 843 6.7 11.4 -13 -73 1G-176BA 740 8 754 845 6.0 12.6 -23 -71 1G-183BA 740 9 746 837 6.6 13.6 -31 -79 1G-155BA 740 10 704 809 6.5 12.3 -73 -107 1D-50BA 1D 675 3 917 964 6.8 13.7 +133 +32 1D-78BA 700 3 892 937 5.9 11.9 +108 +5 1D-106BA 740 3 740 832 7.9 12.8 -44 -100 1H-134BA 1H 700 10 827 900 7.2 13.3 +32 -23 1H-162BA 740 10 724 814 7.8 14.4 -71 -109 * with ΔRp0.2 and ΔRm defined as ΔRp0.2 = Rp0.2(BA) - Rp0.2(HR) and ΔRm = Rm(BA) - Rm(HR) 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 BA settings Cold rolling Structure Tensile properties 30-60 Fit** Application*** Steel Sheet (HR) Sheet (BA) T (°C) t (hrs) CR (%) t (mm) fRex (%) GSRex (µm) Rp0.2 (MPa) Rm (MPa) Ag (%) A50 (%) A50* (%) a (-) b (-) TRB Example 1A 1A-HR600 1A-1BA 675 3 0 3.70 -- -- 748 817 7.5 14.1 10.9 -- -- X Comparative 1A-2BA 10 3.23 <50 -- 871 894 7.1 13.6 10.8 -- -- X Comparative 1A-3BA 20 2.83 <50 -- 925 937 5.1 10.5 8.5 -- -- X Comparative 1A-4BA 30 2.51 <50 -- 945 956 6.2 9.8 8.2 -153 +1480 X Comparative 1A-5BA 40 2.16 <50 -- 936 955 7.2 9.3 8.0 X Comparative 1A-6BA 50 1.78 <50 -- 908 934 7.4 9.3 8.3 X Comparative 1A-7BA 60 1.51 <50 -- 831 872 7.5 14.1 13.0 X Comparative 1B 1B-HR600 1B-8BA 675 3 0 3.70 -- -- 784 848 9.0 15.8 12.2 -- -- X Comparative 1B-9BA 10 3.24 <50 -- 940 961 8.0 15.1 11.9 -- -- X Comparative 1B-10BA 20 2.83 <50 -- 1008 1014 5.4 11.8 9.6 -- -- X Comparative 1B-11BA 30 2.49 <50 -- 1038 1034 5.3 9.8 8.2 -24 +1128 X Comparative 1B-12BA 40 2.14 <50 -- 1051 1051 3.7 7.1 6.1 X Comparative 1B-13BA 50 1.78 <50 -- 1049 1054 5.9 9.7 8.6 X Comparative 1B-14BA 60 1.50 <50 -- 1017 1031 7.1 11.4 10.5 X Comparative 1C 1C-HR600 1C-15BA 675 3 0 3.81 -- -- 896 957 8.4 14.3 10.9 -- -- X Comparative 1C-16BA 10 3.25 <50 -- 1023 1036 7.3 13.6 10.7 -- -- X Comparative 1C-17BA 20 2.88 <50 -- 1069 1073 7.1 12.7 10.3 -- -- X Comparative 1C-18BA 30 2.51 <50 -- 1079 1080 2.3 3.3 2.7 -226 +1858 X Comparative 1C-19BA 40 2.14 <50 -- 1036 1049 7.3 11.0 9.4 X Comparative 1C-20BA 50 1.78 <50 -- 987 1009 8.0 12.4 11.0 X Comparative 1C-21BA 60 1.44 <50 -- 918 952 8.1 11.2 10.4 X Comparative 1D 1D-HR600 1D-22BA 675 3 0 3.2 -- -- 994 1007 7.3 14.1 11.2 -- -- X Comparative 1D-23BA 10 2.85 <50 -- 1027 1035 6.5 12.2 9.9 -- -- X Comparative 1D-24BA 20 2.48 <50 -- 1032 1042 5.8 7.1 5.9 -- -- X Comparative 1D-25BA 30 2.14 <50 -- 1004 1024 7.4 11.1 9.5 -185 +1672 X Comparative 1D-26BA 40 1.79 <50 -- 957 987 7.0 11.0 9.8 X Comparative 1D-27BA 50 1.43 <50 -- 902 942 8.0 11.6 10.8 X Comparative 1D-28BA 60 3.2 <50 -- 994 1007 7.3 14.1 11.2 X Comparative * A50* is the re-calculated A50 value corresponding with a 1 mm thickness (t in mm) according to A50 / t 0.2 ** Fit of the evolution of Rp0.2 as a function of CR% between 30% and 60%: Rp0.2(CR%) = a × ln(CR%) + b with a and b as fit parameters *** Markings: "O" implies suitable as TRB element and "X" means not suitable as TRB element 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 BA settings Cold rolling Structure Tensile properties 30-60 Fit** Application*** Steel Sheet (HR) Sheet (BA) T (°C) t (hrs) CR (%) t (mm) fRex (%) GSRex (µm) Rp0.2 (MPa) Rm (MPa) Ag (%) A50 (%) A50* (%) a (-) b (-) TRB Example 1A 1A-HR540 1A-29BA 675 3 0 3.74 -- -- 782 829 7.3 14.2 10.9 -- -- X Comparative 1A-30BA 10 3.18 <50 -- 935 945 6.9 13.4 10.6 -- -- X Comparative 1A-31BA 20 2.83 <50 -- 978 980 4.7 12.0 9.7 -- -- X Comparative 1A-32BA 30 2.48 <50 -- 977 982 4.5 8.6 7.2 -357 +2190 X Comparative 1A-33BA 40 2.12 <50 -- 870 895 8.3 13.5 11.6 X Comparative 1A-34BA 50 1.80 <50 -- 788 831 9.5 14.4 12.8 X Comparative 1A-35BA 60 1.44 <50 -- 731 778 10.3 15.1 14.0 X Comparative 1B 1B-HR540 1B-36BA 675 3 0 3.92 -- -- 827 878 7.0 15.6 11.9 -- -- X Comparative 1B-37BA 10 3.25 <50 -- 1035 1031 3.9 8.8 7.0 -- -- X Comparative 1B-38BA 20 2.84 <50 -- 1053 1053 3.2 7.4 6.0 -- -- X Comparative 1B-39BA 30 2.48 <50 -- 1051 1057 5.7 8.1 6.8 210 +1769 X Comparative 1B-40BA 40 2.15 <50 -- 985 1003 7.2 11.3 9.7 X Comparative 1B-41BA 50 1.78 <50 -- 978 1000 7.0 11.1 9.9 X Comparative 1B-42BA 60 1.47 <50 -- 889 917 8.5 12.2 11.3 X Comparative 1C 1C-HR540 1C-43BA 675 3 0 3.79 -- -- 948 986 6.4 14.7 11.3 -- -- X Comparative 1C-44BA 10 3.20 <50 -- 1077 1090 -- -- -- -- -- X Comparative 1C-45BA 20 2.83 <50 -- 1107 1108 2.1 7.0 5.7 -- -- X Comparative 1C-46BA 30 2.50 <50 -- 1048 1052 6.7 11.5 9.6 342 +2252 X Comparative 1C-47BA 40 2.10 <50 -- 1066 1073 7.1 12.8 11.0 X Comparative 1C-48BA 50 1.77 <50 -- 911 931 9.0 14.0 12.5 X Comparative 1C-49BA 60 1.44 <50 -- 826 848 10.7 15.5 14.4 X Comparative 1D 1D-HR540 1D-50BA 675 3 0 3.95 -- -- 917 964 6.8 13.7 10.4 -- -- X Comparative 1D-51BA 10 3.22 <50 -- 1052 1046 3.6 8.6 6.8 -- -- X Comparative 1D-52BA 20 2.84 <50 -- 1096 1090 3.0 7.4 6.0 -- -- X Comparative 1D-53BA 30 2.49 <50 -- 1056 1062 4.6 6.0 5.0 -280 +2022 X Comparative 1D-54BA 40 2.16 <50 -- 1013 1028 7.0 11.4 9.8 X Comparative 1D-55BA 50 1.78 <50 -- 929 955 5.8 9.5 8.5 X Comparative 1D-56BA 60 1.43 <50 -- 865 894 7.7 8.3 7.7 X Comparative * A50* is the re-calculated A50 value corresponding with a 1 mm thickness (t in mm) according to A50 / t 0.2 ** Fit of the evolution of Rp0.2 as a function of CR% between 30% and 60%: Rp0.2(CR%) = a × ln(CR%) + b with a and b as fit parameters *** Markings: "O" implies suitable as TRB element and "X" means not suitable asTRB element 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 BA settings Cold rolling Structure Tensile properties 30-60 Fit** App lication*** Steel Sheet (HR) Sheet (BA) T (°C) t (hrs) CR (%) t (mm) fRex (%) GSRex (µm) Rp0.2 (MPa) Rm (MPa) Ag (%) A50 (%) A50* (%) a (-) b (-) TRB Example 1A 1A-HR540 1A-57BA 700 3 0 3.65 -- 775 835 7.0 14.0 10.8 - - X Comparative 1A-58BA 10 3.18 -- 864 888 6.9 12.9 10.2 - - X Comparative 1A-59BA 20 2.82 -- 873 907 7.2 13.2 10.7 - - X Comparative 1A-60BA 30 2.47 -- 793 840 8.3 10.7 8.9 -251 +1639 X Comparative 1A-61BA 40 2.15 - 703 785 7.8 13.7 11.8 X Comparative 1A-62BA 50 1.79 -- 654 737 11.5 17.3 15.4 O Comparative 1A-63BA 60 1.42 -- 619 695 13.4 20.9 19.5 O 1B 1B-HR540 1B-64BA 700 3 0 3.84 -- 854 907 7.4 13.8 10.5 - - X Comparative 1B-65BA 10 3.19 -- 989 992 6.6 13.1 10.4 - - X Comparative 1B-66BA 20 2.84 - 1019 1031 5.2 7.4 6.0 - - X Comparative 1B-67BA 30 2.47 -- 793 840 8.3 10.7 8.9 116 +1227 X Comparative 1B-68BA 40 2.13 - 870 914 8.5 12.3 10.6 X Comparative 1B-69BA 50 1.78 - 764 814 10 13.7 12.2 X Comparative 1B-70BA 60 1.43 -- 729 765 11.7 19.1 17.8 O Comparative 1C 1C-HR540 1C-71BA 700 3 0 3.79 -- 925 967 6.9 13.0 10.0 - - X Comparative 1C-72BA 10 3.22 -- 1051 1052 1.5 2.2 1.7 - - X Comparative 1C-73BA 20 2.87 - 1047 1047 1.1 1.6 1.3 - - X Comparative 1C-74BA 30 2.48 - 1022 1036 6.0 9.5 7.9 112 +2413 X Comparative 1C-75BA 40 2.13 -- 882 918 8.5 13.5 11.6 X Comparative 1C-76BA 50 1.77 - 790 829 10.1 13.1 11.7 X Comparative 1C-77BA 60 1.45 - 740 782 9.6 13.2 12.3 X Comparative 1D 1D-HR540 1D-78BA 700 3 0 3.86 -- 892 937 5.9 11.9 9.1 - - X Comparative 1D-79BA 10 3.19 -- 995 1005 5.2 9.6 7.6 - - X Comparative 1D-80BA 20 2.83 - 1024 1032 6.3 11.5 9.3 - - X Comparative 1D-81BA 30 2.49 -- 964 999 6.5 11.1 9.2 -319 +2032 X Comparative 1D-82BA 40 2.11 -- 829 887 8.5 13.3 11.5 X Comparative 1D-83BA 50 1.77 -- 769 834 8.7 12.4 11.1 X Comparative 1D-84BA 60 1.44 -- 745 807 10.5 13.5 12.6 X Comparative * A50* is the re-calculated A50 value corresponding with a 1 mm thickness (t in mm) according to A50 / t0.2 ** Fit of the evolution of Rp0.2 as a function of CR% between 30% and 60%: Rp0.2(CR%) = a × ln(CR%) + b with a and b as fit parameters *** Markings: "O" implies suitable as TRB element and "X" means not suitable as TRB element 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 BA settings Cold rolling Structure Tensile properties 30-60 Fit** App lication*** Steel Sheet (HR) Sheet (BA) T (°C) t (hrs) CR (%) t (mm) fRex (%) GSRex (µm) Rp0.2 (MPa) Rm (MPa) Ag (%) A50 (%) A50* (%) a (-) b (-) TRB Example 1A 1A-HR540 1A-85BA 740 3 0 3.49 -- 655 774 8.2 15.6 12.1 - - X Comparative 1A-86BA 10 3.19 -- 722 812 6.7 13.5 10.7 - - X Comparative 1A-87BA 20 2.87 - 669 813 7.9 13.4 10.9 - - X Comparative 1A-88BA 30 2.48 - 600 746 8.0 12.9 10.8 154 +1113 X Comparative 1A-89BA 40 2.14 - 529 689 9.8 14.7 12.6 X Comparative 1A-90BA 50 1.81 - 503 658 13.2 19.1 17.0 O Comparative 1A-91BA 60 1.42 -- 494 646 14.1 21.9 20.4 O 1B 1B-HR540 1B-92BA 740 3 0 3.75 -- 676 788 8.2 13.3 10.2 - - X Comparative 1B-93BA 10 3.20 - 800 881 7.6 14.2 11.3 - - X Comparative 1B-94BA 20 2.85 -- 738 847 7.9 13.1 10.6 - - X Comparative 1B-95BA 30 2.47 - 650 778 9.3 14.6 12.2 -117 +1026 X Comparative 1B-96BA 40 2.13 -- 563 699 12.8 20.4 17.5 O Comparative 1B-97BA 50 1.77 -- 555 681 13.7 22.1 19.7 O 1B-98BA 60 1.43 - 570 676 14.3 22.3 20.8 O 1C 1C-HR540 1C-99BA 740 3 0 3.75 -- 756 840 7.0 14.1 10.8 - - X Comparative 1C-100BA 10 3.21 -- 813 922 7.0 12.1 9.6 - - X Comparative 1C-101BA 20 2.85 - 818 900 6.4 10.7 8.7 - - X Comparative 1C-102BA 30 2.48 -- 699 816 5.7 6.2 5.2 -177 +1294 X Comparative 1C-103BA 40 2.15 -- 628 720 12.1 19.7 16.9 O Comparative 1C-104BA 50 1.76 -- 594 712 11.8 20.2 18.0 O 1C-105BA 60 1.43 -- 577 694 13.0 19.7 18.3 O 1D 1D-HR540 1D-106BA 740 3 0 3.72 -- 740 832 7.9 12.8 9.8 - - X Comparative 1D-107BA 10 3.18 -- 841 909 6.9 13.7 10.9 - - X Comparative 1D-108BA 20 2.85 - 802 890 6.7 12.1 9.8 - - X Comparative 1D-109BA 30 2.47 -- 748 853 7.1 12.6 10.5 -215 +1470 X Comparative 1D-110BA 40 2.16 -- 665 782 9.0 12.8 11.0 X Comparative 1D-111BA 50 1.79 -- 618 725 10.0 17.5 15.6 O Comparative 1D-112BA 60 1.43 -- 602 708 12.8 20.0 18.6 O * A50* is the re-calculated A50 value corresponding with a 1 mm thickness (t in mm) according to A50 / t0.2 ** Fit of the evolution of function of CR% between 30% and 60%: b with and b fit Rp0.2 as a Rp0.2(CR%) = a × ln(CR%) + a as parameters *** Markings: "O" implies suitable as TRB element and "X" means not suitable as TRB element 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 BA settings Cold rolling Structure Tensile properties 30-60 Fit** App lication*** Steel Sheet (HR) Sheet (BA) T (°C) t (hrs) CR (%) t (mm) fRex (%) GSRex (µm) Rp0.2 (MPa) Rm (MPa) Ag (%) A50 (%) A50* (%) a (-) b (-) TRB Example 1E 1E-HR540 1E-113BA 700 10 0 3.26 54 -- 799 845 7.5 13.9 11.0 -- -- X Comparative 1E-114BA 10 3.17 35 -- 868 885 6.7 13.0 10.3 -- -- X Comparative 1E-115BA 20 2.83 1 3.88 864 895 6.7 11.4 9.3 -- -- X Comparative 1E-116BA 30 2.47 12 2.82 783 856 7.1 14.4 12.0 309 +1840 X Comparative 1E-117BA 40 2.13 19 2.73 712 805 9.5 14.6 12.6 X Comparative 1E-118BA 50 1.75 61 2.70 619 724 11.2 16.5 14.8 O Comparative 1E-119BA 60 1.51 72 2.77 576 678 12.9 19.0 17.5 O 1F 1F-HR540 1F-120BA 700 10 0 3.33 45 -- 877 911 6.7 12.7 10.0 -- -- X Comparative 1F-121BA 10 3.16 24 -- 891 925 6.5 12.0 9.5 -- -- X Comparative 1F-122BA 20 2.80 0 -- 920 953 6.1 11.2 9.1 -- -- X Comparative 1F-123BA 30 2.49 13 2.29 853 913 7.0 11.4 9.5 -296 +1876 X Comparative 1F-124BA 40 2.10 42 2.62 814 890 7.5 12.5 10.8 X Comparative 1F-125BA 50 1.77 39 2.55 717 808 9.1 12.3 11.0 X Comparative 1F-126BA 60 1.41 70 2.55 654 754 11.4 17.3 16.2 O Comparative 1G 1G-HR540 1G-127BA 700 10 0 3.46 33 -- 832 897 6.7 12.1 9.4 -- -- X Comparative 1G-128BA 10 3.22 18 -- 933 959 5.7 6.4 5.1 -- -- X Comparative 1G-129BA 20 2.85 0 -- 941 977 6.9 12.5 10.1 -- -- X Comparative 1G-130BA 30 2.45 27 2.84 840 916 8.0 13.2 11.0 305 +1862 X Comparative 1G-131BA 40 2.13 40 2.49 721 831 8.9 12.9 11.1 X Comparative 1G-132BA 50 1.74 57 2.43 652 759 11.2 16.5 14.8 O Comparative 1G-133BA 60 1.41 81 2.48 634 724 11.9 16.3 15.2 O 1H 1H-HR540 1H-134BA 700 10 0 3.38 46 -- 827 900 7.2 13.3 10.4 -- -- X Comparative 1H-135BA 10 3.16 32 -- 921 952 6.4 12.0 9.5 -- -- X Comparative 1H-136BA 20 3.47 0 -- 958 979 5.9 10.9 8.5 -- -- X Comparative 1H-137BA 30 2.47 4 2.43 913 957 6.1 10.7 8.9 -366 +2167 X Comparative 1H-138BA 40 2.10 25 2.63 835 906 6.8 11.9 10.3 X Comparative 1H-139BA 50 1.74 58 2.49 729 819 8.2 11.5 10.3 X Comparative 1H-140BA 60 1.41 55 2.46 665 761 10.2 14.4 13.4 X Comparative * A50* is the re-calculated A50 value corresponding with a 1 mm thickness (t in mm) according to A50 / t0.2 ** Fit of the evolution of Rp0.2 as a function of CR% between 30% and 60%: Rp0.2(CR%) = a × ln(CR%) + b with a and b as fit parameters *** Markings: "O" implies suitable as TRB element and "X" means not suitable as TRB element 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 BA settings Cold rolling Structure Tensile properties 30-60 Fit** App lication*** Steel Sheet (HR) Sheet (BA) T (°C) t (hrs) CR (%) t (mm) fRex (%) GSRex (µm) Rp0.2 (MPa) Rm (MPa) Ag (%) A50 (%) A50* (%) a (-) b (-) TRB Example 1E 1E-HR540 1E-141BA 740 10 0 3.45 53 - 650 755 7.5 13.5 10.5 - - X Comparative 1E-142BA 10 3.30 35 - 688 781 7.7 14.7 11.6 - - X Comparative 1E-143BA 20 2.84 0 - 724 805 7.6 13.4 10.9 - - X Comparative 1E-144BA 30 2.53 27 2.81 567 716 7.7 12.7 10.5 117 +955 X Comparative 1E-145BA 40 2.13 85 2.78 512 658 7.7 14.8 12.7 X Comparative 1E-146BA 50 1.79 91 2.41 481 624 13.6 20.5 18.2 O Comparative 1E-147BA 60 1.40 92 2.39 492 629 14.3 21.6 20.2 O 1F 1F-HR540 1F-148BA 740 10 0 3.43 54 - 727 807 7.0 13.1 10.2 - - X Comparative 1F-149BA 10 3.20 29 - 767 844 6.3 9.8 7.8 - - X Comparative 1F-150BA 20 2.87 0 - 746 840 7.3 11.9 9.6 - - X Comparative 1F-151BA 30 2.49 20 3.62 663 793 8.3 12.1 10.1 -174 +1254 X Comparative 1F-152BA 40 2.12 78 2.99 630 760 9.6 14.5 12.5 X Comparative 1F-153BA 50 1.74 87 2.32 548 689 11.9 17.5 15.7 O Comparative 1F-154BA 60 1.40 80 2.62 558 694 12.3 18.4 17.2 O 1G 1G-HR540 1G-155BA 740 10 0 3.52 53 - 704 809 7.2 12.3 9.6 - - X Comparative 1G-156BA 10 3.20 34 - 762 862 7.5 14.0 11.1 - - X Comparative 1G-157BA 20 2.83 4 3.49 670 804 9.0 14.3 11.6 - - X Comparative 1G-158BA 30 2.49 50 3.34 479 667 11.5 17.3 14.4 +48 +317 O Inventive 1G-159BA 40 2.11 73 3.18 497 659 13.1 19.0 16.4 O 1G-160BA 50 1.77 76 3.04 505 658 13.0 18.8 16.8 O 1G-161BA 60 1.44 100 2.79 513 649 13.1 18.5 17.2 O 1H 1H-HR540 1H-162BA 740 10 0 3.39 35 - 724 814 7.8 14.4 11.3 - - X Comparative 1H-163BA 10 3.19 28 - 749 836 7.1 13.8 10.9 - - X Comparative 1H-164BA 20 2.80 8 2.29 775 854 7.6 12.4 10.1 - - X Comparative 1H-165BA 30 2.44 30 3.58 683 797 7.6 12.1 10.1 207 +138 X Comparative 1H-166BA 40 2.09 71 3.03 611 745 9.2 13.5 11.6 X Comparative 1H-167BA 50 1.79 69 2.88 558 693 10.2 12.1 10.8 X Comparative 1H-168BA 60 1.42 88 2.37 544 671 13.2 20.4 19.0 O Comparative * A50* is the re-calculated A50 value corresponding with a 1 mm thickness (t in mm) ccording to A50 / t0.2 ** Fit of the evolution of Rp0.2 as a function of CR% between 30% and 60%: Rp0.2(CR%) = a × ln(CR%) + b with a and b as fit parameters *** Markings: "O" implies suitable as TRB element and "X" means not suitable as TRB element 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 BA settings Cold rolling Structure Tensile properties 30-60 Fit** Application*** Steel Sheet (HR) Sheet (BA) T (°C) t (hrs) CR (%) t (mm) fRex (%) GSRex (µm) Rp0.2 (MPa) Rm (MPa) Ag (%) A50 (%) A50* (%) a (-) b (-) TRB Example 1G 1G-HR540 1G-169BA 740 7 0 3.51 34 -- 764 843 7.1 11.4 8.9 -- -- X Comparative 1G-170BA 10 3.23 23 -- 749 865 7.5 13.3 10.5 -- -- X Comparative 1G-171BA 20 2.82 20 -- 835 908 7.0 12.3 10.0 -- -- X Comparative 1G-172BA 30 2.45 35 3.12 746 849 6.5 10.4 8.7 -281 +1697 X Comparative 1G-173BA 40 2.11 54 2.88 652 785 8.5 12.7 10.9 X Comparative 1G-174BA 50 1.75 605 733 10.8 16.1 14.4 O Comparative 1G-175BA 60 1.45 546 681 11.2 17.0 15.8 O 1G 1G-HR540 1G-176BA 740 8 0 3.63 48 -- 754 845 6.9 12.6 9.7 -- -- X Comparative 1G-177BA 10 3.16 20 -- 781 871 6.4 11.0 8.7 -- -- X Comparative 1G-178BA 20 2.83 20 3.53 807 892 7.1 11.4 9.3 -- -- X Comparative 1G-179BA 30 2.44 19 3.17 639 799 7.0 11.0 9.2 165 +1198 X Comparative 1G-180BA 40 2.10 58 2.92 595 748 8.9 13.2 11.4 X Comparative 1G-181BA 50 1.76 60 2.68 540 688 10.6 14.6 13.0 X Comparative 1G-182BA 60 1.44 74 2.10 532 670 12.5 17.8 16.5 O Comparative 1G 1G-HR540 1G-183BA 740 9 0 3.56 -- 746 837 7.5 13.6 10.5 -- -- X Comparative 1G-184BA 10 3.18 -- 786 873 7.6 13.2 10.5 -- -- X Comparative 1G-185BA 20 2.84 798 879 5.2 8.6 7.0 -- -- X Comparative 1G-186BA 30 2.45 714 844 7.3 10.8 9.0 754 +1544 X Comparative 1G-187BA 40 2.11 561 717 9.4 12.5 10.8 X Comparative 1G-188BA 50 1.78 539 681 11.5 14.1 12.6 X Comparative 1G-189BA 60 1.47 537 672 12.8 17.7 16.4 O Comparative * A50* is the re-calculated A50 value corresponding with a 1 mm thickness (t in mm) according to A50 / t0.2 ** Fit of the evolution of Rp0.2 as a function of CR% between 30% and 60%: Rp0.2(CR%) = a × ln(CR%) + b with a and b as fit parameters *** Markings: "O" implies suitable as TRB element and "X" means not suitable as TRB element - 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.
- (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.
- Tables 2.4 and 2.5 provide the fraction recrystallized ferrite (in %) and the average grain size (in µm) of the recrystallized ferrite based on EBSD measurements.
- 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
- 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 / t0.2 ≥ 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 / t0.2 ≥ 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 / t0.2 < 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 / t0.2 ≥ 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.%) Steel C Mn Si P S Al Nb V Ti Mo N Ratio A Ratio B 2A 0.082 1.849 0.454 0.011 0.0013 0.071 0.044 0.006 0.125 0.156 0.0042 1.89 0.05 2B 0.081 1.873 0.456 0.011 0.0014 0.068 0.046 0.005 0.129 0.147 0.0054 2.08 0.04 2C 0.105 1.858 0.462 0.013 0.0010 0.072 0.046 0.149 0.133 0.245 0.0053 1.28 0.88 2D 0.102 1.850 0.452 0.012 0.0005 0.069 0.046 0.147 0.133 0.150 0.0049 2.09 0.72 2E 0.094 1.877 0.466 0.011 0.0010 0.071 0.046 0.101 0.134 0.248 0.0050 1.27 0.85 2F 0.080 1.422 0.450 0.012 0.0011 0.061 0.045 0.100 0.136 0.151 0.0046 2.11 0.93 2G 0.125 1.848 0.460 0.012 0.0009 0.065 0.064 0.145 0.197 0.297 0.0039 1.55 1.01 with and and and Table 2.2: Tensile properties (longitudinal direction - A50 geometry) of hot-rolled steels coiled at 600 or 540 °C Sheet (HR) Steel CT (°C) Gauge (mm) Rp0.2 (MPa) Rm (MPa) Ag (%) A50 (%) 2A-HR600 2A 600 3.52 852 922 9.8 18.3 2B-HR600 2B 600 3.51 802 861 9.8 18.2 2C-HR540 2C 540 3.33 736 994 7.1 13.2 2D-HR540 2D 540 3.44 775 960 7.8 14.8 2E-HR540 2E 540 3.60 753 973 7.6 14.3 2F-HR540 2F 540 3.33 812 978 7.1 12.8 2G-HR540 2G 540 3.35 687 982 7.0 13.7 Table 2.3: Tensile properties (longitudinal direction) of batch annealed steels coiled at 540 °C (no intermediate cold rolling) and difference (Δ) in Rp0.2 and Rm compared with identical steels but not subjected to batch annealing (see Table 2.2) Sheet (BA) Sheet (HR) T (°C) t (hrs) Rp0.2 (MPa) Rm (MPa) Ag (%) A50 (%) ΔRp0.2* (MPa) ΔRm* (MPa) 2A-1 BA 2A-HR600 675 16 537 684 9.4 17.9 -315 -238 2B-8BA 2B-HR600 720 16 473 630 11.6 21.6 -329 -231 - 2C-HR540 700 16 -- -- -- -- -- -- 2C-20BA 710 16 634 773 7.9 13.9 -102 -221 - 18 -- -- -- -- -- -- 2C-32BA 720 16 612 763 9.2 17.5 -124 -231 2C-39BA 18 609 763 9.1 17.7 -127 -231 2C-46BA 740 10 549 747 8.5 16.9 -187 -247 2C-53BA 12 562 763 9.2 16.3 -174 -231 2D-60BA 2D-HR540 700 16 631 724 8.6 15.8 -144 -236 2D-67BA 720 16 613 704 9.7 18.0 -162 -256 2E-74BA 2E-HR540 700 16 648 737 8.8 15.9 -105 -236 2E-81BA 720 16 617 708 9.1 16.3 -136 -265 2F-88BA 2F-HR540 675 16 700 758 8.7 16.2 -112 -220 2F-95BA 700 16 641 702 9.0 16.8 -171 -276 2F-102BA 720 16 604 663 8.8 15.3 -208 -315 2G-109BA 2G-HR540 740 10 535 766 9.6 18.0 -152 -216 2G-116BA 16 544 772 8.5 13.6 -143 -210 * with ΔRp0.2 and ΔRm defined as ΔRp0.2 = Rp0.2(BA) - Rp0.2(HR) and ΔRm = Rm(BA) - Rm(HR) 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 BA settings Cold rolling Structure Tensile properties 30-60 Fit** Application*** Steel Sheet (HR) Sheet (BA) T (°C) t (hrs) CR (%) t (mm) fRex (%) GSRex (µm) Rp0.2 (MPa) Rm (MPa) Ag (%) A50 (%) A50* (%) a (-) b (-) TRB Example 2A 2A-HR600 2A-1BA 675 16 0 3.48 76 -- 537 684 9.4 15.2 11.8 -- -- X Comparative 2A-2BA 10 3.16 60 -- 635 735 8.0 13.6 10.8 -- -- X Comparative 2A-3BA 20 2.82 3 4.19 701 770 8.0 14.4 11.7 -- -- X Comparative 2A-4BA 30 2.46 17 3.09 657 769 8.6 13.9 11.6 161 +1203 X Comparative 2A-5BA 40 2.09 37 2.69 603 727 9.6 15.3 13.2 X Comparative 2A-6BA 50 1.76 70 2.28 577 695 10.0 13.2 11.8 X Comparative 2A-7BA 60 1.41 77 2.21 542 670 12.0 17.3 16.2 O Comparative 2B 2B-HR600 2B-8BA 720 16 0 3.51 77 -- 473 630 11.6 21.6 16.8 -- -- O Comparative 2B-9BA 10 3.18 58 -- 549 673 8.8 15.3 12.1 -- -- X Comparative 2B-10BA 20 2.81 9 5.46 565 690 7.0 11.9 9.7 -- -- X Comparative 2B-11BA 30 2.48 74 4.89 464 636 8.9 14.6 12.2 -19 +508 X Comparative 2B-12BA 40 2.11 77 3.84 418 605 12.7 20.0 17.2 O Comparative 2B-13BA 50 1.73 89 2.96 406 571 14.3 22.4 20.1 O 2B-14BA 60 1.39 90 2.73 461 611 14.9 22.1 20.7 O * A50* is the re-calculated A50 value corresponding with a 1 mm thickness (t in mm) according to A50 / t0.2 ** Fit of the evolution of Rp0.2 as a function of CR% between 30% and 60%: Rp0.2(CR%) = a × ln(CR%) + b with a and b as fit parameters *** Markings: "O" implies suitable as TRB element and "X" means not suitable as TRB element Table 2.5: Tensile properties (longitudinal direction) of cold-rolled and batch-annealed steel sheets coiled at 540 °C BA settings Cold rolling Structure Tensile properties 30-60 Fit* App lication*** Steel Sheet (HR) Sheet (BA) T (°C) t (hrs) CR (%) t (mm) fRex (%) GSRex (µm) Rp0.2 (MPa) Rm (MPa) Ag (%) A50 (%) A** (%) a (-) b (-) TRB Example 2C 2C-HR540 -- 700 16 0 -- -- -- -- -- -- -- -- -- -- -- -- -- 10 -- -- -- -- -- -- -- -- -- -- -- -- 2C-15BA 20 2.81 26 -- 740 846 7.9 12.8 10.4 -- -- X Comparative 2C-16BA 30 2.42 64 2.99 541 724 10.7 17.3 14.5 +60 +327 O Inventive 2C-17BA 40 2.09 81 1.95 538 680 12.1 18.1 15.6 O 2C-18BA 50 1.75 96 1.88 553 660 14.0 22.7 20.3 O 2C-19BA 60 1.42 96 1.87 585 661 14.2 22.5 21.0 O 2C 2C-HR540 2C-20BA 710 16 0 3.45 -- -- 634 773 7.9 13.9 10.9 -- -- X Comparative 2C-21BA 10 3.16 -- -- 669 751 8.6 14.8 11.8 -- -- X Comparative 2C-22BA 20 2.79 5 2.94 641 818 8.3 13.5 11.0 -- -- X Comparative 2C-23BA 30 2.44 65 3.06 521 682 11.8 18.0 15.1 +26 +425 O Inventive 2C-24BA 40 2.12 91 2.67 512 660 12.9 19.7 17.0 O 2C-25BA 50 1.77 96 1.92 527 647 14.1 22.1 19.7 O 2C-26BA 60 1.40 97 1.79 538 649 14.7 23.1 21.6 O 2C 2C-HR540 -- 710 18 0 -- -- -- -- -- -- -- -- -- -- -- -- -- 10 -- -- -- -- -- -- -- -- -- -- -- -- 2C-27BA 20 2.80 20 4.08 586 763 8.6 12.9 10.5 -- -- X Comparative 2C-28BA 30 2.45 89 3.80 431 649 12.8 20.0 16.7 +172 -161 O Inventive 2C-29BA 40 2.10 84 2.49 469 651 13.2 20.3 17.5 O 2C-30BA 50 1.75 92 2.23 504 648 13.7 22.0 19.7 O 2C-31BA 60 1.42 98 1.80 554 647 14.7 22.9 21.3 O 2C 2C-HR540 2C-32BA 720 16 0 3.42 52 -- 612 763 9.2 17.5 13.7 -- -- X Comparative 2C-33BA 10 3.16 28 -- 675 811 8.3 14.4 11.4 -- -- X Comparative 2C-34BA 20 2.77 37 6.06 641 803 8.2 13.2 10.8 -- -- X Comparative 2C-35BA 30 2.44 66 3.12 485 689 11.0 17.1 14.3 +98 +137 O Inventive 2C-36BA 40 2.08 88 2.37 482 653 13.3 20.8 18.0 O 2C-37BA 50 1.76 -- -- 502 643 14.1 22.6 20.2 O 2C-38BA 60 1.40 -- -- 559 648 14.4 23.3 21.8 O 2C 2C-HR540 2C-39BA 720 18 0 3.42 47 -- 609 763 9.1 16.5 12.9 -- -- X Comparative 2C-40BA 10 3.20 28 -- 647 794 8.6 16.1 12.8 -- -- X Comparative 2C-41BA 20 2.81 7 3.67 649 807 8.7 15.7 12.8 -- -- X Comparative 2C-42BA 30 2.46 70 3.63 524 729 10.7 16.8 14.0 +36 +372 O Inventive 2C-43BA 40 2.13 85 2.48 456 647 13.3 22.7 19.5 O 2C-44BA 50 1.78 94 2.60 505 639 14.6 23.9 21.3 O 2C-45BA 60 1.42 100 1.78 544 640 14.7 22.9 21.3 O 2C 2C-HR540 2C-46BA 740 10 0 3.35 58 -- 549 747 9.3 16.9 13.3 -- -- X Comparative 2C-47BA 10 3.17 44 -- 604 795 8.8 15.6 12.4 -- -- X Comparative 2C-48BA 20 2.81 21 3.77 442 700 11.0 17.6 14.3 -- -- X Comparative 2C-49BA 30 2.49 79 2.62 394 620 14.0 19.5 16.2 +185 +246 O Inventive 2C-50BA 40 2.12 94 2.45 427 610 15.2 25.0 21.5 O 2C-51BA 50 1.78 91 2.19 473 615 15.3 24.2 21.6 O 2C-52BA 60 1.42 99 1.88 524 629 15.2 22.5 21.0 O 2C 2C-HR540 2C-53BA 740 12 0 3.26 45 -- 562 763 9.2 16.3 12.9 -- -- X Comparative 2C-54BA 10 3.17 38 -- 596 784 9.2 16.4 13.0 -- -- X Comparative 2C-55BA 20 2.81 26 5.22 495 737 10.3 16.6 13.5 -- -- X Comparative 2C-56BA 30 2.47 71 3.08 405 646 11.8 21.0 17.5 +164 -172 O Inventive 2C-57BA 40 2.10 96 2.53 401 612 14.6 22.8 19.7 O 2C-58BA 50 1.77 97 2.10 473 613 15.5 26.1 23.3 O 2C-59BA 60 1.41 100 1.81 513 622 15.6 23.5 21.9 O 2D 2D-HR540 2D-60BA 700 16 0 3.47 57 -- 631 724 8.6 15.8 12.3 -- -- X Comparative 2D-61BA 10 3.18 32 -- 692 763 7.5 12.9 10.2 -- -- X Comparative 2D-62BA 20 2.82 9 7.05 545 687 9.9 15.7 12.8 -- -- X Comparative 2D-63BA 30 2.45 86 3.26 466 601 13.5 22.5 18.8 +125 +45 O Inventive 2D-64BA 40 2.11 99 2.38 515 598 14.6 24.8 21.4 O 2D-65BA 50 1.75 100 2.14 531 609 14.4 22.9 20.5 O 2D-66BA 60 1.42 99 1.99 555 619 14.9 24.0 22.4 O 2D 2D-HR540 2D-67BA 720 16 0 3.50 -- -- 613 704 9.6 17.9 13.9 -- -- X Comparative 2D-68BA 10 3.18 -- -- 636 724 8.8 14.7 11.7 -- -- X Comparative 2D-69BA 20 2.79 95 4.10 355 544 14.4 23.3 19.0 -- -- O Inventive 2D-70BA 30 2.45 82 4.99 407 575 13.3 21.2 17.7 +150 -99 O 2D-71BA 40 2.12 96 2.68 458 575 15.2 24.5 21.1 O 2D-72BA 50 1.75 100 2.45 483 564 17.0 28.7 25.7 O 2D-73BA 60 1.40 100 1.98 513 577 17.0 25.6 23.9 O 2E 2E-HR540 2E-74BA 700 16 0 3.44 48 -- 648 737 8.8 15.9 12.4 -- -- X Comparative 2E-75BA 10 3.15 34 -- 704 784 8.5 15.6 12.4 -- -- X Comparative 2E-76BA 20 2.81 16 5.03 652 760 7.9 12.4 10.1 -- -- X Comparative 2E-77BA 30 2.45 57 4.89 496 651 11.1 17.0 14.2 +69 +251 O Inventive 2E-78BA 40 2.11 89 2.92 495 618 13.3 22.7 19.6 O 2E-79BA 50 1.77 98 2.19 505 613 14.2 22.1 19.7 O 2E-80BA 60 1.38 100 1.80 550 619 14.7 23.0 21.6 O 2E 2E-HR540 2E-81BA 720 16 0 3.56 -- -- 617 708 9.1 16.3 12.6 -- -- X Comparative 2E-82BA 10 3.14 -- -- 595 706 9.3 15.4 12.2 -- -- X Comparative 2E-83BA 20 2.78 96 4.06 385 598 9.9 14.7 12.0 -- -- X Comparative 2E-84BA 30 2.46 100 3.38 377 559 15.1 24.8 20.7 +230 -411 O Inventive 2E-85BA 40 2.12 100 2.73 428 565 15.8 25.6 22.0 O 2E-86BA 50 1.77 100 2.10 496 584 15.6 24.7 22.0 O 2E-87BA 60 1.41 100 1.90 532 598 15.5 23.6 22.0 O 2F 2F-HR540 2F-88BA 675 16 0 3.51 -- -- 700 758 8.7 16.2 12.6 -- -- X Comparative 2F-89BA 10 3.13 -- -- 802 836 7.2 13.0 10.3 -- -- X Comparative 2F-90BA 20 2.81 3 4.98 776 822 8.4 13.5 11.0 -- -- X Comparative 2F-91BA 30 2.45 36 3.57 670 741 8.4 12.9 10.8 -125 +1079 X Comparative 2F-92BA 40 2.11 62 3.57 594 667 11.4 16.8 14.5 O Comparative 2F-93BA 50 1.76 100 2.80 580 647 12.3 19.1 17.1 O 2F-94BA 60 1.41 96 2.39 584 640 13.7 21.2 19.8 O 2F 2F-HR540 2F-95BA 700 16 0 3.54 52 -- 641 702 9.0 16.8 13.0 -- -- X Comparative 2F-96BA 10 3.15 33 -- 687 738 8.0 13.7 10.9 -- -- X Comparative 2F-97BA 20 2.80 39 5.66 616 694 9.3 14.4 11.7 -- -- X Comparative 2F-98BA 30 2.45 64 4.80 578 664 11.1 17.6 14.7 -92 +872 O Inventive 2F-99BA 40 2.11 81 3.44 503 601 11.6 17.8 15.3 O 2F-100BA 50 1.75 92 2.98 505 580 14.4 24.0 21.5 O 2F-101BA 60 1.41 94 2.89 513 579 15.2 23.8 22.2 O 2F 2F-HR540 2F-102BA 720 16 0 3.25 -- -- 604 663 8.8 15.3 12.1 -- -- X Comparative 2F-103BA 10 3.13 -- -- 639 686 7.6 12.9 10.3 -- -- X Comparative 2F-104BA 20 2.80 34 5.94 610 684 10.1 15.2 12.4 -- -- X Comparative 2F-105BA 30 2.44 70 5.36 467 578 12.1 19.4 16.2 +29 +359 O Inventive 2F-106BA 40 2.09 92 4.31 451 556 14.0 21.0 18.1 O 2F-107BA 50 1.73 93 3.52 464 549 15.1 24.4 21.9 O 2F-108BA 60 1.40 98 2.95 488 554 15.5 25.1 23.5 O 2G 2G-HR540 2G-109BA 740 10 0 3.53 535 766 9.5 17.9 13.9 -- -- X Comparative 2G-110BA 10 3.16 503 755 7.9 12.6 10.0 -- -- X Comparative 2G-111BA 20 2.82 408 702 10.3 17.1 13.9 -- -- X Comparative 2G-112BA 30 2.46 373 654 12.6 19.9 16.6 +163 -185 O Inventive 2G-113BA 40 2.13 410 623 14.9 23.3 20.0 O 2G-114BA 50 1.77 445 624 15.4 24.5 21.9 O 2G-115BA 60 1.43 488 632 15.4 23.6 22.0 O 2G 2G-HR540 2G-116BA 740 16 0 3.41 -- -- 544 772 8.5 13.6 10.6 -- -- X Comparative 2G-117BA 10 3.17 -- -- 589 808 8.9 14.8 11.8 -- -- X Comparative 2G-118BA 20 2.80 61 5.28 385 691 9.7 14.9 12.1 -- -- X Comparative 2G-119BA 30 2.46 93 3.65 374 658 12.6 19.6 16.4 +139 -97 O Inventive 2G-120BA 40 2.10 91 2.29 408 634 14.7 23.7 20.4 O 2G-121BA 50 1.77 97 2.26 465 631 15.0 23.9 21.3 O 2G-122BA 60 1.39 100 2.03 460 617 16.1 25.8 24.2 O * A50* is the re-calculated A50 value corresponding with a 1 mm thickness (t in mm) according to A50 / t0.2 ** Fit of the evolution of Rp0.2 as a function of CR% between 30% and 60%: Rp0.2(CR%) = a × ln(CR%) + b with a and b as fit parameters *** Markings: "O" implies suitable as TRB element and "X" means not suitable as TRB element
Claims (13)
- Variably cold rolled steel strip, sheet or blank, consisting of the following elements (in wt%):0.05 - 0.20 C0.30 - 0.60 Si0.80 - 2.50 Mn0.01 - 0.10 Al0.07 - 0.25 Ti0.10 - 0.35 V0.05 - 0.40 Mo0.02 - 0.10 Nboptionally 0.01 - 0.80 Crat most 0.06 Pat most 0.01 Sat most 0.01 Nat most 0.005 Cathe 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 / t0.2 ≥ 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 / t0.2 ≥ 17, and 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 further contains cementite and/or pearlite, and wherein the precipitates in said microstructure consist of Ti, V, Mo and Nb.
- 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 C0.90 - 2.30 Mn0.02 - 0.09 Al0.08 - 0.22 Ti0.12 - 0.30 V0.08 - 0.35 Mo0.03 - 0.09 Nboptionally 0.01 - 0.60 Crat most 0.04 Pat most 0.005 Sat most 0.008 Nat most 0.003 Ca.
- 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 C1.20 - 2.00 Mn0.04 - 0.08 Al0.10 - 0.20 Ti0.13 - 0.25 V0.10 - 0.30 Mo0.03 - 0.08 Nboptionally 0.01 - 0.40 Crat most 0.02 Pat most 0.003 S0.002 - 0.007 Nat most 0.001 Ca.
- Variably rolled steel strip, sheet or blank according to any one of the preceding claims, wherein Nb is present in an amount between 0.03 and 0.09%, preferably between 0.03 and 0.08%.
- 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.
- 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.
- 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 / t0.2 ≥ 16, preferably A50 / t0.2 ≥ 18, most preferably A50 / t0.2 ≥ 20, and wherein the steel in the portion with low thickness satisfies the equation: A50 / t0.2 ≥ 18, preferably A50 / t0.2 ≥ 20, most preferably A50 / t0.2 ≥ 22, with A50 / t0.2 in the portion with low thickness equal to or higher than A50 / t0.2 in the portion with high thickness.
- 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 1150 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, 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.
- Method according to claim 10, 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.
- Method according to claim 10 or 11, 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.
- Method according to claim 10,11 or 12, wherein the hot-rolled steel strip or sheet is coiled in the temperature range between 480 and 560° C.
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PCT/EP2019/061654 WO2019215131A1 (en) | 2018-05-08 | 2019-05-07 | Variably rolled steel strip, sheet or blank and production method therefor |
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