Method for manufacturing cold-formable high-strength steel strip and steel strip

Abstract

1. A method for producing steel strip, comprising the steps of: The process of casting molten steel into slabs; reheating the slab to a temperature of at least 1150°C for at least one hour; hot rolling the steel into strip, preferably with an average entry temperature of the F1 slabs above 1000°C; Coiling of hot rolled steel strip; Steel strips under the following conditions: transformation interval temperatures (i.e., Ac1 to Ac3), preferably temperatures below 700°C; Non-oxidizing and non-nitrifying atmosphere; A total annealing time of at least 5 hours, preferably at least 10 hours to enrich the austenite with Mn such that the Mn content of the austenite is at least 1.25 times the bulk Mn content of the steel, and to enrich the austenite with C such that the C content of the austenite is at least 1.2 times the bulk C content of the steel; Cooling the steel after batch annealing by air, forced air or water quenching A method comprising:

Description

Technical Field

[0001] The present invention relates to a method for manufacturing cold-formable coated or uncoated high-strength steel strips for manufacturing steel products and such steel strips.

Background Art

[0002] Cold forming or cold stamping or cold press forming of steel sheets is a method for manufacturing steel parts for various applications in the manufacturing industry, such as automobiles, construction, engineering, infrastructure, etc. It is known that as the strength of the steel sheet increases, the cold formability of the steel sheet decreases. This particularly applies to the first generation of advanced high-strength steels (AHSS) in addition to conventional steel sheets.

[0003] Due to this inverse correlation between the strength and elongation of steel, the characteristic maps of the tensile strength and elongation of these steels are sometimes called "banana diagrams". However, by adopting the concept of so-called second-generation AHSS (2GAHSS), high strength and high formability can be achieved, but these steels are usually highly alloyed and contain expensive alloying elements. An example is twinning-induced plasticity (TWIP) steel with a high manganese content, which is a stainless steel with an Mn content usually higher than 12 wt% and contains large amounts of expensive alloying elements such as chromium, nickel, molybdenum, etc. In addition to being very expensive, another drawback of 2GAHSS is that due to its very high alloy content, it is very difficult to manufacture on a large industrial scale.

[0004] These problems of 2GAHSS are overcome, but in order to achieve high strength and moderately high cold formability, various concepts of third-generation AHSS (3GAHSS), such as quenching and partitioning (Q&P) steels, carbide-free bainite (CFB) steels, and medium Mn steels, have been introduced. These steels are less expensive than 2GAHSS and can be easily processed in existing facilities of steel mills. The present invention focuses on the medium Mn type of 3GAHSS.

[0005] WO16001887 describes a method for manufacturing a high-strength steel sheet, wherein the steel contains 0.1 ≦ C ≦ 0.4, 4.2 ≦ Mn ≦ 8, 1 ≦ Si ≦ 3, 0.2 ≦ Mo ≦ 0.5 in weight %, and the balance is Fe and inevitable impurities, and the method includes a step of continuous annealing above Ac3, a step of quenching to a temperature between martensite start (Ms) and finish (Mf), a step of overaging for more than 10 seconds at 300 to 500 °C, and a step of cooling. This is essentially a quenching and partitioning (Q&P) process, and the C enrichment (and optionally Mn enrichment) of austenite is achieved by the overaging process of a steel containing a certain amount of martensite. The Q&P process is significantly different from the intercritical annealing treatment. In this document, at lower overaging temperatures (300 to 500 °C), the diffusion of Mn in the steel is very slow, so substantial partitioning of Mn into austenite is not expected.

[0006] WO2017021464 describes a high-strength steel in the form of a hot- or cold-rolled strip having the following chemical composition (in wt%): C: 0.005 - 0.6; Mn: 4 - 10; Al: 0.005 - 4; Si: 0.005 - 2; P: 0.001 - 0.2; S: max 0.05; N: 0.001 - 0.3, with the balance being iron and inevitable element inclusions related to steel, which is annealed optionally if flexibly hot-rolled, flexibly cold-rolled, optionally annealed, further flexibly cold-rolled, and then annealed at an annealing temperature of 600°C - 750°C for 1 minute to 48 hours. This patent applies flexible rolling by controlling the roll gap where the shearing conditions vary across the entire width of the strip. Flexible rolling is a different process for parts with varying thicknesses and is different from conventional rolling processes like the present invention where a uniform thickness of the product is obtained across the entire width. The drawback of flexible rolling strips is that non-uniform properties are obtained across the entire width of the strip.

Summary of the Invention

[0007] An object of the present invention is to provide a steel strip that can be highly cold-formed with a cold-rolled thickness under coating or non-coating conditions while maintaining high strength.

[0008] Another object of the present invention is to provide a steel strip that can be highly cold-formed with a hot-rolled thickness range under coating or non-coating conditions while maintaining high strength.

[0009] Both the hot-rolled and cold-rolled steel strips of the present invention have high energy absorption capacity, that is, high impact resistance, are spot weldable, and have resistance to hydrogen embrittlement.

Brief Description of the Drawings

[0010]

Figure 1

DETAILED DESCRIPTION OF THE INVENTION

[0011] The present invention is, firstly, a method for manufacturing a cold-rolled and annealed steel strip, comprising: The steel composition is, by weight, C: 0.05 to 0.3; Mn: 3.0 to 12.0; Al: 0.03 to 3.0; Optionally, the following additional alloying elements: Si: less than 1.5 Cr: less than 2.0 V: less than 0.1 Nb: less than 0.1 Ti: less than 0.1 Mo: less than 0.5 one or more of; inevitable impurities (e.g., S: less than 30 ppm, P: less than 0.04); Fe: the balance and the method comprising the following steps: casting the molten steel into slabs; reheating the slabs and holding at a temperature of 1150 °C or higher for 1 hour or more; hot rolling the steel into strips, preferably with an average inlet temperature of the F1 slab exceeding 1000 °C; coiling the hot-rolled steel strip; pickling the steel strip; intermediate batch annealing the steel strip at a temperature of less than 650 °C for 24 hours or more (for longer than 24 hours), and obtaining at least 60% by volume of ferrite after cooling to room temperature; cold rolling the steel into a cold-rolled steel strip and coiling it; The coiled steel strip is subjected to the following conditions: An intercritical temperature in the Ac1 - Ac3 transformation range that is less than 700 °C; A non - oxidizing and non - nitriding atmosphere; A total annealing time during which the strip is maintained at the intercritical temperature of at least 5 hours, preferably at least 10 hours In a batch annealing step, enriching Mn in austenite such that the Mn content of the austenite is at least 1.25 times the bulk Mn content of the steel, and enriching C in austenite such that the C content of the austenite is at least 1.2 times the bulk C content of the steel; Cooling the steel after batch annealing by air, forced air or water quenching is embodied in the said method.

[0012] Steel containing 0.05 to 0.3 wt% C, 3.0 to 12.0 wt% Mn, 0.03 to 3.0 wt% Al, and optionally other alloying elements and inevitable impurities is processed into a hot-rolled gauge using a specific processing route. The molten steel is cast into slabs, and then the slabs are reheated at a temperature of 1150 °C or higher for 1 hour or more. Next, the slabs are hot-rolled into strips, preferably hot-rolled into strips at a finishing inlet temperature (F1) exceeding 1000 °C. The F1 inlet temperature is the inlet temperature of the strip at the first stand of the finishing rolling mill. The finishing rolling mill is part of the hot rolling mill. After rough rolling or breakdown rolling of the slabs in the roughing mill and before the run-out table is cooled, finishing rolling is carried out. After passing through the run-out table, the hot-rolled strip is wound into coils. When these coils are intercritically annealed at a temperature below 650 °C for at least 24 hours, at least 60% by volume of ferrite is obtained in the steel strip after cooling to room temperature. Next, the steel strip is pickled in an acidic solution, for example, at a temperature of 50 to 90 °C, and cold-rolled to a thinner gauge. The present invention is not limited by the range of the hot-rolled or cold-rolled gauge. However, usually, the hot-rolled gauge is 2 to 10 mm, and the cold-rolled gauge is 0.5 to 2 mm. Then, the cold-rolled steel is batch annealed at a transformation range temperature below 700 °C, preferably below 660 °C, in a non-oxidizing and non-nitriding atmosphere for at least 5 hours, preferably at least 10 hours, so that the Mn content of the intercritical austenite reaches at least 1.25 times the bulk Mn content of the steel, and the C content of the intercritical austenite reaches at least 1.2 times the bulk C content. Since a large amount of Mn can be distributed from ferrite to austenite during annealing, a longer time of 10 hours is preferred. Since Mn is a large substitutional alloying element in iron, it usually takes a long time for diffusion. As the temperature of the batch annealing decreases, the manganese enrichment in the intercritical austenite increases, which results in the austenite in the steel becoming more stable after cooling to room temperature following the batch annealing.The batch annealing time is defined as the time during which the steel strip is maintained at the batch annealing temperature, excluding the time for heating the steel strip to the target temperature.

[0013] The last batch annealing, i.e., the batch annealing of the coiled strip in this document, is carried out for a time as described in the claims, which is long enough to obtain a relatively equiaxed ferrite grain morphology in the steel. The ratio of the grain length to the width is preferably 3 or less. Then, the steel is cooled to room temperature at any cooling rate, for example, by air, forced air or water.

[0014] By implementing the method according to the present invention, the following advantages can be obtained.

[0015] · Medium-Mn steel containing the above alloying elements and 3 to 12 wt% Mn reduces the segregation of Mn. The segregation of Mn is a concern that affects the mechanical properties when Mn is present in a relatively large amount as in the present invention. To minimize segregation and uniformly disperse Mn in the matrix, a relatively high slab reheating temperature of 1150 °C or higher, preferably 1200 °C or higher, more preferably 1250 °C or higher, and a minimum reheating time of 1 hour are selected. Otherwise, the final mechanical properties may be impaired. The selection of the reheating temperature depends on the Mn content of the alloy. When the Mn content of the alloy is close to the lower limit of the Mn range described in the claims, a reheating temperature close to 1150 °C is sufficient to make the Mn distribution uniform, and as the Mn content increases, a higher slab reheating temperature is preferred.

[0016] · The steel can be hot-rolled on an industrial scale with a moderately wide strip width, for example, exceeding 1000 mm. This is achieved by maintaining a high F1 temperature above 1000 °C during hot rolling to keep the required hot rolling load low. When the F1 temperature is low, the hot rolling of the steel strip becomes difficult.

[0017] · The steel is suitable for cold rolling on an industrial scale. This is made possible by using an intermediate batch annealing process for the hot-rolled steel. The intermediate batch annealing is carried out at a temperature in the transformation range of the steel, preferably less than 650 °C, and is selected such that at least 60% by volume of ferrite is obtained in the steel strip, with the balance being retained austenite and martensite.

[0018] · (Additional) batch annealing of the coiled steel strip at less than 700 °C actually produces an appropriate microstructure. This needs to be at least 5 hours, preferably at least 10 hours. During this processing step, the Mn and C of the steel of the present invention are partitioned between the austenite and ferrite in the transformation range, so that the austenite is highly enriched with Mn and C and stabilizes the phase down to room temperature. The austenite is enriched such that the Mn content is at least 1.25 times the bulk Mn content of the steel and the C content is at least 1.2 times the bulk C content of the steel, so that the steel becomes substantially less sensitive to the actual cooling rate, and thus the steel can be cooled with air, forced air or water after batch annealing. The lower the batch annealing temperature is below 700 °C, the richer the Mn content of the austenite in the transformation range becomes.

[0019] A second embodiment of the present invention is a method for manufacturing a hot-rolled and annealed steel strip, wherein the steel composition is, by weight, C: 0.05 to 0.3; Mn: 3.0 to 12.0; Al: 0.03 to 3.0; Optionally, the following additional alloying elements: Si: less than 1.5 Cr: less than 2.0 V: less than 0.1 Nb: less than 0.1 Ti: less than 0.1 Mo: less than 0.5 one or more of; inevitable impurities (e.g., S: less than 30 ppm, P: less than 0.04); Fe: the balance and the method comprising the following steps: casting molten steel into slabs; reheating the slabs at a temperature of 1150 °C or higher for 1 hour or longer; hot rolling the steel into strips, preferably with an average inlet temperature of the F1 slabs exceeding 1000 °C; hot rolling Prolonged strip Steel coiling the strips; pickling the steel strips; batch annealing the coiled steel strips under the following conditions: a transformation interval temperature of less than 700 °C, between Ac1 and Ac3; a non-oxidizing and non-nitriding atmosphere; a total annealing time during which the strips are maintained at the transformation interval temperature of at least 5 hours, preferably at least 10 hours, and enriching Mn in the austenite such that the Mn content in the austenite is at least 1.25 times the bulk Mn content of the steel, and enriching C in the austenite such that the C content in the austenite is at least 1.2 times the bulk C content of the steel; cooling the steel after batch annealing by air, forced air or water quenching is the said method.

[0020] As described above, the steel processed according to the first embodiment up to the hot rolling step is then pickled and then directly fed to the final batch annealing step, skipping the intermediate steps. The batch annealing is carried out according to the claims for a time long enough to obtain a relatively equiaxed ferrite grain morphology in the steel, where the ratio of the length to the width of the grains is preferably 3 or less.

[0021] Thus, the steel is produced as a hot rolled strip instead of a cold rolled strip, but has all the advantages regarding the mechanical properties as a cold rolled strip under the first embodiment.

[0022] The present invention is also embodied in a method of reheating a slab at a temperature of 1200 °C or higher. Thereby, good homogenization of Mn in the as-cast steel slab is achieved and its segregation is reduced.

[0023] The present invention is also embodied in a method of reheating a slab at a temperature of 1250 °C or higher. Thereby, a further reduction in the microsegregation of Mn present in the as-cast steel slab is achieved.

[0024] The present invention is also embodied in a method of batch annealing a coiled steel strip at a transformation range temperature below 660 °C. Thereby, a high enrichment of Mn in the transformation range austenite is achieved, and as a result, the martensite content in the final microstructure can be minimized. The present invention is also embodied in a method in which the obtained strip is coated with any metal coating applied by hot-dip galvanizing, galvanizing, electro-galvanizing, aluminum plating, or other methods (e.g., physical vapor deposition (PVD), chemical vapor deposition (CVD)). Thereby, the corrosion resistance and excellent aesthetic appearance of the steel strip required during application or service are achieved.

[0025] In one embodiment of the method according to the claims, excluding claim 2, the obtained steel strip is subjected to skin pass rolling, also known as temper rolling. Thereby, the cold formability is further improved.

[0026] In one embodiment, the thickness reduction by skin pass rolling is 5% or less. Thereby, the yield point elongation of the steel strip in the tensile test is minimized, and the cold formability and aesthetic appearance of the cold-formed steel strip are improved.

[0027] The present invention is also a steel strip that can be manufactured or is manufactured according to the method of the first or second embodiment of the present invention, wherein the steel strip has the following steel composition by weight%: C: 0.05 to 0.3; Mn: 3.0 to 12.0; Al: 0.03 to 3.0; Optionally, the following additional alloying elements: Si: less than 1.5 Cr: less than 2.0 V: less than 0.1 Nb: less than 0.1 Ti: less than 0.1 Mo: less than 0.5 one or more of; inevitable impurities (e.g., S: less than 30 ppm, P: less than 0.04); Fe: the balance having, In order to obtain metastable retained austenite in the steel strip so as to give the steel a high strain hardening exponent of at least 0.3 measured in a strain range of 7% after yield point elongation in a quasi-static tensile test, the composition of the retained austenite in the steel strip has a Mn content that is at least 1.4 times the bulk Mn content of the steel composition and a C content that is at least 2.3 times the bulk C content of the steel composition, The microstructure after final batch annealing of the coiled steel strip is, in volume %, Ferrite: 30 to 70%; Retained austenite: 20 to 65%; Martensite: less than 20% (including 0 volume%); and is embodied in the steel strip.

[0028] The steel strip has Mn and C enriched to a certain level in the metastable austenite of the final microstructure of the steel under the final use conditions, resulting in a high work hardening rate or strain hardening rate.

[0029] Furthermore, according to the present invention, the cold-rolled or hot-rolled steel strip preferably has a microstructure in which there are 20 to 65% by volume of retained austenite, 30 to 70% by volume of ferrite, and martensite present in less than 20% by volume (including 0% by volume). The ferrite is preferably ultrafine with a crystal grain size of 0.2 to 2 μm. By subjecting the coiled steel strip to sufficiently long batch annealing, this ultrafine ferrite acquires a substantially equiaxed shape with a ratio of grain length / width of 3 or less. In contrast to typical short-time (minutes rather than hours) continuous annealing, which results in elongated grains with a high aspect ratio (length / width ratio), sufficient recrystallization of the ferrite grains occurs during the long batch annealing of the cold-rolled strip used in the present invention.

[0030] Other embodiments of the present invention are in accordance with claims 10 to 14, which result in high mechanical properties and high cold formability achieved in the steel strip when manufactured according to the present invention. These properties are (biaxial) stretchability, bendability, hole expansion performance, yield strength, maximum tensile strength, total elongation, and yield point elongation.

[0031] The present invention is also embodied in a steel strip that has been subjected to skin pass rolling, also known as temper rolling, as described above.

[0032] The present invention is based on the modification of the steel composition and processing it using all the steps described to obtain an optimized microstructure. Due to the obtained microstructure, the cold-rolled and / or hot-rolled steel strip has high cold formability and high mechanical properties.

[0033] The essential elements of steel are Mn, C, and Al. Mn and C are austenite stabilizing elements in steel, and thus, they are added to steel in a predetermined amount to stabilize austenite. Al is a ferrite stabilizing factor and widens the temperature range of Ac1 to Ac3 (Ac1 = the temperature at which austenite transformation starts during heating, Ac3 = the temperature at which austenite transformation is completed during heating). Al is added to enhance the robustness of steel for industrial processing because it reduces the sensitivity of steel to undesirable small temperature changes during the treatment in the transformation range. The present invention is not limited to the presence of other optional elements and inevitable impurities existing in the steel. The range of these optional inevitable alloying elements is shown in the relevant claims.

[0034] A high Mn content (3 - 12 wt.%) in steel causes a large amount of Mn enrichment in austenite during batch annealing of hot-rolled or cold-rolled steel strips in the as-coiled state. This Mn enrichment, together with C enrichment (since C is also an austenite stabilizing element), enhances the thermal stability of austenite in the transformation range by suppressing the Ms temperature of the steel (Ms = the temperature at which martensite transformation starts during cooling). Therefore, during the cooling of the batch-annealed steel strip to room temperature, austenite in the transformation range does not transform much into martensite, and as a result, a large amount of austenite (20% by volume or more) can be retained in the room-temperature microstructure of the steel. The retained austenite with optimal mechanical stability transforms into martensite during loading (during forming or any other deformation), causing the transformation-induced plasticity (TRIP) effect. Due to the TRIP effect that increases the work hardening rate or strain hardening rate, high strength, high elongation, and high cold formability are achieved in the steel strip of the present invention. An Mn content exceeding 12 wt.% makes continuous casting of steel difficult due to extreme segregation and changes the mechanism of plasticity enhancement from TRIP to TWIP (TWIP = twinning-induced plasticity), and a content less than 3 wt.% does not result in sufficient Mn enrichment in austenite and does not achieve a sufficient amount of retained austenite in the room-temperature microstructure.

[0035] Similar to the effect of Mn described above, C is also partitioned into transformation range austenite during final batch annealing, enhancing the thermal stability of austenite and causing stabilization of austenite in the room temperature microstructure. However, C is effective in a smaller amount than Mn. Therefore, the range of C content for modifying the steel chemistry in the present invention is 0.05 to 0.3 wt%. When the C content is less than 0.05 wt%, a sufficient austenite stabilization effect cannot be obtained, and a C content exceeding 0.3 wt% makes post-treatment of the manufactured strip after cold forming, for example, spot welding, difficult. Sol Joint Since welding is essential when assembling automotive parts to the vehicle body, it is very important to consider this aspect. C is also added to the steel of the present invention to increase strength.

[0036] Aluminum is not an austenite stabilizing element of steel but a ferrite stabilizing element. However, it is added to the steel up to 3 wt% to expand the transformation range temperature range (Ac1 - Ac3) of the steel. Due to the high level of Mn, the steel becomes sensitive to slight variations in the processing temperature in industrial scale processing. By adding Al, the robustness of the steel processing is ensured, and as a result, the batch annealing temperature of the steel strip can be selected with slight variations to achieve the desired microstructure. If Al is not intentionally added to the steel as an alloying element (i.e., when the Al content is about 0.03 wt% because it is often added as a deoxidizer to molten steel), a more accurate furnace needs to be used, but the present invention still functions. To suppress the formation of oxide scale during hot rolling and reduce the rolling load during hot and cold rolling, the maximum amount of Al is limited to 3 wt%.

[0037] The combination of the steel composition and the process steps of the method brings about the beneficial effects of the present invention. Mn is an essential alloying element for changing the chemical composition of steel. When its content exceeds about 2% by weight, it tends to segregate after casting. This can affect the product performance by resulting in non-uniform properties and may also cause cracks during the processing steps. Therefore, it is preferable that the casting slab is sufficiently homogenized. Good homogenization of the slab is achieved by using a relatively high slab reheating temperature of 1150 °C or higher, preferably 1200 °C or higher, more preferably 1250 °C or higher, for a sufficiently long time, preferably 60 minutes or more.

[0038] Subsequently, due to the relatively high alloy content of the steel, the rolling load during hot rolling of the strip increases. To roll a moderately wide strip (usually wider than 1000 mm) on an industrial scale, it is preferable to perform hot rolling at a relatively high temperature above Ar3 in the austenite phase of the steel. Ar3 is the temperature at which ferrite begins to form in the steel during cooling. This can be ensured by using a starting - finishing rolling temperature (F1) of about 1000 °C or higher. An F1 temperature lower than this may increase the hot rolling load and lead to hot rolling in the transformation range, which can make hot rolling in industrial-scale large-scale processing difficult. Apart from increasing the rolling load, hot rolling in the transformation range can also cause insufficient recrystallization of the hot-rolled strip.

[0039] Next, when cold rolling is applied to the hot-rolled strip to reduce the gauge of the final steel product, the material cannot be cold rolled unless appropriate pretreatment is employed. In particular, the coiled steel after hot rolling is subjected to an intermediate batch annealing treatment at a low temperature within the transformation temperature range of the steel for 24 hours or more. Since it is a relatively long time, it is batch annealing. The intermediate batch annealing temperature needs to be less than 650 °C because at temperatures higher than this, a large amount of retained austenite is formed after cooling the steel to room temperature. Also, using a higher batch annealing temperature may result in a large amount of martensite appearing in the microstructure. Both martensite and retained austenite make cold rolling difficult by increasing the rolling load. The martensite phase is hard, but the retained austenite transforms into hard martensite during cold rolling, which itself increases the rolling load. Therefore, intermediate batch annealing of the coiled material is part of the method according to this embodiment to maintain the contents of retained austenite and martensite at lower values and increase the amount of ferrite. The ferrite phase does not work-harden as much as retained austenite during cold rolling, keeping the rolling load low and enabling cold rolling. The minimum required amount of the ferrite phase of the hot-rolled steel suitable for cold rolling after this batch annealing is 60% by volume.

[0040] Final batch annealing of hot-rolled steel strip (in the case of hot-rolled products) or cold-rolled steel strip (in the case of cold-rolled products) is very important for obtaining the desired microstructural components in the final product for the present invention to function. This final batch annealing needs to be carried out at a transformation interval temperature (Ac1 - Ac3) below 700°C, preferably below 660°C. The reason is that thermodynamic calculations suggest that for the range of chemical compositions of the steel of the present invention, the peak of C enrichment in the transformation interval austenite occurs below 660°C, while Mn enrichment increases monotonically as the temperature is lowered from 700°C. Therefore, the final batch annealing temperature below 660°C ensures optimal maximum (C + Mn) enrichment in the transformation interval austenite. The annealing temperature needs to be selected so that Mn and C are distributed in the austenite in the maximum amount. Since C and Mn are austenite stabilizing factors, during this final batch annealing at the transformation interval temperature, C and Mn are enriched in the transformation interval austenite. C is a small inter-lattice element in steel and thus diffuses and distributes rapidly, while Mn, being a large substitutional element, diffuses slowly. Therefore, in order to achieve a large amount of Mn in the austenite, a batch annealing time of preferably 5 hours or more, more preferably 10 hours or more is required. Mn enrichment in the austenite should be such that the Mn content is at least 1.25 times, preferably at least 1.4 times the bulk Mn content of the steel. C enrichment should be such that the C content is at least 1.2 times, preferably at least 2.3 times the bulk C content of the steel. These levels of Mn and C enrichment in the transformation interval austenite are required to properly stabilize the austenite at room temperature, whereby at least 20% by volume of retained austenite is obtained in the room temperature microstructure. Also, these levels of Mn and C enrichment are necessary to achieve the optimal mechanical stability (referred to as meta-stability) of the retained austenite, whereby the strain hardening index of the steel can be at least 0.3 during deformation.When the enrichment levels of Mn and C in austenite are lower than the above values, the optimum stability of retained austenite cannot be achieved, and thus the strain hardening index, which is at least 0.3, cannot be achieved either. Therefore, when the annealing period is shorter than 5 hours, these requirements of the present invention are not satisfied. Using a batch annealing temperature above 700 °C results in similar drawbacks. Since there is no necessary C and Mn enrichment in the transformation range austenite, the transformation range austenite lacks sufficient stability and cannot provide at least 20% by volume of retained austenite in the room temperature microstructure after batch annealing and the metastability required for the desired high strain hardening rate. Therefore, the combination of a high proportion of retained austenite and its optimum mechanical stability leads to the desired high strain hardening rate. An annealing temperature above 700 °C also results in more than 20% by volume of martensite, which does not result in the desired strain hardening rate. It is this high strain hardening rate that provides a high combination of cold formability and the strength and ductility of the final product. The high strain hardening rate strengthens the steel sheet while it thins during forming (e.g., forming by stretching), resulting in high cold formability.

[0041] The final batch annealing is carried out in a non-oxidizing and non-nitriding atmosphere in order to minimize the surface deterioration of the steel strip by oxygen and nitrogen. Since a batch annealing time of at least 5 hours is required, if a non-oxidizing atmosphere is not used, the steel surface may be oxidized. Also, decarburization occurs, the C content of the steel decreases, and the effects of the present invention are reduced. For the same reason, nitrogen may react with Al present on the surface of the steel to form nitrides on the surface of the steel. All these forms of surface deterioration have an adverse effect on the mechanical properties and formability of the steel. The preferred annealing atmosphere can be a vacuum, hydrogen or argon atmosphere.

[0042] The above modifications to the steel and its processing result in an appropriate microstructure in the final product for the success of the present invention. A high proportion of retained austenite (20% by volume or more), a low proportion of martensite (less than 20% by volume), and an optimal proportion of ferrite (30 - 70% by volume) result in a combination of high strength, high ductility, and high formability due to a high strain hardening rate. A retained austenite content exceeding 65% by volume cannot be achieved within the composition boundaries of this steel and is not necessary to achieve the minimum required strain hardening rate. Furthermore, a retained austenite exceeding 70% by volume can also cause problems in spot welding and can lead to severe liquid metal embrittlement and a decrease in hydrogen embrittlement resistance during use. Therefore, the composition boundaries of the steel of the present invention were selected taking these factors into account. When the ferrite content exceeds 70% by volume, a high strain hardening index of at least 0.3, mainly due to metastable retained austenite by the TRIP effect, is not achieved. A ferrite proportion of less than 30% by volume is not required to obtain the minimum strain hardening rate. The martensite phase mainly contributes to strength but does not contribute much to the strain hardening rate. Furthermore, a large amount of martensite (more than 20% by volume) can result in weak interfaces with softer phases, such as ferrite and retained austenite. These interfaces function as nucleation sites for damage initiation and thus have an adverse effect on high ductility and high formability. Therefore, the martensite content needs to be kept below 20% by volume (including the case where no martensite is present).

[0043] The ultra-fine crystal grain size is another requirement of the microstructure of the present invention. The crystal grain size of ferrite should be below 2 μm, preferably 0.2 - 2 μm. This ultra-fine crystal grain size gives good ductility to the product and also provides strengthening through the refinement of crystal grains, which contributes to the good mechanical properties of the steel strip of the present invention. This ultra-fine microstructure is also ensured by selecting a low final batch annealing temperature below 700 °C, which limits the growth of crystal grains. Furthermore, due to the steel composition of the present invention and the requirements of the final batch annealing in the transformation range, the phases (ferrite and austenite) at the annealing temperature are restricted from growing with respect to each other. All these factors result in the desired ultra-fine crystal grain size of ferrite. When the crystal grain size of ferrite exceeds 2 μm, it leads to a decrease in strength and ductility. In this specification, the crystal grain size is mainly represented by the length of crystal grains. Through sufficient recrystallization of the steel strip during the final batch annealing, the width of ferrite crystal grains is more than 1 / 3 of the length. As a result, a relatively equiaxed shape of ferrite crystal grains is actually obtained after the final batch annealing.

[0044] The influence of the above ultra-fine ferrite crystal grain size on sheet formability may be the appearance of yield point elongation in the engineering stress-strain curve of the product. This may have an adverse effect on the cold formability of the steel due to strain localization and may deteriorate the aesthetic appearance of cold-formed parts. Therefore, the process variables are selected such that the proportion of ferrite in the final microstructure is at most 70% by volume. Thereby, the yield point elongation, if present, is limited to a maximum engineering strain of 10%, and the best sheet formability and / or aesthetic appearance of cold-formed parts are obtained. However, a yield point elongation exceeding 10% of engineering strain is not a limiting factor for the present invention to function, because potential adverse effects due to yield point elongation can be mitigated by appropriate quality rolling before forming and / or appropriate lubrication during forming.

[0045] In the case where yield point elongation exists, in order to remove the yield point elongation, the steel strip is, in the present invention, optionally slightly reduced in thickness by cold rolling, which is carried out by temper rolling or skin pass rolling, with a maximum thickness reduction of 5%. This slight cold rolling, applied in one or more passes, eliminates the yield point elongation without perceptibly changing the mechanical properties of the steel strip. However, the present invention still functions to provide high cold formability without this temper rolling process even when the yield point elongation exists at an engineering strain of 10% or less.

[0046] Optionally, the hot-rolled or cold-rolled strip after final batch annealing is coated with a metal coating to enhance the aesthetic appearance and corrosion resistance during use. Coating methods include, but are not limited to, hot dip galvanizing, electrogalvanizing, electro-galvanizing, PVD, CVD, etc. The final batch annealed steel strip is very robust in terms of its microstructure and thus the application of a thin coating does not essentially change its properties.

[0047] Cold forming of the steel strip or sheet or blank can be carried out with or without the application of a suitable lubricant to reduce the frictional force between the steel and the tool. In either case, the present invention provides high cold formability. Non-limiting examples of lubrication systems are light oil, Kluber Press Pate, Teflon foil, or combinations thereof.

[0048] The steel used in the method of the present invention is a medium Mn steel containing carbon, manganese and aluminum as main components. Optionally, other alloying elements selected from silicon, chromium, vanadium, niobium, titanium and molybdenum may be present. Inevitable impurities, such as N, P, S, O, Cu, Ni, Sn, Sb, etc. (derived from the starting materials for preparing the steel having this composition), may be present. These are not intentionally added and are not specifically controlled within certain limits. The balance of the steel having this composition is iron.

[0049] Carbon is present in an amount of 0.05 to 0.3% by weight, for example, 0.05 to 0.20% by weight, preferably 0.07 to 0.20% by weight. Carbon is added mainly from the viewpoint of strength, but C also contributes to the stabilization of austenite. In this composition, the austenite stabilizing effect of manganese is much more remarkable because of its high ratio. The preferred range of C is 0.05 to 0.25% by weight, and a more preferred range is 0.08 to 0.21% by weight. If C is excessively low, the desired strength level of 800 MPa cannot be obtained, and if C exceeds 0.21% by weight, the weldability of the formed parts may decrease.

[0050] Manganese is present in an amount of 3.0 to 12.0% by weight. Manganese lowers the Ac1 and Ac3 temperatures, stabilizes austenite, increases strength and toughness, and causes the TRIP effect by stabilizing austenite in the room temperature microstructure. At levels below 3.1% by weight, the desired effect is not achieved, while above 10.5% by weight, problems of casting and segregation occur. The deformation mechanism also changes from transformation-induced plasticity (TRIP) to twinning-induced plasticity (TWIP). If the Mn content is excessively low, sufficient austenite is not retained at room temperature, and the stability of the retained austenite becomes excessively low, and as a result, the benefits of ductility and strain hardening cannot be obtained. Preferably, the Mn content is 3.5 to 10.0% by weight. In one embodiment, Mn reaches 5.0 to 9.0% by weight. In other embodiments, it is 5.5 to 8.5% by weight, for example, 6.0 to 7.5% by weight.

[0051] Aluminum is added to expand the temperature range of Ac1 - Ac3 and improve the robustness of the treatment from the viewpoint of industrial applications. Al is present in an amount of 0.03 to 3.0% by weight, for example, 0.6 to 2.9% by weight, preferably 1.0 to 2.2% by weight.

[0052] When silicon is present, it is added in an amount of less than 1.5 wt% in order to enhance the strength by solid solution strengthening. When present, the amount is usually more than 0.01 wt% and less than 1.5 wt%. Its preferred range is 0.1 - 1.0 wt%.

[0053] Both Al and Si suppress the precipitation of cementite and avoid a decrease in ductility. Furthermore, both Al and Si increase the peak annealing temperature in order to obtain the maximum amount of retained austenite at room temperature after final batch annealing. Therefore, during intercritical annealing, the diffusion of Mn is promoted and effective Mn partitioning in austenite is carried out.

[0054] One or more additional microalloying elements selected from Group V, Nb, Ti, and Mo may optionally be present. These microalloying elements increase the strength by precipitation hardening with carbides, nitrides, or carbonitrides. Cr, another optional element of the present invention, also increases the peak annealing temperature to achieve the maximum amount of retained austenite at room temperature and decreases the sensitivity of the retained austenite content to the annealing temperature. These result in effective Mn partitioning in austenite and an improvement in the robustness of the treatment during annealing. When present, the preferred addition of these optional alloying elements is: V: 0.01 - 0.1 wt%; and / or Nb: 0.01 - 0.1 wt%; and / or Ti: 0.01 - 0.1 wt%; and / or Mo: 0.05 - 0.5 wt%; and / or Cr: 0.1 - 2.0 wt%.

[0055] The composition of the metal coating is not limited. A zinc-based coating can be applied, for example, which essentially contains zinc and contains at least 0.1% by weight of Al, optionally up to 5% by weight of Al and optionally up to 4% by weight of Mg, with the remainder of the coating composition containing, independently of each other, less than 0.3% by weight of additional elements and unavoidable impurities, and is a zinc coating. For example, in order to form spangles and / or prevent dross formation, other additional elements may be present in small amounts of less than 0.3% by weight. The other additional elements may be selected from the group comprising Pb, Sb, Ti, Ca, Mn, Sn, La, Ce, Cr, Ni, Zr and Bi. Pb, Sn, Bi and Sb. Such small amounts of additional elements do not significantly change the properties of the bath or the resulting coating in normal applications. Preferably, if one or more additional elements are present in the coating, each is present in less than 0.02% by weight, and preferably each is present in less than 0.01% by weight. The coating methods can also be various, such as hot dip galvanizing (GI), zinc plating, heat-to-coat cycle, electrogalvanizing, etc. Aluminum-based coatings can also be applied, for example, an Al-Si-X coating (where Si can vary from 0.1 to 10% by weight and X = elements that modify other coatings in the required amounts and unavoidable impurities that do not essentially change the coating properties). Coating methods such as PVD, CVD, etc. are also applicable.

[0056] The final batch annealing procedure is not limited by the type of furnace used or the heating and cooling rates of the strip within the coil. When subjecting the coil to batch annealing, it is understood that the heating rate of the coil can vary from the surface towards the center. However, in the present invention, it is essential that the coiled strip be at the target batch annealing temperature for a minimum of 5 hours, preferably 10 hours or more, such that all parts of the coil undergo sufficient C and Mn enrichment in the austenite. Since the presence of a large amount of Mn improves the hardenability of the steel, the cooling rate after batch annealing is not relevant to the present invention. Thus, the coil can be cooled, air-cooled, forced air-cooled, or even water-quenched within the batch annealing furnace.

[0057] Normalizing rolling or skin pass rolling can be carried out on either bare steel strips or coated steel strips. This can also be carried out in a single pass or multiple passes.

[0058] The resulting steel strip preferably has, in volume %, Ferrite: 30 - 70%; Retained austenite: 20 - 65%; Martensite: less than 20% (including 0%); and has a three-phase or two-phase microstructure, and the ferrite crystal grain size is 0.2 - 2 μm.

[0059] The resulting steel strip is characterized by the following composition for the retained austenite: Mn: 1.25 times, preferably 1.4 times the bulk Mn composition of the steel C: 1.2 times, preferably 2.3 times the bulk C composition of the steel and has.

[0060] Advantageously, the steel strip has the following properties: Yield strength: 600 MPa or more; Maximum tensile strength: 800 MPa or more; Elongation at break: 20% or more; Strain hardening index: 0.3 or more Yield point elongation: Engineering strain preferably 10% or less; Minimum bending angle at a thickness of 1.0 mm: 100° or more.; Hole expansion performance: 20% or more; Minimum stretching strain in bi-axial stretching: 10% or more having.

[0061] The proportion of the above phases was determined using X-ray diffraction (XRD). The amount of retained austenite was determined by XRD at a position where the thickness of the sample was 1 / 4. The XRD pattern was recorded in the range of 45 to 165° (2θ) by a Panalytical Xpert PRO standard powder diffractometer (CoKα line). The quantitative determination of the phase ratio was performed by Rietveld analysis using the Bruker Topas software package for Rietveld refinement. The martensite content was determined from the peak-split at the ferrite diffraction position of the diffraction pattern.

[0062] The crystal grain size of the phase was determined from a scanning electron microscope (SEM) image of the microstructure. The Mn concentration of the retained austenite was determined by an electron probe microanalyzer (EPMA). The C content of the retained austenite was evaluated by a well-known formula proposed by Dyson and Holmes. This formula relates the lattice parameter of austenite that can be determined from XRD data to its C content. This formula can be obtained from the following paper: D.J. Dyson, B. Holmes, Effect of alloying additions on the lattice parameter of austenite. Journal of Iron Steel Institute, vol. 208, year 1970, pages 469-474.

[0063] The yield strength, maximum tensile strength, yield point elongation, and total elongation were determined at room temperature under quasi-static (strain rate 3×10 -4 s -1)It was determined from a tensile test. The shape of the tensile test piece was composed of a gauge length of 80 mm, a width of 30 mm, and a nominal thickness of 1.5 mm in the rolling direction. The strain hardening rate was measured in a range of 7 % strain after the yield point elongation in the tensile curve. The bendability was determined by a three-point bending test conforming to the VDA238-100 standard for test pieces of 40 mm × 30 mm in both the longitudinal and transverse directions with a nominal thickness of 1.5 mm. The bending axis was along the 30 mm dimension, and the bending radius was 0.4 mm. The bending angle obtained from the test piece with a nominal thickness of 1.5 mm was converted to the angle corresponding to a thickness of 1.0 mm using the following formula: Bending angle at a thickness of 1.0 mm = Measured angle × Square root of the actual thickness (mm). From these converted bending angles, for specific heat treatment conditions, the minimum values of the test pieces in the longitudinal and transverse directions were adopted to claim the scope of the present invention. The hole expansion performance (HEC) was determined in accordance with the ISO / TS 16630:2003(E) standard. A test piece with dimensions of 90 mm × 90 mm × 1.5 mm was cut out from a steel strip. A hole with a diameter of 10 mm was drilled in the center of the test piece, and a hole expansion test was carried out. The hole expansion performance (HEC = (expansion of the initial diameter of the hole / initial diameter of the hole) × 100%) was calculated from the measurement data. The biaxial stretching strain was determined from a biaxial stretching test carried out using an Erichsen press that combines a flat punch with a diameter of 75 mm and a die with a diameter of 79.78 mm. The radius of the punch nose was 10 mm, and the radius of the die was 8 mm. The blank holder force was set to the maximum mechanical capacity (about 580 kN) so that no drawing occurred. The test speed was set to 20 mm / min. Strain measurement was performed by applying a 10 mm square grid to the sheet with fine markers.

[0064] In the production of steel strips, by using the final batch annealing process at a temperature in the transformation range of steel below 700°C as described above, Mn distribution from ferrite to austenite occurs, improving the stability of austenite in the transformation range. During cooling after the final batch annealing, since austenite in the transformation range has high stability due to its low Ms, it does not transform much into martensite, resulting in a two-phase microstructure of ferrite and retained austenite. In the case of a low Mn content (e.g., less than 8 wt%), some austenite in the transformation range may transform into martensite, but the martensite content will be less than 20 vol%. Therefore, by increasing the level of Mn and lowering the batch annealing temperature (e.g., below 700°C), a large amount of retained austenite (20 vol% or more) with an optimal metastable state can be ensured. This large amount of retained austenite partially transforms into martensite during deformation in the forming process, causing the transformation-induced plasticity (TRIP) effect and resulting in a high strain hardening index (= high elongation and high formability).

[0065] Due to the steel composition, it is preferable that the total elongation of the steel strip is 20% or more and the strain hardening index is 0.3 or more. Using the transformation range batch annealing process of the medium Mn steel approach, it is preferable to obtain a mixed microstructure of ultrafine ferrite (0.5 - 2.0 microns), martensite, and high retained austenite regions. As a result, high ductility and a high hardening rate are obtained. These lead to high cold formability of the steel strip.

[0066] Preferred steel strips are used as materials for manufacturing automotive parts, especially automotive parts with complex shapes that require the formability of the strip. Parts that require high energy absorption combined with high strength are also suitable for manufacturing from steel strips. Non-limiting examples include automotive interior parts, B-pillars, axial bars, etc.

Examples

[0067] The present invention will be described with reference to FIG. 1 and the examples described below.

[0068] Figure 1 shows the SEM microstructure of steel (Steel A, 650 °C / 10 h) manufactured according to the present invention obtained by final batch annealing. Here, F = ferrite, MA = martensite-austenite.

[0069] Steel ingots A, B, and C having the chemical compositions of three types of the present invention with dimensions of 200 mm × 100 mm × 100 mm were cast by melting the charge in a vacuum induction furnace. The chemical compositions of these steels having the chemical compositions of the present invention are shown in Table 1 together with two reference steels D and E. Steel D is a twinning-induced plasticity (TWIP) steel, and steel E is a DH1000 steel grade. Both were received in their finally cold-rolled and annealed states, and the thicknesses of these steels in the as-received state were 1.7 mm and 1.5 mm, respectively. Then, they were reheated at 1250 °C for 2 hours and rough-rolled to a thickness of 30 mm. Then, the strip was reheated again at 1250 °C for 30 minutes, and hot-rolled to a thickness of 3 mm for Steel A and B and 4 mm for Steel C at a rolling start temperature of 1150 °C and a finish rolling temperature (FRT) of 900 °C (in the austenite phase region for all three steels). A high reheating temperature of 1250 °C and a long time of 2 hours were used for proper homogenization of Mn.

[0070] The transformation temperatures from austenite to ferrite (Ar3) for steels A, B, and C were measured by dilatometry to be 785 °C, 770 °C, and 723 °C, respectively. Then, the hot-rolled steel was subjected to coil cooling simulation in a muffle furnace from 680 °C and thereby cooled to room temperature. Next, the hot-rolled strips of A and B were subjected to intermediate batch annealing at 600 °C for 96 hours in a muffle furnace under a protective atmosphere of argon, and the strip of C was subjected to intermediate batch annealing at 550 °C and air-cooled to room temperature. These annealing temperatures were selected so that a desired proportion of ferrite could be obtained to facilitate subsequent cold rolling. Table 2 shows the phase proportions of steels A, B, and C after this intermediate batch annealing of the hot-rolled strips. The phase proportions were determined by XRD measurement as described above at a position where the strip thickness was 1 / 4. It can be seen that in all three steels, the proportion of ferrite exceeded 60% by volume.

[0071] Next, the strips were pickled with HCl acid at 90 °C to remove oxides, and then all the steels were cold-rolled from their respective hot-rolled gauges to a final thickness of 1.5 mm.

[0072] The cold-rolled strips of A and B were batch annealed at 650 °C for 10 hours, and the cold-rolled strip of steel C was batch annealed at 640 °C for 4 hours and 16 hours using a muffle furnace. An argon atmosphere was used for annealing to ensure that the atmosphere did not contain oxygen and nitrogen, thereby minimizing strip oxidation and unwanted reactions of nitrogen from the atmosphere and aluminum from the steel to form a nitride layer on the surface. After annealing, the samples were air-cooled to room temperature. For comparison, the cold-rolled strip of A was also annealed at 650 °C for 2 minutes, 5 minutes, and 1 hour in the same manner, and the cold-rolled strip of C was annealed at 640 °C for 4 hours. Some test pieces were subjected to skin pass rolling or temper rolling with a maximum thickness reduction of 5%.

[0073] The procedures for material property evaluation and testing are as described above. For the sake of recollection, the microstructure of the samples was characterized using XRD and SEM. The microanalysis of the phase chemistry was performed by EPMA and XRD analysis. The tensile properties were determined by tensile tests on specimens with a gauge length of 80 mm and a width of 30 mm (the shape of A80 specimens). The formability of the strip was evaluated by bending tests, hole expansion tests, and biaxial stretching tests using appropriate lubricants. Regarding bendability, the definitions of L and T specimens are as follows: L = a longitudinal specimen with the bending axis parallel to the rolling direction, T = a transverse specimen with the bending axis perpendicular to the rolling direction.

[0074] Figure 1 shows a typical microstructure obtained after final batch annealing of the cold-rolled strip of Steel A. Here, it is possible to observe ferrite and martensite-austenite regions. Ferrite with an ultrafine crystal grain size can also be recognized. The microstructure characteristics of Steels A, B, and C after subjecting the cold-rolled samples to different final annealing treatments are shown. For all conditions of all steels, the crystal grain size of ferrite is 0.5 - 1.9 μm. For Steels A and C, as the annealing time at each annealing temperature increases, the residual austenite content increases because of the increased Mn partitioning to austenite. The higher the Mn content, the higher the residual austenite content (Steel C has more residual austenite than A and B), indicating the effect of Mn on austenite stabilization. High proportions of residual austenite (more than 33 vol%) were obtained under all conditions except for Steel A annealed at 650 °C for 2 minutes. The Mn and C contents of the annealing conditions of the steels shown in Table 4 indicate that in the residual austenite of all these different conditions of the steels, the Mn enrichment is 1.286 - 2.139 times the bulk Mn content of the steel (except for Steel A under the condition of 650 °C / 2 minutes In steel A under the condition of 650°C / 2 minutes where the Mn content is only 1.09 times the bulk Mn content). Regarding C enrichment in the residual austenite, the C content is 1.17 - 3.085 times the bulk C content of the steel (except for Steel A under the condition of 650 °C / 2 minutes In steel A under the condition of 650°C / 2 minutesThis value is 1.063 times. Due to these low C and Mn enrichments in austenite, the retained austenite content of Steel A under the condition of 650 °C / 2 minutes is also less than 20% by volume. As a result, the martensite content is more than 20% by volume (39.8% by volume). Under all other steels and conditions of the present invention, the martensite content is 16.7% by volume or less (including 0% by volume: for Steel C at 640 °C / 960 minutes). Oh (for Steel C).

[0075] The reason for the low proportion of retained austenite in Steel A under the condition of 650 °C / 2 minutes is that even though the annealing temperature is within the transformation temperature range of Steel A and less than 700 °C, the annealing time of 2 minutes is too short for sufficient Mn diffusion into austenite.

[0076] The significance (consequence) of the above microstructural features is seen in the tensile properties of the steels shown in Table 5. Steel A at 650 °C / 2 min, with a low amount of retained austenite and with Mn and C being less than 1.25 times and less than 2 times their bulk Mn and C contents respectively, showed very high yield strength and maximum tensile strength, but the total elongation was only 3.1%. This is because all of its retained austenite transforms very rapidly to martensite during the tensile test due to the low stability predicted by the low Mn and C enrichment. The small amount of retained austenite shows no yield point elongation and is consumed at a very early stage of deformation. Therefore, the tensile properties of the steel under this condition are low and it cannot be used for cold forming. On the other hand, Steel A under other annealing conditions and Steels B and C under all conditions showed a yield strength above 693 MPa, a maximum tensile strength above 860 MPa and a total elongation above 23.4%. These steels also showed high energy absorption capacity (determined by the product of the maximum tensile strength and the total elongation) and various amounts of yield point elongation. The yield point elongation decreased with the annealing time for Steels A and C, which is due to the increase in the ferrite crystal grain size shown in Table 3. The tensile properties of the steels having the compositions of the present invention are comparable to those of the reference steels described in Table 6. The chemical composition of the steels of the present invention under long-time final batch annealing conditions has a much higher total elongation and energy absorption capacity than the conventional DH1000 steel grade (reference steel E) due to the combination of the chemical composition - treatment - microstructure of the steel. Steel E has a very small amount of retained austenite in its microstructure. Furthermore, the TWIP steel (reference steel D) has a much higher total elongation than the steels of the present invention, but the energy absorption capacity of some of the steels of the present invention is within the range of the TWIP steel having a fully austenitic microstructure.

[0077] The formability parameters of the steel of the present invention are shown in Table 7 in comparison with a reference steel. The formability parameters compared are the biaxial stretchability in terms strains in rolling and transverse directions, the bendability in the longitudinal and perpendicular directions of the sheet, and the flangability represented by the HEC value. Steel A shows that when the final batch annealing is carried out at 650 °C for less than 10 hours, the value of the biaxial elongation strain is 0, but the other parameters are not zero. Steel A annealed at 650 °C for 2 minutes was also very poor in bendability and flangability. The bendability and flangability improve with the increase of the annealing time, but the material has no stretchability until 10 hours of the final batch annealing. Steel B annealed at 650 °C for 10 hours also showed the same formability parameters as Steel A under the same annealing conditions. Steel C annealed at 640 °C for 4 hours showed high bendability and flangability but low stretchability. When Steel C is annealed for 16 hours, the stretchability also improves.

[0078] The cold formability of a steel sheet is a combination of various parameters, such as stretchability, bendability, and flangability. When the steel of the present invention is subjected to final batch annealing at less than 700 °C in the transformation temperature range, since Mn is an element that diffuses slowly in the steel, as seen above, the annealing time for causing the required amount of Mn and C enrichment in the retained austenite is important. High Mn and C enrichment are necessary to achieve a high strain hardening rate. Therefore, the steel annealed for less than 10 hours, which has a low strain hardening index, also has low stretchability, but the other formability parameters are good. To achieve good stretchability in the steel of the present invention, a high strain hardening index of 0.3 or more is required. Otherwise, there is a possibility of premature local fracture. Therefore, from the results, for good cold formability (a combination of stretchability, bendability, and flangability) as described in the claims, it is clear that a final batch annealing of at least 10 hours is necessary for the steel of the present invention to achieve the minimum values of Mn and C enrichment in the retained austenite.

[0079] Comparing the formability of the sample annealed for 10 hours with the steel of the present invention, it can be seen that the formability parameters of the steel of the present invention are in the range of high formability TWIP steel (reference steel E) and are much higher than those of the conventional DH1000 (reference steel D). The biaxial tensile strain of steel C is higher than that of the conventional DH1000 even when annealed for only 4 hours. This high cold formability of the steel of the present invention is due to the high proportion of metastable retained austenite with high Mn and C enrichment achieved in the steel of the present invention through the treatment process of the present invention.

[0080] Table 8 shows the effect of temper rolling on the mechanical properties of Steel B annealed at 650 °C for 10 hours. It is clear that as the temper rolling reduction increases, the yield point elongation decreases. With 2% reduction, the yield point elongation disappeared. The tensile properties did not change much and remained within the scope of the claims of the present invention. Very importantly, the strain hardening index also remained high up to a 5% thickness reduction. Therefore, this disappearance of the yield point elongation by temper rolling up to a maximum of 5%, without significantly changing the mechanical properties, makes the steel strip of the present invention more cold formable, because it reduces the risk of strain localization during stretch-forming and stretcher marks on the surface of the formed product.

[0081]

Table 1

[0082]

Table 2

[0083]

Table 3

[0084]

Table 4

[0085]

Table 5

[0086]

Table 6

[0087]

Table 7

[0088]

Table 8

Claims

1. A method for manufacturing a cold-rolled and annealed steel strip, wherein the steel composition is, by weight%, C: 0.05 to 0.3; Mn: 3.0 to 12.0; Al: 0.03 to 3.0; optionally, the following additional alloying elements: Si: less than 1.5 Cr: less than 2.0 V: less than 0.1 Nb: less than 0.1 Ti: less than 0.1 Mo: less than 0.5 one or more of; inevitable impurities; and Fe: the balance and the method comprising the following steps: casting the molten steel into slabs; reheating the slabs and holding at a temperature of 1150 °C or higher for 1 hour or more; hot-rolling the steel into strips; winding up the hot-rolled steel strips; pickling the steel strips; intercritical batch annealing the steel strips at a temperature of less than 650 °C for 24 hours or more, and obtaining at least 60% by volume of ferrite after cooling to room temperature; cold-rolling the steel into cold-rolled steel strips and winding them up; winding up the wound steel strips under the following conditions: a transformation range temperature between Ac1 and Ac3 that is less than 700 °C; a non-oxidizing and non-nitriding atmosphere; A step of batch annealing for a total annealing time in which the strip is maintained at the transformation interval temperature for at least 5 hours, enriching Mn in austenite such that the Mn content of the austenite is at least 1.25 times the bulk Mn content of the steel, and enriching C in the austenite such that the C content of the austenite is at least 1.2 times the bulk C content of the steel; A step of cooling the steel after batch annealing by air, forced air or water quenching The method including the above.

2. A method for manufacturing a hot-rolled and annealed steel strip, The steel composition is, by weight%, C: 0.05 - 0.3; Mn: 3.0 - 12.0; Al: 0.03 - 3.0; Optionally, the following additional alloying elements: Si: less than 1.5 Cr: less than 2.0 V: less than 0.1 Nb: less than 0.1 Ti: less than 0.1 Mo: less than 0.5 One or more of the above; Inevitable impurities; and Fe: the balance wherein, The method includes the following steps: A step of casting molten steel into a slab; A step of reheating the slab at a temperature of 1150°C or higher for 1 hour or more; A step of hot-rolling the steel into a strip; A step of winding up the hot-rolled steel strip; A step of pickling the steel strip; The wound-up steel strip under the following conditions: The transformation interval temperature of Ac1 - Ac3, which is less than 700°C; A non-oxidizing and non-nitriding atmosphere; A step of batch annealing for a total annealing time in which the strip is maintained at the transformation range temperature for at least 5 hours, enriching Mn in austenite such that the Mn content of the austenite is at least 1.25 times the bulk Mn content of the steel, and enriching C in austenite such that the C content of the austenite is at least 1.2 times the bulk C content of the steel; A step of cooling the steel after batch annealing by air, forced air or water quenching The method including the above.

3. The method according to claim 1 or claim 2, wherein the slab is reheated at a temperature of 1200 °C or higher.

4. The method according to any one of claims 1 to 3, wherein the slab is reheated at a temperature of 1250 °C or higher.

5. The method according to any one of claims 1 to 4, wherein the batch annealing of the coiled steel strip is carried out at a transformation range temperature of less than 660 °C.

6. The method according to any one of claims 1 to 5, wherein the obtained strip is coated with any metal coating applied by hot dip galvanizing, galvanizing, electro-galvanizing, aluminum plating, PVD, or CVD.

7. The method according to claim 1 or any one of claims 3 to 6, wherein the obtained steel strip is subjected to skin pass rolling.

8. The method according to claim 7, wherein the thickness reduction by the skin pass rolling is 5% or less.

9. A steel strip, The steel strip has the following steel composition by weight percentage: C: 0.05 - 0.3; Mn: 3.0 - 12.0; Al: 0.03 - 3.0; Optionally, the following additional alloying elements: Si: less than 1.5 Cr: less than 2.0 V: less than 0.1 Nb: less than 0.1 Ti: less than 0.1 Mo: less than 0.5 one or more of; inevitable impurities; and Fe: the balance having the composition of retained austenite in the steel strip has a Mn content that is at least 1.25 times the bulk Mn content of the steel composition and a C content that is at least 1.2 times the bulk C content of the steel composition, the steel has a high strain hardening index of at least 0.3, measured in the strain range of 7% after the yield point elongation in a quasi-static tensile test, the microstructure after final batch annealing is, in volume %, ferrite: 30 - 70%; retained austenite: 20 - 65%; including, with the balance being martensite: less than 20% (including 0 volume %), the steel strip wherein the length / width ratio of ferrite crystal grains is 3 or less.

10. The steel strip according to claim 9, wherein the ferrite crystal grain size is 0.2 - 2 μm.

11. The steel strip according to claim 9 or 10, wherein the yield point elongation is a maximum of 10% engineering strain measured from the engineering stress-strain curve.

12. The steel strip according to any one of claims 9 - 11, wherein the yield strength is 600 MPa or more, the maximum tensile strength is 800 MPa or more, and the total elongation (A80) is 20% or more.

13. The steel strip according to any one of claims 9 - 12, which has very high formability, characterized in that the elongation strain in an individual direction under biaxial stretching conditions is 10% or more, the VDA bending angle at a thickness of 1.0 mm is 100° or more, and the hole expansion performance is 20% or more.

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