METHOD OF MANUFACTURING A COLD-FORMABLE HIGH-STRENGTH STEEL STRIP AND STEEL STRIP

Abstract

A method for manufacturing a steel strip comprises the steps of: - shaping molten steel into slabs; - reheating the slabs to a temperature of 1150°C or higher for a time of 1 hour or more; hot rolling the steel into a strip, preferably with an average inlet temperature of the plate F1 above 1000°C; - cooling the hot-rolled steel strip; - batch annealing the steel strip: at an intercritical temperature (i.e., between Ac1 and Ac3), preferably below 700°C; - in a non-oxidizing and non-nitrogenous atmosphere; - the total annealing time being at least 5 hours, preferably at least 10 hours to obtain Mn enrichment in austenite so that the Mn content is at least 1.25 times the bulk Mn content of the steel and C enrichment so that the C content is at least 1.2 times the bulk carbon content of the steel: - cool the steel after batch annealing in air, forced air or water quenching.

Description

METHOD FOR MANUFACTURING A COLD-FORMABLE HIGH-STRENGTH STEEL STRIP AND STEEL STRIP FIELD OF INVENTION The present invention relates to a method for manufacturing a coated or uncoated high-strength steel strip that is cold-formed to manufacture a steel article and the steel strip. BACKGROUND OF THE INVENTION Cold forming, also known as cold stamping or cold press forming, is a method for manufacturing steel components in industries such as automotive, construction, engineering, and infrastructure, for various applications. It is known that the cold formability of steel sheets decreases as the strength of the steel increases. This is particularly true for conventional steel sheets, as well as for the first generation of advanced high-strength steel sheets (AHSS). Due to this inverse relationship between the strength and elongation of steel, the strain-strength property map of these steels is sometimes referred to as a "banana diagram." However, high formability at high strength can be achieved using the concepts of so-called second-generation AHSS (2GAHSS), but these steels are usually highly alloyed and also contain expensive alloying elements. Examples include twinned plasticity-induced steels (TWIPs) with high manganese content, where the Mn content is typically greater than 12% by weight, and stainless steels containing high amounts of expensive alloying elements such as chromium, nickel, molybdenum, etc. Besides being very expensive, another disadvantage of 2GAHSS is that, due to their very high alloy content, they are very difficult to manufacture on a large industrial scale. To overcome the problems of 2GAHSS while still achieving reasonably high cold formability at high strength, the concepts of third-generation AHSS (3GAHSS) have been introduced, such as quench and parting steels (Q&P), carbide-free bainitic steels (CFB), and medium-manganese steels. These steels are less expensive than 2GAHSS and can be easily processed in existing steel mill facilities. The present invention focuses on the medium-manganese type of 3GAHSS. WO16001887 describes a method for manufacturing a high-strength steel plate where the steel contains, by weight percent, 0.1 < C < 0.4, 4.2 < Mn < 8, 1 < Si < 3, and 0.2 < Mo < 0.5, the remainder being Fe and unavoidable impurities. The method comprises continuous annealing above Ac3, quenching between the initial (Ms) and final (Mf) temperatures of martensite, overaging at 300–500°C for more than 10 seconds, and cooling. This is essentially a quench and parting (Q&P) process where carbon enrichment (and possibly manganese enrichment) in the austenite is achieved through the overaging step of a steel containing some martensite. The Q&P process is entirely different from an intercritical annealing process. In this document, substantial partitioning of Mn in austenite is not expected because at the lower overaging temperatures (300-500°C) Mn diffusion in steel is extremely low. WO2017021464 describes a high-tensile steel in a hot- or cold-rolled strip, having a chemical composition (in wt. %): C: 0.005 to 0.6; Mn: 4 to 10; Al: 0.005 to 4; Si: 0.005 to 2; P: 0.001 to 0.2; S: up to 0.05; N: 0.001 to 0.3; the remainder being iron with unavoidable elemental inclusions associated with steel, wherein the steel is hot-rolled flexibly, optionally annealed, cold-rolled flexibly, optionally annealed and further cold-rolled flexibly, and subsequently annealed at an annealing temperature of 600°C to 750°C for 1 minute to 48 hours. This patent applies flexible lamination by controlling the lamination gap where cutting conditions vary across the width of the strip.Flexible lamination is a different process for components with variable wall thicknesses, unlike conventional lamination as in the present invention, which produces a uniform product thickness across its entire width. One disadvantage of flexible laminated strip is that it results in heterogeneous properties across the strip's width. BRIEF DESCRIPTION OF THE INVENTION The present invention aims to provide a cold-formable steel strip in cold-rolled thicknesses in a coated or uncoated condition while maintaining high strength. Another objective of the invention is to provide a highly cold-formable steel strip in the cold-rolled thickness range in a coated or uncoated condition while maintaining high strength. Both hot-rolled and cold-rolled steel strips of the present invention have a high energy absorption capacity, i.e., they are highly shock-resistant, spot-weldable, and resistant to hydrogen embrittlement. The invention is first incorporated in a method of manufacturing a cold-rolled and annealed steel strip, the composition of the steel being in % by weight: C: 0.05-0.3; Mn: 3.0-12.0; Al: 0.03-3.0; Optionally, one or more additional alloying elements: Yes: 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; unavoidable impurities, such as S: less than 30 ppm; P: less than 0.04; and the remainder being Fe; The method comprises the following steps: - mold the molten steel into a sheet; - reheat the iron and maintain it at a temperature of 1150°C or more for 1 hour or more; - hot roll the steel into a strip, preferably with an average plate entry temperature F1 above 1000°C; - roll up the hot-rolled steel strip; - pickle the steel strip; - perform intermediate batch annealing of the steel strip at a temperature lower than 650°C for more than 24 hours to achieve at least 60% ferrite by volume after cooling to room temperature; - cold roll the steel into a cold-rolled steel strip and roll this up; - perform batch annealing of the coiled steel strip; - at an intercritical temperature between Ac1 and Ac3 which is less than 700°C; - in a non-oxidizing and non-nitrogenous atmosphere; - the total annealing time for which the strip is held at the critical temperature being at least 5 hours, preferably at least 10 hours to achieve the enrichment of Mn in the austenite such that the Mn content is at least 1.25 times the bulk Mn content of the steel and the enrichment of C such that the C content is at least 1.2 times the bulk C content of the steel; - Cooling the steel after batch annealing in air, forced air, or water cooling. A steel containing essentially 0.05 to 0.3 wt% carbon, 3.0–12.0 wt% manganese, 0.03–3.0 wt% aluminum, and optionally other alloying elements and unavoidable impurities is processed to a hot-rolled gauge using a specific processing route. The molten steel is cast into slabs, which are then reheated to a temperature of 1150°C or higher for one hour or more. The slabs are then hot-rolled into strips, preferably with a finishing entry temperature (F1) above 1000°C. The F1 entry temperature is the temperature at which the strip enters the first station of a finishing mill. The finishing laminator is the part of the hot rolling mill where the finishing lamination takes place after the roughing or breaking of the plates in the roughing laminator and before cooling the finishing table.After passing through the finishing table, the hot-rolled strip is wound into coils, and these coils are then subjected to intermediate batch annealing at a temperature below 650°C for at least 24 hours, so that at least 60% ferrite by volume is achieved in the steel strip after cooling to room temperature. The steel strips are pickled in an acid solution at a temperature of, say, 50–90°C and cold-rolled to thinner gauges. The invention is not limited by the range of hot-rolled or cold-rolled gauges. However, typically the hot-rolled gauge will be in the range of 2 to 10 mm, and the cold-rolled gauge in the range of 0.5 to 2 mm.The cold-rolled steel is then batch-annealed at an intercritical temperature below 700°C, preferably below 660°C, in a non-oxidizing, nitrogen-free 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 critical austenite reaches at least 1.2 times the bulk C content. The longer period of 10 hours is preferred because a larger amount of Mn can be diffused to the austenite from the ferrite during annealing. Mn typically takes longer to diffuse due to its large replacement alloying element in iron.As the annealing temperature decreases, the manganese enrichment in the intercritical austenite increases, making the austenite in the steel more stable at room temperature after cooling following batch annealing. Batch annealing time is defined as the period of time the steel strip is held at the batch annealing temperature, excluding the time required to heat the strip to the target temperature. The final batch annealing, i.e., in this text the batch annealing of the coiled strip, shall take place for a duration mentioned in the claims, which is sufficiently long to obtain in the steel a relatively equiaxed ferrite grain morphology, where the length-to-width ratio of the grain is preferably 3 or less. The steel is cooled at any cooling rate to ambient temperature, such as in air, forced air, or water. By carrying out the method according to the invention, the following advantages are obtained: Medium-manganese steel containing the aforementioned alloying elements along with 3-12 wt% Mn exhibits reduced Mn segregation. This is a concern that affects mechanical properties when Mn is present in a relatively high quantity, as in this invention. A relatively high reheating temperature of the plate, above 1150°C, preferably 1200°C, and more preferably above 1250°C, and a minimum reheating time of 1 hour are chosen to minimize segregation and homogeneously distribute the Mn within the matrix. Consequently, the final mechanical properties may be compromised. The selection of the reheating temperature will depend on the Mn content of the alloy.When the Mn content of the alloy is close to the lower limit of the claimed Mn range, reheating temperatures close to 1150°C will be sufficient to homogenize the Mn distribution, and as the Mn content increases, higher plate reheating temperatures will be preferred. Steel can be hot-rolled on an industrial scale to a reasonably large strip width, for example, over 1000 mm. This is achieved by maintaining a high F1 temperature, above 1000°C, during hot rolling to keep the required hot-rolling strength low. At a lower F1 temperature, hot rolling the steel strip becomes difficult. The steel then becomes suitable for cold rolling on an industrial scale. This is made possible by using an intermediate batch annealing step for the hot-rolled steel. The intermediate batch annealing is carried out at the intercritical temperature of the steel, preferably below 650°C, at a temperature selected such that at least 60% ferrite by volume is achieved in the steel strip, the remainder being retained austenite and martensite. The additional batch annealing of the rolled steel strip below 700°C creates the correct microstructure. This annealing should last at least 5 hours, preferably at least 10 hours. During this process step, the Mn and C in the steel of the invention are distributed between the intercritical austenite and ferrite, so that the austenite is highly enriched by Mn and the C makes the phase stable down to room temperature. Since the austenite is enriched so 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, the steel becomes virtually insensitive to practical cooling rates and can therefore be cooled in air, forced air, or water after batch annealing. The lower the annealing temperature below 700°C, the richer the intercritical austenite is in Mn content. A second embodiment of the invention is a method of manufacturing a hot-rolled and annealed steel strip, the steel composition being % by weight: C: 0.05-0.3; Mn: 3.0-12.0; Al: 0.03-3.0; Optionally one or more additional alloying elements: Yes: 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; unavoidable impurities, such as S: less than 30 ppm; P: less than 0.04; and The rest being Faith; The method comprises the following steps: - mold the molten steel into a sheet; - reheat the iron to a temperature of 1150°C or more, for a time of 1 hour or more; - hot roll the steel into a strip, preferably with an average inlet temperature of plate F1 above 1000°C; - roll up the hot-rolled steel strip; - pickle the steel strip; - Batch annealing of the coiled steel strip: - at an intercritical temperature between Ac1 and Ac3 which is less than 700°C; - in a non-oxidizing and non-nitrogenous atmosphere; - the total annealing time being the time by which the strip is held at the critical temperature for at least 5 hours, preferably at least 10 hours to achieve the enrichment of Mn in the austenite, so that the Mn content is at least 1.25 times the bulk Mn content of the steel and the enrichment of C so that the C content is at least 1.2 times the bulk C content of the steel; - Cool the steel after batch annealing in air, forced air or by water quenching. The steel processed according to the first method up to the hot rolling step, as described above, is then subjected directly to the final batch annealing step after pickling, omitting the intermediate steps. The batch annealing shall take place according to the claims, for a period long enough to obtain in the steel a relatively equiaxed ferrite grain morphology, where the grain length-to-width ratio is preferably 3 or less. The steel is thus manufactured as a hot-rolled strip instead of a cold-rolled strip, however, with all the advantages in terms of mechanical properties compared to the cold-rolled strip under the first method. The invention is also incorporated into a method where the plate is reheated to a temperature of 1200°C or higher. This achieves better homogenization of the Mn in the molten steel plates, reducing its segregation. The invention is also incorporated in a method where the plate is reheated to a temperature of 1250°C or higher. This achieves an even greater reduction in any microsegregation of Mn present in the cast steel plates. The invention also incorporates a method where batch annealing of the coiled steel strip takes place at an intercritical temperature below 660°C. This achieves greater manganese enrichment in the intercritical austenite, so that the martensite content in the final microstructure can be minimized. The invention further incorporates a method where the resulting strip is coated with any metallic coating applied by hot-dip galvanizing, electro-annealing, electrogalvanizing, aluminizing, or any other method such as physical vapor deposition (PVD) or chemical vapor deposition (CVD). This achieves the desired corrosion resistance and good aesthetic appearance of the steel strip in its application or service. In one embodiment of the method according to claims except claim 2, the resulting steel strip is subjected to light cold rolling, also known as temper rolling. This further improves the cold formability properties. In one method, light cold rolling takes place with a thickness reduction of 5% or less. This minimizes the elongation at the yield point or elastic limit of the steel strip during tensile testing, thereby improving cold formability and the aesthetic appearance of the cold-formed steel strip. The invention is also incorporated in a steel strip that can be or is manufactured according to the method of the first or second embodiment of the invention, wherein the steel strip has a steel composition that is % by weight: C: 0.05-0.3; Mn: 3.0-12.0; C 7 bQCn / 77n7 / q / YILI Al: 0.03-3.0; optionally one or more additional alloying elements: Yes: 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; unavoidable impurities, such as S: less than 30 ppm; P: less than 0.04; and the remainder being Fe; wherein the steel strip has a retained austenite composition having a Mn content that is at least 1.4 times the bulk Mn content of the steel composition as well as a C content that is at least 2.3 times the bulk C content of the steel composition, to obtain metastable retained austenite to give the steel a high strain hardening exponent of at least 0.3 measured after yield point elongation for a stress range of 7% in a quasi-static tensile test, wherein the microstructure after final batch annealing of the coiled steel strip comprises in % by volume: ferrite: 30-70%; retained austenite: 20-65%; martensite: < 20%, including 0% by volume. The steel strip has a level of Mn and C enrichment in the metastable austenite of the final microstructure of the steel in the end-use condition to give it a high work hardening or strain hardening rate. Furthermore, according to the invention, the cold-rolled or hot-rolled steel strips preferably have a microstructure where retained austenite is present in the range of 20 to 65% by volume, ferrite in the range of 30 to 70% by volume, and martensite is present below 20% by volume, including 0% by volume. The ferrite is preferably ultrafine with a grain size in the range of 0.2 to 2 µm. This ultrafine ferrite, due to sufficiently prolonged batch annealing of the coiled steel strip, acquires a more or less equiaxed shape, with a grain length-to-width ratio < 3.In contrast to continuous-type annealing, which is typically short (in minutes rather than hours) and produces the elongated shape of the grains with a high aspect ratio (length / width ratio), during the prolonged batch annealing of the cold-rolled strips employed in this invention, sufficient recrystallization of the ferrite grains takes place. The other embodiments of the invention are according to claims 10 to 14, which provide high mechanical properties and cold formability achieved in the steel strips when manufactured according to the invention. These properties are biaxial tensile strength, bending capacity, hole expansion capacity, creep, ultimate tensile strength, total elongation, and elongation at the yield point. The invention is also incorporated into a steel strip that has undergone light cold rolling, also known as temper rolling, as described above. The invention is based on modifying the composition of steel and processing it using all the aforementioned steps to achieve the optimized microstructure. Due to the achieved microstructure, the steel strip in its cold-rolled and / or hot-rolled form has high cold formability and high mechanical properties. The essential elements for steel are Mn, C, and Al. Mn and C are austenite-stabilizing elements in steel and are therefore added to the steel in predetermined amounts to stabilize the austenite. Al is a ferrite stabilizer, but it widens the temperature range between Ac1 and Ac3 (Ac1 = the temperature at which the austenite transformation begins during heating; Ac3 = the temperature at which the austenite transformation is completed during heating). Al is added to increase the steel's robustness for industrial processing, as it makes the steel less sensitive to small, undesirable temperature variations during intercritical processing. The invention is not limited to the presence of other optional elements and unavoidable impurities in the steel. The ranges of these optional and unavoidable alloying elements are provided in the relevant claims. The high manganese (Mn) content in steel, ranging from 3 to 12% by weight, will cause significant manganese enrichment in the austenite during batch annealing of hot-rolled or cold-rolled steel strip. This manganese enrichment, along with carbon enrichment (since carbon is also an austenite-stabilizing element), increases the thermal stability of the intercritical austenite by suppressing the steel's Ms temperature (Ms = the temperature at which martensite transformation begins during cooling). Therefore, during cooling to room temperature after batch annealing of the rolled steel strip, the intercritical austenite does not transform to a greater degree into martensite, allowing a high amount of austenite (>20% by volume) to be retained in the steel's microstructure at room temperature.This retained austenite, with optimal mechanical stability, transforms into martensite during loading (forming or any other deformation), causing a transformation-induced plasticity (TRIP) effect. The high strength, high elongation, and high cold formability of the steel strips of the invention are achieved because the TRIP effect increases the work-hardening rate. A manganese content greater than 12 wt% will make continuous steelmaking difficult due to extreme segregation and will also change the plasticity enhancement mechanism from TRIP to TWIP (twinning-induced plasticity). A manganese content of less than 3 wt% will not provide sufficient manganese enrichment in the austenite to achieve adequate amounts of retained austenite in the microstructure at room temperature. Similar to the effects of Mn as described above, the distribution of C within the intercritical austenite during final batch annealing also increases the thermal stability of the austenite and causes its stabilization at the microstructure at room temperature. However, C is effective in smaller quantities than Mn, and therefore the C content range for modifying the steel chemistry in the present invention is 0.05 to 0.3 wt%. If the C content is below 0.05 wt, sufficient austenite stabilization is not achieved, and a C content above 0.3 wt will make subsequent processing of the strip after cold forming, such as spot welding, difficult. Welding is essential in the assembly of automotive components to the body, and therefore this aspect is very important to consider.Carbon is also added to the steel in the present invention to increase strength. Aluminum is not an austenite-stabilizing element in steel; rather, it is a ferrite-stabilizing element. However, it is added to steel at up to 3 wt% to lengthen the intercritical temperature range (from Ac1 to Ac3). With a high manganese (Mn) content, steel becomes sensitive to small variations in processing temperature during industrial-scale processing. Furthermore, aluminum ensures the robustness of the steelmaking process, so the batch annealing temperature of steel strips can be selected with small variations to achieve the desired microstructures. When aluminum is not deliberately added as an alloying element to steel (i.e., when the aluminum content is approximately 0.03 wt%, as is frequently the case when added as a deoxidizer to molten steel), more precise furnaces are required, but the invention will still function.The highest amount of Al is limited to 3% by weight to reduce oxide scale formation during hot rolling and rolling forces during hot and cold rolling. The combination of the steel composition and the steps of the method leads to the beneficial effects of the invention. Manganese (Mn), an essential alloying element for modifying the chemistry of steel, tends to segregate after melting when its content exceeds approximately 2% by weight. This will affect product performance, resulting in non-homogeneous properties and potentially leading to cracking during processing. Therefore, it is preferable that the cast plates be well homogenized. Good homogenization of the plates is achieved by using relatively high plate reheating temperatures, above 1150°C, preferably above 1200°C, and more preferably above 1250°C, for a sufficiently long time, preferably 60 minutes or more. Therefore, due to the relatively high alloy content of steel, rolling forces are high during the hot rolling of strips. To roll reasonably wide strips on an industrial scale, typically wider than 1000 mm, hot rolling is preferably carried out at relatively high temperatures in the austenitic phase of the steel, above Ar3, where Ar3 is the temperature at which ferrite begins to form in the steel during cooling. This can be ensured by using an initial finish rolling temperature (F1) of around 1000°C or higher. An F1 temperature lower than this can increase the hot rolling force and lead to intercritical hot rolling, which can make hot rolling difficult for large-scale industrial processing.In addition to increasing rolling strength, intercritical hot rolling can also cause insufficient recrystallization of the hot-rolled strip. Therefore, when cold rolling is applied to hot-rolled strips to reduce the gauge of the final steel product, it will not be possible to cold roll the material unless appropriate preprocessing is adopted. In particular, the coiled steel after hot rolling is subjected to an intermediate batch annealing treatment for 24 hours or more at a low temperature within the steel's intercritical temperature range. Since this is for a relatively long period, it is a batch annealing. The intermediate batch annealing temperature should be below 650°C because at higher temperatures, large amounts of retained austenite will form after the steel cools to room temperature. Also, large amounts of martensite may appear in the microstructure if a higher batch annealing temperature is used.Both martensite and retained austenite make cold rolling difficult by increasing the rolling strength. Although the martensite phase is hard, the retained austenite transforms into hard martensite during cold rolling itself, increasing the rolling strength. Therefore, an intermediate batch annealing of the coiled material is part of the method according to this approach, to keep the retained austenite and martensite content lower and increase the amount of ferrite. The ferrite phase does not provide as much work hardening during cold rolling as retained austenite and therefore keeps the rolling strength low, making cold rolling possible. The minimum amount of ferrite phase required after this batch annealing of hot-rolled steel to make it cold-roll sensitive is 60% by volume. A final batch annealing of the hot-rolled steel strip—in the case of a hot-rolled product—or the cold-rolled steel strip—in the case of a cold-rolled product—is critical for obtaining the desired microstructural components in the final product for the invention to function. This final batch annealing must be carried out at an intercritical temperature (in the Ac1 to Ac3 range) below 700°C, preferably below 660°C. This is because thermodynamic calculations suggest that, for the chemical range of the steel of the invention, maximum carbon enrichment in the intercritical austenite occurs below 660°C, while manganese enrichment increases monotonically with decreasing temperature from 700°C. Consequently, with a final batch annealing temperature below 660°C, the optimal maximum enrichment (C + Mn) in the intercritical austenite is ensured.The annealing temperature should be selected to achieve maximum partitioning of Mn and C within the austenite. During this final batch annealing at an intercritical temperature, there will be an enrichment of C and Mn in the intercritical austenite, as C and Mn are austenite stabilizers. Although C diffuses more rapidly into the partition because it is a small interstitial element in steel, Mn, being a large replacement element, diffuses slowly. Therefore, a batch annealing time of 5 hours or more, preferably 10 hours or more, is necessary to achieve a high amount of Mn in the austenite. The Mn enrichment in the austenite should be such that the Mn content is at least 1.25 times the bulk Mn content of the steel, preferably at least 1.4 times. The enrichment of C must be such that the content of C is at least 1.Twice the bulk carbon content of the steel, preferably at least 2.3 times. These levels of manganese and carbon enrichment in the intercritical austenite are necessary to properly stabilize the austenite at room temperature so that at least 20% by volume of retained austenite is obtained in the microstructure at room temperature. These levels of manganese and carbon enrichment are also necessary to achieve optimal mechanical stability of the retained austenite (referred to as metastability) so that during deformation the steel can have a strain-hardening exponent of at least 0.3. If the manganese and carbon enrichments in the austenite are lower than the values ​​mentioned, optimal stability of the retained austenite is not achieved, and therefore the minimum strain-hardening exponent of 0.3 is also not attained.Consequently, when the annealing time is shorter than 5 hours, these requirements of the invention are not met. Similar disadvantages arise if a batch annealing temperature above 700°C is used. C and Mn enrichments in the intercritical austenite are not required, and therefore the critical austenite will not be sufficiently stable to provide a minimum of 20% by volume of retained austenite in the microstructure at room temperature after batch annealing, combined with the metastability required for the high strain-hardening rate needed. Therefore, a combination of high retained austenite fractions and optimal mechanical stability leads to the desired high strain-hardening rate. An annealing temperature above 700°C will also lead to more than 20% by volume of martensite, which will not provide the desired strain-hardening rate.It is this high strain hardening rate that leads to the high cold formability, as well as the high combination of strength and ductility in the final product. The high strain hardening rate strengthens the steel sheet while simultaneously thinning it during forming (e.g., by drawing), resulting in high cold formability. The final batch annealing is carried out in a non-oxidizing, nitrogen-free atmosphere to minimize any surface degradation of the steel strips due to oxygen and nitrogen. Since a minimum of 5 hours is required for batch annealing, the steel surface may oxidize if a non-oxidizing atmosphere is not used. Decarburization will also occur, decreasing the carbon content of the steel and making the invention less effective. For the same reason, nitrogen can react with the aluminum present on the steel surface, forming nitrides. All these forms of surface degradation are detrimental to the mechanical properties and formability of the steel. Preferred annealing atmospheres are vacuum, hydrogen, or argon atmospheres. The modifications to the steel and its processing lead to the appropriate microstructure in the final product for the invention to be successful. High fractions of retained austenite (>20% by volume), with low fractions of martensite (<20% by volume) and optimal fractions of ferrite (30–70% by volume) provide the combination of high strength, ductility, and formability due to the high strain hardening rate. A retained austenite content above 65% by volume cannot be achieved within the limits of the steel composition and is not necessary to achieve the required strain hardness rate. Furthermore, more than 70% retained austenite by volume can also cause problems in spot welding, resulting in severe brittleness of the molten metal, as well as poor resistance to hydrogen embrittlement in service.Consequently, the composition limits of the steels of the invention have been chosen considering these factors. If the ferrite content is above 70% by volume, the minimum high strain hardening exponent of 0.3 will not be achieved, which comes largely from metastable retained austenite through the TRIP effect. A ferrite fraction below 30% by volume is not necessary to obtain the minimum strain hardening rate. The martensite phase contributes primarily to strengthening, but not significantly to the strain hardening rate. Furthermore, high amounts of martensite (>20% by volume) can create weak interfaces with softer phases such as ferrite and retained austenite. These interfaces are detrimental to high ductility and formability, as they act as nucleation sites for the initiation of damage.Therefore, the martensite content must be kept below 20% by volume, including martensite not present. The ultrafine grain size is another microstructural requirement for the invention. The ferrite grain size must be below 2 µm, in the range of 0.2–2 µm. This ultrafine grain size imparts good ductility to the product and also causes strengthening through grain refinement, which contributes to the good mechanical properties of the steel strips of the invention. This ultrafine microstructure is also ensured by selecting a low final batch annealing temperature, below 700°C, which restricts grain growth. Furthermore, due to the composition of the steel of the invention and the requirement for intercritical final batch annealing, the phases at the annealing temperature (ferrite and austenite) are restricted to each other and cannot grow. All these factors lead to the desired ultrafine ferrite grain size. A ferrite grain size larger than 2 µm will result in lower strength and ductility.Grain size is primarily expressed here by grain length. Due to sufficient recrystallization of the steel strip during final batch annealing, the ferrite grain width will be more than 1 / 3 of the length. This will result in the relatively equiaxed shape of the ferrite grains after final batch annealing. One effect on the formability of the sheet metal with the aforementioned ultrafine ferrite grain size could be the appearance of yield point elongation in the engineering or design stress-strain curve of the products. This can be detrimental to the cold formability of the steel due to the localization of the strain, as well as to the degradation of the aesthetic appearance of the cold-formed components. Therefore, the process variables were selected such that the ferrite fraction in the final microstructure is a maximum of 70% by volume. This will ensure that yield point elongation, if present, is limited to a maximum of 10% of the design strain to obtain the best sheet formability and / or aesthetic appearance of the cold-formed components. However, a yield point elongation greater than 10% of the design strain is not a limiting factor for the operation of this invention.This is because, through temper rolling before forming and / or with appropriate lubrication during forming, the potential negative effects of yield point elongation can be mitigated. To eliminate yield point elongation, if present, the steel strips are optionally subjected in this invention to a small amount of reduction by cold rolling, temper rolling, or light cold rolling, with a thickness reduction of up to 5%. This small amount of rolling applied in one or more processes will eliminate the yield point elongation without altering the mechanical properties of the steel strips to any appreciable degree. However, the invention still provides high cold formability without the temper rolling step, even if yield point elongation is present, achieving up to less than 10% of the design strain. Optionally, hot-rolled or cold-rolled strips, after final batch annealing, are coated with a metallic coating to enhance their aesthetic appearance and corrosion resistance in service. The coating process may include, but is not limited to, hot-dip galvanizing, electro-annealing, electrogalvanizing, PVD, CVD, etc. The final batch annealed steel strip is very robust in terms of its microstructure and therefore its characteristics are not significantly altered by the application of a thin coating. The cold forming of steel strips, sheets, or parts can be carried out with or without the application of suitable lubrication to reduce friction between the steel and the tools. In both cases, the invention provides high cold formability. Non-limiting examples of lubrication systems include light oil, Klüber press paste, Teflon sheet, or combinations thereof. The steel used in the method according to the invention is a medium-manganese steel comprising carbon, manganese, and aluminum as its principal constituents. Optionally, other alloying elements selected from silicon, chromium, vanadium, niobium, titanium, and molybdenum may be present. Unavoidable impurities such as nitrogen, phosphorus, sulfur, oxygen, copper, nickel, tin, antimony, etc. (originating from the starting materials used to prepare the steel composition) may be present. These impurities are not intentionally added or specifically controlled within predetermined limits. The remainder of the steel composition is iron. Carbon is present in an amount of 0.05–0.3 wt%, as 0.05–0.20 wt%, preferably 0.07–0.20 wt%. It is added primarily for strength, although carbon also contributes to stabilizing the austenite. In the present composition, the austenite-stabilizing effect of manganese is much more pronounced due to its higher proportion. A preferred range for carbon is 0.05–0.25 wt%, and a more preferred range is 0.08–0.21 wt%. Too little carbon will not yield the desired strength level of 800 MPa, and if carbon exceeds 0.21 wt%, the weldability of the formed parts may be compromised. Manganese is present in amounts ranging from 3.0 to 12.0 wt%. Manganese reduces Ac1 and Ac3 temperatures, stabilizes austenite, increases strength and toughness, and causes the TRIP effect by stabilizing austenite in the microstructure at room temperature. At levels below 3.1 wt, the desired effects are not achieved, while amounts above 10.5 wt will cause casting problems and segregation. Furthermore, the deformation mechanism would change from transformation-induced plasticity (TRIP) to twinning-induced plasticity (TWIP). If the Mn content is too low, insufficient austenite will be retained at room temperature, and the stability of the retained austenite will be too low, preventing the desired benefits of ductility and strain hardening. Preferably, the Mn content is in the range of 3.5–10.0 wt%. In one formulation, the Mn content reaches 5.0–9.0 wt. In another, it is 5.5–8.5 wt, and in still others, 6.0–7.5 wt. Aluminum is added to expand the temperature range from Ac1 to Ac3 to increase the robustness of the process for industrial applications. Al is present in amounts from 0.03 to 3.0% by weight, as 0.6–2.9% by weight, preferably in the range of 1.0–2.2% by weight. Silicon, if present, is added in an amount less than 1.5 wt% to increase strength by solid solution strengthening. If present, the amount is typically greater than 0.01 wt% and less than 1.5 wt%, with a preferred range of 0.1–1.0 wt%. Both aluminum (Al) and silicon (Si) contribute to suppressing cementite precipitation, thus preventing ductility deterioration. Furthermore, both Al and Si increase the maximum annealing temperature, resulting in the highest amount of retained austenite at room temperature after the final batch annealing. Therefore, during intercritical annealing, manganese (Mn) diffusion is facilitated, leading to effective Mn partitioning within the austenite. One or more additional microalloying elements, selected from the group of V, Nb, Ti, and Mo, are optionally present. These microalloying elements increase strength through precipitation hardening by their carbides, nitrides, or carbonitrides. Cr, another optional element of this invention, also increases the maximum annealing temperature for achieving the highest amount of retained austenite at room temperature and reduces the sensitivity of the retained austenite content to the annealing temperature. This results in the effective partitioning of Mn in the austenite and increased process robustness during annealing. If present, the preferred additions of these optional alloying elements are: 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% by weight. The composition of the metallic coating is not limited. Zinc-based coatings, such as zinc containing essentially zinc, at least 0.1 wt% of aluminum, and optionally up to 5 wt% of aluminum and up to 4 wt% of magnesium, may be used. The remainder of the coating composition comprises additional elements, all individually in less than 0.3 wt%, and unavoidable impurities. Other additional elements that may be present in a small amount of less than 0.3 wt%, for example, to form flake and / or prevent slag formation, could be selected from the group comprising lead, antimony, titanium, calcium, manganese, tin, ladium, ce, chromium, nickel, zinc, and bismuth. Small amounts of these additional elements do not alter the properties of the resulting batch or coating to any significant degree for usual applications. Preferably, when one or more additional elements are present in the coating, each is present in an amount of less than 0.3 wt%.0.02% by weight, preferably each present in an amount of <0.01% by weight. The coating method may also vary from hot-dip galvanizing (Gl), electro-annealing, heat cycling to coat, electro-galvanizing. An aluminum-based coating, such as Al-Si-X coatings, may also be applied, where Si may vary from 0.1 to 10% by weight and X = any other coating-modifying element present in any required amount, plus unavoidable impurities that do not substantially alter the coating characteristics. Coating methods such as PVD, CVD, etc., are also applicable. The final batch annealing procedure is not limited by the type of furnace used or even the heating and cooling rates of the strip in the coil. It should be understood that the heating rate of a coil can vary from the surface to the center when subjected to batch annealing. However, for this invention, it is essential that the coiled strip be held at the target batch annealing temperature for a minimum of 5 hours, preferably more than 10 hours, so that each part of the coil experiences sufficient enrichment of C and Mn in austenite. The cooling rate after batch annealing is irrelevant to the invention since the presence of high amounts of Mn increases the hardenability of the steel. Thus, the coil can be cooled inside the batch annealing furnace, air-cooled, forced-air cooled, or even water-cooled. Tempered or light cold rolling can be performed on uncoated or coated steel strip. It can also be carried out in a single pass or in multiple passes. The steel strip obtained preferably has a triple or double microstructure comprising (in % by volume): ferrite: 30-70%; retained austenite: 20-65%; martensite: less than 20%, including 0%; and ferrite grain size: 0.2-2 pm. The resulting steel strip has the following characteristics of the retained austenite composition: Mn: 1.25 times the bulk Mn composition of the steel, preferably 1.4 times C: 1.2 times the bulk C composition of steel, preferably 2.3 times Advantageously, the steel strip has the following properties: flow: 600 MPa or more; ultimate breaking strength: 800 MPa or more; total elongation: 20% or more; strain hardening exponent: 0.3 or more; elongation at yield point: preferably 10% of the design strain or less; minimum bend angle at a thickness of 1.0 mm: 100° or more; hole expansion capacity: 20% or more; minimum stretching deformation in biaxial stretching: 10% or more. The phase fractions mentioned above were determined using X-ray diffraction (XRD). The amount of retained austenite was determined by XRD as a percentage of the sample thickness. XRD parameters were recorded in the 45–165° (2Θ) range using a standard Panalytical Xpert PRO powder diffractometer (CoKa radiation). Quantitative determination of the phase proportions was performed by Rietveld analysis using the Bruker Topas software package for Rietveld refinement. Martensite content was determined from peak shift at ferrite diffraction sites in the diffractograms. The grain size of the phases was determined from scanning electron microscopy (SEM) images of the microstructure. The Mn concentration of the retained austenite was determined using an electron probe microanalyzer (ERMA). The C content of the retained austenite was determined using the well-known formula developed by Dyson and Holmes. This formula relates the austenite lattice parameter, which can be determined from XRD data, to its C content. This formula can be found in the following article: DJ 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. Yield strength, ultimate tensile strength, yield elongation, and total elongation were determined from near-static tensile tests (strain rate of 3 x 10⁴ s⁻¹) at room temperature according to NEN 10002. The tensile specimen geometry consisted of an 80 mm gauge length in the rolling direction, 30 mm width, and a nominal thickness of 1.5 mm. The strain hardening rate was measured at a 7% strain interval after yield elongation on the tensile curve. Bendability was determined by three-point bend tests following VDA 238-100 on 40 mm x 30 mm specimens nominally 1.5 mm thick in both the longitudinal and transverse directions. The bend angles were along the 30 mm dimension, and the bend radius was 0.4 mm. The bending angles obtained from nominally 1 specimens.The 5 mm bend angles were converted to the corresponding 1.0 mm thickness using the following formula: bend angle at 1.0 mm thickness = measured angle × the square root of the actual thickness in mm. From these converted bend angles, for a specific heat treatment condition, the lowest value from the longitudinal and transverse specimens was taken to claim the ranges in this invention. The hole expansion capacity (HEC) was determined according to ISO / TS 16630:2003(E). Specimens with dimensions of 90 mm × 90 mm × 1.5 mm were cut from the steel strips. A 10 mm diameter hole was drilled in the middle of the specimens, and hole expansion tests were performed. The hole expansion capacity (HEC = (expansion of the initial hole diameter / initial hole diameter) × 100%) was calculated from the measured data.Biaxial strain was determined from biaxial stretching tests performed on the Erichsen press using a 75 mm diameter flat punch in combination with a 79.78 mm diameter die. The punch tip had a radius of 10 mm and the die, 8 mm. The workpiece clamping force was set to the machine's maximum capacity (-580 kN) to ensure no inward stretching occurred. The test speed was set to 20 mm / min. Strain was measured by applying a 10 mm square grid to the sheets with a fine marker. In the manufacture of steel strip, using the final batch annealing step at the steel's intercritical temperature below 700°C, as explained previously, Mn partitioning from ferrite to austenite occurs, increasing the stability of the intercritical austenite. During cooling after the final batch annealing, the intercritical austenite does not significantly transform to martensite due to its high stability resulting from the low Ms content, yielding a duplex microstructure of ferrite and retained austenite. For low Mn content, for example, below 8 wt%, some of the intercritical austenite may transform to martensite, but the martensite content will be 20% by volume or less. Therefore, due to the increased levels of Mn and the low batch annealing temperature (e.g., less than 700°C), a high amount of retained austenite (20% by volume or more) with optimal metastability can be guaranteed.This high amount of retained austenite is partially transformed into martensite during deformation in the forming step, causing a transformation-induced plasticity (TRIP) effect that results in a high strain hardening exponent (= high elongation and high conformability). The total elongation of the steel strip is preferably 20% or more, and the strain hardening exponent is 0.3 or more due to the steel composition. An intercritical batch annealing step of a medium-manganese approach steel is preferably used to obtain a mixed microstructure of ultrafine ferrite (0.5–2.0 micrometers) and areas of martensite and high retained austenite. Therefore, high ductility and a high hardening rate are achieved, leading to high cold formability of the steel strip. A preferred steel strip is used as a material for manufacturing automotive components, particularly those with complex shapes requiring strip formability. Components requiring high energy absorption combined with high strength are also suitable for manufacturing from steel strip. Non-limiting examples include interior car parts, B-pillars, and longitudinal bars. BRIEF DESCRIPTION OF THE FIGURE The invention will be clarified with reference to Figure 1 and the examples described below. Figure 1 shows an SEM microstructure of a steel manufactured according to this invention obtained by final batch annealing (steel A, 650°C / 10 hours) where F = ferrite, MA = martensite-austenite. DETAILED DESCRIPTION OF THE INVENTION Steel ingots of the three steel chemistries of the invention, A, B, and C, measuring 200 mm x 100 mm x 100 mm, were melted by melting the charges in a vacuum induction furnace. The chemical compositions of these steel chemistries of the invention, along with two reference steels, D and E, are provided in Table 1. Steel D is a twin-induced plasticity (TWIP) steel, and steel E is a DH1000 grade steel, both received in their final annealed cold-rolled condition. The thicknesses of these steels as received were 1.7 mm and 1.5 mm, respectively. They were reheated for 2 hours at 1250°C and as-rolled to a thickness of 30 mm.The strips were then reheated to 1250°C for 30 minutes and hot-rolled to a thickness of 3 mm for steels A and B and 4 mm for steel C, with a starting rolling temperature of 1150°C and a finish rolling temperature (FRT) of 900°C, which falls within the austenitic phase range for all three steels. The higher reheating temperature of 1250°C and the longer duration of 2 hours were used to ensure proper Mn homogenization. The austenite-to-ferrite (Ar3) transformation temperature for steels A, B, and C, measured by dilatometry, was 785, 770, and 723°C, respectively. The hot-rolled steels were then subjected to simulated coil quenching at 680°C in a muffle furnace and subsequently cooled to room temperature. Hot-rolled strips of A and B were then intermediate-batch annealed for 96 hours at 600°C, while strips of C were intermediate-batch annealed at 550°C in a muffle furnace under a protective argon atmosphere and air-cooled to room temperature. These annealing temperatures were selected to achieve the desired ferrite fractions to facilitate subsequent cooling-rolling in the process. The fractions of the phases of steels A, B and C after this intermediate batch annealing of the hot-rolled strips are provided in Table 2.Phase fractions were determined from the percentage of strip thickness by XRD measurements as described above. Ferrite fractions in all three steels were found to be greater than 60% by volume. Next, the strips were pickled in HCI acid at 90°C to remove the oxides, and then all the steels were cold rolled to a final thickness of 1.5 mm of their respective hot-rolled gauges. Cold-rolled strips of steel A and B were batch-annealed at 650°C for 10 hours, and those of steel C at 640°C for 4 and 16 hours, using a muffle furnace. An argon atmosphere was used for annealing, ensuring that the atmosphere was free of oxygen and nitrogen to minimize oxidation of the strips and any undesirable reaction between atmospheric nitrogen and the aluminum in the steel to form nitride layers on the surface. After annealing, the samples were air-cooled to room temperature. For comparison, cold-rolled strips of steel A were similarly annealed at 650°C for 2 minutes, 5 minutes, and 1 hour, and those of steel C at 640°C for 4 hours. Some specimens underwent light cold rolling or temper rolling to reduce thickness by up to 5%. The procedures for characterizing and testing the material have been described above. As a reminder, the microstructure of the samples was characterized using XRD and SEM. Microanalysis of phase chemistry was performed by ERMA and XRD analysis. Tensile properties were determined by tensile testing of specimens with a gauge length of 80 mm and a width of 30 mm (specimen geometry A80). The formability of the strips was evaluated by bending, hole expansion, and biaxial stretching tests using lubrication. For bending capability, the definitions of specimens L and T are as follows: L = longitudinal specimen where the bend axis is parallel to the rolling direction, T = transverse specimen where the bend axis is perpendicular to the rolling direction. A typical microstructure obtained after final batch annealing of cold-rolled strips of steel A is shown in Figure 1, where areas of ferrite and austenite martensite can be observed. The ultrafine ferrite grain size is also evident. The microstructural characteristics of steels A, B, and C are provided after different final annealing treatments of the cold-rolled samples. For all steels under all conditions, the ferrite grain size ranges from 0.5 to 1.9 pm. For steels A and C, the retained austenite content increases with increasing annealing time at their respective annealing temperatures due to greater Mn partitioning in the austenite. The retained austenite content was also higher with higher Mn content (steel C has more retained austenite than A and B), demonstrating the effects of Mn on austenite stabilization.Under all conditions, high fractions of retained austenite (above 33% by volume) were obtained except for steel A annealed for 2 minutes at 650°C. The Mn and C content of the annealing conditions of the steels provided in Table 4 shows that in the retained austenite of these three conditions, the Mn enrichment of the steels ranges from 1.286 to 2.139 times the bulk Mn content of the steels, except for steel A—condition 650°C / 2 min—where the Mn content is only 1.09 times the bulk Mn content. For C enrichment in retained austenite, the C content fluctuates from 1.17 to 3.085 times the bulk C content of the steels, except for A-condition steel at 650°C / 2 min where this value is 1.063 times.Due to these low C and Mn enrichments in the austenite, the retained austenite content of the 650°C / 2 min conditioning steel is also less than 20% by volume, and consequently, the martensite content is greater than 20% by volume (39.8% by volume). In all other steels and conditions of the invention, the martensite content is 16.7% by volume or less, including 0% by volume (C-640°C / 960 min steel). The lower retained austenite fraction for the A-steel condition of 650°C / 2 min is clearly due to the fact that the 2 min annealing time was too short for sufficient diffusion of Mn into the austenite even though the annealing temperature was in the intercritical temperature range of A-steel and below 700°C. The consequence of the aforementioned microstructural characteristics can be observed in the tensile properties of the steels provided in Table 5. The A-650°C / 2 min steel exhibited the lowest amounts of retained austenite. At 1.25 times Mn and 2 times C, based on their bulk Mn and C contents, respectively, it displayed very high yield strength and ultimate tensile strength, but a total elongation of only 3.1%. This is because, during the tensile test, all of its retained austenite transforms very rapidly to martensite due to its low stability resulting from the low Mn and C enrichment. The small amount of retained austenite is consumed very early during deformation without yet showing any yield point elongation. Thus, the tensile properties of this steel condition are poor and unsuitable for cold forming.On the other hand, steel A under other annealed conditions and steels B and C under all conditions showed yield strengths greater than 693 MPa, ultimate tensile strength greater than 860 MPa, and total elongation greater than 23.4%. These steels also exhibited high energy absorption capacity (determined by the product of ultimate tensile strength and total elongation) and varying amounts of yield point elongation. The yield point elongation decreased with annealing time for steels A and C due to the increase in ferrite grain size, as shown in Table 3. The tensile properties of these steel compositions of the invention can be compared with the reference steels listed in Table 6.The steels of the invention, under extended final batch annealing conditions, exhibit significantly higher total elongation and energy absorption capacity than conventional DH1000 grade steel (reference steel E) due to the combination of steel chemistry, processing, and microstructure. Steel E has a very low amount of retained austenite in its microstructure. Furthermore, although TWIP steel (reference steel D) has a much higher total elongation than the steels of the invention, the energy absorption capacity of some of the steels of the invention falls within the range of TWIP steel, which has a completely austenitic microstructure. The formability parameters of the steels of the invention are shown in Table 7 in comparison with the reference steels. The formability parameters compared are biaxial stretchability in terms of rolling and transverse strains, bending capacity in the longitudinal and perpendicular directions of the sheet, and flange capacity according to the HEC values. Steel A shows that when the final batch annealing is carried out for less than 10 hours at 650°C, the biaxial stretch strain values ​​are 0, although the other parameters are not zero. Steel A annealed at 650°C for 2 minutes also showed very low bending and flange capacity. Although bending and flange capacity improve with increasing annealing time, the material is not stretchable until 10 hours after the final batch annealing.Steel B, annealed at 650°C for 10 hours, also showed formability parameters similar to those of steel A under the same annealing conditions. Steel C, annealed at 640°C for 4 hours, showed high bending and flangeability but low elongation. When steel C was annealed for 16 hours, its elongation also improved. The cold formability of steel sheets is a combination of several parameters, such as elongation, bendability, and flangeability. When the sheets of the invention are batch-annealed below 700°C within their intercritical temperature range, the annealing time is important to provide the required amounts of Mn and C enrichment in the retained austenite, as noted previously, since Mn is an element that diffuses slowly from steel. High Mn and C enrichment is necessary to achieve a high strain hardening exponent. Therefore, steels annealed in less than 10 hours, exhibiting the lowest strain hardening exponent, also exhibited low elongation, although other formability parameters were good. A high strain hardening exponent above 0 is required.3. To obtain good elongation capacity in the steels of the invention; otherwise, premature local fracture may occur. Consequently, it appears from the results that for good cold formability (combination of elongation capacity, bending capacity, and flange capacity) as mentioned in the claims, a minimum of 10 hours of final batch annealing is necessary for the steels of the invention to obtain the minimum values ​​of Mn and C enrichment in the retained austenite. When the formability of the samples annealed in 10-hour batches is compared with the steels of the invention, it is observed that the formability parameters of the steels of the invention fall within the range of highly formable TWIP steel (reference steel E) and are significantly higher than those of conventional DH1000 (reference steel D). The biaxial stretching deformations of steel C, even when annealed for only 4 hours, are greater than those of conventional DH1000. This high cold formability of the steels of the invention is due to the high fractions of retained metastable austenite with high Mn and C enrichment obtained in the steels of the invention through the processing steps of the invention. The effects of quench rolling on the mechanical properties of B steel annealed for 10 hours at 650°C are shown in Table 8. It is evident that the elongation at the yield point decreases with increasing reduction by quench rolling. With a 2% reduction, the elongation at the yield point disappears. The tensile properties did not change much, remaining within the claimed range of this invention. More importantly, the strain hardening exponent also remains high even up to a 5% thickness reduction.Therefore, this elimination of the yield point elongation without appreciable change in mechanical properties with temper rolling up to 5% will make the steel strips of this invention even more cold-formable, since this will reduce the risk of localized deformation during stretch forming as well as the marks of the stretcher on the surface of the formed articles. Table 1: Steel composition in % by weight. Steel C Mn Si Al PSB Cr Mo Ni Cu A 0.094 7.15 0.20 1.54 0.001 0.0014 0.0001 0.003 0.01 0.0015 0.02 B 0.13 7.32 0.22 1.57 0.001 0.0011 0.0002 0.004 0.001 0.002 0.03 C 0.16 9.81 0.19 1.40 0.002 0.0018 0.0001 0.024 0.001 0.014 0.03 D 0.72 14.5 0.25 0.05 0.002 0.0012 0.0003 0.030 0.002 0.003 0.03 E 0.15 2.24 1.0 0.033 0.001 0.0001 0.0002 0.002 0.001 0.016 0.02 Table 1 (continued) Steel Nb Ti VWN Sn Co Fe Observation A 0.0007 0.001 0.00014 0.001 0.005 0.0010 0.001 Bal. Invention B 0.0008 0.001 0.00013 0.001 0.004 0.0008 0.001 Bal. Invention C 0.0005 0.002 0.001 0.001 0.003 0.0007 0.0005 Bal. Invention D 0.0004 0.0001 0.002 0.002 0.006 0.0009 0.0004 Bal. Reference (TWIP) E 0.001 0.0016 0.001 0.001 0.0033 0.0008 0.0003 Bal. Reference (DH1000) Table 2: Phase fractions after intermediate batch annealing of hot-rolled strips Steel Ferrite (% by volume) Retained Austenite (% by volume) Martensite (% by volume) A 83.3 14.4 2.3 B 79.5 18.4 2.1 C 71.6 28.4 0 Table 3: Phase fractions and average grain sizes of austenite and ferrite after final batch annealing of cold-rolled strips Steel Annealing Temperature (°C) Annealing Time (minutes) Phase Fractions (% by volume) Average Grain Size (qm) Ferrite Retained Austenite Martensite Ferrite Retained Austenite A 650 2 49.5 10.7 39.8 0.50 0.20 5 50.1 33.2 16.7 0.70 0.34 60 49.3 36.9 13.8 0.85 0.38 600 49.1 40.2 10.7 1.20 0.40 B 650 600 48.9 42.7 8.4 1.30 0.41 C 640 240 41.2 56.3 2.5 1.01 0.50 960 39.5 60.5 0 1.90 0.89 Table 4: Manganese and carbon enrichments in the retained austenite after final annealing of the cold-rolled material Steel Annealing temperature (°C) Annealing time (minutes) Mn content (% by weight) C content (% by weight) A 650 2 7.8 0.10 5 9.2 0.11 60 10.1 0.18 600 13.0 0.29 B 650 600 13.8 0.31 C 640 240 11.1 0.22 960 15.3 0.29 Table 5: Tensile properties of steels after annealing of cold-rolled steels Steel Annealing Temperature (°C) Annealing Time (minutes) Lower Yield Strength (MPa) Ultimate Tensile Strength (MPa) Total Elongation (%) Elongation at Yield Point (%) Strain Hardening Exponent Energy Absorption (% of MPa) A 650 2 1258 1323 3.1 NANA 4101.3 5 939 959 23.4 14.3 0.15 22440.6 60 844 860 28.5 13.2 0.15 24510 600 693 887 29.3 4.6 0.36 25989.1 B 650 600 702 910 32.1 5.2 0.37 29211 C 640 240 837 1149 35.1 14.2 0.37 40329.9 960 781 1114 42.3 8.7 0.42 47122.2 Table 6: Tensile properties and microstructure of the reference steels Steel Yield Strength (MPa) Ultimate Tensile Strength (MPa) Total Elongation (%) Energy Absorption (% of MPa) Microstructure D 632 979 47.3 46306.7 100% by volume of Austenite E 798 1005 10.5 10552.5 8% by volume of retained Austenite + 45% by volume of Martensite + 47% by volume of Ferrite Table 7: Formability parameters of steels after final annealing Steel Annealing Temperature (°C) Annealing Time (minutes) Biaxial Stretch Deformation (%) Bendability (°) Hole Expansion Capacity (%) Rolling Direction Cross Direction LTA 650 2 0 0 21.4 44.0 1 5 0 0 81.2 116.0 15 60 0 0 159.9 155.6 41 600 17.5 10.0 153.9 154.1 28 B 650 600 18.1 10.3 151.2 153.6 32 C 640 240 6.1 5.2 145.6 147.1 21 960 16.2 10.5 133.7 135.6 27 D NA 2.0 3.0 96.5 100.2 15 E 17.5 19.0 143.1 147.8 41 Table 8: Tensile properties of steel B in quenched rolled conditions applied after heat treatment to the best formable condition Batch Annealing Final Reduction by Rolling Quenching (%) Elongation at Yield Point Yield Strength (MPa) Ultimate Tensile Strength (MPa) Total Elongation (%) Strain Hardening Exponent 650°C, 10 hours 0 5.2 702 910 32.1 0.37 0.5 3.5 695 878 32.1 0.37 1 0.5 703 881 32.0 0.37 2 0 693 883 31.5 0.37 5 0 768 927 30.4 0.36

Claims

CLAIMS 1. A method of manufacturing a cold-rolled and annealed steel strip, the steel composition being in wt%: C: 0.05-0.3; Mn: 3.0-12.0; Al: 0.03-3.0; optionally one or more 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; unavoidable impurities, such as S: less than 30 ppm; P: less than 0.04; the remainder being Fe; the method being characterized in that it comprises the steps of: - molding the molten steel into a plate; - reheating the plate and holding it at a temperature of 1150°C or more for a time of 1 hour or more; - hot roll the steel into a strip, preferably with an average plate entry temperature F1 above 1000°C; - coil the hot-rolled steel strip; - pickle the steel strip;- Carry out batch intermediate annealing of the steel strip at a temperature below 650°C for more than 24 hours to achieve at least 60% ferrite by volume after cooling to room temperature; - cold roll the steel into a cold-rolled steel strip and coil this; - carry out batch annealing of the coiled steel strip; - at an intercritical temperature between Ac1 and Ac3 which is below 700°C; - in a non-oxidizing and nitrogen-free atmosphere; - the total annealing time for which the strip is held at the critical temperature being at least 5 hours, preferably at least 10 hours to achieve Mn enrichment in the austenite such that the Mn content is at least 1.25 times the bulk Mn content of the steel and C enrichment such that the C content is at least 1.2 times the bulk C content of the steel;- Cooling the steel after batch annealing in air, forced air, or water cooling.

2. A method of manufacturing a hot-rolled and annealed steel strip, the steel composition being in % by weight: C: 0.05-0.3; Mn: 3.0-12.0; Al: 0.03-3.0; optionally one or more 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; unavoidable impurities, such as S: less than 30 ppm; P: less than 0.04; and the remainder being Fe; the method is characterized in that it comprises the steps of: - molding the molten steel into a plate; - reheating the plate to a temperature of 1150°C or more, for a time of 1 hour or more; - hot rolling the steel into a strip, preferably with an average entry temperature of the plate F1 above 1000°C; - coiling the hot-rolled steel strip; - pickling the steel strip; - batch annealing the coiled steel strip: - at an intercritical temperature between Ac1 and Ac3 which is less than 700°C; - in a non-oxidizing and non-nitrogenous atmosphere; - the total annealing time being the time for which the strip is held at the critical temperature for at least 5 hours, preferably at least 10 hours to achieve the enrichment of Mn in the austenite so that the Mn content is at least 1.25 times the bulk Mn content of the steel and the enrichment of C so that the C content is at least 1.2 times the bulk carbon content of the steel; - cool the steel after batch annealing in air, forced air or by water quenching.

3. The method according to claim 1 or claim 2, further characterized in that the reheating of the iron is to a temperature of 1200°C or more.

4. The method in accordance with any of claims 1 to 3, further characterized in that the reheating of the iron is to a temperature of 1250°C or more.

5. The method in accordance with any of claims 1 to 4, further characterized in that the batch annealing of the coiled steel strip takes place at an intercritical temperature below 660°C.

6. The method in accordance with any of claims 1 to 5, further characterized in that the resulting strip is coated with any metallic coating applied by hot-dip galvanizing, electro-annealing, electro-galvanizing, aluminizing or any other method such as PVD, CVD.

7. The method according to claim 1 or any of claims 3 to 6, further characterized in that the resulting steel strip is subjected to light cold rolling.

8. The method according to claim 7, further characterized in that the light cold rolling takes place with a thickness reduction of 5% or less.

9. A steel strip obtainable by the method according to any one of claims 1 to 8, characterized in that the steel strip has a steel composition which is in wt%: C: 0.05-0.3; Mn: 3.0-12.0; Al: 0.03-3.0; optionally containing one or more 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; unavoidable impurities, such as S: less than 30 ppm; P: less than 0.04; and the remainder being Fe; where the steel strip has a retained austenite composition having a Mn content that is at least 1.4 times the bulk Mn content of the steel composition as well as a C content that is at least 2.3 times the bulk C content of the steel composition, to obtain metastable retained austenite to give the steel a high strain hardening exponent of at least 0.3 measured after yield point elongation for a 7% stress range in a quasi-elastic tensile test, wherein the microstructure after final batch annealing of the coiled steel strip comprises in % by volume: ferrite: 30-70%; retained austenite: 20-65%; martensite: < 20%, including 0% by volume.

10. The steel strip according to claim 9, further characterized in that the size of the ferrite grains is 0.2-2 pm.

11. The steel strip according to any of claims 9 or 10, further characterized in that the length / width ratio of the ferrite grains is 3 or less.

12. The steel strip according to any of claims 9 to 11, further characterized in that it has a yield point elongation greater than 10% of the design strain measured from its engineering stress-strain curve.

13. The steel strip according to any of claims 9 to 12, further characterized in that it has a yield strength of 600 MPa or more and an ultimate tensile strength of 800 MPa or more and a total elongation (A80) of 20% or more.

14. The steel strip according to any of claims 9 to 13, having very high formability, is further characterized in that it has a single-direction stretch deformation in a biaxial strain condition of 10% or more, a bend angle VDA of 10°, a thickness of 1.0 mm, and a hole expansion capacity of 20% or more.