Method of manufacturing a cold formable high strength steel strip and steel strip

Abstract

Method of manufacturing a steel strip comprising the steps of: - casting a molten steel into a slab; - reheating the slab at a temperature of 1150°C or more for a time of 1 hour or more; - hot rolling the steel into a strip, preferably using an average F1 slab entry temperature of more than 1000°C; - coiling the hot rolled steel strip; - batch annealing the steel strip: - at a temperature in the intercritical range (i.e. between Ac1 and Ac3), preferably less than 700°C; - in a non-oxidizing and non-nitrogenizing atmosphere; - for a total annealing time of at least 5 hours, preferably at least 10 hours so as to obtain: a Mn enrichment in austenite such that the Mn content is at least 1.25 times the overall Mn content of the steel; and a C enrichment such that the C content is at least 1.2 times the overall C content of the steel; - cooling the steel after batch annealing in air, forced air or water quenching.

Description

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

[0002] Cold forming, cold stamping, or cold pressing of steel sheets is a method for producing steel components for various applications in manufacturing industries such as automotive, construction, engineering, and infrastructure. It is known that the cold formability of steel sheets decreases with increasing steel strength. This is especially true for conventional steel sheets and first-generation advanced high-strength steels (AHSS).

[0003] Due to this inverse relationship between strength and elongation in steel, the distribution of tensile strength and elongation properties in these steels is sometimes referred to as a "banana diagram." However, high formability at high strength can be achieved using the so-called second-generation AHSS (2GAHSS) concept, but these steels are typically highly alloyed and contain expensive alloying elements. Examples include high-manganese twinned induced plasticity (TWIP) steels with a typical Mn content exceeding 12 wt% and stainless steels containing significant amounts of expensive alloying elements such as chromium, nickel, and molybdenum. Besides being very expensive, another drawback of 2GAHSS is that their very high alloy content makes them very difficult to manufacture on a large industrial scale.

[0004] To overcome the problems of 2GAHSS while still achieving reasonably high cold formability at high strength, various third-generation AHSS (3GAHSS) concepts have been introduced, such as quenched and partitioned (Q&P) steels, carbide-free bainitic (CFB) steels, and medium-Mn steels. These steels are less expensive than 2GAHSS and can be easily processed in existing steel mill equipment. This invention focuses on medium-Mn type 3GAHSS.

[0005] WO16001887 discloses a method for manufacturing high-strength steel sheets, wherein the steel contains, by weight percentage, 0.1 ≤ C ≤ 0.4, 4.2 ≤ Mn ≤ 8, 1 ≤ Si ≤ 3, 0.2 ≤ Mo ≤ 0.5, with the balance being Fe and unavoidable impurities. The method includes continuous annealing with a temperature greater than Ac3, quenching to a temperature between the start (Ms) and end (Mf) of martensite, over-aging for more than 10 s between 300-500 °C, and cooling. This is essentially a quench and partition (Q&P) method, in which C enrichment (and possibly some Mn enrichment) in austenite is achieved through an over-aging step in steel containing a certain amount of martensite. The Q&P method is entirely different from the critical zone annealing method. In this document, a large amount of Mn partitioning into the austenitic steel is not expected because Mn diffusion in the steel is extremely slow at the lower over-aging temperatures (300-500 °C).

[0006] WO2017021464 discloses high-tensile steel in hot-rolled or cold-rolled strip form, having a chemical composition (in weight %): 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 balance being iron and unavoidable elemental inclusions related to steel. If the steel is flexibly hot-rolled, optionally annealed, or flexibly cold-rolled, optionally annealed, it is further flexibly cold-rolled and subsequently annealed at an annealing temperature of 600°C to 750°C for 1 minute to 48 hours. This patent applies flexible rolling by controlling the roll gap, wherein shear conditions vary across the width of the strip. Flexible rolling is a different method for parts with varying wall thicknesses, unlike the conventional rolling method of this invention, in which a uniform product thickness is obtained across the entire width. The disadvantage of flexible rolled strip is that it becomes uneven across the width of the strip.

[0007] The present invention aims to provide a highly cold-formable steel strip with a cold-rolled thickness that maintains high strength under both coated and uncoated conditions.

[0008] Another object of the present invention is to provide a highly cold-formable steel strip with a range of hot-rolled thicknesses that maintains high strength under both coated and uncoated conditions.

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

[0010] This invention is primarily embodied in a method for manufacturing cold-rolled and annealed steel strip, wherein the steel composition, in weight percent, is:

[0011] C: 0.05-0.3;

[0012] Mn: 3.0-12.0;

[0013] Al: 0.03-3.0;

[0014] One or more additional alloying elements may be selected:

[0015] Si: less than 1.5;

[0016] Cr: less than 2.0;

[0017] V: less than 0.1;

[0018] Nb: less than 0.1;

[0019] Ti: less than 0.1;

[0020] Mo: less than 0.5;

[0021] Unavoidable impurities, such as

[0022] S: Less than 30ppm;

[0023] P: less than 0.04; and

[0024] The balance is Fe;

[0025] The method includes the following steps:

[0026] - Cast molten steel into slabs;

[0027] - Preheat the slab and hold it at 1150°C or higher for 1 hour or more;

[0028] - Hot rolling steel into strip, preferably using an average F1 slab entry temperature greater than 1000°C;

[0029] - Winding hot-rolled steel strip;

[0030] - Pickled steel strip;

[0031] - The steel strip is annealed in intermediate batches at a temperature below 650°C for more than 24 hours to achieve at least 60% ferrite by volume after cooling to room temperature;

[0032] - Cold rolling steel into cold-rolled steel strip and then winding it;

[0033] - Anneal the wound steel strip in batches:

[0034] -Critical temperature range of less than 700°C between Ac1 and Ac3;

[0035] -In a non-oxidizing and non-nitriding atmosphere;

[0036] - The total annealing time for maintaining the strip at the critical zone temperature is at least 5 hours, preferably at least 10 hours, to obtain: Mn enrichment in austenite such that the Mn content is at least 1.25 times the overall Mn content of the steel; and C enrichment such that the C content is at least 1.2 times the overall C content of the steel;

[0037] - Cooling the steel after batch annealing in air, in forced air, or by water quenching.

[0038] Steel containing substantially 0.05 to 0.3 wt% C, 3.0–12.0 wt% Mn, 0.03–3.0 wt% Al, and optional other alloying elements and unavoidable impurities is processed into hot-rolled specifications using a specific processing route. The molten steel is cast into a slab, which is then reheated to 1150°C or higher for 1 hour or more. The slab is then hot-rolled into strip, preferably at a finishing mill inlet temperature (F1) greater than 1000°C. The F1 inlet temperature is the entry temperature of the strip into the first pedestal of the finishing mill. The finishing mill is the part of the hot rolling mill where finishing rolling takes place after roughing or slab rolling of the slab in the roughing mill and before cooling on the output roll table. After passing through the output roll table, the hot-rolled strip is wound into coils, and these coils are then intermediately annealed in batches at a temperature below 650°C for at least 24 hours to achieve at least 60 vol% ferrite in the strip after cooling to room temperature. The steel strip is then pickled in an acid solution at, for example, 50-90°C and cold-rolled to thinner specifications. This invention is not limited to hot-rolled or cold-rolled specifications. However, typically hot-rolled specifications will be in the range of 2 to 10 mm and cold-rolled specifications in the range of 0.5 to 2 mm. The cold-rolled steel is then batch-annealed at a critical temperature below 700°C, preferably below 660°C, in a non-oxidizing and non-nitriding atmosphere for a duration of at least 5 hours, preferably at least 10 hours, such that the Mn content in the critical austenite reaches at least 1.25 times the overall Mn content of the steel and the C content in the critical austenite reaches at least 1.2 times the overall C content. The longer duration of 10 hours is preferred because a greater amount of Mn can partition from ferrite into austenite during annealing. Mn typically takes a long time to diffuse because it is a large substitutional alloying element in iron. As the batch annealing temperature decreases, the manganese enrichment in the critical austenite increases, which makes the austenite in the steel more stable upon cooling to room temperature after batch annealing. Batch annealing time is defined as the time period during which the steel strip is held at the batch annealing temperature, excluding the time for heating the strip to the target temperature.

[0039] A final batch annealing will occur, i.e., the batch annealing of the wound strip herein continues for the duration mentioned in the claims, long enough to obtain a relatively equiaxed ferrite grain morphology in the steel, wherein the ratio of grain length to width is preferably 3 or less. The steel is then cooled to room temperature at any cooling rate, for example in air, forced air, or water.

[0040] By performing the method according to the present invention, the following advantages are obtained:

[0041] - Medium Mn steel containing the mentioned alloying elements, along with 3-12 wt% Mn, reduces Mn segregation. This is a consideration affecting mechanical properties when Mn is present in a relatively high amount, as in this invention. A relatively high slab reheating temperature of greater than 1150°C, preferably 1200°C, and more preferably greater than 1250°C, and a minimum reheating time of 1 hour are chosen to minimize segregation and ensure uniform Mn distribution in the matrix. However, the final mechanical properties may be compromised. The choice of 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, a reheating temperature close to 1150°C will be sufficient to ensure uniform Mn distribution, and higher slab reheating temperatures are preferred as the Mn content increases.

[0042] Steel can be hot-rolled on an industrial scale to achieve strip widths of considerable size, such as greater than 1000 mm. This is achieved by maintaining a high F1 temperature (greater than 1000°C) during hot rolling, thereby reducing the required hot rolling force. Using lower F1 temperatures, hot rolling of steel strip becomes difficult.

[0043] - Steel becomes suitable for cold rolling on an industrial scale. This is made possible by using an intermediate batch annealing step in hot-rolled steel. Intermediate batch annealing is preferably carried out at a temperature below 650°C in the critical zone of the steel, at which such a selected temperature results in at least 60% vol% ferrite in the steel strip, with the balance being retained austenite and martensite.

[0044] - Further batch annealing of the wound steel strip at temperatures below 700°C actually produces a suitable microstructure. This should last for at least 5 hours, preferably at least 10 hours. During this process step, the distribution of Mn and C in the steel of the present invention between the critical region austenite and ferrite results in austenite becoming highly enriched by Mn and C, stabilizing the phases until room temperature. Because the austenite enrichment results in an Mn content at least 1.25 times the overall Mn content of the steel and a C content at least 1.2 times the overall C content of the steel, the steel is practically insensitive to the actual cooling rate and therefore can be cooled in air, forced air, or water after batch annealing. The lower the batch annealing temperature, below 700°C, the more abundant the Mn content is in the critical region austenite.

[0045] A second embodiment of the present invention is a method for manufacturing hot-rolled and annealed steel strip, wherein the steel composition, in weight percent, is:

[0046] C: 0.05-0.3;

[0047] Mn: 3.0-12.0;

[0048] Al: 0.03-3.0;

[0049] One or more additional alloying elements may be selected:

[0050] Si: less than 1.5;

[0051] Cr: less than 2.0;

[0052] V: less than 0.1;

[0053] Nb: less than 0.1;

[0054] Ti: less than 0.1;

[0055] Mo: less than 0.5;

[0056] Unavoidable impurities, such as

[0057] S: Less than 30ppm;

[0058] P: less than 0.04; and

[0059] The balance is Fe;

[0060] The method includes the following steps:

[0061] - Cast molten steel into slabs;

[0062] - Reheat the slab at 1150°C or higher for 1 hour or more;

[0063] - Hot rolling steel into strip, preferably using an average F1 slab entry temperature greater than 1000°C;

[0064] - Winding hot-rolled steel strip;

[0065] - Pickled steel strip;

[0066] - Anneal the wound steel strip in batches:

[0067] -Critical temperature range of less than 700°C between Ac1 and Ac3;

[0068] -In a non-oxidizing and non-nitriding atmosphere;

[0069] - The total annealing time for maintaining the strip at the critical zone temperature is at least 5 hours, preferably at least 10 hours, to obtain: Mn enrichment in austenite such that the Mn content is at least 1.25 times the overall Mn content of the steel; and C enrichment such that the C content is at least 1.2 times the overall C content of the steel;

[0070] - Cooling the steel after batch annealing in air, in forced air, or by water quenching.

[0071] As described above, steel processed up to the hot rolling step according to the first embodiment is then skipped after pickling and directly subjected to the final batch annealing step. The batch annealing is carried out for a sufficient time according to the claims to obtain a relatively equiaxed ferrite grain morphology in the steel, wherein the ratio of grain length to width is preferably 3 or less.

[0072] Therefore, steel is manufactured as hot-rolled strip instead of cold-rolled strip, yet it has all the advantages in terms of mechanical properties as the cold-rolled strip of the first embodiment.

[0073] The invention is also embodied in the method of reheating the slab at a temperature of 1200°C or higher. This achieves better homogenization of Mn in the cast steel slab, thereby reducing its segregation.

[0074] The invention is also embodied in the method of reheating the slab at a temperature of 1250°C or higher. This achieves, and even further reduces, any microsegregation of Mn present in the cast steel slab.

[0075] The invention is further embodied in a method of batch annealing of the wound steel strip at a critical temperature below 660°C. This achieves a high enrichment of Mn in the austenite of the critical region, thereby minimizing the martensite content in the final microstructure. The invention is also embodied in a method of coating the resulting strip with any metallic coating applied by hot-dip galvanizing, zinc annealing, electroplating, aluminizing, or any other method such as physical vapor deposition (PVD), chemical vapor deposition (CVD). This achieves a good aesthetic appearance and the required corrosion resistance for the steel strip in application or use.

[0076] In embodiments according to the method of claim excluding claim 2, the resulting steel strip undergoes a final cold rolling process, also known as leveling. This further improves its cold formability.

[0077] In this implementation, the finishing cold rolling is performed with a thickness reduction of 5% or less. This minimizes the yield point elongation of the steel strip during tensile testing, thereby improving the cold formability and aesthetic appearance of the cold-formed steel strip.

[0078] The invention is also embodied in steel strips that may be manufactured by a method according to a first or second embodiment of the invention, wherein the steel strip has a steel composition, said steel composition being, in weight percent:

[0079] C: 0.05-0.3;

[0080] Mn: 3.0-12.0;

[0081] Al: 0.03-3.0;

[0082] One or more additional alloying elements may be selected:

[0083] Si: less than 1.5;

[0084] Cr: less than 2.0;

[0085] V: less than 0.1;

[0086] Nb: less than 0.1;

[0087] Ti: less than 0.1;

[0088] Mo: less than 0.5;

[0089] Unavoidable impurities, such as

[0090] S: Less than 30ppm;

[0091] P: less than 0.04; and

[0092] The balance is Fe;

[0093] The steel strip has a retained austenite composition having a Mn content at least 1.4 times that of the overall Mn content of the steel composition and a C content at least 2.3 times that of the overall C content of the steel composition, in order to obtain metastable retained austenite so that the steel has a high strain hardening index of at least 0.3 as measured after yield point elongation for a strain range of 7% in a quasi-static tensile test, wherein the microstructure of the wound steel strip after final batch annealing comprises, by volume %:

[0094] Ferrite: 30-70%;

[0095] Residual austenite: 20-65%;

[0096] Martensite: <20%, including 0% by volume.

[0097] The steel strip has a level of Mn and C enrichment in the metastable austenite of the final microstructure of the steel under final service conditions, which enables it to have a high work hardening or strain hardening rate.

[0098] Furthermore, according to the invention, the cold-rolled or hot-rolled steel strip preferably has a microstructure containing 20 to 65 vol% of retained austenite, 30 to 70 vol% of ferrite, and less than 20 vol% (including 0 vol%) of martensite. The ferrite is preferably ultrafine, with grains ranging from 0.2 to 2 μm. This ultrafine ferrite achieves a more or less equiaxed shape due to sufficiently long batch annealing of the coiled steel strip, wherein the length / width ratio of the grains is ≤3. Unlike the typically short (in minutes rather than hours) discontinuous type of annealing that produces 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 in the present invention.

[0099] Other embodiments of the invention, according to claims 10 to 14, impart high mechanical and cold-formable properties to the steel strip when manufactured according to the invention. These properties are (biaxial) stretchability, bendability, hole-expanding ability, yield strength, ultimate tensile strength, total elongation, and yield point elongation.

[0100] As described above, the invention is also embodied in steel strips that have undergone smoothing and cold rolling (also known as leveling).

[0101] This invention is based on modifying the steel composition and processing it using all the mentioned steps to achieve an optimized microstructure. Due to the achieved microstructure, cold-rolled and / or hot-rolled steel strips exhibit high cold formability and high mechanical properties.

[0102] The essential elements for steel are Mn, C, and Al. Mn and C are austenite stabilizers in steel and are therefore added to steel in predetermined amounts to stabilize 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 completes during heating). Adding Al improves the robustness of steel to industrial processing because it makes the steel less sensitive to undesirable small temperature changes during critical processing. This invention does not limit the presence of other optional elements and unavoidable impurities in steel. The scope of these optional and unavoidable alloying elements is given in the relevant claims.

[0103] An increase in the amount of Mn in steel, ranging from 3 to 12 wt%, will cause a significant Mn enrichment in austenite during batch annealing of hot-rolled or cold-rolled steel strip in its coiled state. Since C is also an austenite-stabilizing element, this Mn enrichment, along with C enrichment, enhances the thermal stability of the critical region austenite by suppressing the Ms temperature of the steel (Ms = the temperature at which martensitic transformation begins during cooling). Therefore, during cooling to room temperature after batch annealing of the coiled steel strip, the critical region austenite does not transform to martensite to a large extent, allowing a significant amount of austenite (>20 vol%) to remain in the room-temperature microstructure of the steel. This retained austenite, possessing optimal mechanical stability, transforms into martensite during loading (forming or any other deformation), thereby inducing a transformation-induced plasticity (TRIP) effect. Due to the TRIP effect, which increases the rate of work or strain hardening, high strength, high elongation, and high cold formability are achieved in the steel strip of this invention. Due to extreme segregation and the change in the mechanism of plastic enhancement from TRIP to TWIP (TWIP = twin-induced plasticity), a Mn content higher than 12% by weight will make continuous casting of steel difficult, and a content lower than 3% by weight will not produce enough Mn enrichment in austenite to achieve sufficient retained austenite in the room temperature microstructure.

[0104] Similar to the effects of Mn described above, C also distributes into the critical region austenite during the final batch annealing, improving the thermal stability of austenite and causing austenite stabilization in the room temperature microstructure. However, C is more effective than Mn in smaller amounts, and therefore the C content used in this invention for modifying the chemical properties of steel ranges from 0.05 to 0.3 wt%. If the C content is less than 0.05 wt%, sufficient austenite stabilization effect is not obtained, and if the C content is greater than 0.3 wt%, it will cause difficulties in spot welding, for example, in strips manufactured after cold forming and post-processing. Welding is essential in assembling automotive parts into vehicle bodies, so this aspect is very important to consider. C is also added to the steel of this invention to improve strength.

[0105] Aluminum is not an austenite-stabilizing element in steel, but rather a ferrite-stabilizing element. However, it is added to steel at most 3% by weight to extend the critical temperature range (Ac1 to Ac3). With high levels of Mn, steel becomes sensitive to small variations in processing temperature during industrial-scale processing. The addition of Al ensures the processing robustness of steel, allowing for small variations in the batch annealing temperature of the selected steel strip to achieve the desired microstructure. When Al is intentionally not added to steel as an alloying element (i.e., when the Al content is about 0.03% by weight, as it is typically added to molten steel as a deoxidizer), a more precise furnace is necessary, but the present invention still applies. The maximum amount of Al is limited to 3% by weight to reduce oxide scale formation during hot rolling and rolling forces during both hot and cold rolling.

[0106] The combination of steel composition and method steps results in the beneficial effects of the invention. Mn, an essential alloying element for modifying the chemical properties of steel, tends to segregate after casting when its content exceeds about 2% by weight. This will affect product properties by producing inhomogeneous properties and can also lead to cracking during processing steps. Therefore, it is preferable that the cast slab be highly homogeneous. Good homogenization of the slab is achieved by using a relatively high slab reheating temperature greater than 1150°C, preferably greater than 1200°C, more preferably greater than 1250°C, and sustaining it for a sufficiently long time, preferably 60 minutes or more.

[0107] Then, due to the relatively high alloy content of the steel, the rolling force is high during hot rolling of the strip. To roll fairly wide strips on an industrial scale, typically wider than 1000 mm, hot rolling is preferably carried out at a relatively high temperature of the austenitic phase in the steel, greater than Ar3, where Ar3 is the temperature at which ferrite begins to form in the steel during cooling. This can be ensured by using a starting finishing temperature (F1) of approximately 1000 °C or higher. A lower F1 temperature would increase the hot rolling force and result in critical zone hot rolling, which would make large-scale industrial processing difficult. In addition to increasing the rolling force, critical zone hot rolling can also cause insufficient recrystallization of the hot-rolled strip.

[0108] Then, when cold rolling is applied to hot-rolled strip to reduce the dimensions of the final steel product, the material cannot be cold-rolled unless appropriate pre-processing is employed. Specifically, the steel in the coiled state after hot rolling undergoes intermediate batch annealing at low temperatures within the critical temperature range of the steel for 24 hours or longer. Because it lasts for a relatively long period, it is intermittent annealing. The intermediate batch annealing temperature should be below 650°C, because at temperatures above this, a large amount of retained austenite will form after the steel is cooled to room temperature. Additionally, if a higher batch annealing temperature is used, a large amount of martensite can appear in the microstructure. Both martensite and retained austenite make cold rolling difficult by increasing the rolling force. Although the martensitic phase is hard, the retained austenite transforms into hard martensite during the cold rolling process itself, thus increasing the rolling force. Therefore, intermediate batch annealing of the coiled material is part of the method according to this embodiment to maintain low levels of retained austenite and martensite and increase the amount of ferrite. The ferrite phase does not produce the same level of work hardening as the retained austenite during cold rolling, and therefore maintains low rolling forces, making cold rolling possible. Following this intermittent annealing of hot-rolled steel, the minimum amount of ferrite phase required to make it suitable for cold rolling is 60% by volume.

[0109] Final batch annealing of hot-rolled steel strip in the case of hot-rolled products and cold-rolled steel strip in the case of cold-rolled products are crucial for obtaining the desired microstructure in the final product for the purpose of this invention. This final batch annealing must be performed at a critical zone temperature (within the range of Ac1 to Ac3) below 700°C, preferably below 660°C. This is because thermodynamic calculations suggest that for the chemical properties of the steel of the invention, peak enrichment of C occurs in the critical zone austenite below 660°C, while Mn enrichment monotonically increases with decreasing temperature from 700°C. Therefore, using a final batch annealing temperature below 660°C ensures optimal maximum (C+Mn) enrichment in the critical zone austenite. The annealing temperature must be selected in such a way that the maximum amount of Mn and C is distributed into the austenite. During this final batch annealing at the critical zone temperature, C and Mn enrichment will be present in the critical zone austenite because C and Mn are austenite stabilizers. While carbon (C) diffuses and distributes rapidly in steel as a small interstitial element, mn (Mn), as a large substitutional element, diffuses slowly. Therefore, a batch annealing time of 5 hours or longer, more preferably 10 hours or longer, is required to achieve a high concentration of Mn in austenite. Mn enrichment in austenite should result in an Mn content at least 1.25 times, preferably at least 1.4 times, the overall Mn content of the steel. C enrichment should result in a C content at least 1.2 times, preferably at least 2.3 times, the overall C content of the steel. These levels of Mn and C enrichment in the critical region austenite are necessary to adequately stabilize the austenite to room temperature, resulting in at least 20% by volume of retained austenite in the room temperature microstructure. Furthermore, these levels of Mn and C enrichment are essential for achieving optimal mechanical stability (referred to as metastability) of the retained austenite, allowing the steel to have a strain hardening exponent of at least 0.3 during deformation. If Mn and C enrichment in austenite is below the aforementioned values, optimal stability of the retained austenite is not achieved, and therefore, the minimum strain hardening exponent of 0.3 is not achieved. Therefore, when the annealing duration is less than 5 hours, these requirements of the present invention are not met. Similar drawbacks occur if batch annealing temperatures above 700°C are used. The desired enrichment of C and Mn will not occur in the critical austenite region, and therefore the critical austenite region will not be sufficiently stable to produce a minimum of 20 vol% retained austenite in the room temperature microstructure after batch annealing, along with the metastability required by the desired high strain hardening rate. Thus, the combination of a high retained austenite fraction and its optimal mechanical stability results in the desired high strain hardening rate. Annealing temperatures above 700°C will also result in more than 20 vol% martensite, which will not produce the desired strain hardening rate. It is this high strain hardening rate that leads to high cold formability and a high combination of strength and ductility in the final product. High strain hardening rate strengthens the steel sheet while simultaneously thinning it during forming (e.g., by stretching), which results in high cold formability.

[0110] The final batch annealing is performed in a non-oxidizing and non-nitriding atmosphere to minimize any surface degradation of the steel strip caused by oxygen and nitrogen. Because the batch annealing process requires a minimum of 5 hours, the steel surface can oxidize if a non-oxidizing atmosphere is not used. Additionally, decarburization will occur, reducing the carbon content of the steel and making the invention less effective. For the same reason, nitrogen can react with Al present on the steel surface, forming its nitrides on the steel surface. All these forms of surface degradation are detrimental to the mechanical properties and formability of the steel. Preferred annealing atmospheres may be vacuum, hydrogen, or argon.

[0111] The above modifications to the steel and its processing result in a suitable microstructure in the final product, thus enabling the success of this invention. A high retained austenite fraction (>20 vol%), combined with a low martensite fraction (<20 vol%) and an optimal ferrite fraction (30-70 vol%), produces a combination of high strength, ductility, and formability due to a high strain hardening rate. A retained austenite content greater than 65 vol% is impossible to achieve within the compositional boundary of the steel and is unnecessary for achieving the minimum required strain hardening rate. Furthermore, retained austenite above 70 vol% can cause problems in spot welding, leading to severe liquid metal embrittlement and poor resistance to hydrogen embrittlement in use. Therefore, the compositional boundary of the steel of this invention is selected considering these factors. If the ferrite content is greater than 70 vol%, a high strain hardening index of 0.3 will not be achieved, primarily due to the metastable retained austenite via the TRIP effect. A ferrite fraction less than 30 vol% is not required to obtain the minimum strain hardening rate. The martensitic phase primarily contributes to strength without significantly contributing to the strain hardening rate. Furthermore, a large amount of martensite (>20 vol%) can form weak interfaces with softer phases such as ferrite and retained austenite. These interfaces are detrimental to high ductility and formability because they act as nucleation sites for damage initiation. Therefore, the martensite content should be kept below 20 vol%, or even absent altogether.

[0112] Ultrafine grain size is another microstructure requirement of this invention. Ferrite grains should be less than 2 μm, in the range of 0.2-2 μm. This ultrafine grain size imparts good ductility to the product and also contributes to the good mechanical properties of the steel strip in this invention by strengthening through grain refinement. 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 steel composition of this invention and the requirement of final batch annealing in the critical region, the phases (ferrite and austenite) at the annealing temperature are constrained and cannot grow. All these factors lead to the desired ultrafine grain size of ferrite. Ferrite grain sizes greater than 2 μm will result in lower strength and ductility. Grain size here is primarily represented by grain length. Due to sufficient recrystallization of the steel strip during final batch annealing, the width of the ferrite grains will be greater than 1 / 3 of their length. This will effectively produce a relatively equiaxed shape of the ferrite grains after final batch annealing.

[0113] The sheet formability effect of the above-mentioned ultrafine ferrite grain size can be seen in the presence of yield point elongation in the engineered stress-strain curve of the product. This can be detrimental to the cold formability of steel due to strain concentration and deterioration of the aesthetic appearance of cold-formed parts. Therefore, process variables are selected in such a way that the ferrite fraction in the final microstructure is a maximum of 70% by volume. This will ensure that the yield point elongation (if present) will be limited to a maximum of 10% of the engineered strain to obtain the best sheet formability and / or aesthetic appearance of the cold-formed parts. However, a yield point elongation greater than 10% of the engineered strain is not a limiting factor for the operation of this invention. This is because the potential negative effects of yield point elongation can be mitigated by leveling before forming and / or by using appropriate lubrication during forming.

[0114] To remove yield point elongation (if present), the steel strip in this invention is optionally subjected to a small cold rolling reduction, with a thickness reduction of up to 5%, via leveling or finishing cold rolling. This small amount of cold rolling applied in one or more passes removes the yield point elongation without altering the mechanical properties of the steel strip to any perceptible degree. However, the invention still functions without this leveling step, producing high cold formability even with a total yield point elongation of less than 10% of the engineering strain.

[0115] Optionally, hot-rolled or cold-rolled strips are coated with a metallic coating after their final batch annealing to improve their aesthetic appearance and corrosion resistance in use. Coating methods may include, but are not limited to, hot-dip galvanizing, zinc annealing, electro-galvanizing, PVD, CVD, etc. The steel strips that are ultimately batch annealed are very stable in terms of their microstructure and therefore their properties are not substantially altered by the application of a thin coating.

[0116] Cold forming of steel strips, sheets, or blanks can be performed with or without the application of suitable lubrication to reduce friction between the steel and the tool. In both cases, the present invention provides high cold formability. Non-limiting examples of lubrication systems are light oils, Klüber Press Pate, Teflon foil, or combinations thereof.

[0117] The steel used in the method according to the invention is a medium-Mn steel containing carbon, manganese, and aluminum as main components. Optional alloying elements selected from silicon, chromium, vanadium, niobium, titanium, and molybdenum may be present. Unavoidable impurities such as N, P, S, O, Cu, Ni, Sn, Sb, etc. (derived from the starting materials used to prepare the steel composition) may be present. They are not intentionally added or deliberately controlled within predetermined limits. The balance in the steel composition is iron.

[0118] Carbon is present in an amount of 0.05-0.3 wt%, for example 0.05-0.20 wt%, preferably 0.07-0.20 wt%. Although C also helps stabilize austenite, it is added primarily for strength considerations. In this composition, the austenite-stabilizing effect of manganese is much more pronounced due to its higher proportion. The preferred range of C is 0.05-0.25 wt%, and a more preferred range is 0.08-0.21 wt%. Too little C will not produce the desired strength level of 800 MPa, and if C is above 0.21, the weldability of the formed parts may deteriorate.

[0119] Manganese is present in an amount of 3.0-12.0 wt%. Manganese reduces the Ac1 and Ac3 temperatures, stabilizes austenite, improves strength and toughness, and induces the TRIP effect by stabilizing the austenite in the room-temperature microstructure. The desired effect is not achieved at levels below 3.1 wt%, while amounts above 10.5 wt% will cause casting and segregation problems. Furthermore, the deformation mechanism will change from transformation-induced plasticity (TRIP) to twinning-induced plasticity (TWIP). If the Mn content is too low, insufficient austenite will remain at room temperature, and the stability of the retained austenite will be too low, resulting in the inability to obtain the ductility and strain hardening benefits. Preferably, the Mn content is in the range of 3.5-10.0 wt%. In embodiments, the total Mn content is 5.0-9.0 wt%. In other embodiments, it is 5.5-8.5 wt%, for example, 6.0-7.5 wt%.

[0120] Adding aluminum expands the temperature range Ac1 to Ac3, thereby improving the robustness of the method in industrial applications. Al is present in amounts ranging from 0.03 to 3.0 wt%, for example 0.6-2.9 wt%, preferably 1.0-2.2 wt%.

[0121] Silicon (if present) is added in an amount less than 1.5 wt% to improve strength through solid solution strengthening. If present, the amount is typically greater than 0.01 wt% and less than 1.5 wt%. Its preferred range is 0.1–1.0 wt%.

[0122] Both Al and Si help suppress cementite precipitation to avoid deterioration of ductility. Furthermore, both Al and Si increase the peak annealing temperature to obtain the highest amount of retained austenite at room temperature after final batch annealing. Therefore, annealing in the critical zone promotes Mn diffusion to achieve efficient Mn distribution within the austenite.

[0123] Optionally, one or more additional microalloying elements, selected from V, Nb, Ti, and Mo, are present. These microalloying elements improve strength through precipitation hardening of their carbides, nitrides, or carbonitrides. Cr (another optional element of the invention) also increases the peak annealing temperature to achieve the highest amount of retained austenite at room temperature and reduces the sensitivity of retained austenite content to annealing temperature. This results in efficient Mn distribution in austenite and improved method robustness during annealing. If 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%.

[0124] There are no restrictions on the composition of the metallic coating. A Zn-based coating can be applied, for example, a zinc coating primarily containing zinc, at least 0.1 wt% Al, and optionally up to 5 wt% Al and up to 4 wt% Mg, with the remainder of the coating composition comprising all other elements individually less than 0.3 wt% and unavoidable impurities. Small amounts of other additional elements, such as those for forming zinc flowers and / or preventing slag formation, may be present, for example, in amounts less than 0.3 wt%, selected from the group consisting of: Pb, Sb, Ti, Ca, Mn, Sn, La, Ce, Cr, Ni, Zr, and Bi. Small amounts of these additional elements do not significantly alter the properties of the bath or the resulting coating for general applications. Preferably, when one or more additional elements are present in the coating, each is present in an amount of <0.02 wt%, more preferably in an amount of <0.01 wt%. Coating methods may also range from hot-dip galvanizing (GI), zinc annealing, heated coating cycles, and electroplating variations. Aluminum-based coatings, such as Al-Si-X coatings, can also be applied, where Si can vary from 0.1 to 10% by weight and X = any other coating-modifying element present in any desired amount, plus unavoidable impurities that substantially do not alter the coating properties. Coating methods such as PVD, CVD, etc., are also applicable.

[0125] The final batch annealing process is not limited to the type of furnace used or even the heating and cooling rates of the coil. It should be understood that the heating rate of the coil can vary from the surface to the center as it undergoes batch annealing. However, for the purposes of this invention, it is necessary to maintain the wound strip at the target batch annealing temperature for at least 5 hours, preferably more than 10 hours, so that each portion of the coil undergoes sufficient enrichment of C and Mn in austenite. The cooling rate after batch annealing is irrelevant to this invention, as the presence of a large amount of Mn increases the hardenability of the steel. Therefore, the coil can be cooled in a batch annealing furnace, air-cooled, forced-air-cooled, or even water-quenched.

[0126] The leveling or finishing cold rolling can be completed either on bare steel strip or on coated steel strip. This can also be done in a single pass or multiple passes.

[0127] The obtained steel strip preferably has a three-phase or two-phase microstructure, comprising (by volume%):

[0128] Ferrite: 30-70%;

[0129] Residual austenite: 20-65%;

[0130] Martensite: less than 20%, including 0%; and

[0131] Ferrite grain size: 0.2-2μm.

[0132] The obtained steel strip has the following compositional characteristics of retained austenite:

[0133] Mn: The overall Mn content of the steel is 1.25 times, preferably 1.4 times.

[0134] C: The overall C content of steel is 1.2 times, preferably 2.3 times.

[0135] Advantageously, the steel strip has the following properties:

[0136] Yield strength: 600 MPa or greater;

[0137] Ultimate tensile strength: 800 MPa or greater;

[0138] Total elongation: 20% or greater;

[0139] Strain hardening index: 0.3 or greater

[0140] Yield point elongation: preferably 10% of engineering strain or less

[0141] Minimum bending angle at 1.0mm thickness: 100° or greater;

[0142] Hole enlargement capacity: 20% or greater;

[0143] Minimum tensile strain in biaxial tension: 10% or greater.

[0144] The phase fractions mentioned above were determined using X-ray diffraction (XRD). The amount of retained austenite was determined by XRD at a position of 1 / 4 thickness of the sample. The determination was performed using a Panalytical Xpert PRO standard powder diffractometer (CoK). α XRD patterns were recorded on the surface of the ferrite (radio) in the range of 45 to 165° (2Θ). Phase-to-phase quantitative determinations were performed using the BrukerTopas software package for Rietveld refinement. Martensite content was determined by peak splitting at the ferrite diffraction positions in the diffraction patterns.

[0145] 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 by electron probe microanalysis (EPMA). The C content of the retained austenite was estimated using a well-known formula proposed by Dyson and Holmes. This formula relates the lattice parameter of austenite, which can be determined from XRD data, to its C content. This formula is available in the following article: DJ Dyson, B. Holmes, Effect of alloying additions on the lattice parameter of austenite. Journal of Iron and Steel Institute, Vol. 208, 1970, pp. 469-474.

[0146] From room temperature according to NEN 10002 standard quasi-static (strain rate 3×10) -4 s -1Tensile testing was used to determine yield strength, ultimate tensile strength, elongation at yield point, and total elongation. The geometry of the tensile specimen was an 80 mm nominal length, 30 mm width, and a nominal thickness of 1.5 mm along the rolling direction. The strain hardening rate was measured in the tensile curve after the elongation at yield point within a range of 7% strain. Bendability was determined in both longitudinal and transverse directions on a nominal 1.5 mm thick, 40 mm × 30 mm specimen by a three-point bend test according to VDA 238-100. The bending axis was along the 30 mm dimension and the bending radius was 0.4 mm. The bending angle obtained from the nominal 1.5 mm specimen was converted to the angle corresponding to a 1.0 mm thickness using the following formula: Bending angle at 1.0 mm thickness = Measured angle × Square root of actual thickness in mm. From these converted bending angles, the lowest values ​​of the longitudinal and transverse specimens were taken for a specific heat treatment condition to claim the scope of the invention. Hollow expansion capacity (HEC) was determined according to ISO / TS 16630:2003(E). Specimens measuring 90mm × 90mm × 1.5mm were cut from steel strip. A 10mm diameter hole was punched in the center of the specimen, and a hole reaming test was performed. The hole reaming capacity (HEC = (increase in initial hole diameter / initial hole diameter) × 100%) was calculated from the measured data. Biaxial tensile strain was determined by a biaxial tensile test performed on an Erichsen press using a 75mm diameter flat punch and a 79.78mm diameter die. The punch nose had a radius of 10mm, and the die was 8mm. The blank holder force was set to the maximum machine capacity (~580kN) to ensure no shrinkage occurred. The test speed was set to 20mm / min. Strain measurements were performed by applying a 10mm square grid to a sheet with fine markings.

[0147] In steel strip manufacturing, as explained above, by using a final batch annealing step at the critical zone temperature of the steel (below 700°C), Mn partitions from ferrite to austenite, thereby improving the stability of the critical zone austenite. During cooling after the final batch annealing, the critical zone austenite does not significantly transform into martensite due to its high stability caused by low Ms, resulting in a dual-phase microstructure of ferrite and retained austenite. For low Mn contents, such as less than 8% by weight, some critical zone austenite may transform into martensite, but the martensite content will be 20% by volume or less. Therefore, by increasing the Mn level and using a low batch annealing temperature (e.g., below 700°C), a large amount of retained austenite (20% by volume or more) with optimal metastability can be guaranteed. During deformation in the forming step, this large amount of retained austenite partially transforms into martensite, causing a transformation-induced plasticity (TRIP) effect, resulting in a high strain hardening exponent (= high elongation and high formability).

[0148] Due to the steel composition, the total elongation of the steel strip is preferably 20% or greater, and the strain hardening index is 0.3 or greater. A critical-division batch annealing step using a medium-Mn steel approach is preferred to obtain a mixed microstructure of ultrafine ferrite (0.5-2.0 μm) with regions of martensite and high retained austenite. This results in high ductility and a high hardening rate. These lead to high cold formability of the steel strip.

[0149] Preferred steel strips are used as materials for manufacturing automotive parts, particularly those with complex shapes where strip formability is required. Components requiring high energy absorption combined with high strength are also suitable for manufacture from steel strips. Non-limiting examples include interior automotive parts, B-pillars, and longitudinal ribs.

[0150] Reference Figure 1 The invention is illustrated by the examples described below.

[0151] Figure 1 The SEM microstructure of the steel manufactured according to the present invention, obtained by final batch annealing (steel A, 650°C / 10 hours), is shown, where F = ferrite and MA = martensite-austenite.

[0152] Three ingots of the present invention, chemical compositions A, B, and C, with dimensions of 200 mm × 100 mm × 100 mm, were cast by melting the feed in a vacuum induction furnace. The chemical compositions of these steels, along with those of two reference steels, D and E, are given in Table 1. Steel D is a twin-induced plasticity (TWIP) steel, and steel E is a DH1000 grade, both received under their final cold-rolled and annealed conditions. The thicknesses of these received steels were 1.7 mm and 1.5 mm, respectively. They were then reheated at 1250 °C for 2 hours and rough-rolled to a thickness of 30 mm. The strip was then reheated again at 1250 °C for 30 minutes, and steels A and B were hot-rolled to a thickness of 3 mm using a rolling start temperature of 1150 °C and a finishing rolling temperature (FRT) of 900 °C (which is within the austenitic phase region for all three steels) and a finishing rolling temperature (FRT) of 900 °C to a thickness of 4 mm. A high reheating temperature of 1250 °C and a long duration of 2 hours were used for suitable homogenization of Mn.

[0153] The austenite-to-ferrite transformation temperatures (Ar3) of steels A, B, and C, measured by dilatational analysis, were 785, 770, and 723 °C, respectively. The hot-rolled steels were then subjected to coil cooling simulation in a muffle furnace from 680 °C, followed by cooling to room temperature. Hot-rolled strips of A and B were then intermediately batch annealed at 600 °C for 96 hours, while strip C was intermediately batch annealed in a muffle furnace at 550 °C under an argon protective atmosphere and air-cooled to room temperature. These annealing temperatures were chosen in this manner to achieve the desired ferrite fraction to facilitate the later stages of cold rolling in this process. The phase fractions of steels A, B, and C after this intermediate batch annealing of the hot-rolled strips are given in Table 2. The phase fractions were determined by XRD measurements from the quarter-thickness position of the strip as described previously. It can be seen that the ferrite fraction in all three steels is higher than 60% by volume.

[0154] Next, the strip is pickled in HCl acid at 90°C to remove oxides, and then all the steel is cold rolled from their respective hot-rolled specifications to a final thickness of 1.5 mm.

[0155] Cold-rolled strips of steel A and B were batch annealed at 650°C for 10 hours using a muffle furnace, and cold-rolled strips of steel C were batch annealed at 640°C for 4 and 16 hours. An argon atmosphere was used for annealing, ensuring the atmosphere was free of oxygen and nitrogen to minimize oxidation of the strip and any unwanted reactions between atmospheric nitrogen and aluminum from the steel to form a nitride layer on the surface. After annealing, the samples were cooled to room temperature. For comparison, cold-rolled strips of steel A were also annealed in a similar manner at 650°C for 2 minutes, 5 minutes, and 1 hour, and steel C was annealed at 640°C for 4 hours. Some samples underwent finishing cold rolling or leveling with a reduction of up to 5% in thickness.

[0156] The material characterization and testing procedures have been described above. Recall that the microstructure of the samples was characterized using XRD and SEM. Trace analysis of the phase chemical composition was performed using EPMA and XRD analysis. Tensile properties were determined by tensile testing of specimens with an 80 mm length and 30 mm width (A80 specimen geometry). The formability of the strip was evaluated by biaxial tensile testing with bending, reaming, and appropriate lubrication. For bendability, the L and T specimens are defined as follows: L = longitudinal specimen, where the bending axis is parallel to the rolling direction; T = transverse specimen, where the bending axis is perpendicular to the rolling direction.

[0157] Figure 1The typical microstructure obtained after final batch annealing of cold-rolled strip of steel A is shown, in which ferrite and martensite-austenite regions are observed. Furthermore, ultrafine ferrite grain sizes are achieved. Microstructural characteristics of steels A, B, and C after different final annealing treatments of cold-rolled samples are provided. For all steels under all conditions, the ferrite grain size ranges from 0.5 to 1.9 μm. For steels A and C, the retained austenite content increases with increasing annealing time at their respective annealing temperatures, as more Mn is allocated to the austenite. Higher Mn content results in higher retained austenite content (steel C has higher retained austenite content than A and B), demonstrating the effect of Mn on austenite stability. Under all conditions, except for steel A annealed at 650 °C for 2 minutes, a high fraction of retained austenite (greater than 33% by volume) was obtained. The Mn and C contents of the steels under the annealing conditions given in Table 4 show that the Mn enrichment in the retained austenite of all these different conditions ranges from 1.286 to 2.139 times the overall Mn content of the steel, except for steel A at -650°C / 2 min, where the Mn content is only 1.09 times the overall Mn content. For C enrichment in the retained austenite, the C content ranges from 1.17 to 3.085 times the overall C content of the steel, except for steel A 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 steel A at -650°C / 2 min is also less than 20% by volume, and therefore the martensite content is greater than 20% by volume (39.8% by volume). In all other steels and conditions of the present invention, the martensite content is 16.7% by volume or less, including 0% by volume (steel C at -640°C / 960 min).

[0158] The lower residual austenite fraction of steel A at 650℃ / 2min is clearly due to the fact that the annealing time of 2min is too short for sufficient Mn to diffuse into the austenite, even though the annealing temperature is in the critical temperature range of steel A and is less than 700℃.

[0159] The results of the above microstructure characteristics can be observed in the tensile properties of the steels given in Table 5. Steel A, with its bulk Mn and C contents of less than 1.25 times the amount of retained austenite and less than 2 times the amount of C (at 650℃ / 2min), exhibits 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 into martensite due to its low stability as expected from the enrichment of low Mn and C. The small amount of retained austenite is consumed very early during deformation without even showing any yield point elongation. Therefore, the poor tensile properties of this steel condition make it unsuitable for cold forming. On the other hand, steel A under other annealing conditions and steels B and C under all conditions show yield strengths above 693 MPa, ultimate tensile strengths above 860 MPa, and total elongation above 23.4%. These steels also exhibit high energy absorption capacity (determined by the product of ultimate tensile strength and total elongation) and varying amounts of yield point elongation. As can be seen from Table 3, the yield elongation decreases with annealing time for steels A and C due to the increased ferrite grain size. The tensile properties of these steel compositions of the present invention are comparable to those of the reference steels listed in Table 6. Due to the combination of steel chemical composition, processing, and microstructure, the steel compositions of the present invention exhibit significantly higher total elongation and energy absorption capacity under long final batch annealing conditions than the conventional DH1000 grade (reference steel E). Steel E has a very small 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 present invention, some of the steels of the present invention have energy absorption capacities within the range of TWIP steels with a fully austenitic microstructure.

[0160] Table 7 shows the formability parameters of the steels of the present invention compared to reference steels. The formability parameters compared are biaxial stretchability in terms of strain in the rolling and transverse directions, bendability in the longitudinal and perpendicular directions of the sheet, and stretchable flanging as expressed by HEC values. Steel A shows a biaxial tensile strain of 0 when subjected to final batch annealing at 650°C for less than 10 hours, although other parameters are non-zero. Steel A annealed at 650°C for 2 minutes shows very poor bendability and stretchable flanging. While bendability and stretchable flanging improve with increasing annealing time, the material remains unstretchable until the final batch annealing of 10 hours. Steel B annealed at 650°C for 10 hours also shows similar formability parameters to Steel A under the same annealing conditions. Steel C annealed at 640°C for 4 hours shows high bendability and stretchable flanging but low stretchability. Stretchability also improves when steel C is annealed for 16 hours.

[0161] The cold formability of steel sheets is a combination of various parameters such as stretchability, bendability, and stretchable flanging. When the steels of the present invention are ultimately batch annealed at less than 700°C in their critical temperature range, the annealing time, as previously observed, is important for producing the required amounts of Mn and C enrichment in the retained austenite, since Mn is a slow-diffusing element in steel. High Mn and C enrichment is necessary to achieve a high strain hardening rate. Therefore, steels exhibiting a lower strain hardening index and annealing time of less than 10 hours also exhibit low stretchability, but other formability parameters are good. A high strain hardening index greater than 0.3 is necessary to achieve good stretchability in the steels of the present invention; otherwise, premature local cracking may occur. Therefore, from the results, a minimum of 10 hours of final batch annealing is necessary for the steels of the present invention to achieve a minimum amount of Mn and C enrichment in the retained austenite, as the good cold formability (combination of stretchability, bendability, and stretchable flanging) mentioned in the claims appears.

[0162] When the formability of samples annealed in batches over 10 hours was compared with that of the steel of the present invention, the formability parameters of the steel of the present invention were observed to be in the range of highly formable TWIP steel (reference steel E) and much higher than those of conventional DH1000 (reference steel D). Even when annealed for only 4 hours, the biaxial tensile strain of steel C was higher than that of conventional DH1000. This high cold formability of the steel of the present invention is due to the high fraction of metastable retained austenite and high enrichment of Mn and C, which is achieved in the steel of the present invention through the processing steps of the present invention.

[0163] Table 8 shows the effect of leveling on the mechanical properties of steel B annealed at 650°C for 10 hours. It is evident that the elongation at yield point decreases with increasing leveling reduction. With a 2% reduction, the elongation at yield point disappears. The tensile properties are not significantly altered and remain within the scope of the claims of this invention. Most importantly, even with a reduction of up to 5% thickness, the strain hardening exponent remains high. Therefore, this elimination of elongation at yield point without significant alteration of mechanical properties at a maximum leveling reduction of 5% will make the steel strip of this invention even more cold-formable, as this will reduce the risk of strain concentration during stretching and the presence of stretcher marks on the surface of the formed article.

[0164] Table 1: Composition of steel (by weight %)

[0165] steel C Mn Si Al P S B 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

[0166]

[0167]

[0168] Table 2: Phase fraction after intermediate batch annealing of hot-rolled strip

[0169] steel Ferrite (volume %) Residual austenite (volume %) Martensite (volume %) A 83.3 14.4 2.3 B 79.5 18.4 2.1 C 71.6 28.4 0

[0170] Table 3: Phase fraction and average austenite and ferrite grain size of cold-rolled strip after final batch annealing

[0171]

[0172] Table 4: Enrichment of manganese and carbon in retained austenite after final annealing of cold-rolled materials

[0173]

[0174] Table 5: Tensile properties of cold-rolled steel after annealing

[0175]

[0176] Table 6: Tensile properties and microstructure of reference steels

[0177]

[0178] Table 7: Formability parameters of steel after final annealing

[0179]

[0180] Table 8: Tensile properties of steel B under flattening conditions after heat treatment to optimal formability.

[0181]

Claims

1. A method for manufacturing cold-rolled and annealed steel strip, wherein the steel composition, in weight percent, is: C：0.05-0.3； Mn: 3.0-12.0; Al：0.03-3.0； 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; and The balance is Fe; The method includes the following steps: - Cast molten steel into slabs; - Reheat the slab and hold it at 1150°C or higher for 1 hour or more; - Use an average F1 slab entry temperature greater than 1000℃ to hot roll the steel into strip; - Winding hot-rolled steel strip; - Pickled steel strip; - The steel strip is annealed in intermediate batches 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 rolling steel into cold-rolled steel strip and then winding it; - Anneal the wound steel strip in batches: -Critical temperature range of less than 700°C between Ac1 and Ac3; -In a non-oxidizing and non-nitriding atmosphere; - The total annealing time for maintaining the strip at the critical zone temperature is at least 10 hours to obtain: Mn enrichment in austenite such that the Mn content is at least 1.25 times the overall Mn content of the steel; and C enrichment such that the C content is at least 1.2 times the overall C content of the steel; - Cooling the steel after batch annealing in air, in forced air, or by water quenching.

2. A method for manufacturing hot-rolled and annealed steel strip, wherein the steel composition, in weight percent, is: C：0.05-0.3； Mn: 3.0-12.0; Al：0.03-3.0； 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; and The balance is Fe; The method includes the following steps: - Cast molten steel into slabs; - Reheat the slab at 1150°C or higher for 1 hour or more; - Use an average F1 slab entry temperature greater than 1000℃ to hot roll the steel into strip; - Winding hot-rolled steel strip; - Pickled steel strip; - Anneal the wound steel strip in batches: -Critical temperature range of less than 700°C between Ac1 and Ac3; -In a non-oxidizing and non-nitriding atmosphere; - The total annealing time for maintaining the strip at the critical zone temperature is at least 10 hours to obtain: Mn enrichment in austenite such that the Mn content is at least 1.25 times the overall Mn content of the steel; and C enrichment such that the C content is at least 1.2 times the overall C content of the steel; - Cooling the steel after batch annealing in air, in forced air, or by water quenching.

3. The method according to claim 1 or claim 2, wherein the unavoidable impurities comprise: S: less than 30 ppm; P: less than 0.

04.

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

5. The method according to claim 1 or claim 2, wherein the slab is reheated at a temperature of 1250°C or higher.

6. The method according to claim 1 or claim 2, wherein the batch annealing of the wound steel strip occurs at a critical temperature of less than 660°C.

7. The method according to claim 1 or claim 2, wherein the resulting strip is coated with any metallic coating applied by hot-dip galvanizing, zinc annealing, electroplating, aluminizing, PVD or CVD.

8. The method according to claim 1, wherein the resulting steel strip undergoes a light-finish cold rolling process.

9. The method of claim 8, wherein the finishing cold rolling is performed with a thickness reduction of 5% or less.

10. A steel strip obtainable by the method of any one of claims 1 to 9, wherein the steel strip has a steel composition, said steel composition being, in weight percent: C：0.05-0.3； Mn: 3.0-12.0; Al：0.03-3.0； 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; and The balance is Fe; The steel strip has a retained austenite composition, with a Mn content at least 1.25 times the overall Mn content of the steel composition and a C content at least 1.2 times the overall C content of the steel composition, in order to obtain metastable retained austenite, thereby giving the steel a high strain hardening index of at least 0.3 as measured after yield point elongation at 7% strain in a quasi-static tensile test, wherein the microstructure after final batch annealing comprises, by volume %,: Ferrite: 30-70%; Residual austenite: 20-65%; The balance is martensite: < 20%, including 0% by volume; wherein the length / width ratio of the ferrite grains is 3 or less.

11. The steel strip according to claim 10, wherein the size of the ferrite grains is 0.2-2 μm.

12. The steel strip according to claim 10 or claim 11, having a yield point elongation of 10% of the maximum engineering strain as measured by its engineering stress-strain curve.

13. The steel strip according to claim 10 or claim 11, having a yield strength of 600 MPa or greater, an ultimate tensile strength of 800 MPa or greater, and a total elongation of 20% or greater.

14. The steel strip according to claim 10 or claim 11 has very high formability, characterized in that the tensile strain in a single direction is 10% or greater under biaxial tension, the VDA bending angle is 100° or greater at a thickness of 1.0 mm, and the hole expansion capacity is 20% or greater.

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