Patent Application
20250188563
The present disclosure relates to a steel sheet and a method of manufacturing the same, and more particularly, to an ultra-high-strength steel sheet with excellent elongation and a method of manufacturing the same.
Automotive steel sheets have been developed with a focus on increasing strength for user safety and vehicle weight reduction, as well as ensuring elongation for workability. Currently, the commonly used ultra-high-strength steels include dual-phase steels which consist of ferrite and martensite phases to secure elongation, and transformation-induced plasticity (TRIP) steels which achieve strength and elongation through the phase transformation of retained austenite in the final structure during plastic deformation. However, the development based on dual-phase steels which may not overcome the limitations of the rule of mixtures and TRIP steels which exhibit relatively low strength due to the main matrix of bainite has reached its limits. Therefore, the development of next-generation ultra-high-strength automotive steel sheets capable of ensuring ultra-high strength and high formability by improving the microstructure of TRIP steels is attracting the attention of steelmakers.
The present disclosure provides an ultra-high-strength steel sheet with excellent elongation and a method of manufacturing the same.
According to an aspect of the present disclosure, there is provided an ultra-high-strength steel sheet comprising, consisting essentially of or consisting of: carbon (C): 0.1 wt % to 0.30 wt %, silicon (Si): 1.0 wt % to 2.0 wt %, manganese (Mn): 1.5 wt % to 3.0 wt %, aluminum (Al): more than 0 wt % and not more than 0.05 wt %, a sum of one or more elements selected from niobium (Nb), titanium (Ti), and vanadium (V): more than 0 wt % and not more than 0.05 wt %, phosphorus (P): more than 0 wt % and not more than 0.02 wt %, sulfur (S): more than 0 wt % and not more than 0.005 wt %, nitrogen (N): more than 0 wt % and not more than 0.006 wt %, and a balance of iron (Fe) and other unavoidable impurities, wherein a final microstructure of the ultra-high-strength steel sheet consists of tempered martensite, martensite, ferrite, and retained austenite, and wherein the ultra-high-strength steel sheet has a yield strength (YP) of 850 MPa or more, a tensile strength (TS) of 1180 MPa or more, and an elongation (El) of 14% or more.
In certain aspects, the final microstructure may have a martensite matrix rather than bainite.
In certain aspects, the ferrite suitably may occupy an area fraction of 10% to 20%, the retained austenite suitably may occupy an area fraction of 10% to 20%, and the tempered martensite and martensite may occupy a remaining area fraction.
In aspects, the ferrite may include carbides, and a density of the carbides with an average size of 50 nm or more is 20 particles/μm2 or less.
In aspects, a product of TS and El suitably may be no less than 15,000 MPa·% and in aspect suitably may be no less than 15,500 MPa·% or no less than 16,000 MPa %.
According to another aspect of the present disclosure, there is provided a method of manufacturing an ultra-high-strength steel sheet, the method comprising, consisting essentially of or consisting of: producing a hot-rolled steel sheet including carbon (C): 0.1 wt % to 0.30 wt %, silicon (Si): 1.0 wt % to 2.0 wt %, manganese (Mn): 1.5 wt % to 3.0 wt %, aluminum (Al): more than 0 wt % and not more than 0.05 wt %, a sum of one or more elements selected from niobium (Nb), titanium (Ti), and vanadium (V): more than 0 wt % and not more than 0.05 wt %, phosphorus (P): more than 0 wt % and not more than 0.02 wt %, sulfur (S): more than 0 wt % and not more than 0.005 wt %, nitrogen (N): more than 0 wt % and not more than 0.006 wt %, and a balance of iron (Fe) and other unavoidable impurities; softening the hot-rolled steel sheet by maintaining a temperature of 500° C. to 650° C.; producing a cold-rolled steel sheet by cold rolling the softened steel sheet; annealing the cold-rolled steel sheet at a temperature of 820° C. to 860° C.; multi-stage cooling the cold-rolled steel sheet to a temperature of 200° C. to 260° C.; and partitioning the cooled cold-rolled steel sheet at a temperature of 370° C. to 460° C.
In aspects, preferably the hot-rolled steel sheet is softened by constantly maintaining a temperature 500° C. to 650° C. during the softening step.
The producing of the hot-rolled steel sheet may include preparing a steel slab having the above alloy composition; reheating the steel slab at 1150° C. to 1250° C.; producing a hot-rolled steel sheet by hot rolling the reheated steel slab at a finishing delivery temperature (FDT) of 850° C. to 1000° C.; and coiling the hot-rolled steel sheet at 500° C. to 700° C.
In aspects, the annealing of the cold-rolled steel sheet may be performed by heating the cold-rolled steel sheet at a heating rate of 1° C./s to 10° C./s and maintaining a temperature of 820° C. to 860° C. for 40 sec. to 120 sec.
In aspects, the multi-stage cooling of the cold-rolled steel sheet may include a primary cooling step for slowly cooling the cold-rolled steel sheet at a cooling rate of 1° C./s to 10° C./s to an end temperature of 550° C. to 750° C.; and a secondary cooling step for rapidly cooling the cold-rolled steel sheet at a cooling rate of 50° C./s or more to an end temperature of 200° C. to 260° C.
In aspects, the partitioning may be performed by heating the cooled cold-rolled steel sheet at a heating rate of 20° C./s or more and maintaining a temperature of 370° C. to 460° C. for 10 sec. to 240 sec.
As referred to herein, the term area fraction refers to percentage of a phase or composition in a two-dimensional plane of steel based on the total amount (area or volume) of the steel composition. Area fraction of a particular phase or composition may be determined by known methods including image analysis, point counting or aerial analysis.
As referred to herein, yield point (YP) and tensile stress (TS) and elongation (EL) can be measured using a commercially available tensile tester and according to the ISO standard ISO 6892-1, published in October 2009. As referred to herein, yield point refers to a lower yield point (MPa).
According to an embodiment of the present disclosure, an ultra-high-strength steel sheet with excellent elongation and a method of manufacturing the same may be implemented.
However, the scope of the present disclosure is not limited to the above effect.
An ultra-high-strength steel sheet with excellent elongation and a method of manufacturing the same, according to an embodiment of the present disclosure, will now be described in detail. The terms used herein are selected based on their functions in the present disclosure, and their definitions should be made in the context of the entire specification.
Automotive steel sheets have been developed with a focus on increasing strength for user safety and vehicle weight reduction, as well as ensuring elongation for workability. Currently, the commonly used ultra-high-strength steels include dual-phase steels which consist of ferrite and martensite phases to secure elongation, and transformation-induced plasticity (TRIP) steels which achieve strength and elongation through the phase transformation of retained austenite in the final structure during plastic deformation. However, the development based on dual-phase steels which may not overcome the limitations of the rule of mixtures and TRIP steels which exhibit relatively low strength due to the main matrix of bainite has reached its limits. Therefore, the development of next-generation ultra-high-strength automotive steel sheets capable of ensuring ultra-high strength and high formability by improving the microstructure of TRIP steels is attracting the attention of steelmakers.
The present disclosure provides a next-generation ultra-high-strength automotive coated steel sheet capable of ensuring high strength and appropriate elongation by substituting the bainite matrix of TRIP steels with martensite to overcome the mechanical property limitations of existing TRIP steels.
An ultra-high-strength steel sheet with excellent elongation, according to an embodiment of the present disclosure, will now be described in detail.
An ultra-high-strength steel sheet with excellent elongation, according to an embodiment of the present disclosure, includes a steel sheet consisting of carbon (C): 0.1 wt % to 0.30 wt %, silicon (Si): 1.0 wt % to 2.0 wt %, manganese (Mn): 1.5 wt % to 3.0 wt %, aluminum (Al): more than 0 wt % and not more than 0.05 wt %, a sum of one or more elements selected from niobium (Nb), titanium (Ti), and vanadium (V): more than 0 wt % and not more than 0.05 wt %, phosphorus (P): more than 0 wt % and not more than 0.02 wt %, sulfur (S): more than 0 wt % and not more than 0.005 wt %, nitrogen (N): more than 0 wt % and not more than 0.006 wt %, and a balance of iron (Fe) and other unavoidable impurities. The functions and contents of the components included in the steel sheet will now be described.
C is the most important alloying element in steelmaking, and is primarily intended for basic strengthening and austenite stabilization in the present disclosure. A high concentration of C in austenite may increase the stability of austenite and thus appropriate austenite for enhancing material properties may be easily ensured. When the content of C is less than 0.1 wt % of the total weight, the above-described effects may not be achieved and a sufficient strength may not be ensured. On the other hand, when the content of C is greater than 0.3 wt % of the total weight, the carbon equivalent may be increased to cause reduced weldability and decreased workability.
Si is an element for suppressing the formation of carbides and, in particular, prevents degradation in material properties caused by the formation of Fe3C. Si is well known as a ferrite stabilizing element and thus may improve ductility by increasing the fraction of ferrite during cooling. Si is also known as an element capable of ensuring strength by promoting the formation of martensite through carbon enrichment in austenite. Meanwhile, Si may be added together with Al as a deoxidizer for removing oxygen from steel during a steelmaking process, and have a solid solution strengthening effect.
Si may be added at a content ratio of 1.0 wt % to 2.0 wt % of the total weight in the steel sheet for forming the ultra-high-strength steel sheet with excellent elongation, according to an embodiment of the present disclosure. When the content of Si is less than 1.0 wt % of the total weight, ductility may not be ensured and the above-described effects of Si addition may not be properly realized. On the other hand, when the content of Si is greater than 2.0 wt % of the total weight due to excessive addition, oxide (SiO2) may be formed on the surface of the steel sheet to deteriorate wettability and reduce coatability, red scale may be formed during reheating and hot rolling to degrade the surface quality, toughness and plasticity may be reduced, and the weldability of steel may also be reduced.
Mn is a major element for stabilizing austenite. When Mn is added, the martensite start temperature, Ms, may be gradually lowered and thus the fraction of retained austenite may be increased during a continuous annealing process. Mn is also an element for facilitating low-temperature transformation and increasing strength through solid solution strengthening. Mn reduces the acid resistance and oxidation resistance of steel, but increases yield strength by refining pearlite and solid-solution-strengthening ferrite.
Mn may be added at a content ratio of 1.5 wt % to 3.0 wt % of the total weight in the steel sheet for forming the ultra-high-strength steel sheet with excellent elongation, according to an embodiment of the present disclosure. When the content of Mn is less than 1.5 wt %, the above-described effect of strength enhancement may not be sufficiently realized. When the content of Mn is greater than 3.0 wt %, the carbon equivalent may be increased to reduce weldability, oxide (MnO) may be formed on the surface of the steel sheet during processing to deteriorate wettability and reduce coatability, and segregation zones may be formed inside and outside the continuously casted slab and the steel sheet and the formation and propagation of cracks may be caused to reduce bendability.
Al functions similarly to Si and mainly serves to strengthen a solid solution and suppress the formation of carbides. Al is an element commonly used as a deoxidizer, and promotes the formation of ferrite, increases elongation, and stabilizes austenite by increasing the concentration of C in austenite. Al serves as a layer between Fe and zinc (Zn) coating to enhance coatability, and effectively suppresses the formation of Mn bands in a hot-rolled coil.
Al may be added at a content ratio of more than 0 wt % and not more than 0.05 wt % of the total weight in the steel sheet for forming the ultra-high-strength steel sheet with excellent elongation, according to an embodiment of the present disclosure. When the content of Al is greater than 0.05 wt % due to excessive addition, Al inclusions may be increased to reduce continuous castability, the enrichment of Al may occur on the surface of the steel sheet to deteriorate coatability, and AlN may be formed in the slab to cause hot rolling cracks.
One or More Elements Selected from Niobium (Nb), Vanadium (V), and Titanium (Ti)
Nb, V, and/or Ti are major elements precipitated in the form of carbides inside steel, and are intended in the present disclosure to ensure the stability of retained austenite and enhance strength by refining the initial austenite grains through the formation of precipitates, and to enable precipitation hardening through ferrite grain refinement and the presence of precipitates in ferrite.
The sum of one or more elements selected from Nb, V, and Ti may be added at a content ratio of more than 0 wt % and not more than 0.05 wt % of the total weight in the steel sheet for forming the ultra-high-strength steel sheet with excellent elongation, according to an embodiment of the present disclosure. When the content of the sum of one or more elements selected from Nb, V, and Ti is greater than 0.05 wt %, degradation in material properties and increase in production costs may occur, grain coarsening may result from the formation of coarse precipitates, and the recrystallization temperature may be excessively increased to cause a non-uniform structure.
P functions similarly to Si and may serve to increase strength through solid solution strengthening and suppress the formation of carbides. P may be added at a content ratio of more than 0 wt % and not more than 0.02 wt % of the total weight in the steel sheet for forming the ultra-high-strength steel sheet with excellent elongation, according to an embodiment of the present disclosure. When the content of P is greater than 0.02 wt %, problems such as embrittlement of welded joints, brittleness, reduced press formability, and decreased impact resistance may occur. Thus, the content of P needs to be controlled as low as possible in the present disclosure.
S is an element that combines with Mn or Ti to increase the machinability of steel and forms fine MnS precipitates to increase workability, but generally hinders ductility and weldability. S may be added at a content ratio of more than 0 wt % and 0.005 wt % of the total weight in the steel sheet for forming the ultra-high-strength steel sheet with excellent elongation, according to an embodiment of the present disclosure. When the content of S is greater than 0.005 wt %, the number of MnS inclusions may be increased to reduce workability and segregation may occur during continuous casting solidification to cause high-temperature cracks. Thus, the content of S needs to be controlled as low as possible in the present disclosure.
N is a solid solution strengthening element capable of increasing the strength of the steel sheet, and is generally introduced from the atmosphere. The content of N needs to be controlled through a degassing process during steelmaking. When the content of N is greater than 0.006 wt %, problems such as embrittlement of welded joints, low-temperature brittleness, reduced press formability, and decreased impact resistance may occur. Thus, the content of N needs to be controlled as low as possible in the present disclosure.
In the ultra-high-strength steel sheet with excellent elongation having the above-described composition of alloying elements according to an embodiment of the present disclosure, the steel sheet may have a yield strength (YP) of 850 MPa or more, a tensile strength (TS) of 1180 MPa or more, and an elongation (El) of 14% or more. For example, the steel sheet may have a YP of 850 MPa to 1070 MPa, a TS of 1180 MPa to 1250 MPa, and an El of 14% to 20%. Furthermore, the product of TS and El may be no less than 15,000 MPa·%. In the present disclosure, the product of TS and El is 15,000 MPa·% or more, which is more than twice the value of 7,000 MPa·%, typically representing the product of TS and total El at a similar strength level, and thus it may be understood that superior formability may be achieved compared to existing ultra-high-strength steels at the same strength.
Referring to
A method of manufacturing an ultra-high-strength steel sheet with excellent elongation having the above-described composition, properties, and microstructure, according to an embodiment of the present disclosure, will now be described.
Referring to
The step of producing the hot-rolled steel sheet (S100) may include preparing a steel slab having the above alloy composition; reheating the steel slab at 1150° C. to 1250° C.; producing a hot-rolled steel sheet by hot rolling the reheated steel slab at a finishing delivery temperature (FDT) of 850° C. to 1000° C.; and coiling the hot-rolled steel sheet at 500° C. to 700° C.
The step of producing the hot-rolled steel sheet (S100) may include reheating the steel material at 1150° C. to 1250° C. The steel slab having the composition of the present disclosure is reheated to a temperature of Ac3 or more to redissolve the as-cast components. When the reheating temperature is low, the hot rolling load may increase, and when the reheating temperature is high, the slab may warp so as not to be easily inserted into or discharged from the furnace. When the steel material is reheated at the above-mentioned temperature, the components segregated during continuous casting may be redissolved. To enhance strength through precipitation and solid solution strengthening, strengthening elements need to be sufficiently dissolved in austenite before hot rolling and thus the steel material needs to be heated to 1150° C. or above. When the reheating temperature is lower than 1150° C., the hot rolling load may increase, various carbides may not be dissolved sufficiently, and the components segregated during continuous casting may not be uniformly distributed.
However, a reheating temperature higher than 1250° C. may cause adverse effects such as austenite coarsening and decarburization and prevent the desired strength from being achieved. That is, when the reheating temperature is higher than 1250° C., very coarse austenite grains may be formed and thus strength may not be easily ensured. In addition, when the reheating temperature is higher than 1250° C., the slab may warp so as not to be easily inserted into or discharged from the furnace, and heating costs and process time may be increased to increase production costs and reduce productivity.
The step of producing the hot-rolled steel sheet (S100) may include performing hot rolling under conditions of a FDT of 850° C. to 1000° C., a cooling rate of 10° C./s to 30° C./s, and a coiling temperature (CT) of 500° C. to 700° C. Because high-alloy steels require minimizing edge cracking and rolling load to ensure producibility, the FDT and CT may be set in high-temperature ranges. The FDT is an important factor affecting the final material properties, and rolling at 850° C. to 1000° C. may refine austenite.
However, when the hot rolling temperature is lower than 850° C., the rolling load may increase and a mixed grain structure may occur at the edge. Rolling at a temperature higher than 1000° C. may cause grain coarsening and prevent the desired mechanical properties from being achieved. After hot rolling, cooling is performed at a cooling rate of 10° C./s to 30° C./s. The higher the cooling rate, the smaller the average grain size.
Meanwhile, when the CT is lower than 500° C., the hot-rolled coil may have a non-uniform shape and the cold rolling load may be increased. When the CT is higher than 700° C., a non-uniform microstructure may occur due to the difference in cooling rate between the center and edge of the steel sheet, and the inside of the grain boundaries may be oxidized.
In the ultra-high-strength steel sheet manufacturing method, the step of annealing the cold-rolled steel sheet may be performed by heating the cold-rolled steel sheet at a heating rate of 1° C./s to 10° C./s and maintaining a temperature of 820° C. to 860° C. for 40 sec. to 120 sec.
In the ultra-high-strength steel sheet manufacturing method, the step of multi-stage cooling the cold-rolled steel sheet may include a primary cooling step for slowly cooling the cold-rolled steel sheet at a cooling rate of 1° C./s to 10° C./s to an end temperature of 550° C. to 750° C.; and a secondary cooling step for rapidly cooling the cold-rolled steel sheet at a cooling rate of 50° C./s or more to an end temperature of 200° C. to 260° C.
In the ultra-high-strength steel sheet manufacturing method, the partitioning may be performed by heating the cooled cold-rolled steel sheet at a heating rate of 20° C./s or more and maintaining a temperature of 370° C. to 460° C. for 10 sec. to 240 sec.
After the hot-rolled steel sheet is produced (S100) and before the cold-rolled steel sheet is produced by cold rolling the softened steel sheet (S300), the hot-rolled steel sheet may be softened by constantly maintaining a temperature of 500° C. to 650° C. (S200).
The component systems utilized in the present disclosure correspond to high-alloy steels with a high content of an alloying element such as Mn, and thus a delay in hardening of the material occurs. While cooling to room temperature after hot rolling, a hard phase of martensite is formed instead of soft phases such as ferrite and pearlite and thus the strength of the hot-rolled material is very high. As such, softening is performed after hot rolling. Because problems such as thickness hunting and shape defects may occur during cold rolling when the strength of the hot-rolled plate is high, for efficient cold rolling, cold rollability may be ensured by reducing the hardness of the hot-rolled material through softening.
Table 1 shows specific process conditions for the softening step (S200). Referring to Table 1, the softening step (S200) may include, for example, heating to 500° C. to 650° C. at a heating rate of 30° C./hr to 80° C./hr, holding for 2 hours to 16 hours, and then cooling to room temperature at a cooling rate of 30° C./hr to 60° C./hr.
For softening, heat treatment is performed in a temperature range of 500° C. to 650° C. for a sufficient time. When the softening temperature is lower than 500° C., martensite formed after hot rolling does not undergo recrystallization but only tempering such that C oversaturated in the microstructure forms and spheroidizes cementite (0) in the microstructure. In this case, the brittleness of martensite may occur to cause safety accidents such as sheet fracture during cold rolling. When the softening temperature is higher than 650° C., excessive formation of austenite may occur during softening, and martensite may be formed during cooling, preventing an effective reduction in strength. Thus, the strength of the hot-rolled material may be reduced by controlling the temperature range to 500° C. to 650° C. The step of producing the cold-rolled steel sheet by cold rolling the softened steel sheet (S300) is performed to reduce the hot-rolled material to the final thickness of the steel sheet. Pickling may be performed before cold rolling. The reduction ratio in the cold rolling process may be controlled according to the final product specifications, for example, to 40% to 60%. A higher reduction ratio in cold rolling enhances formability due to the refining effect on the microstructure. When the reduction in cold rolling is less than 40%, a uniform microstructure is not easily achieved, and when the reduction in cold rolling is greater than 60%, the roll force is increased to raise the process load.
The step of annealing the cold-rolled steel sheet (S400) may include heating the cold-rolled steel sheet at a heating rate of 1° C./s to 10° C./s and maintaining an annealing temperature of 820° C. to 860° C. for 40 sec. to 120 sec. The annealing process is performed under conditions of the annealing start temperature and the annealing holding time within a dual-phase region of austenite and ferrite. In the current embodiment, heat treatment is performed within the dual-phase region to achieve the desired final material properties of the steel sheet by ensuring an appropriate fraction of ferrite for an ideal composition of ferrite, tempered martensite, and retained austenite in the final microstructure.
The annealed cold-rolled steel sheet is multi-stage-cooled to a temperature of 200° C. to 260° C. (S500). The step of multi-stage cooling the cold-rolled steel sheet (S500) may include a primary cooling step for slowly cooling the cold-rolled steel sheet at a cooling rate of 1° C./s to 10° C./s to an end temperature of 550° C. to 750° C.; and a secondary cooling step for rapidly cooling the cold-rolled steel sheet at a cooling rate of 50° C./s or more to an end temperature of 200° C. to 260° C.
The primary cooling step is a step of performing slow cooling after annealing to achieve the plasticity of the final microstructure by attempting to ensure a certain amount of ferrite in the final microstructure during heat treatment. The secondary cooling step may include a step of performing rapid cooling, for example, to a rapid cooling end temperature of 200° C. to 260° C. at a cooling rate of 70° C./s or more after the primary cooling step. The final material properties may be easily ensured by transforming austenite into martensite in the microstructure after slow cooling through the control of the rapid cooling end temperature, and a cooling rate of 70° C./s or more is required to suppress phase transformation during the rapid cooling process.
The multi-stage-cooled cold-rolled steel sheet is partitioned at a temperature of 370° C. to 460° C. (S600). The partitioning step (S600) may be performed by heating the cooled cold-rolled steel sheet at a heating rate of 20° C./s or more and holding a temperature of 370° C. to 460° C. for 10 sec. to 240 sec. The partitioning step S600 aims to ensure strength and elongation through carbon enrichment in retained austenite and martensite tempering, and lastly maintain the final microstructure configuration. The partitioning step (S600) may include, for example, a process of maintaining a reheating temperature of 400° C. to 460° C. within 60 sec.
After the partitioning step (S600) is performed, the steel sheet may be immersed in a Zn—Mg—Al coating bath and then cooled to room temperature to produce a final steel sheet.
The ultra-high-strength steel sheet of the present disclosure manufactured by performing the above-described steps may have a YP of 850 MPa or more, a TS of 1180 MPa or more, and an El of 14% or more. For example, the steel sheet may have a YP of 850 MPa to 1070 MPa, a TS of 1180 MPa to 1250 MPa, and an El of 14% to 20%. Furthermore, the product of TS and El may be no less than 15,000 MPa·%. In the present disclosure, the product of TS and El is 15,000 MPa·% or more, which is more than twice the value of 7,000 MPa·%, typically representing the product of TS and total El at a similar strength level, and thus it may be understood that superior formability may be achieved compared to existing ultra-high-strength steels at the same strength.
In the existing process of forming vehicle body parts, the fracture of components made from ultra-high-strength materials during forming may be explained by evaluation criteria such as drawability and bi-axial stretchability, which are represented in the general forming limit diagram, as well as hole expansion ratio, which is not shown in the forming limit diagram. Generally, sheet materials with better formability evaluation results may be used for more complex forming structures in the processing of vehicle body parts. These formability indices are important factors in vehicle part forming processes carried out through pressing operations. Typically, ultra-high-strength materials show a tendency for reduced elongation as their strength increases. To improve the formability of these ultra-high-strength materials, materials are being developed by applying specialized forming processes or deformation mechanisms designed to enhance formability.
The present disclosure aims to improve formability compared to existing ultra-high-strength steels by utilizing the microstructure of ferrite, retained austenite, and martensite/tempered martensite, and seeks to develop cold-rolled steel sheets that maintain the microstructure of ferrite, retained austenite, and martensite/tempered martensite while satisfying the material properties through additional component system and heat treatment control. Based on the results of simulation testing, the present disclosure proposes component system and heat treatment ranges.
In the case of existing ultra-high-strength steels with a dual-phase structure of ferrite and martensite, when the steel material undergoes plastic deformation, dislocations are formed and move in the microstructure. This movement of dislocations leads to plastic deformation through the fundamental deformation mechanism in which fracture occurs as defects form and grow. To ensure strength under this deformation mechanism, hard phases such as martensite and bainite are formed. However, increasing the fraction of the hard phases to ensure strength inevitably leads to a reduction in elongation. To compensate for the reduced elongation, a soft phase such as ferrite is formed in the microstructure. In ultra-high-strength steels with the above final microstructure, strength and elongation follow the rule of mixture (ROM) and thus enhancing material properties beyond the ROM is inherently limited.
Steels developed to improve the ultra-high-strength steels with a dual-phase structure of ferrite and martensite include TRIP steels capable of ensuring strength and elongation through the phase transformation of retained austenite during plastic deformation by ensuring retained austenite in the final structure. However, TRIP steels may not significantly enhance formability due to the low area fraction of retained austenite contained in the final microstructure.
Therefore, the present disclosure aims to enhance the formability of ultra-high-strength steel by ensuring retained austenite in the final microstructure, and the final microstructure of the produced steel sheet consists of ferrite, retained austenite, and tempered martensite (see
Accordingly, the methods proposed in the present disclosure to ensure yield strength, tensile strength, and elongation by implementing the above microstructures may be summarized as follows.
In the present disclosure, the final microstructure is set based on the following principles to ensure high formability, and process optimization for implementing the same was performed through simulation testing.
To configure the above-described microstructure, in the present disclosure, heat treatment conditions for cold-rolled steel sheets with a final microstructure of ferrite, retained austenite, and tempered martensite are established by applying intercritical annealing, rapid cooling, reheating, and continuous coating processes.
In relation to Design Directions 1 and 2, the application of intercritical annealing, rapid cooling, and reheating is required to configure the final microstructure.
Firstly, it is important to ensure a sufficient amount of ferrite in the final microstructure through intercritical annealing. Ferrite in the final microstructure may occur in a total of two steps of annealing and slow cooling among the heat treatment processes. When the annealing region is set as a single-phase region, it enters a dual-phase region during slow cooling and a phase transformation of austenite→ferrite is enabled. However, the amount is slight due to a very short time and thus a sufficient amount of ferrite desired in the present disclosure may not be formed. On the other hand, when the annealing region is set within a dual-phase region, ferrite may be ensured during slow cooling in addition to ferrite ensured in the dual-phase region, and thus the desired fraction of ferrite may be ensured and easily configured in the final microstructure. In the present disclosure, considering the control capability of a general continuous annealing line, the annealing temperature may be set within the dual-phase region, and the slow cooling end temperature may be set to 550° C. to 750° C. More specifically, the annealing temperature may be set to 820° C. to 860° C., and the slow cooling end temperature may be set to 600° C. to 700° C.
Secondly, it is important to ensure retained austenite through rapid cooling after slow cooling. When the steel sheet is cooled at a high rate without hardenability interference after slow cooling, austenite, which is present in addition to ferrite, is transformed into martensite and austenite due to rapid cooling. As such, the martensite structure capable of easily ensuring strength and the fraction of retained austenite capable of easily ensuring formability in the final microstructure may be determined before reheating. Therefore, when determining the fraction, it is important to find an appropriate temperature range between the cooling rate and the martensite start and finish temperatures. In the present disclosure, considering the control capability of a continuous annealing line, the rapid cooling rate is set to 70° C./s or more and the rapid cooling end temperature is set to 200° C. and 260° C. More specifically, the rapid cooling end temperature may be set to 210° C. to 250° C.
Thirdly, stabilization of the formed austenite and softening of the martensite structure are simultaneously caused by inducing carbon redistribution between phases in the formed microstructure through reheating (partitioning). After the reheating process, the stabilized austenite does not undergo further phase transformation even when cooled to room temperature, and this austenite is referred to as retained austenite. The retained austenite serves as a main base for simultaneously ensuring the formability and strength of the steel sheet by causing the TRIP phenomenon during subsequent plastic deformation. The amount of redistributed carbon generally increases with increasing reheating temperature and time, and the reheating process may be flexibly changed depending on the desired final material properties. In the present disclosure, the reheating temperature may be set to 370° C. and 460° C., and more specifically, to 400° C. and 430° C.
Lastly, the effects of the annealing heating rate and holding time, the slow cooling rate, and the reheating rate and holding time were checked. When the total annealing time is excessively long or short, the phase transformation in the corresponding unit process may proceed excessively to cause recrystallization and grain growth and prevent the formation of the desired microstructure, or proceed insufficiently so as not to satisfy the desired material properties. In the present disclosure, the annealing heating rate and holding time are set to 1° C./s to 10° C./s and 40 sec. to 120 sec., the slow cooling rate is set to 1° C./s to 10° C./s, and the reheating rate and holding time are set to 20° C./s or more and 10 sec. or more. More specifically, the annealing heating rate and holding time may be set to 1° C./s to 5° C./s and 50 sec. to 110 sec., the slow cooling rate may be set to 5° C./s to 10° C./s, and the reheating rate and holding time may be set to 30° C./s or more and 10 sec. to 240 sec.
In relation to Design Direction 3, the control of carbides formed in the final microstructure is required to ensure the optimal stability of retained austenite.
The present disclosure includes the final microstructure ensuring concept of Design Directions 1 and 2 and, at the same time, aims to ensure the desired material properties by controlling phase transformation based on plastic deformation behavior through the stability control of retained austenite which is a main structure for ensuring the material properties. An appropriate amount of an austenite stabilizing element needs to be included in the austenite structure during a series of heat treatment processes, particularly annealing and reheating. To this end, the reduction in diffusion of the austenite stabilizing element due to the formation of unnecessary inclusions needs to be suppressed to finally form adequately stable retained austenite in the microstructure after heat treatment.
In the present disclosure, C is specifically utilized as an austenite stabilizing element to form retained austenite in the final microstructure and, in this case, the formation of unnecessary carbide needs to be suppressed for the diffusion of C into austenite. In general, carbides are predominantly observed in martensite and ferrite phases. Particularly, in the present disclosure, the formation and growth of carbides occur prominently during reheating after annealing and cooling. The formation of carbides is primarily driven by two factors: 1) the phase transformation from austenite after rapid cooling into some bainite during reheating, and 2) the tempering of martensite during reheating. In this case, the formation of carbides is minimized by adding Si into the steel. However, although the above method is effective in limiting the growth of carbides previously formed during hot rolling and softening, controlling the amount of previously formed carbides is ultimately required to promote the diffusion behavior of C into austenite together with the component control effect.
Therefore, the present disclosure aims to ensure the stability of retained austenite by minimizing the amount of carbides formed in the final microstructure. Test results of the present disclosure show that optimal material properties may be ensured when a density of carbides (with an average size of 50 nm or more) in ferrite is 20 particles/μm2 or less. Additionally, these carbides in ferrite (soft phase) may act as crack initiation sites under deformation, thereby adversely affecting formability.
Based on the ultra-high-strength steel sheet and the method of manufacturing the same, according to the technical features of the present disclosure, strength and elongation were ensured by utilizing the strengthening mechanism and carbide control based on TRIP steels. To ensure the desired strength and elongation, unlike existing ultra-high-strength steels that ensure elongation by ensuring soft phases in the final microstructure, the present disclosure has ensured enhanced elongation by ensuring more retained austenite utilized for TRIP steels, in the final microstructure. Furthermore, the present disclosure has established carbide conditions for controlling the diffusion of an austenite stabilizing element such as C to optimally control the stability of retained austenite and satisfy the desired material properties.
Test examples will now be described for better understanding of the present disclosure. However, the following test examples are merely to promote understanding of the present disclosure, and the present disclosure is not limited to thereto.
The present test examples provide samples with the alloying element composition of Table 2.
al.
indicates data missing or illegible when filed
The component system of Table 2 represents a composition for configuring an ultra-high-strength steel sheet with excellent elongation, according to an embodiment of the present disclosure, and satisfies the composition of C: 0.1 wt % to 0.30 wt %, Si: 1.0 wt % to 2.0 wt %, Mn: 1.5 wt % to 3.0 wt %, Al: more than 0 wt % and not more than 0.05 wt %, a sum of one or more elements selected from Nb, Ti, and V: more than 0 wt % and not more than 0.05 wt %, P: more than 0 wt % and not more than 0.02 wt %, S: more than 0 wt % and not more than 0.005 wt %, N: more than 0 wt % and not more than 0.006 wt %, and a balance of Fe.
Table 3 shows process conditions applied to the present test examples with the alloy composition of Table 2. In Table 3, softening corresponds to the softening step (S200) shown in
Table 4 shows microstructures and properties achieved by the present test examples according to Tables 2 and 3. In Table 4, the criteria for determining whether the material properties are achieved include a YP of 850 MPa or more, a TS of 1180 MPa or more, an El of 14% or more, an area fraction of ferrite of 10% to 20%, an area fraction of retained austenite of 10% to 20%, an area fraction of tempered martensite/martensite of the remainder, and a density of carbides with an average size of 50 nm or more in ferrite of 20 particles/μm2 or less.
low
apid
nnealing
nnealing
nnealing
low
apid
eheating
eheating
eheating
00
30
00
40
00
30
00
40
00
30
00
40
00
30
30
40
75
30
30
40
00
50
00
40
00
70
00
40
00
30
30
40
00
30
70
40
00
30
00
40
00
30
00
40
indicates data missing or illegible when filed
errite
indicates data missing or illegible when filed
Referring to Tables 3 and 4, Test Examples 1, 2, 4, 6, and 8 correspond to embodiments of the present disclosure and satisfy process conditions such as a softening temperature of 500° C. to 650° C., an annealing heating rate of 1° C./s to 10° C./s, an annealing temperature of 820° C. to 860° C., an annealing holding time of 40 sec. to 120 sec., a slow cooling rate of 1° C./s to 10° C./s, a slow cooling end temperature of 550° C. to 750° C., a rapid cooling rate of 50° C./s or more, a rapid cooling end temperature of 200° C. to 260° C., a reheating rate of 20° C./s or more, a reheating temperature of 370° C. to 460° C., and a reheating holding time of 10 sec. to 240 sec. As such, the final microstructure consists of tempered martensite, martensite, ferrite, and retained austenite, the ferrite occupies an area fraction of 10% to 20%, the retained austenite occupies an area fraction of 10% to 20%, and the tempered martensite and martensite occupy the remaining area fraction. The ferrite includes carbides, and a density of the carbides with an average size of 50 nm or more is 20 particles/μm2 or less. A YP of 850 MPa to 1070 MPa, a TS of 1180 MPa to 1250 MPa, an El of 14% to 20%, and a product of TS and El of 15,000 MPa·% or more are all satisfied. On the other hand, Test Example 3 corresponds to a comparative example of the present disclosure and does not satisfy but exceeds the annealing temperature range of 820° C. to 860° C. As such, the area fraction of ferrite does not satisfy but falls below the range of 10% to 20%, the YP does not satisfy but exceeds the range of 850 MPa to 1070 MPa, the El does not satisfy but falls below the range of 14% to 20%, and the product of TS and El does not satisfy but falls below the range of 15,000 MPa·% or more. That is, when the annealing region is set as a single-phase region, it enters a dual-phase region during slow cooling and a phase transformation of austenite→ferrite is enabled. However, the amount is slight due to a very short time and thus a sufficient amount of ferrite desired in the present disclosure may not be formed, thereby decreasing elongation.
Test Example 5 corresponds to a comparative example of the present disclosure and does not satisfy but exceeds the slow cooling end temperature range of 550° C. to 750° C. As such, the area fraction of ferrite does not satisfy but falls below the range of 10% to 20%, the YP does not satisfy but exceeds the range of 850 MPa to 1070 MPa, the El does not satisfy but falls below the range of 14% to 20%, and the product of TS and El does not satisfy but falls below the range of 15,000 MPa·% or more.
Test Example 7 corresponds to a comparative example of the present disclosure and does not satisfy but exceeds the rapid cooling end temperature range of 200° C. to 260° C. As such, the area fraction of retained austenite does not satisfy but falls below the range of 10% to 20%, and the YP does not satisfy but falls below the range of 850 MPa to 1070 MPa.
Test Example 9 corresponds to a comparative example of the present disclosure and does not satisfy but exceeds the reheating temperature range of 370° C. to 460° C. As such, the YP does not satisfy but exceeds the range of 850 MPa to 1070 MPa, the El does not satisfy but falls below the range of 14% to 20%, and the product of TS and El does not satisfy but falls below the range of 15,000 MPa·% or more.
Test Examples 10 and 11 correspond to comparative examples of the present disclosure and do not satisfy but fall below the softening temperature range of 500° C. to 650° C. As such, the density of carbides with an average size of 50 nm or more in ferrite does not satisfy but exceeds the range of 20 particles/μm2 or less, and the El does not satisfy but falls below the range of 14% to 20%.
While the present disclosure has been particularly shown and described with reference to embodiments thereof, it will be understood by one of ordinary skill in the art that various changes in form and details may be made therein without departing from the scope of the present disclosure as defined by the following claims.
| Number | Date | Country | Kind |
|---|---|---|---|
| 10-2022-0125652 | Sep 2022 | KR | national |
This application is a continuation of International Application No. PCT/KR2022/019631 filed on Dec. 5, 2022, which claims under 35 U.S.C. § 119(a) the benefit of Korean Patent Application No. 10-2022-0125652 filed on Sep. 30, 2022, the entire contents of which applications are incorporated by reference herein.
| Number | Date | Country | |
|---|---|---|---|
| Parent | PCT/KR2022/019631 | Dec 2022 | WO |
| Child | 19058558 | US |
