Patent Application
20210071278
The present disclosure relates to a high-strength steel sheet and a method of manufacturing the same, and more particularly, to a high-strength steel sheet having a high yield ratio and excellent plating wettability and stretch flangeability, and a method of manufacturing the same.
Recently, the use of high-strength steel sheets is required to improve fuel efficiency or improve durability by various environmental regulations and energy use regulations. In particular, as impact stability regulations of automobiles have recently been expanded, high-strength steel sheets having excellent yield strength have been adopted in structural members such as members, seat rails, pillars and the like to improve impact resistance of the vehicle body. The structural member has a characteristic that is advantageous in absorbing impact energy as the yield strength compared with tensile strength, that is, the yield ratio (yield strength/tensile strength) is higher than the tensile strength. However, in general, as the strength of the steel sheet increases, elongation decreases, resulting in a problem in which workability deteriorates. Therefore, there is a need to develop a material that may compensate therefor.
Typically, methods for strengthening steel include solid solution strengthening, precipitation strengthening, strengthening by grain refinement, transformation strengthening, and the like. However, the solid solution strengthening and the strengthening by grain refinement among the above methods have a disadvantage that it is very difficult to manufacture high strength steel having a tensile strength of 490 MPa or higher.
On the other hand, the precipitation-strengthened high-strength steel is the technique to strengthen the steel by adding carbon and nitride forming elements such as Cu, Nb, Ti, V, etc. to precipitate carbon and nitride or to secure the strength by refinement of grains by suppressing grain growth by fine precipitates. The above technique has an advantage of easily obtaining high strength compared to low manufacturing costs, but has a disadvantage in that since the recrystallization temperature rapidly rises by fine precipitates, high temperature annealing must be performed to secure sufficient ductility by causing sufficient recrystallization. In addition, precipitation-strengthened steel, which is strengthened by precipitating carbon and nitride on a ferrite matrix, has a problem in that it is difficult to obtain high-strength steel of 600 MPa or more.
On the other hand, as the transformation reinforced high strength steel, ferrite-martensitic dual-phase steel containing hard martensite in a ferrite matrix, Transformation Induced Plasticity (TRIP) steel using the transformation induced plasticity of residual austenite, or Complex Phase (CP) steel which consists of low-temperature structure steel of ferrite and hard bainite or martensite, have been variously developed. However, the tensile strength that may be implemented in the improved high-strength steel (of course, the strength may be increased by increasing the amount of carbon, but considering practical aspects such as spot weldability or the like) has a limit of about 1200 Mpa level. In addition, hot press forming steel, which secures final strength by rapid cooling through direct contact with a die that is water-cooled after molding at high temperature, has been spotlighted for application to structural members to secure safety against collision. However, the expansion of application is not large due to excessive investment in facilities and high heat treatment and process costs.
In recent years, to further improve the stability of a passenger in a collision, high-strength and light-weighting of a vehicle's seat parts have been simultaneously conducted. These parts are manufactured by two methods, not only roll forming but also press forming. The seat part is a part that connects the passenger and the vehicle body, and must be supported with high stress to prevent the passenger from being bounced out during a collision. For this, high yield strength and yield ratio are required.
In addition, most of the parts to be processed require stretch flangeability, and therefore, it is required to apply a steel material having excellent hole expandability.
A representative manufacturing method for increasing yield strength is to use water cooling during continuous annealing. That is, a steel sheet having a tempered martensite structure may be produced by tempering process of soaking the steel in an annealing process and immersing the soaked steel sheet in water. However, this method may cause inferior shape quality due to temperature variations in the width and length directions during water cooling, so that has very serious disadvantages such as deterioration of workability and material deviation by location when roll forming is applied.
As one of the prior art references, Patent Document 1 may be provided. In this technique, after continuously annealing steel having 0.18% or more of carbon, water cooling is performed to room temperature, and then, the steel is subjected to an overaging treatment at a temperature of 120 to 300° C. for 1 to 15 minutes, thereby producing the steel containing 80 to 97% by volume of martensite. When the ultra-high strength steel is manufactured by the method of tempering after water cooling as in Patent Document 1, the yield ratio is very high, but a problem may arise in that the shape quality of the coil deteriorates due to the temperature deviation in the width direction and the length direction. Therefore, during roll forming, problems such as material defects and workability deterioration may occur depending on parts.
In addition, as one of the other prior art references, Patent Document 2 can be cited. In this case, a method of manufacturing a cold rolled steel sheet that obtains high strength and high ductility at the same time by utilizing tempered martensite and also has excellent plate shape after continuous annealing is disclosed. But there is concern about the inferiority of weldability due to the high carbon content of 0.2% or higher, and the possibility of causing dents in the furnace due to the high content of Si.
In addition, in the prior art, there is a problem that plating peeling occurs due to poor plating wettability depending on the components and operating conditions.
Therefore, there is a need for a high-strength steel material exhibiting high yield ratio and hole workability and excellent plating wettability by solving the above-mentioned problems.
(Patent Document 1) Japanese Patent Laid-Open Publication No. 1992-289120
(Patent Document 2) Japanese Patent Laid-Open Publication No. 2010-090432
An aspect of the present disclosure is to provide a high yield ratio-type high strength steel sheet having excellent plating wettability and stretch flangeability.
Another aspect of the present disclosure is to provide a method of manufacturing a high yield ratio-type high strength steel sheet having excellent plating wettability and stretch flangeability.
According to an aspect of the present disclosure, a high yield ratio-type high strength steel sheet includes:
in weight %, 0.04 to 0.09% of C, 0.2% or less (excluding 0%) of Si, 2.0 to 3.5% of Mn, 0.3 to 1.2% of Cr, 0.03 to 0.08% of Ti, 0.01 to 0.05% of Nb, 0.0010 to 0.0050% of B, 0.01 to 0.10% of Sol.Al, 0.001 to 0.10% of P, 0.010% or less (including 0%) of S, 0.010% or less (excluding 0%) of N, a remainder of Fe and unavoidable impurities, the contents of Si, Cr and C satisfying the following relationship 1,
(Si+Cr)/2C≥5 [Relationship 1]
where each component indicates a weight content; and
as a microstructure, in an area fraction, 50% or more of bainite, 20 to 40% of tempered martensite, and residual ferrite, in which a hardness ratio of a bainite phase and a tempered martensite phase at a ¼t point of a thickness of a steel sheet (where t is the thickness of the steel sheet (mm)), represented by the following relationship 2, is 1.3 or less,
(HB/HTM)≤1.3 [Relationship 2]
where HB is a bainite hardness, and HTM is a tempered martensite hardness.
The tempered martensite may have an average particle diameter of 2 μm or less, the bainite may have an average particle diameter of 3 μm or less, and an area fraction of a bainite structure having a particle size of more than 3 μm may be 5% or less.
A distribution density of a nano precipitate having a diameter of 10 nm or less among the bainite and tempered martensite structures may be 150 pieces/μm2 or more.
The steel sheet may have a tensile strength of 980 MPa or higher and a yield strength of 780 MPa or higher.
The steel sheet may have a yield ratio of 0.8 or higher and a HER value of 50% or higher.
The steel sheet may be provided with a zinc-based plating layer disposed on at least one surface thereof.
According to another aspect of the present disclosure, a method of manufacturing a high yield ratio-type high strength steel sheet, includes:
preparing a slab containing, in weight %, 0.04 to 0.09% of C, 0.2% or less (excluding 0%) of Si, 2.0 to 3.5% of Mn, 0.3 to 1.2% of Cr, 0.03 to 0.08% of Ti, 0.01 to 0.05% of Nb, 0.0010 to 0.0050% of B, 0.01 to 0.10% of Sol.Al, 0.001 to 0.10% of P, 0.010% or less (including 0%) of S, 0.010% or less (excluding 0%) of N, a remainder of Fe and unavoidable impurities, the contents of the Si, the Cr and the C satisfying the following relationship 1,
(Si+Cr)/2C≥5 [Relationship 1]
where each component indicates a weight content;
hot rolling the slab such that a temperature at a finish rolling outlet side is Ar3 to Ar3+50° C.° C., so as to obtain a hot rolled steel sheet;
coiling the hot rolled steel sheet at a temperature in a range of 600 to 750° C.;
cold rolling the hot rolled steel sheet at a reduction ratio of 40 to 70% so as to obtain a cold rolled steel sheet;
continuously annealing the cold rolled steel sheet at a temperature in a range of Ac3-60 to Ac3° C.;
performing a primary cooling to a temperature of 600 to 700° C. at a cooling rate of 1 to 20° C./sec after the continuously annealing;
after the primary cooling, performing a secondary cooling to a temperature of Ms-100 to Ms ° C. at a cooling rate of 5 to 20° C./second; and
after the secondary cooling, performing an overaging treatment,
wherein contents of the Mn and the B and an annealing temperature (a, ° C.) satisfy the following relationship 3,
19000 [B]+4.5 [Mn]+0.06a≤104 [Relationship 3]
where [B] and [Mn] indicate weight percent of a corresponding element.
An aging time in the overaging treatment may be controlled such that in an entire structure, an area fraction of bainite is 50% or more, and an area fraction of tempered martensite is 20 to 40%.
The method of manufacturing a high yield ratio-type high strength steel sheet may further include a zinc-based plating operation of forming a zinc-based plating layer on at least one surface of a steel sheet after the overaging treatment
The method of manufacturing a high yield ratio-type high strength steel sheet may further include an operation of skin-pass rolling at an elongation of 0.1 to 1.0% after the overaging treatment or after the zinc-based plating.
According to an exemplary embodiment of the present disclosure, a high yield ratio high strength steel sheet having not only high tensile strength of 980 MPa or more but also excellent plating wettability and stretch flangeability.
Hereinafter, exemplary embodiments of the present disclosure will be described.
However, embodiments of the present disclosure are provided to more fully describe exemplary embodiments of the present disclosure to those skilled in the art.
In addition, embodiments of the present disclosure may be modified in various other forms, and the scope of the present disclosure is not limited to the embodiments described below.
In addition, ‘comprising’ certain components throughout the specification indicates that other components may be further included, rather than excluding other components, unless otherwise specified.
Hereinafter, a high yield ratio-type high strength steel sheet according to an exemplary embodiment of the present disclosure will be described in detail.
A high yield ratio-type high strength steel sheet according to an exemplary embodiment of the present disclosure, includes: in weight %, 0.04 to 0.09% of C, 0.2% or less (excluding 0%) of Si, 2.0 to 3.5% of Mn, 0.3 to 1.2% of Cr, 0.03 to 0.08% of Ti, 0.01 to 0.05% of Nb, 0.0010 to 0.0050% of B, 0.01 to 0.10% of Sol.Al, 0.001 to 0.10% of P, 0.010% or less (including 0%) of S, 0.010% or less (excluding 0%) of N, a remainder of Fe and unavoidable impurities, the contents of Si, Cr and C satisfying the following relationship 1,
(Si+Cr)/2C≥5 [Relationship 1]
where each component indicates a weight content; and
as a microstructure, in an area fraction, 50% or more of bainite, 20 to 40% of tempered martensite, and residual ferrite, in which a hardness ratio of a bainite phase and a tempered martensite phase at a ¼t point of a thickness of a steel sheet (where t is the thickness of the steel sheet (mm)), represented by the following relationship 2, is 1.3 or less,
(HB/HTM)≤1.3 [Relationship 2]
where HB is a bainite hardness, and HTM is a tempered martensite hardness.
Hereinafter, components and component ranges will be described.
C: 0.04 to 0.09% by weight (hereinafter also referred to as “%”)
Carbon (C) is a very important element added to strengthen the transformation structure. Carbon promotes high strength and promotes the formation of martensite in the transformation structure steel. When the carbon content increases, the amount of martensite in steel increases. However, if the amount exceeds 0.09%, the strength of martensite increases, but the difference in strength with ferrite having a relatively low carbon concentration increases. This strength difference deteriorates stretch flangeability because breakage easily occurs at the interphase interface when stress is applied. In addition, due to poor weldability, welding defects occur when machining parts. On the other hand, if the carbon content is lowered to be less than 0.04%, it is difficult to secure the strength of martensite presented in the present disclosure. Therefore the content of C may be preferably limited to 0.04 to 0.09%, and the content of C may be, in more detail, 0.05 to 0.09%.
Si: 0.2% or less (excluding 0%)
Silicon (Si) is an element that promotes ferrite transformation and increases the carbon content in non-transformed austenite to form a composite structure of ferrite and martensite, thereby interfering with an increase in the strength of martensite. In addition, since it not only causes surface scale defects in relation to surface properties, but also reduces chemical conversion treatment, it may be desirable to limit the Si addition as much as possible.
Therefore, it may be preferable to limit the content of Si to 0.2% or less (excluding 0%).
Mn: 2.0 to 3.5%
Manganese (Mn) is an element that refines particles without damaging ductility and prevents hot embrittlement by the formation of FeS by precipitating sulfur into MnS, and at the same time, lowers the critical cooling rate at which martensite phase is obtained, thereby forming martensite more easily. If the content is less than 2.0%, it is difficult to secure the target strength in the present disclosure, and if the content exceeds 3.5%, since there is a high possibility of occurrence of a problem such as deterioration of weldability, hot rolling properties and plating wettability, it may be preferable to limit to the range of the manganese content to 2.0 to 3.5%, and in more detail, the manganese content may be 2.3 to 3.2%.
Cr: 0.3 to 1.2%
Chromium (Cr) is an element added to improve the hardenability of the steel and ensure high strength, and in the present disclosure, is an element that plays a very important role in forming the low-temperature transformation phase, martensite. If the content is less than 0.3%, it is difficult to secure the above-mentioned effect, and if the content exceeds 1.2%, the effect is not only saturated, but an excessive increase in hot rolling strength also causes the problem of cold rolling deterioration. Therefore, it may be desirable to limit the Cr content to 0.3 to 1.2%. The Cr content may be, in more detail, 0.3 to 1.0%.
Ti: 0.03 to 0.08% and Nb: 0.01 to 0.05%
Titanium (Ti) and niobium (Nb) are elements effective for increasing the strength of the steel sheet and refining the grains by nano precipitates. When these elements are added, the elements are combined with carbon to form very fine nano precipitates. These nano-precipitates serve to reduce the difference in hardness between phases by strengthening the matrix structure.
In the present disclosure, the content of titanium may be preferably limited to 0.03 to 0.08%, and the content of niobium may be preferably limited to 0.01 to 0.05%. If the content of each element is less than a lower limit value of the content of each component, the distribution density of the nano-precipitates decreases and the difference in the hardness ratio between phases increases. If the content of each element exceeds the lower limit value of the content of each component, ductility may be significantly reduced due to an increase in manufacturing costs and excessive precipitates. On the other hand, Ti may be, in more detail, 0.03 to 0.06%.
B: 0.0010 to 0.0050%
Boron (B) is an element that delays the transformation of austenite into pearlite in a process of cooling during annealing, and is added as an element that inhibits ferrite formation and promotes the formation of martensite. However, if the content is less than 0.0010%, it may be difficult to obtain the above-described effect, and if the content exceeds 0.0050%, a large amount of annealing oxide is generated in annealing, resulting in poor plating wettability and an increase in cost due to excessive ferroalloy content. Therefore, it may be desirable to limit the content thereof to 0.0010% to 0.0050%.
Sol.Al: 0.01 to 0.10%
Soluble aluminum (Sol.Al) is combined with oxygen to cause deoxidation and is a component effective to improve the martensitic hardenability by distributing carbon in ferrite to austenite. If the content is less than 0.01%, it may be difficult to secure the effect, and if the content exceeds 0.1%, the effect is not only saturated, but also increases the manufacturing costs, and thus, it may be preferable to limit the content of Sol.Al to 0.01 to 0.1%.
P: 0.001 to 0.10%
Phosphorus (P) is a substituted alloy element with a greatest solid solution strengthening effect, and serves to improve in-plane anisotropy and improve strength. If the content is less than 0.001%, not only the effect may not be secured, but also causes a problem in terms of manufacturing costs, whereas if added excessively, press formability may deteriorate and steel embrittlement may occur. Therefore, it may be desirable to limit the content to 0.001 to 0.10%.
S: 0.010% or less (including 0%)
Sulfur (S) is an impurity element in steel and is an element that inhibits the ductility and weldability of the steel sheet. If the content exceeds 0.010%, since the possibility of inhibiting the ductility and weldability of the steel sheet is high, it may be preferable to limit the content of S to 0.010% or less.
N: 0.010% or less (excluding 0%)
Nitrogen (N) is an element that is effective in stabilizing austenite, and if the content exceeds 0.010%, the risk of cracking during continuous casting through AlN formation or the like is greatly increased, and thus, the upper limit may be preferably limited to 0.01%.
It may be preferable that the contents of Si, Cr and C satisfy the following relationship 1.
(Si+Cr)/2C≥5, [Relationship 1]
in this case, each component indicates the weight content.
In a case in which the contents of Si, Cr and C do not satisfy the relationship 1, there is a problem that the strength is lowered.
In the relationship 1, the value of (Si+Cr)/2C should be 5 or more to secure the strength, and the C content of a secondary phase may be reduced to reduce the difference in phase strength. In detail, in the relationship 1, the solid solution concentration of Si and Cr, in detail, in ferrite, is increased to effectively reduce the difference in hardness between phases. The value of (Si+Cr)/2C may be, in more detail, 5.2 or more.
A high yield ratio-type high strength steel sheet according to an exemplary embodiment of the present disclosure includes, as a microstructure, in an area fraction, 50% or more of bainite, 20 to 40% of tempered martensite, and residual ferrite, in which a hardness ratio of a bainite phase and a tempered martensite phase at a ¼t point of a thickness of a steel sheet (where t is the thickness of the steel sheet (mm)), represented by the following relationship 2, is 1.3 or less.
(HB/HTM)≤1.3 [Relationship 2]
where HB is a bainite hardness, and HTM is a tempered martensite hardness.
Control of microstructures and precipitates is very important to manufacture a steel material having stretch flangeability of, e.g., HER 50% or more and a high yield ratio, e.g., a yield ratio of 0.8 or higher. As a method of simultaneously increasing the stretch flangeability and yield ratio, a technique for securing a low-temperature single-phase structure is required. In general, the structure having the highest strength among the low-temperature structure is martensite, and as is well known, an easiest method for preparing martensite is to perform water cooling and tempering treatment after annealing which is maintained for a sufficient time to form austenite.
However, the present disclosure is intended to secure martensite by controlling alloy elements, since the water cooling method may cause productivity deterioration due to problems such as material deviation, shape defects or the like. For example, it is the technique to secure martensite even at a low cooling rate by adding a certain amount or more of a hardenable element such as Mn, Cr and the like. However, this method may cause problems such as deterioration of weldability due to high addition of the alloying element. Therefore, in the present disclosure, it was intended to significantly reduce the carbon content that has the greatest influence on the weldability. Accordingly, in the present disclosure, the carbon content is limited to 0.09% or less.
To secure a relatively high yield ratio under the cooling conditions as in the present disclosure, it is necessary to add as many alloy elements as possible. However, these attempts additionally cause problems such as deterioration of weldability and an increase in hot rolled strength, and thus, the problems need to be resolved. However, when controlling nanoprecipitates and the size of martensite without adding excessive alloying elements according to exemplary embodiment, the stretch flangeability and yield ratio proposed in the present disclosure may be satisfied.
The high-strength steel sheet according to an exemplary embodiment of the present disclosure includes transformation structure and contains 50% or more of bainite and 20 to 40% of tempered martensite in an area fraction as a microstructure of a transformation structure. To increase the HER and yield ratio, a 100% transformation structure (bainite and tempered martensite) may be advantageous. To increase strength, it may be desirable to ensure that the size of the transformation structure is as small as possible. When the average particle diameter of martensite and bainite was outside the scope of the present disclosure, the bending workability, stretch flangeability and yield ratio could not be satisfied.
If the fraction of bainite in the microstructure is less than 50%, the required strength may not be satisfied.
If the content of the tempered martensite in the microstructure is less than 20%, the required strength may not be satisfied, and if the content exceeds 40%, the elongation may be lowered.
The average particle diameter of the tempered martensite may be 2 μm or less, the average particle diameter of bainite may be 3 μm or less, and the area fraction of the bainite structure having a particle size exceeding 3 μm may be 5% or less.
If the average particle diameter of the martensite is increased to 2 μm or more, or if the average particle diameter of bainite is 3 μm or more, the bending workability, stretch flangeability and yield ratio to be obtained in the present disclosure may not be obtained.
The distribution density of nano precipitates having a diameter of 10 nm or less among the bainite and tempered martensite structures may be 150/μm2 or more.
According to the research of the present inventors, it may be preferable to make the size of the precipitate as small as possible to improve the strength of the microstructure, and in detail, when the precipitation of 10 nm or less is secured to be 150 pieces/μm2 or more, a high yield ratio of 0.8 or more suggested by the present disclosure may be secured. In addition, the strength of the matrix structure is increased by the fine precipitates in the steel, so that the hardness ratio between phases is 1.3 or less, thereby manufacturing a high-strength steel sheet having a HER value of 50% or more and excellent stretch flange properties and yield strength.
The nano precipitates may be (Ti, Nb, Cr)C precipitates.
In terms of stretch flangeability, it is very important to control the strength ratio between phases, and for example, to secure 50% or more of HER, it is necessary to control the hardness ratio of the soft phase and the hard phase in the microstructure to be within 1.3. In the case in which the hardness value of a transformation phase and the hardness ratio between phases were not satisfied, it was impossible to secure a HER value of 50% or more and a yield ratio of 0.8 or more.
If the hardness ratio of the bainite phase and the tempered martensite phase represented by the following relationship 2 at the point of ¼ t of the thickness of the steel sheet (where t indicates the thickness of the steel sheet (mm)) is greater than 1.3, as the phase-to-phase strength deviation increases, the HER and bending workability may be deteriorated.
(HB/HTM)≤1.3 [Relationship 2]
where HB indicates bainite hardness, and HTM indicates tempered martensite hardness.
The steel sheet may have a tensile strength of 980 MPa or more and a yield strength of 780 MPa or more.
The steel sheet may have a yield ratio of 0.8 or more and a HER value of 50% or more.
A zinc-based plating layer may be formed on at least one surface of the steel sheet.
Hereinafter, a method of manufacturing a high yield ratio-type high strength steel sheet according to another exemplary embodiment of the present disclosure will be described in detail.
A method of manufacturing a high yield ratio-type high strength steel sheet according to another exemplary embodiment of the present disclosure includes: preparing a slab including, in weight %, 0.04 to 0.09% of C, 0.2% or less (excluding 0%) of Si, 2.0 to 3.5% of Mn, 0.3 to 1.2% of Cr, 0.03 to 0.08% of Ti, 0.01 to 0.05% of Nb, 0.0010 to 0.0050% of B, 0.01 to 0.10% of Sol.Al, 0.001 to 0.10% of P, 0.010% or less (including 0%) of S, 0.010% or less (excluding 0%) of N, a remainder of Fe and unavoidable impurities, the contents of the Si, the Cr and the C satisfying the following relationship 1,
(Si+Cr)/2C≥5 [Relationship 1]
where each component indicates a weight content;
hot rolling the slab such that a temperature at a finish rolling outlet side is Ar3 to Ar3+50° C.° C., so as to obtain a hot rolled steel shee;
coiling the hot rolled steel sheet at a temperature in a range of 600 to 750° C.;
cold rolling the hot rolled steel sheet at a reduction ratio of 40 to 70% so as to obtain a cold rolled steel sheet;
continuously annealing the cold rolled steel sheet at a temperature in a range of Ac3-60 to Ac3° C.;
performing a primary cooling to a temperature of 600 to 700° C. at a cooling rate of 1 to 20° C./sec after the continuously annealing;
after the primary cooling, performing a secondary cooling to a temperature of Ms-100 to Ms ° C. at a cooling rate of 5 to 20° C./second; and
after the secondary cooling, performing an overaging treatment,
wherein contents of the Mn and the B and an annealing temperature (a, ° C.) satisfy the following relationship 3.
19000 [B]+4.5 [Mn]+0.06a≤104 [Relationship 3]
where [B] and [Mn] indicate weight percent of a corresponding element.
Hot Rolling
The hot-rolled steel sheet is obtained by hot rolling the slab having the above-described steel composition such that the temperature at the outlet side of the finish rolling is Ar3 to Ar3+50° C.
If the temperature at the outlet side is less than Ar3 during hot rolling, the hot deformation resistance is likely to increase rapidly, and top and tail portions and edges of the hot rolled coil become single-phase areas, thereby increasing in-plane anisotropy and deteriorating formability. On the other hand, if the temperature at the outlet side exceeds Ar3+50° C., not only too large oxidation scales occur, but also there is a possibility that the microstructure of the steel sheet may become coarse.
Coiling
The hot-rolled steel sheet, having been subjected to hot rolling as described above, is coiled at a temperature in a range of 600 to 750° C.
If the coiling temperature is less than 600° C., excessive martensite or bainite is generated, resulting in excessive strength increase of the hot rolled steel sheet, which may cause manufacturing problems such as shape defect or the like due to load during cold rolling. On the other hand, if the temperature exceeds 750° C., the pickling property deteriorates due to an increase in the surface scale, and thus, it may be preferable to limit the coiling temperature to 600 to 750° C.
Cold Rolling
After the coiling operation, the hot-rolled steel sheet is cold-rolled at a reduction ratio of 40 to 70% to obtain a cold-rolled steel sheet. It may be preferable to perform a pickling process before cold rolling the hot rolled steel sheet.
If the rolling reduction is less than 40% during the cold rolling, the recrystallization driving force is weakened, and there is a great possibility of problems in obtaining a good recrystallized grain, and shape correction is very difficult.
On the other hand, if the rolling reduction ratio exceeds 70%, there is a high possibility of cracking at the edge of the steel sheet, and the rolling load increases rapidly.
Continuous Annealing
The cold rolled steel sheet having subjected to cold rolling as described above is continuously annealed at a temperature in the range of Ac3-60 to Ac3° c.
The cold rolled steel sheet undergoes a continuous annealing process, thereby providing a base for a microstructure required in the present disclosure. The continuous annealing may be preferably performed in the temperature range of Ac3-60 to Ac3. If the annealing temperature is less than Ac3-60° C., it is difficult to secure a yield ratio of 0.8 or higher since the ferrite is generated in a relatively large amount and the yield strength is thus lowered. In detail, the difference in hardness between phases with the transformation phase increases due to the production of large amounts of ferrite, and thus, the condition of 1.3 or less may not be satisfied. On the other hand, if the annealing temperature exceeds Ac3, the size of the martensite phase produced upon cooling increases due to the increase in the austenitic grain size due to the high temperature annealing, and thus, the particle size of the preferred microstructure of the present disclosure may not be satisfied. In addition, as the annealing temperature increases, the amount of annealing oxide increases on a surface layer portion, resulting in poor plating adhesion. In addition, even if the annealing temperature condition is satisfied, if the annealing temperature does not satisfy the following relationship 3, the plating wettability is inferior and the plating adhesion deteriorates.
19000[B]+4.5[Mn]+0.06a≤104 [Relationship 3]
where [B] and [Mn] indicate the weight percent of the corresponding element, and a represents the annealing temperature (° C.)).
Primary Cooling
After the continuous annealing, primary cooling is performed at a cooling rate of 1 to 20° C./sec to a temperature of 600 to 700° C.
This primary cooling is a process of inhibiting ferrite transformation to transform most of austenite to martensite.
In the case of the primary cooling, the ferrite fraction may increase if the cooling end temperature is less than 600° C., and the ferrite fraction may increase if the cooling end temperature exceeds 700° C.
In the case of the primary cooling, if the cooling rate is less than 1° C./sec, the ferrite fraction may increase, and there is no problem even in the case in which the ferrite fraction exceeds 20° C./sec, but the upper limit of the cooling rate may be limited to 20° C./sec in consideration of the cooling capacity of the cooling equipment or the like.
Secondary Cooling
After the primary cooling, secondary cooling is performed at a cooling rate of 5 to 20° C./sec to a temperature of Ms-100° C. to Ms° C.
In the secondary cooling, austenite begins to transform into martensite, but it does not cool down to below the temperature at which the martensite transformation ends, thereby obtaining a tempered martensite structure.
By performing the secondary cooling to the temperature range of Ms-100 to Ms ° C., and then performing overaging while maintaining the temperature, high yield strength and a high hole expanding ratio (HER) may be secured, as well as securing the shape in the width and length directions of the coil.
In the case of the secondary cooling, if the cooling end temperature is less than Ms-100° C., the amount of martensite increases excessively during the overaging treatment, resulting in simultaneous increase in yield strength and tensile strength and causing very deteriorating ductility. In detail, shape deterioration occurs due to rapid cooling, and thus workability deterioration may occur when machining automotive parts. On the other hand, if cooling to a temperature exceeding Ms is performed, the austenite produced during annealing may not be transformed into martensite and may be produced as high temperature transformation phases, such as bainite, granular bainite, or the like, causing a problem in which the yield strength is rapidly deteriorated. The occurrence of such a structure is accompanied by a decrease in yield ratio and deterioration of hole expandability, and thus, a high yield ratio-type high strength steel sheet having excellent stretch flangeability proposed in the present disclosure may not be manufactured.
In the case of the secondary cooling, if the cooling rate is less than 5° C./sec, the ferrite fraction may increase, and if the cooling rate exceeds 20° C./sec, there is a concern that quality problems such as twisting of the coil shape or the like due to rapid cooling may occur.
Overaging
After the secondary cooling, overaging is performed.
In the overaging, the aging time may be preferably adjusted such that the entire structure becomes 50% or more of bainite and 20 to 40% of tempered martensite in an area fraction.
After the overaging, a zinc-based plating operation of forming a zinc-based plating layer on at least one surface of the steel sheet may be further included.
Skin Pass Rolling
After the overaging, or after the zinc-based plating operation, an operation of skin pass rolling with an elongation of 0.1 to 1.0% may further be included.
In general, in case of skin-pass rolling the transformed structure steel, a yield strength increase of at least 50 MPa or more occurs with little increase in tensile strength. If the elongation is less than 0.1%, it is difficult to control the shape, and if the elongation exceeds 1.0%, the workability is greatly unstable by the high elongation operation.
Therefore, it may be preferable to limit the elongation to 0.1 to 1.0% at the time of performing skin pass rolling.
According to a method of manufacturing a high yield ratio-type high strength steel sheet according to another exemplary embodiment of the present disclosure, a high yield ratio-type high strength steel sheet having, not only a tensile strength of 980 MPa or more, a yield strength of 780 MPa or more, a yield ratio of 0.8 or more and an HER value of 50% or more, but also excellent plating wettability, may be obtained.
Hereinafter, an exemplary embodiment of the present disclosure will be described in more detail through examples. However, it should be noted that the embodiments described below are only intended to exemplify the present disclosure and are not intended to limit the scope of the present disclosure. This is because the scope of the present disclosure is determined by the items described in the claims and the items reasonably inferred therefrom.
The steel slab having the steel composition of Table 1 and satisfying the conditions of Table 2 below was reheated for 1 hour at a temperature of 1200° C. in a heating furnace, and was hot rolled under the conditions of Table 3, followed by being coiled.
After pickling the hot-rolled steel sheet, cold rolling was performed at a cold rolling reduction of 45% to prepare a cold-rolled steel sheet.
The cold-rolled steel sheet was subjected to continuous annealing, primary and secondary cooling under the conditions illustrated in Table 3 below, and then, subjected to skin pass rolling at a rolling reduction rate of 0.2%. In this case, the primary cooling was performed to 650° C.
In Table 3, FDT is a hot finish rolling temperature, CT is a coiling temperature, SS is a continuous annealing temperature, and RCS is a secondary cooling end temperature.
For the cold-rolled steel sheet having been subjected to the skin pass rolling as described above, the microstructure fraction, the average particle diameter of martensite (TM) and bainite (B), the hardness ratio between phases, and the distribution density of nano precipitates of 10 nm or less in the steel were investigated, and the results are illustrated in Table 4 below.
In this case, the hardness of the transformation structure was measured by using a Nano-Indenter (NT110) device to measure 100 points in a square shape with a load of 2 g to utilize values excluding maximum and minimum values. In addition, bainite, martensite and nano precipitates were evaluated by FE-TEM. In detail, the size and distribution density of nano precipitates were evaluated with an image of precipitate structure measured by FE-TEM, by using an image analyzer (image analysis) equipment. In addition, the fraction of the microstructure was observed by SEM and then the image analyzer (image analysis) was used.
In addition, a JIS 5 tensile specimen was prepared to measure yield strength (YS), tensile strength (TS), elongation (T-El), HER, yield ratio (YR) and plating wettability, and the results are illustrated in Table 5 below.
As illustrated in
On the other hand, the microstructure of the inventive example (5-1) was observed, and the results are illustrated in
As illustrated in Tables 1 to 5, since the inventive examples satisfying the steel composition and manufacturing conditions of the present disclosure exhibit tensile strength of 980 MPa or higher, yield strength of 780 MPa or higher, the yield ratio of 0.8 or higher, and the HER value of 50% or higher; it can be seen that plating wettability and elongation as well as the tensile strength, the yield strength, the yield ratio, and hole expandability (HER) are excellent.
As illustrated in
Meanwhile, in the case of comparative examples that do not satisfy at least one of the steel composition and manufacturing conditions proposed in the present disclosure, it can be seen that at least one property of tensile strength, yield strength, yield ratio, HER value and plating wettability is inferior or insufficient.
In the case of Comparative Examples (2-1 and 5-2), the composition satisfy the scope of the present disclosure, but the secondary cooling conditions, RCS temperatures, are 440° C. and 450° C., which do not satisfy the cooling temperature of the Ms temperature or less suggested in the present disclosure. Therefore, the austenite produced during annealing could not be transformed into martensite due to high temperature overaging, and was generated as high temperature transformation phases such as bainite, granular bainite and the like, such that coarse transformation phases occurred. These coarse transformation phases have high hardness ratios between phases of microstructures, resulting in low yield ratios and deterioration of HER values.
In Comparative Example (6-2), the annealing temperature was very low and annealing was performed in a two-phase region of austenite and ferrite, whereby the bainite fraction was 41% and the tempered martensite transformation fraction was 10%, which did not satisfy the microstructure fraction according to an exemplary embodiment of the present disclosure. The production of such ferrite caused an increase in deviation of a hardness ratio between phases, resulting in a low yield ratio and deterioration of the HER value.
In Comparative Example (6-3), the annealing temperature was very high temperature of 890° C., resulting in an increase in the austenitic grain size due to high temperature annealing to increase the size of martensite structure produced during cooling. Thus, martensite having the average particle diameter of 2 μm or less and bainite having the average particle diameter of 3 μm or less could not be manufactured, and thus, the yield ratio and HER value were deteriorated.
Comparative Examples 10 to 15 and 17 did not satisfy the relationship 1. This increase in carbon serves to increase the strength of martensite produced in the quenching process after annealing. However, when overaging after quenching, all martensite cannot be tempered and remain in lath form. In the case of the tempered martensite generated at this time, the strength decreases due to the precipitation of carbon, but the lath-type martensite which is not tempered has a very high strength due to the added carbon as a very stable martensite. Therefore, if the carbon content exceeds the composition suggested in the present disclosure steel, the HER value and yield ratio do not satisfy the criteria suggested in the inventive steel due to the increase in the strength difference between the lath martensite and the tempered martensite produced in the overaging treatment.
Comparative Examples 1-3, 5-2, 5-3, 6-3, 10, 11, 12, 15, 16 and 17 did not satisfy the relationship 3, resulting in poor plating wettability. As the B and Mn contents increase, Mn—B-based oxides are formed on the surface during annealing, thereby weakening the wettability of molten zinc and composite oxide to deteriorate plating wettability. Even in the case in which the contents of B and Mn are increased, when the annealing temperature is relatively low, Mn—B oxide is relatively less generated, and thus the plating adhesion is good.
In Comparative Examples 12 to 17, the carbon content or Mn, Cr, Ti and Nb contents did not satisfy the component range according to an exemplary embodiment of the present disclosure. In Comparative Examples 12 and 15, the C content was high, and thus, the hardness ratio between phases was high, and the yield ratio and HER value were deteriorated by the generation of coarse martensite.
In Comparative Examples 13 and 14, due to the low Mn or Cr contents, sufficient transformation of martensite did not occur, and Comparative Example 16 has a very high Si content.
In general, Si is a ferrite forming element, and when the addition amount is increased, the formation of ferrite is promoted upon cooling. In Comparative Example 16, the amount of martensite produced due to high Si addition was 11%, which did not satisfy the criteria presented in the present disclosure, and the yield ratio was low and the HER value deteriorated due to an increase in the phase-to-phase hardness ratio.
Comparative Example 17 is a case in which Ti and Nb contents did not satisfy the composition range proposed in the present disclosure. As mentioned above, Ti and Nb are combined with carbon to form nano-precipitates, and these nano-precipitates strengthen the matrix structure to reduce the difference in hardness between phases. However, in Comparative Example 17, Ti and Nb were very small, so that sufficient precipitates could not be formed, and as a result, yield ratios and HER values were deteriorated due to the distribution of nanoprecipitates and an increase in the hardness ratio between phases.
Although the embodiments of the present disclosure have been described above, those skilled in the art to which the present disclosure pertains will understand that the present disclosure may be implemented in other specific forms without changing its technical spirit or essential features.
Therefore, it should be understood that the above-described embodiments are illustrative in all respects and not restrictive. The scope of the present disclosure is indicated by the following claims rather than the above detailed description, and all changes or modified forms derived from the meaning and scope of the claims and equivalent concepts thereof should be interpreted to be included in the scope of the present disclosure.
| Number | Date | Country | Kind |
|---|---|---|---|
| 10-2017-0178944 | Dec 2017 | KR | national |
| Filing Document | Filing Date | Country | Kind |
|---|---|---|---|
| PCT/KR2018/014272 | 11/20/2018 | WO | 00 |
