KR102131527B1 - High-strength steel sheet with excellent durability and method for manufacturing thereof - Google Patents

Abstract

The present invention relates to steel materials used for members of chassis parts of commercial vehicles and wheel discs, and more particularly, to high-strength steel materials having excellent durability and a method for manufacturing the same.

Description

High-strength steel with excellent durability and its manufacturing method {HIGH-STRENGTH STEEL SHEET WITH EXCELLENT DURABILITY AND METHOD FOR MANUFACTURING THEREOF}

The present invention relates to a high strength steel material having excellent durability and a method for manufacturing the same.

Conventionally, members of chassis parts and wheel disks of commercial vehicles have used a high-strength steel sheet having a thickness of 5 mm or more and a yield strength of 450 to 600 MPa in order to secure high rigidity due to the characteristics of the vehicle, but recently, the weight and strength of automobiles have been improved. In order to apply a high-strength steel material with a tensile strength of 650 MPa or more.

When manufacturing a part using high-strength steel, it undergoes a step of manufacturing by pressing and forming a sheet material that has undergone shear molding and punching. During the above shear molding and punching molding, fine cracks are generated in the front end portion of the steel sheet to finalize it. There is a disadvantage that the durability life of the product (part) is shortened.

As a method for solving this problem, in Patent Document 1, the ferrite phase is formed as a matrix structure by winding at a high temperature after hot rolling in a normal austenite region to form a precipitate finely. In addition, Patent Document 2 proposes a technique for cooling the coiling temperature to a temperature at which the bainite phase is formed as a matrix structure so that a coarse pearlite structure is not formed. Patent Document 3 discloses a technique for miniaturizing austenite grains by rolling with a rolling reduction of 40% or more in an unrecrystallized region during hot rolling using titanium (Ti), niobium (Nb), or the like.

For the manufacture of high-strength steel, alloy components such as Si, Mn, Al, Mo, and Cr are mainly used. In this case, it is effective in improving the strength of the hot-rolled steel sheet, but when a large amount of alloy components is added, some components are segregated in the steel ( segregation), or non-uniformity of the microstructure, resulting in poor shear formability, and micro-cracks on the shear surface are easily propagated in the fatigue environment, resulting in component damage.

Particularly, as the thickness of the steel material increases, the microstructure non-uniformity between the thickness surface layer portion and the center portion increases, resulting in increased cracking of the shear surface, and a rapid propagation rate of cracks in a fatigue environment, resulting in poor durability.

However, the previous techniques (Patent Documents 1 to 3) do not consider the fatigue properties of thick steel materials having high strength.

In addition, in the case of using precipitate forming elements such as Ti, Nb, V, etc. in order to refine the grains of the thick steel material and obtain a precipitation strengthening effect, it is wound at a high temperature of about 500 to 700° C., where precipitates are easily formed, or cooled after hot rolling. If the cooling rate of the steel sheet is not controlled, coarse carbides are formed at the center of the thickness of the thick steel material, and the quality of the shear surface is inferior. In addition, applying a large pressure of 40% in the non-recrystallized zone during hot rolling degrades the shape quality of the rolled sheet and brings a load on equipment, making it difficult to apply in practice.

Japanese Patent Application Publication No. 2002-322541 Korean Registered Publication No. 10-1528084 Japanese Patent Application Publication No. 1997-143570

One aspect of the present invention is to provide a steel material having a certain thickness, and not only has high strength, but also has excellent durability and a method for manufacturing the same.

The subject of this invention is not limited to the above-mentioned content. Anyone having ordinary knowledge in the technical field to which the present invention pertains will have no difficulty in understanding additional problems of the present invention from the contents throughout the present specification.

One aspect of the present invention, by weight, carbon (C): 0.05 to 0.15%, silicon (Si): 0.01 to 1.0%, manganese (Mn): 1.0 to 2.3%, aluminum (Al): 0.01 to 0.1% , Chromium (Cr): 0.005 to 1.0%, Phosphorus (P): 0.001 to 0.05%, Sulfur (S): 0.001 to 0.01%, Nitrogen (N): 0.001 to 0.01%, Niobium (Nb): 0.005 to 0.07% , Titanium (Ti): 0.005 ~ 0.11%, contains the balance of Fe and other inevitable impurities,

As a microstructure, the fraction of ferrite and bainite phases is 90% or more, and the proportion of crystal grains (ratio of short side/long side) in the central part (thickness t/4 to t/2 point) is less than 50%. , And provides a high strength steel material having excellent durability at a grain boundary length of 700 mm or more observed in a unit area (1 cm ² ) at the center.

Another aspect of the present invention, heating the steel slab having the above-described alloy composition in a temperature range of 1200 ~ 1350 ℃; Hot-rolling the heated steel slab to produce a hot-rolled steel sheet; Cooling the hot-rolled steel sheet to a temperature range of 400 ~ 500 ℃ and then winding (CT); And after the winding, including the step of air cooling to a temperature range of room temperature ~ 200 ℃,

The hot rolling is finished hot rolling at a temperature (FDT (°C)) satisfying the following [Relational Formula 1], the cooling is performed by the following primary cooling and secondary cooling, and the primary cooling is [Relational Formula 2] at a cooling rate (CR ₁₎ is satisfied, there is provided a method of manufacturing a high-strength steel material having excellent durability, characterized in that for performing a cooling rate (CR ₂₎ satisfying the following the secondary cooling [Expression 3].

[Relationship 1]

Tn-50 ≤ FDT (Hot rolling end temperature (℃)) ≤ Tn

Tn = 730 + 92×[C] + 70×[Mn] + 45×[Cr] + 650×[Nb] + 410×[Ti]-80×[Si]-1.4×(t-5) (where Each element means weight content (%), and t means thickness (mm) of the final hot rolled steel sheet)

[Relationship 2]

CR ₁ ≥ 196-300 × [C] + 4.5 × [Si]-71.8 × [Mn]-59.6 × [Cr] + 187 × [Ti] + 852 × [Nb] (where each element has a weight content (% Means ))

[Relationship 3]

CR _Min ≤ CR ₂ ≤ CR _Max

(CR _Max = 76.6-157 × [C]-25.2 × [Si]-14.1 × [Mn]-27.3 × [Cr] + 61 × [Ti] + 448 × [Nb], CR _Min = 27.4-45.3 × [C] + 5.28×[Si]-11×[Mn]-7.33×[Cr] + 42.3×[Ti] + 82×[Nb], each element means weight content (%))

According to the present invention, it is possible to provide a thick steel material having high strength and excellent cross-section quality during molding, thereby ensuring excellent ratio of fatigue strength and yield strength of the steel after molding.

The steel material of the present invention has an effect that can be suitably applied to members of a chassis component of a vehicle and wheel disk.

1 is a graph showing the ratio of fatigue strength and yield strength according to the thickness of an invention steel and a comparative steel in an embodiment of the present invention.
Figure 2, in one embodiment of the present invention, shows the results of observing the microstructure of invention steel 4 (in Figure 2, Min, Max are the minimum and maximum values of the aspect ratio (short side length of crystal grains/long side length of crystal grains), respectively. Means, Total fraction refers to the area fraction of crystal grains that fall within the range of Min (excess) to Max (below).

The present inventors have studied in depth to solve the problem of deterioration in durability when forming a steel material for automobiles.

In particular, the present inventors investigated the change in the distribution of cracks and durability in the shear surface after molding according to the components and microstructures of existing thick steel materials, and confirmed that the durability characteristics were changed depending on the shape control of the grains in the center of the thickness of the steel material.

From this, the present inventors have confirmed that it is possible to provide a steel material having high strength and a targeted durability due to excellent cross-section quality during molding, and to complete the present invention.

Hereinafter, the present invention will be described in detail.

The high strength steel material having excellent durability according to an aspect of the present invention is in weight%, carbon (C): 0.05 to 0.15%, silicon (Si): 0.01 to 1.0%, manganese (Mn): 1.0 to 2.3%, aluminum (Al ): 0.01 to 0.1%, chromium (Cr): 0.005 to 1.0%, phosphorus (P): 0.001 to 0.05%, sulfur (S): 0.001 to 0.01%, nitrogen (N): 0.001 to 0.01%, niobium (Nb) ): 0.005 to 0.07%, titanium (Ti): may include 0.005 to 0.11%.

Hereinafter, the reason for limiting the alloy composition of the hot-rolled steel sheet provided in the present invention as described above will be described in detail.

On the other hand, unless specifically stated in the present invention, the content of each element is based on weight, and the proportion of tissue is based on area.

Carbon (C): 0.05~0.15%

Carbon (C) is the most economical and effective element for reinforcing steel, and as the amount added increases, the precipitation strengthening effect increases or the fraction of bainite phase increases, thereby improving tensile strength. In addition, the thicker the thickness of the hot-rolled steel material, the slower the cooling rate in the center of the thickness during cooling after hot rolling, so that when the content of C is large, coarse carbide or pearlite is likely to be formed.

In the present invention, if the content of C is less than 0.05%, it is difficult to sufficiently obtain the strengthening effect of steel, whereas if it exceeds 0.15%, pearlite phase or coarse carbide is formed in the center of the thickness, resulting in poor shear formability and reduced durability. there is a problem.

Therefore, in the present invention, the C may be included as 0.05 to 0.15%, and more advantageously as 0.06 to 0.12%.

Silicon (Si): 0.01~1.0%

Silicon (Si) deoxidizes molten steel and has a solid solution strengthening effect, and is advantageous in improving moldability by delaying the formation of coarse carbide.

If the Si content is less than 0.01%, the solid solution strengthening effect is small, and the effect of delaying carbide formation is also low, making it difficult to improve moldability. On the other hand, when the content exceeds 1.0%, a red scale formed by Si is formed on the surface of the steel sheet during hot rolling, so that the surface quality of the steel sheet is very poor, and there is a problem that the ductility and weldability are also deteriorated.

Therefore, in the present invention, the Si may be included in an amount of 0.01 to 1.0%, and more advantageously, in an amount of 0.2 to 0.7%.

Manganese (Mn): 1.0-2.3%

Manganese (Mn), like Si, is an effective element for solid solution strengthening of steel, and increases the hardenability of steel to facilitate formation of a bainite phase during cooling after hot rolling.

If the Mn content is less than 1.0%, the above-described effect cannot be sufficiently obtained. On the other hand, when the content exceeds 2.3%, the hardenability is greatly increased, so that the transformation of martensite is likely to occur, and the segregation part is greatly developed at the center of the thickness during slab casting in the playing process, and when cooled after hot rolling, fine in the thickness direction The tissue is formed non-uniformly, resulting in poor shear formability and durability.

Therefore, in the present invention, the Mn may be included in an amount of 1.0 to 2.3%, and more advantageously, in an amount of 1.1 to 2.0%.

Aluminum (Al): 0.01~0.1%

Aluminum (Al) is an element mainly added for deoxidation. If the content is less than 0.01%, the additive effect cannot be sufficiently obtained. On the other hand, when the content exceeds 0.1%, by forming AlN in combination with nitrogen (N) in steel, corner cracks are easily generated in the slab during continuous casting, and there is a risk of defects due to inclusion formation.

Therefore, in the present invention, Al may be included as 0.01 to 0.1.

Meanwhile, in the present invention, aluminum means soluble aluminum (Sol. Al).

Chromium (Cr): 0.005~1.0%

Chromium (Cr) solidifies and strengthens the steel, and helps to form bainite at the coiling temperature by retarding the transformation of the ferrite phase upon cooling. In order to obtain the above-described effect, it is preferable to contain Cr in an amount of 0.005% or more, but if the content exceeds 1.0%, the elongation is inferior as the martensite phase is formed by excessively delaying ferrite transformation. In addition, similar to Mn, the segregation part is greatly developed in the center of the thickness, and the shearing formability and durability are deteriorated by making the microstructure in the thickness direction non-uniform.

Therefore, in the present invention, the Cr may be included at 0.005 to 1.0%, and more advantageously at 0.3 to 0.9%.

Phosphorus (P): 0.001~0.05%

Phosphorus (P) is an element that simultaneously has the effect of strengthening solid solution and promoting ferrite transformation. In order to manufacture the P content of less than 0.001%, the manufacturing cost is excessive, which is economically disadvantageous, and it is difficult to secure the target level of strength. On the other hand, when the content of P exceeds 0.05%, brittleness by grain boundary segregation occurs, fine cracking is likely to occur during molding, and shear formability and durability are greatly deteriorated.

Therefore, in the present invention, the P may be included as 0.001 to 0.05%.

Sulfur (S): 0.001 to 0.01%

Sulfur (S) is an impurity present in the steel, and when its content exceeds 0.01%, it forms a non-metallic inclusion by combining with Mn, etc., thereby easily generating fine cracks during cutting of steel and greatly improving shear formability and durability. There is a problem of deterioration. On the other hand, in order to manufacture the content of S to less than 0.001%, it takes too much time during steel-making operation, and productivity decreases.

Therefore, in the present invention, S may be included as 0.001 to 0.01%.

Nitrogen (N): 0.001 to 0.01%

Nitrogen (N) is a typical solid solution strengthening element with C, and combines with Ti, Al, etc. to form coarse precipitates. In general, the solidification effect of N is superior to carbon, but there is a problem that the toughness of steel decreases as the amount of N increases in steel. In consideration of this, it is preferable to include the N in an amount of 0.01% or less, but in order to manufacture the content to less than 0.001%, it takes a lot of time during steelmaking operation, resulting in a decrease in productivity.

Therefore, in the present invention, the N may be included as 0.001 to 0.01%.

Niobium (Nb): 0.005 to 0.07%

Niobium (Nb) precipitation strengthening element. It is effective in improving the strength and impact toughness of steel due to the grain refinement effect due to recrystallization delay by precipitation during hot rolling. In order to sufficiently obtain the above-described effect, it may be included in an amount of 0.005% or more, while if the content exceeds 0.07%, formability and durability are formed by formation of stretched grains due to excessive recrystallization delay during hot rolling and formation of coarse composite precipitates. You become inferior.

Therefore, in the present invention, the Nb may be included as 0.005 to 0.07%, and more advantageously as 0.01 to 0.06%.

Titanium (Ti): 0.005~0.11%

Titanium (Ti) is a representative precipitation strengthening element together with the Nb, and forms a coarse TiN in the steel with strong affinity with N. The TiN has an effect of inhibiting the growth of crystal grains during the heating process for hot rolling. In addition, the Ti remaining after reacting with N is dissolved in the steel to form TiC precipitates by bonding with carbon, which is useful for improving the strength of the steel.

In order to sufficiently obtain the above-described effect, it is necessary to contain Ti in an amount of 0.005% or more, but when the content exceeds 0.11%, there is a problem of inferior impact resistance during molding by generation of coarse TiN and coarsening of precipitates. have.

Therefore, in the present invention, the Ti may be included in an amount of 0.005 to 0.11%, and more advantageously, in an amount of 0.01 to 0.1%.

The remaining component of the invention is iron (Fe). However, in the normal manufacturing process, impurities that are not intended from the raw material or the surrounding environment can be inevitably mixed, and therefore cannot be excluded. Since these impurities are known to anyone skilled in the ordinary manufacturing process, they are not specifically mentioned in this specification.

The steel material of the present invention having the above-described alloy composition may be composed of a microstructure is a composite structure of ferrite and bainite phase.

In this case, the sum of the fractions of the ferrite and bainite phases is preferably 90% or more of the area fraction, and the bainite phase may be 50% or more of the area fraction.

If the fraction of the bainite phase is less than 50% of the area fraction, it is difficult to secure the target strength, and when the coarse ferrite phase increases, it has a non-uniform microstructure, and thus, it is easy to generate fine cracks during shear deformation or punching deformation.

Here, the ferrite phase means a polygonal ferrite phase, which is a high-temperature ferrite phase, and the bainite phase collectively refers to both a needle-shaped ferrite and a bainitic ferrite phase, which are low-temperature ferrite phases.

Residual tissues other than the composite tissue may include a MA phase (a mixture of martensite and austenite) and a martensite phase. At this time, the two phases may be combined to include an area fraction of 1 to 10%, of which the MA phase is preferably less than 3%.

If, when the combined fraction of the MA phase and the martensite phase exceeds 10%, the tensile strength increases, while the hardness value is high compared to the surrounding microstructure, at the interface between the MA phase and the martensite phase during shear deformation or punching deformation. Cracking occurs, and fatigue characteristics are deteriorated accordingly. In particular, the MA phase has an average size of 1/10 of that of the martensite phase, but the tendency to crack at the interface of the phase is similar to that of the martensite phase. It is preferred.

Thus, it is possible to obtain the effect of resolving the tissue unevenness by minimizing the fraction of the coarse MA phase and martensite phase in the matrix tissue.

On the other hand, even if the steel material of the present invention contains 3% or less (including 0%) of the pearlite phase in addition to the above-described structure, there is no great difficulty in securing the intended physical properties.

Particularly, the steel material of the present invention has a shape ratio (ratio of short side length (short axis)/long side length (long axis), aspect ratio) of crystal grains of 0.3 or less in the center corresponding to the point t/4 from the thickness direction t/2. It is preferable that the fraction is less than 50%, and the length of the grain boundaries observed in the unit area (1 mm ² ) at the center is 700 mm or more.

If the proportion of crystal grains having a shape ratio of crystal grains of 0.3 or less in the central portion is 50% or more, the growth of cracks is facilitated when cracks are generated, and durability is inferior. In addition, if the length of the grain boundary in the central portion is less than 700 mm, the strength of the central portion decreases, and cracks are easily propagated, resulting in poor durability.

The method for analyzing the shape ratio of the crystal grains and the length of the grain boundaries is not particularly limited, but for example, it can be analyzed by using Back Scattered Diffraction (EBSD). Specifically, in the EBSD measurement result of the rolled section, the grain size having a large-diameter grain boundary of 15° or more is determined by the length of the grain boundary per unit area (1 mm ² ), and the shape ratio can be determined by the ratio of the shortening of the grain size and the long axis.

The steel of the present invention having the above-described alloy composition and microstructure is a thick steel material having a thickness of 5 mm or more and a maximum of 12 mm, the tensile strength is 650 MPa or more, and the ratio of fatigue limit and yield strength (fatigue limit/yield strength) is 0.25 or more. As a result, it is possible to secure excellent durability along with high strength.

Hereinafter, a method of manufacturing a high strength steel material having excellent durability, which is another aspect of the present invention, will be described in detail.

The high-strength steel according to the present invention can be produced by performing a series of processes of [heating-hot rolling-winding-cooling] a steel slab satisfying the alloy composition proposed in the present invention.

Hereinafter, each of the process conditions will be described in detail.

Steel slab heating

In the present invention, it is preferable to undergo a process of heating and homogenizing a steel slab before performing hot rolling, and it is preferable to perform a heating process at 1200 to 1350°C.

If the heating temperature is less than 1200°C, the precipitate is not sufficiently re-used, and thus the formation of the precipitate decreases in the process after hot rolling, and there is a problem that coarse TiN remains. On the other hand, when the temperature exceeds 1350°C, the strength is lowered due to abnormal grain growth of austenite grains, which is not preferable.

Hot rolled

It is preferable that the reheated steel slab is hot-rolled to produce a hot-rolled steel sheet. At this time, it is performed in a temperature range of 800 to 1150°C, and it is preferable to perform hot-finishing under conditions satisfying the following [Relational Formula 1].

When the hot rolling is performed at a temperature higher than 1150°C, the temperature of the hot rolled steel sheet becomes high, the grain size becomes coarse, and the surface quality of the hot rolled steel sheet becomes inferior. On the other hand, when hot rolling is performed at a temperature lower than 800° C., crystal grains elongated due to excessive recrystallization delay develop, anisotropy increases, moldability deteriorates, and unevenness occurs due to rolling at a temperature below the austenite temperature range. Microstructure develops more severely.

Particularly, in the hot rolling process of the present invention, when rolling is terminated at a temperature higher than the temperature range suggested in the following relation 1 (temperature above Tn), the microstructure of the steel is coarse and non-uniform, and phase transformation is delayed to coarse. Due to the formation of the MA phase and martensite phase, fine cracks are excessively formed during shear forming and knitting forming, resulting in poor durability. On the other hand, when rolling is finished at a temperature lower than the temperature range suggested by the following relational formula 1 (a temperature less than Tn-50), in thick steel materials having a thickness of 5 mm or more, the thickness direction t/4 under the surface layer where the temperature is relatively low. Ferrite phase transformation is promoted at the point, and the phase fraction of fine ferrite increases, but it has an elongated grain shape, which causes cracks to spread rapidly, and uneven microstructure can remain in the center of thickness, which is disadvantageous in securing durability. do.

[Relationship 1]

Tn-50 ≤ FDT (Hot rolling end temperature (℃)) ≤ Tn

Tn = 730 + 92×[C] + 70×[Mn] + 45×[Cr] + 650×[Nb] + 410×[Ti]-80×[Si]-1.4×(t-5) (where Each element means weight content (%), and t means thickness (mm) of the final hot rolled steel sheet)

Cooling and Winding

The hot-rolled steel sheet manufactured by performing hot rolling as described above can be cooled to a temperature range of 400 to 500° C., and then a winding process can be performed at that temperature.

The cooling is performed by primary cooling and secondary cooling, and the primary cooling is a cooling rate (CR ₁ ) that satisfies [Relational Formula 2], and a cooling rate (CR) that satisfies [Relational Formula 3] for the secondary cooling. It is preferable to perform with ₂ ).

Specifically, the primary cooling is preferably terminated in a temperature section in which phase transformation of ferrite occurs during cooling, but the temperature at which phase transformation of ferrite occurs may vary according to the alloy composition proposed in the present invention. More specifically, the primary cooling is preferably performed to a temperature at which transformation of hard phases such as bainite phase, MA phase and martensite phase does not occur. Even more preferably, the primary cooling may be performed until the temperature of the hot-rolled steel sheet obtained by hot rolling reaches 600°C.

When the thickness of the rolled sheet is 5 mm or more as in the present invention during the first cooling in the temperature section, the cooling rate in the center of the thickness of the rolled sheet is slower than the cooling rate in the area directly under the surface layer ~ t/4, so adjust it from the center of thickness. For the ferrite phase, it may have a non-uniform microstructure.

Therefore, in the present invention, it is preferable to cool at a cooling rate faster than a specific cooling rate (CR ₁ ) represented by Equation 2 below so that excessive ferrite phase is not formed during the primary cooling or the ferrite phase is not coarsened.

[Relationship 2]

CR ₁ ≥ 196-300 × [C] + 4.5 × [Si]-71.8 × [Mn]-59.6 × [Cr] + 187 × [Ti] + 852 × [Nb] (where each element has a weight content (% Means ))

It is preferable to perform the secondary cooling immediately after the primary cooling is completed under the above-described conditions, and the secondary cooling is preferably terminated at a coiling temperature (CT (° C.)).

In the second cooling in the temperature section, the untransformed phase is transformed into the bainite phase over the entire thickness of the steel material, so that 90% (area fraction) of the matrix structure is formed into the ferrite and bainite phases. It is preferable to perform cooling at a specific cooling rate (CR ₂ ). At this time, if the cooling rate is slower than CR _Min , carbides are formed rather than the bainite phase to grow coarse, which mainly exists at the grain boundary of the ferrite phase, and when the cooling rate is slower, a pearlite phase is formed to crack during shear molding or punching molding. It is easy to form and there is a problem that cracks propagate along grain boundaries even with a small external force. On the other hand, when the cooling rate exceeds CR _Max , the MA phase or the martensite phase, which increases the difference in hardness between phases, is excessively formed, resulting in poor durability.

Therefore, it is necessary to perform cooling at a cooling rate that satisfies the following relational expression (3) during secondary cooling in the temperature section.

[Relationship 3]

CR _Min ≤ CR ₂ ≤ CR _Max

(CR _Max = 76.6-157 × [C]-25.2 × [Si]-14.1 × [Mn]-27.3 × [Cr] + 61 × [Ti] + 448 × [Nb], CR _Min = 27.4-45.3 × [C] + 5.28×[Si]-11×[Mn]-7.33×[Cr] + 42.3×[Ti] + 82×[Nb], each element means weight content (%))

On the other hand, when the coiling temperature exceeds 500°C when winding after completing the above-described cooling process, the pearlite phase is formed and the strength of the steel becomes insufficient, whereas when it is less than 400°C, the martensite phase is excessively formed, resulting in shear formability and punching formability. Durability is inferior.

In the present invention, as the process conditions are controlled to satisfy the above-described relational expressions 1 to 3 in manufacturing the intended steel material, the area ratio of the grains having an aspect ratio of 0.3 or less in the central portion of the thickness of the steel material is less than 50% While securing it, the length of the grain boundary observed within the unit area (1 mm ² ) can be secured to 700 mm or more.

When producing thick steel materials having a thickness of 5 mm or more, it is difficult to uniformly secure the microstructure in the center of the thickness when it is performed by ordinary hot rolling. In particular, when hot rolling is performed at an excessively low temperature in order to obtain a retarding effect of recrystallization in the center of thickness, the deformed structure develops strongly in the region up to t/4 directly below the surface layer in the thickness direction of the rolled plate, and rather, the phase with the center of thickness Due to the increase in non-uniformity, fine cracks are likely to occur in the non-uniform part during shear deformation or punching deformation, and the durability of parts is also poor. Therefore, as shown in the relational expression 1 above, it is necessary to complete the hot rolling between Tn (°C) temperature and Tn-50 (°C), which is the temperature at which the recrystallization delay starts.

When hot rolling is terminated at a temperature at which the temperature suggested in the relational expression 1 is high, a coarse ferrite phase and a polygonal ferrite phase are formed, so that the area fraction of the crystal grains having a shape ratio of 0.3 or less is greatly reduced, while the size of the grain boundaries is also significantly reduced, and the strength of the core is significantly reduced. There is a fear that the fall, there is a problem that facilitates the growth of the crack when the crack is formed. In addition, when hot rolling is finished at a temperature lower than the temperature suggested in the relational expression 1, the severely stretched crystal grains increase, and the area fraction of grains having a formation ratio of 0.3 or less increases significantly, and the formation of coarse carbide or martensite phases on grain boundaries As a result, cracks formed during shear molding are easily propagated by external force, resulting in poor durability.

In addition, the relational expressions 2 and 3 correspond to cooling conditions for optimizing the microstructure so as to improve the strength and durability of the steel through a phase transformation process during cooling. That is, since not only the type and fraction of the tissue phase, but also the shape ratio of the crystal grains and the size of the grain boundaries vary depending on the cooling conditions, it will be preferable to perform cooling under the conditions satisfying the above relational expressions 2 and 3.

Air cooling

The coil obtained by completing the cooling and coiling process according to the above can be air-cooled to a temperature range of room temperature to 200°C. At this time, the air cooling process of the coil has a cooling rate of 0.001 to 10°C/hour, which means cooling in the air. At this time, when the cooling rate exceeds 10°C/hour, some untransformed phases of the steel are easily transformed into the MA phase, thereby deteriorating the shear formability and punching formability and durability of the steel. On the other hand, in order to control the cooling rate to less than 0.001°C/hour, separate heating and heat preservation facilities are required, which is economically disadvantageous.

On the other hand, as described above, the air-cooled steel can be pickled and oiled, and then heated to a temperature range of 450 to 740°C to perform a hot dip galvanizing process.

The hot dip galvanizing process may use a zinc-based plating bath, the alloy composition in the zinc-based plating bath is not particularly limited, for example, magnesium (Mg): 0.01 to 30% by weight, aluminum (Al): 0.01 It may be a plating bath containing ˜50% by weight and the balance Zn and unavoidable impurities.

Hereinafter, the present invention will be described in more detail through examples. However, it is necessary to note that the following examples are only intended to illustrate the present invention in more detail and are not intended to limit the scope of the present invention. This is because the scope of the present invention is determined by the items described in the claims and the items reasonably inferred therefrom.

( Example )

A steel slab having an alloy composition of Table 1 below was prepared. At this time, the content of the alloy composition is weight%, and the rest includes Fe and unavoidable impurities. Prepared steel slabs were prepared according to the manufacturing conditions in Table 2, respectively. At this time, when cooling after hot rolling, the primary cooling was completed at 600°C, and the secondary cooling was completed at the coiling temperature.

In Table 2, FDT is the temperature at the time of finishing hot rolling (the end temperature of hot rolling), CT is the winding temperature, and after completion of winding, the cooling rate at air cooling was constant at 1°C/hour.

Steel Alloy composition (% by weight) C Si Mn Cr Al P S N Ti Nb Comparative steel 1 0.06 0.3 1.8 0.2 0.03 0.01 0.004 0.004 0.005 0.02 Comparative steel 2 0.06 0.3 1.8 0.2 0.03 0.008 0.004 0.004 0.005 0.02 Comparative steel 3 0.07 0.04 1.7 0.6 0.03 0.01 0.005 0.004 0.05 0.005 Comparative River 4 0.06 0.5 2.0 0.007 0.03 0.01 0.004 0.005 0.04 0.03 Comparative steel 5 0.06 0.5 2.0 0.007 0.03 0.005 0.004 0.005 0.04 0.03 Comparative steel 6 0.07 0.5 1.6 0.008 0.03 0.01 0.003 0.004 0.08 0.03 Comparative River 7 0.07 0.5 1.6 0.008 0.03 0.01 0.003 0.004 0.08 0.03 Comparative Steel 8 0.07 0.4 2.2 0.012 0.03 0.007 0.004 0.004 0.1 0.02 Comparative Steel 9 0.07 0.4 2.2 0.012 0.03 0.01 0.004 0.004 0.1 0.02 Comparative River 10 0.08 0.4 1.6 0.8 0.05 0.01 0.003 0.006 0.04 0.045 Comparative River 11 0.08 0.4 1.6 0.8 0.05 0.01 0.003 0.006 0.04 0.045 Comparative steel 12 0.04 0.5 1.8 0.3 0.03 0.01 0.002 0.004 0.065 0.03 Comparative steel 13 0.16 0.55 1.6 0.2 0.03 0.01 0.003 0.004 0.07 0.03 Comparative steel 14 0.08 1.2 2.0 0.3 0.03 0.01 0.003 0.004 0.06 0.025 Comparative steel 15 0.08 0.5 0.8 0.8 0.03 0.01 0.003 0.004 0.05 0.035 Comparative River 16 0.07 0.5 2.5 0.01 0.03 0.01 0.003 0.004 0.07 0.03 Comparative Steel 17 0.08 0.5 1.7 1.1 0.03 0.01 0.004 0.004 0.05 0.03 Invention Steel 1 0.06 0.05 1.5 0.05 0.03 0.005 0.003 0.005 0.095 0.03 Invention Steel 2 0.06 0.3 1.2 0.9 0.03 0.01 0.003 0.005 0.04 0.04 Invention Steel 3 0.08 0.5 1.7 0.5 0.03 0.01 0.003 0.005 0.06 0.05 Invention Steel 4 0.07 0.3 1.6 0.8 0.03 0.008 0.003 0.005 0.07 0.06 Invention Steel 5 0.09 0.3 1.6 0.9 0.03 0.01 0.002 0.004 0.07 0.04 Invention Steel 6 0.09 0.1 1.85 0.8 0.03 0.01 0.003 0.004 0.05 0.04 Invention Steel 7 0.11 0.5 1.95 0.7 0.03 0.01 0.003 0.004 0.06 0.045

(Comparative steels 1 to 11 in Table 1, but the alloy composition satisfies the scope of the present invention, but in Table 2 below, the manufacturing conditions are denoted as comparative steel according to the present invention.)

Steel
thickness
(mm)
FDT
(℃) CT
(℃) CR ₁
(℃/s)
CR ₂
(℃/s) Relation 1 Relation 2 Relation 3 Tn Satisfaction
Whether CR ₁ Satisfaction
Whether CR _Max CR _Min Satisfaction
Whether Comparative steel 1 2.9 890 455 78 30 865 × 56.2 ○ 38.0 6.9 ○ Comparative steel 2 10 885 460 65 25 855 × 56.2 ○ 38.0 6.9 ○ Comparative steel 3 7 835 470 72 22 900 × 31.0 ○ 29.5 3.9 ○ Comparative River 4 3.2 870 443 54 38 874 ○ 69.3 × 42.1 9.4 ○ Comparative steel 5 8 858 485 51 25 868 ○ 69.3 × 42.1 9.4 ○ Comparative steel 6 3.3 906 475 88 35 863 × 102.4 × 48.6 15.1 ○ Comparative River 7 9 895 430 80 42 856 × 102.4 × 48.6 15.1 ○ Comparative Steel 8 3.8 850 450 35 30 915 × 53.9 × 39.2 7.9 ○ Comparative Steel 9 8 845 448 33 35 909 × 53.9 × 39.2 7.9 ○ Comparative River 10 9 880 450 78 75 893 ○ 57.1 ○ 32.2 7.8 × Comparative River 11 9 875 465 85 3 893 ○ 57.1 ○ 32.2 7.8 × Comparative steel 12 8 840 450 120 35 875 ○ 76.8 ○ 41.6 11.4 × Comparative steel 13 8 848 450 95 22 866 ○ 62.3 ○ 27.3 9.4 ○ Comparative steel 14 8 820 455 105 12 832 ○ 48.4 ○ 12.3 10.5 × Comparative steel 15 8 816 450 125 28 828 ○ 108.3 ○ 37.1 16.7 ○ Comparative River 16 8 910 445 80 20 916 ○ 35.8 ○ 35.2 4.7 ○ Comparative Steel 17 8 904 462 75 9 902 ○ 21.5 ○ 13.9 4.2 ○ Invention Steel 1 8 855 442 120 42 893 ○ 110.9 ○ 62.6 14.6 ○ Invention Steel 2 7 850 450 125 31 876 ○ 81.1 ○ 38.5 11.4 ○ Invention Steel 3 9 870 443 98 18 890 ○ 76.2 ○ 39.9 10.7 ○ Invention Steel 4 8 880 455 85 20 924 ○ 78.0 ○ 44.8 10.2 ○ Invention Steel 5 9 872 466 105 28 916 ○ 49.0 ○ 30.0 7.0 ○ Invention Steel 6 10 890 452 89 25 935 ○ 32.4 ○ 33.0 3.0 ○ Invention Steel 7 11 880 440 102 18 914 ○ 33.1 ○ 23.9 4.7 ○

For each steel sheet prepared according to the above, the mechanical properties and durability of tensile strength (TS), yield strength (YS), and elongation (T-El) were evaluated, and microstructure was observed and the results were as follows. It is shown in Table 3.

Specifically, the yield strength and the elongation each indicate 0.2% off-set yield strength and the elongation at break, and the tensile strength was measured by taking specimens of JIS 5 standard specimens in a direction perpendicular to the rolling direction.

The durability was evaluated by performing a high-cycle fatigue test (bending fatigue test) on the test piece having the punching molded part. At this time, the test piece for the fatigue test was produced by punching a hole of 10 mm in diameter with a clearance of 12% by punching molding in the center of a bending fatigue test piece with a length of 40 mm and a width of 20 mm, and tested under a stress ratio of -1 and a frequency of 15 Hz. The fatigue strength (S _Fatigue ) was expressed as the strength ratio (S _Fatigue /YS) compared to the yield strength, from which it was possible to confirm the change in cross-section quality and durability of the punched area.

On the other hand, the microstructure of each steel material showed the results observed from the central portion (t/2) in the thickness direction. The aspect ratio (AR), which is the ratio of the length of the grain boundaries per unit area (1 mm ² ) corresponding to the area of the grain boundaries, and the aspect ratio (AR) of the grain boundaries, are applied to back-scattered electron diffraction (EBSD) for grains having a large-angle grain boundary of 15° or more. It was measured using. After the area fraction on the MA was etched by the Lepera etching method, the results of analysis at 1000 magnification using an optical microscope and an image analyzer were shown. In addition, the phase fractions of martensite (M), ferrite (F), bainite (B), and pearlite (P) were measured from the results analyzed at 3000 magnification and 5000 magnification using an electron scanning microscope (SEM).

In Table 3 below, F denotes a polygonal ferrite having an equiaxed crystal shape, and B denotes a sum of all the fractions of the ferrite phase observed in a low temperature region such as a bainite phase, acicular ferrite, and bainitic ferrite.

In addition, in Table 3, AR0.3 represents a ratio (area fraction) of crystal grains having an aspect ratio of 0.3 or less, and shows results obtained by observing at 1000 magnification.

Steel Mechanical properties Microstructure durability YS
(MPa) TS
(MPa) T-El
(%) Grain boundary
Length (mm) AR0.3
(%) F
(%) B
(%) P
(%) M
(%) MA
(%) S _Fatigue
(MPa) S _Fatigue
/YS Comparative steel 1 518 632 19 855 56 54 43 0 One 2 169 0.33 Comparative steel 2 470 565 27 608 35 47 45 2 One 5 113 0.24 Comparative steel 3 539 651 26 785 53 66 30 One 2 One 128 0.24 Comparative River 4 632 775 17 880 58 43 53 0 3 One 198 0.31 Comparative steel 5 510 623 18 730 45 59 32 4 One 4 121 0.24 Comparative steel 6 560 686 17 826 48 41 54 One 2 2 172 0.31 Comparative River 7 538 652 26 688 29 45 46 4 One 4 129 0.24 Comparative Steel 8 715 872 14 1028 62 61 30 2 5 2 204 0.29 Comparative Steel 9 602 738 22 840 39 67 22 6 One 4 130 0.22 Comparative River 10 625 764 22 963 31 40 48 0 8 4 142 0.23 Comparative River 11 587 708 24 765 40 56 38 4 One One 136 0.23 Comparative steel 12 499 604 26 670 45 70 28 0 One One 117 0.23 Comparative steel 13 762 915 16 966 41 29 50 8 9 4 180 0.24 Comparative steel 14 568 690 25 799 67 36 56 One One 6 137 0.24 Comparative steel 15 504 611 27 882 72 59 38 0 One 2 123 0.24 Comparative River 16 715 866 19 1036 28 13 72 2 8 5 168 0.23 Comparative Steel 17 721 882 20 1015 35 17 76 0 6 One 172 0.24 Invention Steel 1 554 675 25 750 25 47 51 0 One One 169 0.31 Invention Steel 2 601 724 22 882 30 35 62 0 One 2 187 0.31 Invention Steel 3 674 816 22 977 28 30 65 0 4 One 193 0.29 Invention Steel 4 690 821 21 1014 30 32 64 0 3 One 195 0.28 Invention Steel 5 762 943 19 1020 35 28 67 0 3 2 213 0.28 Invention Steel 6 771 924 18 1155 38 20 76 One 2 One 208 0.27 Invention Steel 7 780 955 17 1084 42 9 85 One 3 2 220 0.28

As shown in Tables 1 to 3, the inventive steels 1 to 7 satisfying both the alloy composition and the manufacturing conditions proposed in the present invention were formed of a ferrite and bainite composite structure. In addition, the fraction of the crystal grains having a grain shape ratio in the center of the thickness direction of the steel material of 0.3 or less was less than 50% (see FIG. 2), and since the grain boundary lengths were all formed to be 700 mm or more, the intended high strength and excellent durability were secured. .

On the other hand, comparative steels 1 to 11 are satisfactory in the alloy composition proposed in the present invention, but the manufacturing conditions are outside the present invention, and the intended physical properties could not be secured.

Comparative steels 1 to 3 are cases in which the hot-rolled finishing temperature does not satisfy the relational expression 1 proposed in the present invention, and the comparative steel 1 has a final steel thickness of 2.9 mm, and the ferrite phase stretched from the center is excessively formed, but fatigue The characteristics were not significantly inferior. This is due to the fact that, when hot-rolled to a thickness of 2.9 mm, the amount of rolling in the unrecrystallized temperature range greatly increased, and elongated microstructure developed, but the quality of the cross-section of the punched area was good as the microstructure in the thickness direction was uniform. Was judged. On the other hand, Comparative Steels 2 and 3 are thick steels with thicknesses of 10 mm and 7 mm, respectively, and Comparative Steel 2 has micro-cracks formed on the cross section when exposed to a fatigue environment as the MA phase develops in the central microstructure and the grain boundary length is less than 700 mm. It has been shown to grow easily and have poor fatigue properties. In addition, in Comparative Steel 3, due to hot rolling in a low-temperature region, excessively formed crystal grains stretched in the center of the thickness, and it was judged that fatigue fracture occurred along the weak grain boundaries. That is, it is due to the development of fine cracks in the center of the thickness during punching molding along the drawn ferrite grain boundaries.

Comparative steels 4 and 5 have the same components, and the condition of primary cooling when cooling after hot rolling does not satisfy relational expression 2, comparative steel 4 has a thickness of 3.2 mm, and comparative steel 5 has a thickness of 8 mm. To have. Among them, comparative steel 4 having a thickness of less than 5 mm formed many elongated grains similar to comparative steel 1, but even when the cooling rate was slow during the primary cooling, coarse carbides were hardly formed in the grain boundaries, so fatigue characteristics were not significantly inferior. . On the other hand, the comparative steel 5 having a thick thickness has a slow cooling rate during the first cooling, so that pearlite is formed in the center of the thickness, the fraction of ferrite phase is somewhat excessive, and the MA phase is also observed in the crystal grains, indicating that the fatigue properties are deteriorated. .

Comparative steels 6 and 7 have the same components as each other, but have thicknesses of 3.3 mm and 9 mm, respectively, and are not satisfied with both relations 1 and 2. Comparative steel 6 is a thin material, and it is judged that the effect of delaying recrystallization can be secured in the entire thickness even when the hot rolling temperature is high. It was good. On the other hand, the thick comparative steel 7 has a large microstructure due to a high rolling temperature and a slow cooling rate during primary cooling, and a grain boundary length of less than 700 mm is formed, and the MA phase and pearlite phase are also formed, resulting in poor fatigue properties.

Comparative steels 8 and 9 are cases in which the finishing temperature at the time of hot pressure is lower than the range suggested by the present invention, and the cooling rate at the time of primary cooling is slow. These also have the same components but different thicknesses. As for the comparative steel 8, which is a thin material, many fine and elongated ferrite phases were formed over the entire thickness, but the fatigue properties were not inferior, whereas the comparative steel 9, a thick material, was MA at the center of the thickness. The phase and pearlite phase were excessively formed, resulting in poor fatigue properties.

Comparative steel 10 is a relational formula 3, that is, when the cooling rate during secondary cooling is outside the present invention, the cooling rate during secondary cooling is too fast, so that the martensite phase is excessively formed in the center of the thickness, and thus, when exposed to a fatigue environment, the surrounding phase (phase It was judged that the fracture proceeded easily in the region where the hardness difference with) was large.

Comparative steel 11 also does not satisfy the relational expression 3, and the cooling rate is too slow during the secondary cooling, resulting in excessive pearlite phase formation and inferior fatigue properties.

On the other hand, the comparative steel 12 to 17 is the case where the alloy composition is out of the present invention, all of the relational formulas 1 to 3 are satisfied at the time of manufacture, and they are all manufactured to have the same thickness (8 mm), but the fatigue properties are inferior.

Specifically, the comparative steel 12 is a case where the C content is insufficient, the ferrite phase is excessively formed in the center of the thickness, and the bainite phase is not sufficiently formed. Due to this, the microstructure was coarse and the fatigue strength was low.

Comparative steel 13 is a case in which the C content is excessively added, and the pearlite and martensite phases are excessively formed due to the high C concentration in the untransformed phase during the phase transformation process, thereby exhibiting lower fatigue strength than the yield strength.

Comparative steel 14 is a case where the Si content is too high, and the MA phase is formed together with the bainite phase, and a lot of stretched microstructures are observed. Due to this, the fatigue properties were inferior, which is thought to be due to the formation of many cracks around the relatively hard phase MA phase.

Comparative steel 15 is a case where the Mn content is insufficient, and despite the fact that it was prepared by satisfying the relations 1 to 3 to obtain a recrystallization delay effect and a uniform microstructure, the ferrite phase is excessively formed in the center of the thickness, resulting in strength and fatigue strength All appeared low.

In Comparative Steel 16, when Mn content was excessively added, the martensite phase was excessively developed along the Mn segregation zone developed at the center of the thickness, resulting in poor cross-section quality and fatigue characteristics.

In addition, in Comparative Steel 17, when the Cr content was excessive, similarly to Comparative Steel 16 above, a lot of martensite phases formed locally at the center of the thickness were observed, and thus the fatigue properties were inferior.

Claims (10)

In weight percent, carbon (C): 0.05 to 0.15%, silicon (Si): 0.01 to 1.0%, manganese (Mn): 1.0 to 2.3%, aluminum (Al): 0.01 to 0.1%, chromium (Cr): 0.005 ~1.0%, phosphorus (P): 0.001 to 0.05%, sulfur (S): 0.001 to 0.01%, nitrogen (N): 0.001 to 0.01%, niobium (Nb): 0.005 to 0.07%, titanium (Ti): 0.005 ~0.11%, contains the balance Fe and other unavoidable impurities,
As a microstructure, the sum of fractions of ferrite and bainite phase is 90% or more,
The fraction of crystal grains with a shape ratio (ratio of short side/long side) of grains in the center (thickness t/4 to t/2) of 0.3 or less is less than 50%, and the grain boundary length observed within a unit area (1 cm ² ) in the center High-strength steel with excellent durability over 700mm.
According to claim 1,
The steel sheet is a high strength steel material having excellent durability with a fraction of the MA phase (martensite and austenite mixed structure) of less than 3%.
According to claim 1,
The steel sheet is a high-strength steel material having excellent durability in which the sum of the MA phase (a mixture of martensite and austenite) and the martensite phase is 1-10% by area.
According to claim 1,
The steel sheet is a high-strength steel material having excellent tensile strength of 650 MPa or more, and a ratio of fatigue limit and yield strength (fatigue limit/yield strength) of 0.25 or more.
In weight percent, carbon (C): 0.05 to 0.15%, silicon (Si): 0.01 to 1.0%, manganese (Mn): 1.0 to 2.3%, aluminum (Al): 0.01 to 0.1%, chromium (Cr): 0.005 ~1.0%, phosphorus (P): 0.001 to 0.05%, sulfur (S): 0.001 to 0.01%, nitrogen (N): 0.001 to 0.01%, niobium (Nb): 0.005 to 0.07%, titanium (Ti): 0.005 Heating a steel slab containing ˜0.11%, balance Fe and other inevitable impurities in a temperature range of 1200 to 1350° C.;
Hot-rolling the heated steel slab to produce a hot-rolled steel sheet;
Cooling the hot-rolled steel sheet to a temperature range of 400 ~ 500 ℃ and then winding (CT); And
After the winding, including the step of air cooling to a temperature range of room temperature to 200 ℃,
The hot rolling is finished hot rolling at a temperature (FDT (°C)) satisfying the following [Relational Formula 1],
The cooling is performed by the following primary cooling and secondary cooling, and the primary cooling is a cooling rate (CR ₁ ) that satisfies [Relational Formula 2], and a cooling rate (which satisfies [Relational Formula 3]) that is performed by the secondary cooling. CR ₂ ) method for producing high strength steel with excellent durability.

[Relationship 1]
Tn-50 ≤ FDT (Hot rolling end temperature (℃)) ≤ Tn
Tn = 730 + 92×[C] + 70×[Mn] + 45×[Cr] + 650×[Nb] + 410×[Ti]-80×[Si]-1.4×(t-5) (where Each element means weight content (%), and t means thickness (mm) of the final hot rolled steel sheet)

[Relationship 2]
CR ₁ ≥ 196-300 × [C] + 4.5 × [Si]-71.8 × [Mn]-59.6 × [Cr] + 187 × [Ti] + 852 × [Nb] (where each element has a weight content (% Means)

[Relationship 3]
CR _Min ≤ CR ₂ ≤ CR _Max
(CR _Max = 76.6-157 × [C]-25.2 × [Si]-14.1 × [Mn]-27.3 × [Cr] + 61 × [Ti] + 448 × [Nb], CR _Min = 27.4-45.3 × [C] + 5.28×[Si]-11×[Mn]-7.33×[Cr] + 42.3×[Ti] + 82×[Nb], each element means weight content (%))
The method of claim 5,
The primary cooling is a method of manufacturing a high-strength steel having excellent durability, which is terminated at 600°C.
The method of claim 5,
The secondary cooling is a method of manufacturing a high-strength steel having excellent durability, which is terminated at a coiling temperature (CT (°C)).
The method of claim 5,
Method of manufacturing a high-strength steel excellent in durability further comprising the step of pickling and oiling the steel sheet after the cooling.
The method of claim 8,
A method of manufacturing a high strength steel material having excellent durability, further comprising heating the steel sheet after the pickling and oiling to a temperature range of 450 to 740°C, and then hot-dip galvanizing.
The method of claim 9,
The hot-dip galvanizing method using a plating bath containing magnesium (Mg): 0.01 to 30% by weight, aluminum (Al): 0.01 to 50%, and residual Zn and inevitable impurities.

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