JP4854333B2 - High strength steel plate, unannealed high strength steel plate and method for producing them - Google Patents

Description

  The invention according to the claims relates to a high-strength steel sheet having high workability while having high tensile strength, an unannealed high-strength steel sheet as an intermediate product therefor, and a method for producing them.

  The recent demand for high-strength steel sheets with excellent workability will be described by taking the case of automobiles as an example. From the viewpoint of global environmental conservation, it is essential to reduce the amount of exhaust gas such as CO2 in the automobile field. For that purpose, further weight reduction of the automobile body becomes indispensable. In order to reduce the weight of the car body, it is necessary to increase the strength of steel plates used in automobiles and reduce the thickness. At the same time, passengers must ensure the safety of passengers. For this purpose, it is necessary to further increase the strength of the steel sheet.

  As a method for increasing the strength, solid solution strengthening, precipitation strengthening, crystal grain refinement and the like are basic methods. In order to ensure weldability while saving resources and reducing manufacturing costs, it is extremely important to reduce the amount of alloy addition as much as possible. Therefore, it is impossible to produce a steel sheet that requires extremely high strength only by applying a strengthening mechanism that requires a large amount of alloy addition such as solid solution strengthening or precipitation strengthening. Even if a strengthening mechanism by grain refinement is applied, the crystal grain size that can be produced commercially is 1 to 3 μm, so even if the strength can be increased to some extent, the ductility can be improved to meet the increase in strength. Cannot be expected.

When the strength of the steel plate increases, the workability deteriorates, and at the same time, the dimensional accuracy also deteriorates due to the spring back after press forming. It is difficult to apply high-strength steel sheets by cold working methods such as ordinary press forming.
In the hot pressing method, since the hot pressing is performed, the amount of spring back is extremely small and the shape freezing property is good. And the part with very high intensity | strength can be provided with high precision by the hardening effect in the case of a press. However, it is necessary to heat the steel plate before pressing, and an operation to drop the scale is necessary after pressing. Therefore, the work efficiency is very poor. Furthermore, since the mold contacts the heated steel plate, the short life of the mold is also a drawback, which increases the manufacturing cost.
The steel sheet after hot pressing has a small elongation value, and when the member is deformed, it breaks even if it is slightly deformed. Therefore, it is evaluated that the impact absorbing ability is small. Therefore, it is very difficult to use a hot press part as an important safety part for an automobile or the like.

  One method of improving the dimensional accuracy of a molded product by cold pressing is a method of suppressing the occurrence of springback by providing a bead on the molded product. However, this processing method requires extremely high workability for the steel sheet.

Generally, when the strength of a steel plate is increased, ductility is reduced and workability is reduced. Conventional techniques for increasing the ductility of high-strength steel sheets are called composite phase (Dual Phase) steel sheets composed of ferrite and martensite structures, and TRIP (Transformation Induced Plasticity) steel sheets composed of ferrite and retained austenite structures.
Examples of composite structure steel sheets are shown in Patent Documents 1 and 2. In such a steel sheet, hard martensite is finely dispersed in the ferrite. However, this hard martensite causes a large work hardening at the time of deformation and brings high ductility to the steel sheet.
Examples of the TRIP steel sheet are shown in Patent Documents 3 and 4. This type of steel sheet containing retained austenite has very good ductility and formability due to transformation induced plasticity, depending on its amount and stability to deformation.
JP-A-56-133423 JP 2005-230896 JP JP 62-182225 A JP 2005-76078

  As a conventional technique for improving the elongation characteristics in cold working while having high strength, the composite structure steel plate and the TRIP steel plate can be mentioned.

  With a composite structure steel plate, high strength can be obtained even with a relatively low alloy addition amount, and at the same time, good uniform elongation characteristics can be obtained by work hardening. Among them, the steel sheets of Patent Document 1 and Patent Document 2 are further improved in strength and elongation characteristics by finely dispersing martensite or bainite in the ferrite matrix. However, in both documents, the amount of retained austenite is scarce or the amount is insufficiently controlled, so the elongation properties of TRIP steel sheets are not obtained, and the elongation properties are high enough to withstand the processing of parts with complex shapes. Not obtained.

The TRIP steel plate proposed in Patent Document 3 and the TRIP steel plate described in Patent Document 4 in which retained austenite is finely dispersed have high strength, high ductility, and high deep drawability. is there. Therefore, application to a member that requires high workability with a complicated shape and requires high strength is directed.
However, in the annealing cooling process, in order to stabilize austenite, it is necessary to maintain a low temperature near 400 ° C. just above the Ms point at which martensitic transformation starts. Furthermore, due to the low temperature retention, there is a decisive disadvantage that the steel sheet strength is low for the carbon content and the alloy addition amount compared to the composite structure steel plate. The application of TRIP steel sheets was expected in the strength range of 780MPa or more where good elongation characteristics of the steel sheets were difficult to obtain. However, if the same strength level is to be ensured, the required amount of carbon and the amount of alloy addition increase, so that good spot bondability cannot be ensured. For this reason, this steel sheet is not widely used at present.

  Accordingly, the inventors of the present application have studied a manufacturing process in which cooling is performed at around 400 ° C. without cooling in the vicinity of 400 ° C. in the cooling process of final annealing in order to obtain high steel sheet strength even with a low alloy. Through such a process, we have developed a new low-alloy / high-strength steel sheet that has both high strength and good workability, and its manufacturing method.

  As a result of diligent research, the inventors have found that a preferable high-strength steel sheet can be obtained by adopting an appropriate component composition and devising a structure before final annealing. That is, even if it is a low alloy comparable to a composite structure steel plate, high strength and excellent ductility can be simultaneously imparted to the steel plate. Details are shown below.

The high-strength steel sheet described in claim 1 is a lath austenite having a volume ratio of 3% or more and a carbon concentration of 0.9% or more by mass% in the grains and boundaries of ferrite having an average particle diameter of 10 μm or less, and a volume ratio. The high-strength steel sheet is characterized by the presence of 10% or more of lath or granular martensite having an average particle diameter of 10 μm or less.
Since martensite is present in a volume ratio of 10% or more, it has high strength even if the amount of alloy addition is small. Residual austenite has a carbon concentration of 0.9% by mass or more, and therefore can exist metastable at room temperature, and can induce transformation due to deformation. Furthermore, since it is effectively finely dispersed at a volume ratio of 3% or more, good work hardening can be continuously obtained and good elongation characteristics can be obtained. Since both retained austenite and martensite are present in lath or fine granular form, this steel sheet is also excellent in ductility and stretch flangeability. Since the retained austenite is finely dispersed, there is an advantage that delayed fracture does not occur.
As used herein, the lath shape refers to an austenite or martensite crystal grain having an aspect ratio of 3 or more, and less than 3 is referred to as granular.

As for the above-mentioned high-strength steel sheet, in particular, C: 0.05 to 0.25%, Si: 0.01 to 1.50%, Mn: 0.6 to 2.2%, Al: 0.01 to 1.50% in mass%, or Ti: 0.02 to It is preferable that one or both of 0.22% and Nb: 0.02 to 0.10% are contained.
Furthermore, you may contain at least 1 sort (s) selected from Cu: 0.01-1.0%, Ni: 0.01-1.0%, Cr: 0.01-1.0%, Mo: 0.01-1.0%. In each case, the balance is Fe and inevitable impurities.
If such an appropriate kind and amount of chemical components are included, it is easy to obtain a high-strength steel sheet having the above-described structure and exhibiting desirable mechanical properties. In addition, the effect | action of each component is mentioned later.

When the high-strength steel sheet contains a large amount of elements that enhance hardenability, preferable characteristics may not be obtained. The hardenability of the steel sheet was experimentally derived with the components contained. The hardenability index was arranged by the following formula based on the mass% of each component element of Mn, Cr, P, Mo, B as Mn equivalent.
Mn equivalent = Mn + 1.4Cr + 3P + 1.5Mo + 180B (1)
The high strength steel plate preferably has an experimentally derived Mn equivalent of 2.35 or less.
When the Mn equivalent is suppressed to 2.35 or less, the ratio of transformation of austenite having a high carbon concentration to martensite is lowered, the amount of retained austenite is increased, and ductility is ensured, thereby improving material properties.

The high-strength steel sheet preferably has the above-described structure and has a product TS × EL of tensile strength TS (MPa) and elongation value EL (%) of 20000 (MPa ·%) or more.
This is because such a steel sheet has the above-described structure and has both high strength and good elongation characteristics.

  The unannealed high-strength steel sheet described in the claims becomes a high-strength steel sheet as described above through annealing and subsequent cooling. That is, such a steel sheet becomes a high-strength steel sheet having the above-described characteristics by subsequent appropriate annealing and cooling, and exhibits high strength and elongation characteristics.

As an unannealed high-strength steel sheet, in mass%, C: 0.05 to 0.25%, Si: 0.01 to 1.50%, Mn: 0.6 to 2.2%, Al: 0.01 to 1.50%, or Ti: 0.02 to 0.22%, One or both of Nb: 0.02 to 0.10%, or Cu: 0.01 to 1.0%, Ni: 0.01 to 1.0%, Cr: 0.01 to 1.0%, Mo: 0.01 to 1.0% Examples thereof include those having a lower bainite structure in which a prior austenite particle size is 20 μm or less and cementite having a particle size of 1 μm or less is finely dispersed in a chain train. Such an unannealed high-strength steel sheet becomes a high-strength steel sheet according to the above claims by annealing and cooling as described later, for example. Here, the chain form means a state in which two or more cementite particles are formed in a line.
The hot-rolled sheet having coarse austenite grains cannot improve the elongation characteristics even if a certain amount of retained austenite is obtained after the final annealing. In addition, if martensite and retained austenite are unevenly distributed, good work-hardening characteristics cannot be obtained. In that respect, if the prior austenite grain size is 20 μm or less as described above and the structure mainly comprises lower bainite, in addition to making the ferrite grain size obtained by subsequent heat treatment fine, martensite and residual austenite Can be finely dispersed. Since this steel also has an appropriate amount of the above-mentioned elements, this steel sheet can be made into the above-described high-strength steel sheet by appropriate heat treatment. A mechanism for producing a high-strength steel sheet having the above structure from a lower bainite structure in which cementite is finely dispersed in a chain train in the grain will be described later (see FIG. 5). The former austenite grain size referred to here refers to the lower bainite grain size.

As a similar unannealed high-strength steel sheet, a steel having a lower bainite structure in which the Mn equivalent is 2.35 or less, the prior austenite grain size is 20 μm or less, and cementite is finely dispersed in a chain form in the grain is also suitable. .
Since the Mn equivalent is 2.35 or less, the ratio of austenite transforming to martensite is reduced, and the amount of retained austenite is increased. As a result, the material properties can be improved.

The method for producing an unannealed high-strength steel sheet according to the claim is:
1) By mass% C: 0.05-0.25%, Si: 0.01-1.50%, Mn: 0.6-2.2%, Al: 0.01-1.50%, or Ti: 0.02-0.22%, Nb: 0.02-0.10% Any one or both of Cu: 0.01 to 1.0%, Ni: 0.01 to 1.0%, Cr: 0.01 to 1.0%, Mo: 0.01 to 1.0% , A steel material having an Mn equivalent of 2.35 or less,
2) With a hot rolling mill having multiple stands, hot rolling is performed so that the accumulated strain is 0.4 or more, or the rolling reduction in the final stand to be used is 15% or more.
3) Cooling is started from a temperature of Ar3 or higher at the end of rolling, and winding is performed in a temperature range of 150 ° C or higher and 400 ° C or lower.
According to the inventors' production test, as described later, the above-described unannealed high-strength steel sheet could be obtained under these conditions.

The manufacturing method of the high-strength steel sheet according to the claim is the unannealed high-strength steel sheet obtained in 1) to 3) above.
4) Heat to Ac1 or higher and Ac3 + 20 ° C or lower for annealing.
5) It is characterized by continuous cooling to 350 ° C or less at a rate (ΔCR) of 10 ° C / sec or more.
If it does in this way, said high-strength steel plate can be obtained as mentioned later. A mechanism for producing the steel sheet having the above structure by such a process will be described later (see FIG. 5).
In addition, after performing cold rolling to the said unannealed high strength steel plate, the said high strength steel plate can also be obtained similarly by performing appropriate annealing and cooling of 4) and 5).

The cooling rate ΔCR in the above 5) is between the above-mentioned Mn equivalents,
Mn equivalent ≤ -0.005 (ΔCR) + 2.4 (2)
It is better to have a relationship of
Then, the transformation from austenite to martensite can be suppressed during cooling, and the amount of retained austenite can be secured to improve the material properties.

  The high-strength steel sheet described in the claims is a steel sheet having both strength and processing characteristics, which are mutually contradictory characteristics, because retained austenite and martensite are mixed in a state of being finely dispersed in a large amount.

  The unannealed high-strength steel sheet described in the claim is complicated by the fact that, by performing an appropriate heat treatment later, a very high strength is obtained while being a low alloy, and abundant residual austenite is included. It exhibits good processing characteristics that can withstand processing.

  According to the manufacturing method described in the claims, the above-described high-strength steel plate or unannealed high-strength steel plate can be manufactured smoothly.

  Hereinafter, embodiments of a thin steel plate used for a processed part that requires excellent workability while having a tensile strength of about 600 to 1500 MPa and a manufacturing method thereof will be described. As a component system of the steel sheet, C: 0.05 to 0.25%, Si: 0.01 to 1.50%, Mn: 0.6 to 2.2%, Al: 0.01 to 1.50% in mass%, and Ti: 0.02 to 0.22%, Nb: Contains at least one of 0.02 to 0.10%, or contains a component selected from Cu: 0.01 to 1.0%, Ni: 0.01 to 1.0%, Cr: 0.01 to 1.0%, Mo: 0.01 to 1.0% Based on possible compositions. In addition, the thin steel plate described here is a steel plate having a thickness of about 4 mm or less. The steel sheet to be produced can be used mainly for parts that require high workability and strength, such as automobiles, home appliances, and electronic equipment products. In addition, it can be applied as a material for steel pipes.

First, the components of the steel sheet will be described.
As carbon (C), an amount in the range of 0.05 to 0.25% is required. If it is less than 0.05%, a steel plate strength of about 600 MPa or more cannot be obtained, so a carbon content of 0.05% or more is necessary. On the other hand, when the carbon content is 0.25% or more, the welded portion is excessively hardened and easily breaks from the welded portion. This is a limitation in use for thin steel plates, so an upper limit was set for the carbon content. And it has been found that when the carbon content is 0.05 to 0.25%, a composite structure in accordance with the gist of the present invention can be obtained.

  The amount of silicon (Si) is in the range of 0.01 to 1.50%. Silicon is used to improve strength by solid solution strengthening and to stabilize retained austenite. If the amount of silicon is 0.01% or more, the composite structure and material characteristics of the present invention can be obtained. As the amount of silicon increases, the amount of retained austenite can be increased and at the same time the stability thereof is promoted. However, when the silicon content is 1.50% or more, silicon scale is likely to be generated on the surface of the steel sheet, and a good surface property cannot be obtained as a thin steel sheet. Therefore, the upper limit of the silicon content is 1.50%.

Manganese (Mn) content is in the range of 0.6-2.2%. When the manganese content is 0.6% or less, the hardenability is insufficient and the composite structure of the present invention cannot be obtained. As a result, the strength and ductility of the steel sheet are lowered, so the manganese content is 0.6% or more.
On the other hand, the gist of the present invention is that a sufficient amount of austenite to improve the material properties is obtained even if austenite having a high carbon concentration is formed in the heating soaking process of the final annealing and then continuously cooled to 350 ° C. or lower. Is to remain. In order to obtain high strength, it is preferred to add a large amount of manganese. However, if it is too high, the quenching effect of manganese will transform even austenite with a high carbon concentration into martensite. Therefore, the upper limit of manganese content is set to 2.2%.
In addition, since elements that enhance hardenability, including inevitable impurity elements, have the same effect, the relationship with the cooling rate of the steel sheet after final annealing may satisfy formula (2) (above and below). is important.

  If aluminum (Al) is 0.01% or more, the composite structure and material characteristics of the present invention can be obtained. As the amount of aluminum increases, the amount of retained austenite can be increased, and at the same time, the stabilization thereof is promoted. However, when the amount of aluminum is 1.50% or more, it becomes easy to cause clogging of the casting nozzle in the continuous casting process, and it becomes extremely difficult to manufacture the slab by continuous casting. Therefore, the upper limit of the aluminum content is 1.50%.

Titanium (Ti) has an effect of suppressing recrystallization and crystal grain growth in the final annealing process as well as refinement of crystal grains in the hot rolling process.
Since the annealing temperature range in the final process is high, recrystallization or crystal grain growth occurs in a steel sheet having extremely fine crystal grains or a steel sheet subjected to cold rolling. If recrystallization or crystal grain growth occurs easily, the fine structure created in the hot rolling process is eliminated, and the structure is changed to a coarse structure that does not meet the gist of the present invention.
In the present invention, it is important to maintain a fine structure formed by the hot rolling process. Therefore, it is important to suppress recrystallization and crystal grain growth as much as possible in the annealing process. Titanium having a large effect of suppressing recrystallization and grain boundary migration is an effective element for the steel of the present invention. If the amount thereof is in the range of 0.02 to 0.22%, the effects according to the gist of the present invention can be exhibited. Even if the amount of titanium exceeds 0.22%, the effect does not increase so much, so the upper limit is made 0.22%.

  Niobium (Nb) also has the effect of suppressing recrystallization and crystal grain growth, similar to titanium. In order to obtain a structure in accordance with the gist of the present invention, the amount of niobium is required to be in the range of 0.02 to 0.10%. When it is 0.02% or less, the effect of suppressing recrystallization and crystal grain growth is lost. Moreover, even if the amount of niobium increases beyond 0.10%, the effect does not increase so much, so the upper limit was made 0.10%.

  Copper (Cu) improves the strength of the steel by solid solution strengthening, but if it is intentionally added beyond the unavoidable addition in production, the upper limit is set to 1.0%.

  Nickel (Ni) can increase the amount of retained austenite as well as silicon and aluminum while improving the strength of the steel by solid solution strengthening. However, the content is inevitably added in production, and if it is intentionally added, the cost increases, so the upper limit was made 1.0%.

  Chromium (Cr) is an element that enhances hardenability like manganese, and can improve the strength of steel, but it affects the formula (1), so the content is adjusted in a complex manner including Mn. There is a need. In addition, the content is inevitably contained in the production, and if it is intentionally added, the cost increases. Therefore, the upper limit is set to 1.0%.

  Molybdenum (Mo) is an element that enhances hardenability like manganese and can improve the strength of steel, but it affects the formula (1), so the content is adjusted in a complex manner including Mn. There is a need. In addition, the content is inevitably contained in the production, and if it is intentionally added, the cost increases. Therefore, the upper limit is set to 1.0%.

  Phosphorus (P) is not a necessary component in the steel of the present invention, but is inevitably contained in production. Phosphorus is an element that remarkably inhibits spot weldability and enhances hardenability like manganese, and affects the formula (1). Therefore, it is basically preferable to reduce it as much as possible.

  Since boron (B) is not a necessary component in the steel of the present invention but affects the formula (1), it is basically preferred to reduce it as much as possible.

The slab adjusted to the above-described reference component is either hot-rolled after reheating or hot-rolled immediately after casting.
In performing hot rolling, hot rolling is performed by a hot rolling mill having a plurality of stands so that the cumulative strain is 0.4 or more, or the rolling reduction in the final stand to be used is 15% or more.
In order to perform such high-pressure reduction, it is necessary to use a small-diameter roll mill having a work roll diameter of 600 mm or less or a different-diameter roll mill having an average work roll diameter of 600 mm or less for at least a plurality of subsequent stages. preferable. Moreover, in order to suppress the temperature rise of the steel sheet due to the heat generated by processing, it is also preferable to use a curtain wall type cooling device having a high cooling capacity arranged on each outlet side of these high pressure mills.
It is important to complete the rolling in the austenite region of Ar3 or higher. When the rolling completion temperature falls below the austenite region, a ferrite phase not suitable for the present invention is generated in addition to the lower bainite phase. The ferrite phase produced here is abnormally coarsened and significantly inhibits the material. Therefore, the temperature at which hot rolling is completed is Ar3 or higher.
Here, “strain” means the difference between the thickness h ₀ of the steel sheet at the entrance side of each stand (each stage) and the thickness h ₁ at the exit side divided by the average thickness of both ε = (h ₀ −h ₁ ) / {(h ₀ + h ₁ ) / 2}
“Cumulative strain” is a weighted integration of strain at each stage of the latter three stands, taking into account the strength of the influence on the metal structure, and the distortion at the last stage and its previous and previous stages. When ε _n , ε _n-1 and ε _n-2 respectively,
ε _C = ε _n + ε _n-1 / 2 + ε _n-2 / 4
Ε _C represented by

After completion of hot rolling, it is necessary to start cooling from Ar3 or higher, perform cooling at a rate of Δ40 ° C./sec or higher, continuously cool to a temperature range of 150 ° C. or higher and 400 ° C. or lower, and then wind up. When the cooling start temperature is lower than Ar3, ferrite transformation partially occurs. Similarly, when the cooling rate is lower than Δ40 ° C./sec, ferrite transformation occurs. Therefore, in order to obtain a desirable phase mainly composed of the lower bainite of the present invention, after completion of hot rolling, cooling is started from Ar3 or higher, and cooled to 150 ° C or higher and 400 ° C or lower at a cooling rate of Δ40 ° C / sec or higher. It is important to wind up later.
If the cooling start temperature after completion of hot rolling is too high, the form of carbide formed in the lower bainite is not preferable, and the cooling start temperature is more preferably Ar 3 or more and Ar 3 + 50 ° C. or less.

The structure of the hot-rolled steel sheet treated under appropriate conditions is a low-temperature transformation phase of the lower bainite phase, and cementite is precipitated in a chain form in the grains.
In addition, as described above, the rolling is performed so that the cumulative strain is 0.4 or more by hot rolling or the reduction ratio in the final stand to be used is 15% or more, so that the grain size of the prior austenite is 20 μm or less. I can do it.

Although cold rolling can be performed after hot rolling, excessive dislocations are introduced into the grains when rolling is performed at an excessive reduction rate, and recrystallization is likely to occur during annealing. The local high carbon concentration region around the cementite formed in the grains is combined with the grain boundaries. At this time, carbon can easily diffuse.
In the present invention, high carbon concentration austenite is formed in the ferrite grains during annealing at a temperature of Ac1 to Ac3 + 20 ° C., which is in a non-equilibrium state. When carbon diffuses easily, carbon falls to an equilibrium state. In order to increase the amount of retained austenite in the final product, it is important to suppress the decrease in concentration due to the diffusion of carbon, and excessive cold rolling that promotes diffusion is not preferable. When excessive cold rolling is performed, it is preferable to add an element that suppresses the movement of grain boundaries such as titanium and niobium.
Moreover, since the steel plate after hot rolling in the present invention has a hard structure composed of a lower bainite phase, a large load is applied to the cold rolling mill when performing cold rolling.
Considering the above, when performing cold reduction, the reduction rate is preferably 50% or less.
In addition, in the manufacturing example demonstrated here as embodiment, cold rolling is not performed.

In the final annealing, it is heated in the range of Ac1 to Ac3 + 20 ° C and kept isothermal for 1 to 120 seconds to form ferrite and austenite, then gradually cooled to a temperature of 650 to 700 ° C, and then cooled to 10 ° C / sec or more. Cool to 350 ° C or lower at speed.
At this time, in order to obtain high strength with a low alloy composition, it is necessary not to stop cooling to 350 ° C. or lower.
In the conventional TRIP steel plate manufacturing process, in order to concentrate carbon in austenite, cooling was stopped at a temperature just above the Ms point, which is the temperature at which martensitic transformation starts, and it was necessary to hold for several tens of seconds. However, according to the present invention, since austenite having a high carbon concentration is formed during annealing of Ac1 to Ac3 + 20 ° C., residual austenite having a high carbon concentration can be formed without holding at a temperature just above the Ms point during the cooling process. Is possible. Of course, in the steel sheet of the present invention, it is possible to concentrate carbon in the austenite by holding at a temperature just above the Ms point, but it is most important to increase the strength with a low alloy and cooling. It is the main point of the production process of the present invention that the temperature is not stopped below 350 ° C.

  The structure of the high-strength steel sheet manufactured on the basis of the above conditions has lath-like retained austenite in and within the ferrite grain with an average grain size of 10 μm or less, and the average grain size is on the ferrite grain boundary Granular martensite of 10 μm or less is finely dispersed. Some martensite has the same form as partially retained austenite. This coexistence of hard martensite and retained austenite gives rise to properties with high strength characteristics even in low alloys similar to composite steel sheets, and also produces the same elongation and processing characteristics as TRIP steel sheets.

FIG. 1 shows the concept of temperature history in hot rolling and annealing in the manufacturing process of the embodiment of the present invention, in which the horizontal axis represents time and the vertical axis represents temperature. From the left side of the figure, the range a indicates that a rough rolling process is performed, b indicates a finish rolling process, and c indicates a winding process.
2 (a) and 2 (b) show the coiling temperature (horizontal axis) after completion of hot rolling in the manufacturing process shown in FIG. 1 and the steel sheet material properties after final annealing (tensile strength TS, elongation value EL, etc. The relationship between the product TS × EL) and the form of retained austenite after final annealing (volume fraction of retained austenite Vf [γ], carbon concentration C [γ]) for two different steel types 1 and 2. is there. The final annealing was performed under the condition of cooling to near room temperature as described above without holding at around 400 ° C.
Although there is a difference depending on the steel type, the amount of retained austenite after final annealing increases with a decrease in the coiling temperature, and further decreases at 150 ° C. or less. In particular, the amount of retained austenite at a coiling temperature of 150 ° C. or higher and 400 ° C. or lower (preferably 150 ° C. to 350 ° C.) increases remarkably.
The material properties after the final annealing show no significant change in tensile strength, but the ductility is greatly improved as the amount of retained austenite increases. It can be said that the elongation characteristics are improved by the effect of the processing-induced transformation by retained austenite.
The results of measuring the carbon concentration in the retained austenite are also shown in the figure. Although the maximum carbon concentration is seen around the coiling temperature of 350 ° C, a carbon concentration of 0.9% or higher is obtained at temperatures of 150 ° C or higher and 400 ° C or lower, and austenite is stable even at room temperature. The carbon concentration necessary to exist is obtained.

However, the effect of FIG. 2 that the amount of retained austenite is large and the ductility is high when a specific hot rolling coiling temperature is taken is not an effect obtained in any component range. As a result of analyzing many experimental results, it was found that there is an appropriate component range in order to obtain the effect of FIG. The material properties of the steel of the present invention can be organized by elements that enhance the hardenability, and the index was experimentally derived with the Mn equivalent as the index. Further, the inventors have found that the threshold value of Mn equivalent for obtaining good material properties depends on the cooling rate in the final annealing step. The experimental results are shown in FIG.
That is, in order to obtain the effect shown in FIG. 2, the component adjustment shown in equations (1) and (2) is important.
Mn equivalent (mass%) = Mn + 1.4Cr + 3P + 1.5Mo + 180B (1)
Mn equivalent ≤ -0.005 (ΔCR) + 2.4 (2)
ΔCR (° C./sec): Cooling rate in the final annealing For example, when the cooling rate in the final annealing step is set to 10 ° C./sec or more as described above to cool to 350 ° C. or less, in order to obtain the effect of FIG. Therefore, it is necessary to keep the Mn equivalent to 2.35 or less.

In addition, the measurement of the retained austenite in FIG. 2 was calculated | required by the X ray diffraction method using the K alpha ray of Cu. After surface electropolishing finish at 1 / 2t thickness, the integrated strength of the (200) (220) (311) face of the austenite phase and the (200) (211) face of the ferrite phase is measured and calculated from the respective combinations. The average value of the retained austenite volume fraction was used.
The carbon concentration in the retained austenite is also calculated by X-ray diffraction, using the diffraction angle indicating the peak intensity of the (111) (200) (220) (311) plane of the austenite phase to obtain the lattice constant of the austenite phase. did.

FIG. 4 shows a micrograph of a typical cross-sectional structure of the steel of the present invention. Residual austenite (white portion in the photograph) is present in the form of lath in the ferrite grains and at the grain boundaries. In addition, martensite (the black part in the photograph) has a similar texture, and a martensite is present in a granular shape at the grain boundary triple point. The structural properties of the steel of the present invention are characterized by a large amount of both retained austenite and martensite.
In FIG. 4, the structure of the ferrite phase and the austenite phase is identified using the EBSP method.

The mechanism leading to the structure formation of the steel of the present invention will be described below. FIG. 5 shows the mechanism concept.
The structure of the hot-rolled sheet rolled up at 150 ° C. or more and 400 ° C. or less (FIG. 5A) is a lower bainite structure, and cementite is mainly dispersed in a chain form in the grains. When the temperature is raised to Ac1 to Ac3 + 20 ° C., austenite is preferentially formed in the vicinity of cementite in the grains ((b)). Since this austenite formation field is a state where carbides are densely packed, it becomes austenite having a high carbon concentration. When elements such as silicon and aluminum are present, they suppress the diffusion of carbon into the ferrite, resulting in a further increase in the carbon concentration in the austenite. Further, since carbon diffuses at grain boundaries at a high speed, austenite grains having a carbon concentration close to an equilibrium state are formed at grain boundary triple points and the like. When it is cooled to near room temperature, high carbon concentration austenite is transformed into retained austenite and part martensite, and low carbon concentration austenite is transformed into martensite ((c)).
At this time, if the Mn equivalent in the formula (1) becomes higher than the threshold, the ratio of transformation to martensite increases even with a high carbon concentration austenite, the amount of retained austenite decreases, and the material properties cannot be improved. .
In addition, when ac1 to ac3 + 20 ° C annealing is accompanied by the movement of grain boundaries such as recrystallization and grain growth, the high carbon concentration austenite formed in the ferrite grains and the like is combined with the grain boundaries, The carbon concentration decreases to an equilibrium state. Therefore, suppression of recrystallization and crystal grain growth is an important metallurgical element in the present invention. Steel sheets with extremely fine crystal grains and cold-rolled steel sheets are prone to recrystallization and grain growth during annealing in Ac1 to Ac3 + 20 ° C. In the steel plate, it is a desirable component element to be contained.
The present invention has been developed based on the above findings.

Examples of the invention will be described below.
The molten steel having chemical components shown in Table 1 was used as a slab (rolled material) by a continuous casting method or a forging method. Subsequently, these slabs were reheated and hot-rolled to obtain hot-rolled steel sheets.

In steel types A, E, G, J, and K adjusted to the above-described reference components, the amounts of carbon, silicon, manganese, and titanium are 0.05 to 0.20%, 0.20 to 1.20%, 1.55 to 1.85%, and 0.06 to 0.15%, respectively. Steel grade H is further 0.50% aluminum.
Steel type I has no addition of titanium and niobium, steel type C has titanium and niobium content of 0.04% and 0.03%, steel type D has no addition of titanium, and 0.04% of niobium is added as a grain boundary migration inhibiting element. .

Steel types B, F, and L, which are comparative examples, have carbon, silicon, and manganese amounts within the standard range, but the Mn equivalent is high due to other additions and inevitable elements, and is outside the standard range.
In the examples, since the cooling rate in the final annealing process was set to 100 ° C./sec, steel type F that would not satisfy the formula (2) in relation to the Mn equivalent was also used as a comparative example.

  About hot rolling, each steel type of an Example and a comparative example was rolled by sheet thickness 1.4-1.6mm on the conditions shown in Table 2, and it was set as the hot-rolled steel plate. The numbers underlined in Table 2 are outside the standard conditions. For steel grades B, F, and L, the chemical composition deviates from the standard, and the steel grade indication is underlined.

Except for the hot rolling conditions of No. 2 in Table 2, the final rolling reduction is 15% or more or the cumulative strain is 0.4 or more, finish rolling is finished with Ar3 or more, and cooling is started within 1 second. did. In the test conditions of No. 2, the final stage rolling reduction is 15% or less, the cumulative strain is also less than 0.4, and the rolling conditions are out of the standard.
Thereafter, the coiling temperature was controlled in the range of 100 to 610 ° C. The structure of each hot-rolled steel sheet after winding (in the hot rolling stage) is shown in the same table. The ones that have been subjected to low-temperature winding at 150 ° C or higher and 400 ° C or lower are structures mainly composed of lower bainite, whereas those below 150 ° C are completely martensite in which no precipitation of chain-like cementite is observed in the grains. The site phase, which is wound at 400 ° C. or higher, has a structure mainly composed of ferrite, pearlite, or upper bainite. In addition, the structure | tissue was observed with the optical microscope about the rolling direction cross section of the steel plate.

  For continuous annealing, it was heated to 800 ° C. and kept isothermal for 30 seconds, then allowed to cool to about 690 ° C., and then continuously cooled to about room temperature at Δ100 ° C./sec by the air-water cooling method. In steel types G and K, the cooling was stopped at 420 to 430 ° C., and then the condition of holding at the same temperature for several seconds was performed as a comparative example.

Table 2 also shows the martensite volume fraction and retained austenite volume fraction after continuous annealing, and the results of the tensile test and the hole expansion test. In Table 2, TS is the tensile strength, EL is the elongation value, and λ is the hole expansion rate.
In the examples of steel types A, C, and D (No. 1, 4, and 5), since all the steel types have a relatively small amount of carbon, the amount of martensite and the amount of retained austenite are small and the strength level is low. It exhibits relatively good elongation characteristics and good hole expansion characteristics.
However, since the steel type B of the comparative example (No. 3) has a high Mn equivalent, a sufficient amount of retained austenite cannot be obtained at the time of annealing and cooling, and the elongation characteristics are poor as compared with the examples of the steel type A.
In steel type A, the temperature conditions in hot rolling and continuous annealing are within the standard range, but the final rolling reduction in hot rolling is less than 15% and the cumulative strain is less than 0.4% (No. 2) ). Under these conditions, retained austenite can be obtained without performing low-temperature holding under continuous annealing conditions, but since the structure is coarse, the material properties are inferior compared to the examples of the same steel type.
In the examples of steel types E and K, the level of the coiling temperature in hot rolling is changed, and the influence on the structure and material after annealing is investigated. Examples (Nos. 9 and 21) in which the coiling temperature is higher than 400 ° C. or lower than 150 ° C. (No. 6, 7, 8, 10, 18, 19) and higher than 150 ° C. and 400 ° C. or lower. ), The amount of retained austenite tends to increase relatively, and the amount of martensite decreases accordingly. In each steel type, the strength level does not change greatly due to the difference in the coiling temperature level of hot rolling, but the elongation values in the examples are improved as the amount of retained austenite increases. Furthermore, the hole expansion value also increases as the winding temperature decreases.
In steel types G and K, the coiling temperature in hot rolling is controlled to 150 ° C. or more and 400 ° C. or less, and Comparative Examples (No. 12, 13, 20) in which 420 to 430 ° C. is maintained in the cooling process after annealing. Examples (Nos. 14 and 21) in which the retention was not performed at around 400 ° C. are shown. When holding at around 400 ° C. after annealing, the amount of retained austenite increases, but the amount of martensite decreases significantly compared to the case of not holding at around 400 ° C. As a result, although a good elongation value is obtained, the strength level is significantly reduced. On the other hand, even when annealing is not performed at around 400 ° C., abundant retained austenite is obtained, and abundant martensite is obtained, so that the strength level is high. In addition, TS × EL (product of tensile strength TS and elongation value EL), which shows a balance between strength and ductility, is as high as TRIP steel (20,000MPa ·% or more), and good elongation characteristics are obtained for high strength. It has been.

FIG. 6 shows the results of arranging examples and comparative examples of steel types E, G, and K by tensile strength and elongation value. In contrast to the comparative example in which the coiling temperature in the hot rolling is out of the reference range, the examples in the reference range can obtain elongation characteristics that may be at the same strength level.
Further, in contrast to the comparative example in which the holding is performed at around 400 ° C. after annealing, the example in which the holding is not performed has a high strength level and a good balance between strength and ductility.

The volume ratio of martensite is determined by polishing the cross section in the rolling direction of the steel sheet and then etching with a solution of 4% picric alcohol and 2% sodium pyrosulfate in a one-to-one relationship. And martensite etched in white by image analysis processing was measured and determined.
The volume fraction of retained austenite was calculated by the X-ray diffraction method.
Tensile properties (tensile strength TS, elongation value EL) were measured by tensile testing with the shape of JIS No. 5 test piece.
The hole expansion test is a measurement and evaluation of the hole expansion rate based on the Japan Iron and Steel Federation standard JFS T 1001-1996 hole expansion test method.

As described above, the high-strength steel sheets of the examples showing high strength and high ductility characteristics in a low alloy composition are suitable for use as automobile structural members and the like.
For example, the center pillar of an automobile requires the tensile strength necessary to prevent deformation at the time of collision as well as the support of the door, as well as bending, drawing workability for press molding, etc., to form mounting holes for related equipment It can be said that it is a very preferable steel plate as a member that requires a high level of hole expansion workability and further weldability for joining with other body parts.

It is a diagram which shows the concept of the temperature history in the hot rolling and annealing in the manufacturing process of embodiment of this invention. 2 (a) and 2 (b) are diagrams showing the relationship between the coiling temperature after completion of hot rolling, the steel sheet material properties after final annealing, and the form of retained austenite for two different steel types 1 and 2. . It is a figure which shows an experimental result about the threshold value of Mn equivalent for obtaining a favorable material characteristic depending on the cooling rate in the last annealing process. It is a typical cross-sectional structure | tissue photograph about the high strength steel plate by this invention. FIGS. 5A, 5B, and 5C are conceptual diagrams for explaining a mechanism that leads to formation of a structure of a high-strength steel sheet according to the present invention. It is a figure which shows the result which arranged the Example of the high strength steel plate by this invention, and the comparative example which is not so with the tensile strength and the elongation value.

Claims (12)

C: 0.05 to 0.25% by mass, Si: 0.01 to 1.50%, Mn: 0.6 to 2.2%, Al: 0.01 to 1.50%, the balance is Fe and impurities, Mn equivalent is 2.35 or less,
Lath-like austenite with a volume fraction of 3% or more and a carbon concentration of 0.9% or more in mass% and a lath-like or average with a volume ratio of 10% or more in the grains and at the grain boundaries of ferrite having an average grain size of 10 μm or less A high-strength steel sheet characterized by the presence of granular martensite having a particle size of 10 µm or less.
However, Mn equivalent is a value calculated by the following formula from the mass% of each component element to contain.
Mn equivalent = Mn + 1.4Cr + 3P + 1.5Mo + 180B   The high-strength steel sheet according to claim 1, further comprising any one or both of Ti: 0.02 to 0.22% and Nb: 0.02 to 0.10%.   It further contains any one or more of Cu: 0.01-1.0%, Ni: 0.01-1.0%, Cr: 0.01-1.0%, Mo: 0.01-1.0%, It described in Claim 1 or 2 characterized by the above-mentioned High strength steel plate. The high strength steel sheet according to any one of claims 1 to 3 , wherein a product of tensile strength (MPa) and elongation value (%) is 20000 (MPa ·%) or more. % By weight C: 0.05~0.25%, Si: 0.01~1.50 %, Mn: 0.6~2.2%, Al: 0.01~1.50% only contains, consists balance of Fe and impurities, Mn eq is at 2.35 or less,
An unannealed high-strength steel sheet characterized by having a lower bainite structure in which cementite having a prior austenite grain size of 20 μm or less and having a grain size of 1 μm or less dispersed in a chain form.
However, Mn equivalent is a value calculated by the following formula from the mass% of each component element to contain.
Mn equivalent = Mn + 1.4Cr + 3P + 1.5Mo + 180B The unannealed high-strength steel sheet according to claim 5, further comprising any one or both of Ti: 0.02 to 0.22% and Nb: 0.02 to 0.10%. It further contains any one or more of Cu: 0.01-1.0%, Ni: 0.01-1.0%, Cr: 0.01-1.0%, Mo: 0.01-1.0%, The claim 5 or 6 characterized by the above-mentioned. Unannealed high strength steel sheet. 0.05~0.25%, Si:: C mass% 0.01~1.50%, Mn: 0.6~2.2% , Al: 0.01~1.50% only contains, consists balance Fe and impurities, the steel Mn eq is 2.35 or less ,
With a hot rolling mill having multiple stands, hot rolling is performed so that the cumulative strain is 0.4 or more, or the rolling reduction in the final stand to be used is 15% or more, and from the temperature of Ar3 or more at the end of rolling. A method for producing an unannealed high-strength steel sheet, which starts cooling and winds in a temperature range of 150 ° C or higher and 400 ° C or lower.
However, Mn equivalent is a value calculated by the following formula from the mass% of each component element to contain.
Mn equivalent = Mn + 1.4Cr + 3P + 1.5Mo + 180B The method for producing an unannealed high-strength steel sheet according to claim 8, wherein the steel material further contains any one or both of Ti: 0.02 to 0.22% and Nb: 0.02 to 0.10%. The steel material further contains at least one of Cu: 0.01 to 1.0%, Ni: 0.01 to 1.0%, Cr: 0.01 to 1.0%, Mo: 0.01 to 1.0%, or The manufacturing method of the unannealed high-strength steel plate described in 9. An unannealed high-strength steel sheet obtained by the method for producing an unannealed high-strength steel sheet according to any one of claims 8 to 10 ,
A high-strength steel sheet that is heated to a temperature of Ac1 or higher and Ac3 + 20 ° C or lower for annealing, held at that temperature, and then continuously cooled to 350 ° C or lower at a rate ΔCR of 10 ° C / sec or higher. Production method. The relationship between the Mn equivalent described above and the cooling rate ΔCR after the annealing described above,
Mn equivalent ≤ -0.005 (ΔCR) + 2.4
The method for producing a high-strength steel sheet according to claim 11 .