JP7376771B2 - High-strength hot-rolled steel sheet and its manufacturing method - Google Patents

Description

The present invention relates to a high-strength hot-rolled steel sheet with excellent workability and cold rollability, and a method for manufacturing the same.

The strength of steel plates used for automobile parts and the like has been increasing in recent years. For example, steel plates with a tensile strength of 590 MPa or 780 MPa are commonly used in frame members, and high-strength steel plates with a tensile strength of 980 MPa or higher have begun to be used in some areas.
In order to manufacture such high-strength steel sheets, it is necessary to combine multiple strengthening mechanisms (methods) such as solid solution strengthening, precipitation strengthening, and even transformation structure strengthening, which inevitably increases the concentration of added elements. It's coming.

As a result, when manufactured under the same conditions as low-strength steel, when processed into parts as hot-rolled steel sheets, the load required for forming (for example, press) increases, the number of forming steps increases, and productivity decreases, or it becomes difficult to form. A problem arises in that the shape is restricted.
In addition, when hot-rolled steel sheets are further cold-rolled and used as cold-rolled steel sheets, the load required for cold rolling increases, resulting in restrictions on the cold-rolling rate, or after intermediate annealing and then re-cooling. Occasionally, there are situations where it is necessary to perform inter-rolling.

One possible way to deal with this situation is to perform slow cooling after hot rolling, and to set the winding temperature higher. By doing this, a microstructure consisting of a soft ferrite phase (hereinafter referred to as F phase) and a hard structure (pearlite, hereinafter referred to as P structure) is obtained, and press forming and cold rolling are relatively easy. become.

However, with a steel sheet composed of the F phase and P structure as described above, another problem remains that it is not easy to uniformly increase the strength of molded products and cold rolled steel sheets, which was the original purpose.
In the first place, the reason why the concentration of added elements was increased to the extent that the strength of the hot-rolled steel sheet affected the formability and cold rollability was specifically to take advantage of transformation structure reinforcement.

In order to utilize transformation structure strengthening, carbides containing P structure must be removed during heat treatment after forming when used as a hot-rolled steel sheet, and during heat treatment after cold rolling when used as a high-strength cold-rolled steel sheet. It is necessary to increase the carbon concentration in the austenite phase (hereinafter referred to as phase A) by dissolving as much as possible, but this is not so easy to achieve. It is necessary to either raise the heating temperature or hold it for a long time, but due to equipment constraints and the need to ensure productivity, there is not much flexibility in heating temperature and holding time. Furthermore, even if it were possible to sufficiently dissolve carbides by processing without restrictions, there would be a large difference in carbon concentration between the region that was in the F phase before heating and the region that was in the P structure or carbide. In order to reduce and achieve uniformity, a considerably long diffusion process is required, so this is not realistically selected.

On the other hand, it is also possible to manufacture a hot rolled steel sheet with priority given to making the carbon concentration uniform and making the size of carbides relatively small to make them easier to melt.
Instead of the above method, by increasing the cooling rate after hot rolling and lowering the coiling temperature, it is possible to obtain a steel sheet with a microstructure mainly composed of bainite structure (hereinafter referred to as B structure). However, in this case, it is difficult to ensure formability as a hot-rolled steel sheet, and the required load during cold rolling increases.

As mentioned above, hot-rolled steel sheets that achieve both ease of forming and cold rollability (low rolling load), uniformity of carbon concentration, and ease of dissolving carbides have been proposed. The reality is that cooling conditions and winding conditions after rolling have not been found.

For example, Non-Patent Document 1 discloses a method for prototyping a hot-rolled steel sheet with good cold rolling properties, in which a hot-rolled coil is wound at a high temperature, air-cooled for a certain period of time to promote ferrite transformation, and then immersed in water-cooled. ing. For steel having mass percentages of C: 0.17%, Si: 1.3%, and Mn: 2.0%, hot rolling was completed at 920°C, coiled at 630°C, and immersed in water after 30 or 60 minutes. It has been shown that by cooling, it was possible to achieve softening over the entire length of the hot rolled coil and to reduce the occurrence of grain boundary oxidation.
However, since the microstructure estimated from the temperature history is said to be composed of F phase, P structure, and low-temperature products, it is assumed that the nonuniformity of carbon concentration when heated to A phase will not be improved at all. It's presumed and never actually mentioned.

Patent Document 1 discloses that after hot rolling is performed at a reduction rate of 50% or more in a temperature range of Ar ₃ or higher, then preliminary cooling is performed from a temperature of Ar ₃ or higher at a cooling rate of 2°C/sec or higher. After that, pre-cooling is temporarily interrupted at a temperature of Ar ₃ or below and Ar ₃ -100°C or above, and after waiting for a predetermined period of time, the temperature is increased again from 400 to 600°C at a cooling rate of 3°C/s or more and 15°C/s or less. A manufacturing method is disclosed in which a steel plate with a low yield ratio is obtained by heating and cooling the steel sheet to a temperature range of 1 to 100 nm to make the microstructure a mixture of F phase and B structure or martensitic phase (hereinafter referred to as M phase). .
However, since there is no description of material properties other than yield ratio, it is not clear about the uniformity of the material or the ease with which carbides can be dissolved, which facilitates increasing the strength by heat treatment.

Patent Document 2 discloses that rolling of a steel billet having a predetermined chemical composition is completed at a temperature equal to or higher than the Ar ₃ transformation point, the coil is once coiled at a temperature of 750 to 600°C, and the coil is unrolled after being held for 10 to 30 minutes. However, a method for producing a hot rolled steel sheet is disclosed in which the steel sheet is cooled at a cooling rate of 20° C./sec or higher and then re-rolled at a temperature of 550° C. or lower.
According to this method, hot-rolled steel sheets are soft and grain boundary oxidation is suppressed, so a high yield can be achieved in cold rolling, but what kind of microstructure can be obtained? Further, it is not clear about the uniformity of the material or the ease with which carbides can be dissolved, which facilitates increasing the strength by heating or annealing. Furthermore, this technology requires equipment that has a winding device and an unwinding device located in the middle of common hot rolling cooling equipment (called a run-out table or hot run table), so even if desirable characteristics cannot be obtained, However, there is a problem in that it cannot be easily implemented industrially.

Japanese Patent Application Publication No. 2000-087138 Japanese Patent Application Publication No. 2013-253301

Kobayashi, Doi, Kimura, Koizumi, Kimijima, Sano, Sekisui, Morimoto, Ishikawa: Tetsu to Hagane, vol. 100 (2014) No. 5, 616-624

In view of the above circumstances, the present invention aims to improve the good formability and cold ductility of a steel plate composed of F phase and P structure, the material uniformity of a steel plate mainly composed of B structure, and the high strength achieved by heating and annealing. The objective is to obtain a steel sheet that is compatible with the ease of dissolving carbides.

The inventors of the present invention conducted extensive studies in order to obtain a steel sheet that is compatible with good formability, good cold rollability, material uniformity, and ease of dissolving carbides.
Specifically, the microstructure is made mainly of the B structure to first ensure uniformity of the material and ease of dissolving carbides, and then it is softened to have excellent formability and cold rollability. We conducted research to see if there was any. As a result, it has been found that such a steel plate can be obtained by cooling again at a certain cooling rate or higher within a very short time after the completion of bainite transformation, and then winding it at a predetermined temperature or lower.
The gist of the invention thus completed is as follows.

(1) Contains C: 0.05 to 0.50%, Si: 0.01 to 2.0%, Mn: 0.5 to 3.0%, and P: 0.03% or less in mass % , S: 0.02% or less, N: 0.05% or less, Al: 0.05% or less, with the remainder consisting of Fe and unavoidable impurities, and the microstructure is Contains a bainitic structure with an area ratio of 50% or more and a ferrite phase with an area ratio of 15 to 30%, the average crystal grain size of the ferrite phase is 30 μm or less, and the absolute value of local strain in the ferrite phase determined by the fine lattice marker method. A high-strength hot rolled steel sheet having an average of 0.050 or more and a tensile strength of 440 MPa or more.

(2) The above steel plate further includes, in mass %, Ti: 0.1% or less, Nb: 0.1% or less, B: 0.01% or less, Cr: 1.5% or less, Cu: 1 .0% or less, Ni: 1.0% or less, and the high-strength hot-rolled steel sheet according to (1) above.
(3) The above steel plate further contains, in mass%, Mo: 0.01 to 1.0%, W: 0.01 to 0.5%, and V: 0.01 to 0.5%. The high-strength hot-rolled steel sheet according to (1) or (2) above, characterized in that the high-strength hot-rolled steel sheet contains one or more of the following.

(4) A method for producing a high-strength hot-rolled steel sheet according to any one of (1) to (3) above, in which steel having a chemical composition according to any one of (1) to (3) above is cast. After that, hot rolling is carried out directly or by reheating to 1300°C or less, and in the hot rolling, finish rolling is performed with a cumulative reduction rate of 50% or more, which is completed at Ar ₃ points or more, and after hot rolling, As the first stage of cooling, cooling is performed to 400 to 550 °C at an average cooling rate of 15 to 35 °C/s, and then as the second stage of cooling, at a rate of 15 °C/s or less (bainite structure is 50% or more and austenite). cooling to any point between 100 seconds after the bainite structure becomes 50% or more and the austenite phase becomes 20%), and then a third stage of cooling. A method for producing a high-strength hot-rolled steel sheet, the method comprising: cooling to 300°C or less at an average cooling rate of 50°C/second or more until the temperature reaches 300°C, and then winding the sheet.

If the steel sheet of the present invention is used, hot-rolled steel sheets that can be formed with low loads and easily heat-treated after forming to make high-strength members, and hot-rolled steel sheets that can be cold-rolled with low loads and can be heat-treated after cold-rolling to have high strength. It is possible to obtain a cold-rolled steel plate material that can be easily converted into a raw material for cold-rolled steel sheets. Furthermore, since it can be manufactured using general-purpose equipment, it can contribute to a wide range of industries.

It is a schematic diagram explaining a fine lattice marker. FIG. 3 is a schematic diagram showing a grid marker after heat treatment. It is a figure which shows the relationship between the average absolute value of local strain in F phase, and yield strength.

The steel plate of the present invention and its manufacturing method will be explained in detail.
[Steel plate]
First, the chemical composition of the steel plate will be explained.

<C: 0.05-0.50%>
C is an essential element to obtain a high-strength steel plate. In order to obtain a steel plate having a tensile strength of at least 440 MPa or more, it is necessary to contain 0.05% or more. On the other hand, the upper limit is set at 0.50% due to the need to ensure weldability and the toughness of the welded part.

<Si: 0.01-2.0%>
Si is an element useful in designing the strength of steel sheets because it has the effect of strengthening the F phase. Since it also has the effect of suppressing the formation of cementite, it is particularly useful for microstructural design of cold-rolled steel sheets. However, if it is contained in excess, it will have an adverse effect on surface treatment properties such as pickling properties and chemical conversion treatments. Therefore, the upper limit is set at 2.0%. On the other hand, reducing the content to less than 0.01% places an excessive burden on the steel manufacturing process, so the lower limit is set at 0.01%.

<Mn: 0.5-3.0%>
In addition to solid solution strengthening, Mn has high hardenability, and is therefore an extremely important element for increasing the strength of steel sheets through strengthening the transformation structure. Therefore, the lower limit is set at 0.5%, at which a clear effect is produced. On the other hand, if the content exceeds 3.0%, the mechanical properties may deteriorate due to solidification segregation, so the upper limit is set at 3.0%.

<P: 0.03% or less>
P is an impurity and must be limited to 0.03% or less since it has an adverse effect on hot workability. On the other hand, although there is no particular lower limit set, 0.001% may be used as a guideline since reducing the content more than necessary will put a heavy burden on the steel manufacturing process.

<S: 0.02% or less>
S is an impurity and must be limited to 0.02% or less because it has an adverse effect on hot workability and mechanical properties such as ductility and toughness. On the other hand, there is no particular lower limit set, but reducing it more than necessary will put a heavy load on the steel manufacturing process, so 0.0001% may be used as a guide.

<N: 0.05% or less>
N forms nitrides with B, reducing the contribution of B to hardenability, and forms nitrides with Ti, deteriorating mechanical properties, so it is desirable to reduce it as much as possible. It is acceptable if it is 0.05% or less. On the other hand, reducing the content more than necessary places a heavy burden on the steel manufacturing process, so 0.0010% may be used as a guideline.

<Al: 0.05% or less>
Al is used as a deoxidizing element, but its oxide affects the surface quality and the oxide film also affects the surface treatment characteristics, so it needs to be kept at 0.05% or less. On the other hand, since it places a heavy burden on the steel manufacturing process, 0.01% may be used as a lower limit.

The steel sheet of the present invention is based on the above-mentioned chemical components containing C, Si, and Mn, and with limited content of P, S, N, and Al, but furthermore, the following chemical components are included: Ti, Nb, and B. , Cr, Cu, and Ni, as well as one or more of Mo, W, and V, if necessary.

<Ti: 0.1% or less>
Ti combines with N and suppresses N from reducing the contribution of B to hardenability. Therefore, when containing B and utilizing its hardenability to increase strength, Ti is 0.01 It is desirable to add % or more as a guideline. On the other hand, as already mentioned, TiN may impair mechanical properties, and excessive addition may suppress recrystallization after cold rolling and impair productivity. The upper limit is set at 0.1%, since there is a risk that the hardenability may be lowered by decreasing the content.

<Nb: 0.1% or less>
Like Ti, Nb combines with N and suppresses N from reducing the contribution of B to hardenability. Therefore, when incorporating B and utilizing its hardenability to increase strength, , it is desirable to add 0.01% or more. On the other hand, if it is added in excess of 0.1%, the effect will be saturated, and there is a risk that it will further suppress recrystallization after cold rolling and impair productivity, and that it will combine with C and reduce hardenability. Therefore, the upper limit is set at 0.1%.

<B: 0.01% or less>
B has the effect of increasing hardenability when added in an amount of 0.0001% or more, so it can be added as necessary. On the other hand, excessive addition leads to deterioration of hot workability and ductility, so the upper limit is set at 0.01%.

<Cr: 1.5% or less>
Since Cr is an element that has hardenability, it can be used as appropriate. In order to obtain this effect, it is preferable to contain 0.01% or more. However, if it is added in an amount exceeding 1.5%, the effect will be saturated and the manufacturing cost will only increase, so the upper limit is set at 1.5%.

<Cu: 1.0% or less>
Cu has the effect of increasing strength, so it can be added as necessary. In order to obtain this effect, it is preferable to contain 0.01% or more. However, if it exceeds 1.0%, the surface quality of the hot rolled steel sheet will be impaired, so the upper limit is set at 1.0%.

<Ni: 1.0% or less>
Since Ni is an element that improves hardenability, it can be added as necessary. In order to obtain this effect, it is preferable to contain 0.01% or more. On the other hand, since it is an expensive element, the upper limit is set at 1.0% at which the addition effect is saturated. Further, since Ni has the effect of suppressing deterioration of the surface quality of hot rolled steel sheets due to Cu, it is desirable to include Ni at the same time as Cu.

<Mo: 0.01-1.0%>
<W: 0.01-0.5%>
<V:0.01~0.5%>
All of these elements are elements that improve hardenability. In order to obtain the effect of addition, 0.01% or more of each can be added as necessary. On the other hand, since these elements are expensive, the upper limit for addition is set at the point where the effect of addition is saturated. The upper limit of Mo is 1.0%, and the upper limit of W and V is 0.5%.

<Remainder>
In the present invention, the component other than the above (the remainder) is Fe, but unavoidable impurities mixed in from melted raw materials such as scrap, refractories, etc. are allowed.

Next, the microstructure and other requirements of the steel plate of the present invention will be explained.
<Bainitic structure with an area ratio of 50% or more (hereinafter referred to as B structure) and ferrite phase with an area ratio of 15 to 30% (hereinafter referred to as F phase)>
The steel sheet of the present invention has a microstructure in which the B structure occupies the largest area ratio, followed by the F phase.
By setting the B structure to 50% or more, the uniformity of the carbon concentration is ensured, and a layered non-uniform structure of pearlite and ferrite, called a so-called band structure, is avoided. When this non-uniform structure is formed, it becomes a starting point for cracking and destruction, so the B structure should be 50% or more. Preferably it is 60% or more, more preferably 70% or more.
Furthermore, in order to cause local strain of 0.050 or more to remain in the F phase adjacent to the B structure, as described later, the area ratio of the F phase must be 30% or less. If it exceeds 30%, the local strain that can remain in the F phase will be dispersed and become small, which is not preferable. If the F phase is less than 15%, the material may lack ductility when used as a high tensile strength material for automobiles, etc., so the F phase should be 15% or more.

In addition, as microstructures other than those mentioned above, 10% or less of a martensite phase (hereinafter referred to as M phase) and 10% or less of a retained austenite phase (hereinafter sometimes referred to as A phase) are allowed.
Note that each area ratio is determined by polishing a cross section of the steel plate parallel to the rolling direction and corroding it with a nital solution, and then observing the 1/4 position of the plate thickness.

<Average grain size of F phase is 30 μm or less>
The average grain size of the F phase is a factor that greatly affects the strength and toughness of a steel sheet. Generally, the finer the crystal grain size, the better the tensile strength. In order to maintain such strength, the average crystal grain size of the F phase needs to be 30 μm or less.

<The average absolute value of local strain in the ferrite phase determined by the fine lattice marker method is 0.050 or more>
These conditions must be met in order to obtain a steel sheet that has a microstructure consisting of a B structure and an F phase, but exhibits formability and cold rollability equivalent to a steel sheet that has a microstructure consisting of an F phase and a P structure. There is a need.
The steel sheet of the present invention is rolled in the A phase, cooled to 400 to 550 °C at a cooling rate of 15 to 35 °C/s as the first stage cooling, and further cooled as the second stage cooling (B structure is 50% or more and the austenite phase is 35%) to (100 seconds after the B structure is 50% or more and the austenite phase is 20%), the cooling rate is Cool or maintain temperature within 15° C./s to transform into bainite. At this time, the crystal grains of the F phase adjacent to the B structure are subjected to an external compressive or tensile force, and strain remains within the crystal grains. Then, by cooling from the above point to 300° C. or lower, including room temperature, at an average cooling rate of 50° C./s or higher as the third stage of cooling, it becomes possible to keep those strains remaining.

In the present invention, the strain that remains (exists) in the crystal grains of the F phase is defined as local strain in each crystal grain.
Local strain in the F phase has the effect of facilitating the movement of dislocations when a steel plate undergoes deformation such as press forming or cold rolling, regardless of whether it is caused by an external force such as compression or tension. Therefore, it is thought that this can be enjoyed as a reduction in forming load and cold rolling load.
The local strain differs for each crystal grain of the F phase, and is determined by the size of the adjacent B structure, the crystal orientation relationship with the same structure, the formation temperature of the F phase, and the like. The larger the absolute value, the greater the effect. Therefore, the formability and cold rollability of a steel sheet can be evaluated by measuring the local strain for as many F phases as possible within a practical range and averaging their absolute values. On the other hand, such effects can be quantified as (reduction in) yield strength in a tensile test.

As a result of experiments conducted by the present inventors, when the average absolute value of local strain is 0.050 or more, the yield strength is equivalent to (or lower than) that of a steel plate having a microstructure consisting of an F phase and a P structure. In other words, it has become clear that a steel plate exhibiting formability and cold rollability can be obtained. This provision was made based on this.

The local strain within the F phase was determined by the fine lattice marker method. The procedure will be explained below.
First, grid markers are formed on the surface of a steel plate as follows.
The steel plate whose local strain is to be measured is sampled, its surface chemically polished, and then corroded with nital solution to reveal the microstructure. Next, a positive photoresist (photosensitive material) is applied. As the photoresist, for example, ZEP520A manufactured by Zeon Corporation can be used.
A photoresist film formed on a steel plate is scanned with an electron beam to expose it to light in the form of a square lattice. The width of the grating was set to 100 nm, and the distance between the centers of the width of the grating was set to 500 nm (see FIG. 1). The grid was fabricated in an area of 500 μm×500 μm.
The photoresist irradiated with electron beams was developed and the exposed areas, that is, the areas corresponding to the sides of the grid, were removed. For development, ZED-N50 manufactured by Nippon Zeon Co., Ltd. or the like can be used.
Gold is deposited on the developed surface. The gold is deposited directly on the surface of the steel plate in the areas that are removed by development (the sides of the grid), while it is deposited on the resist in the areas that are not exposed to light and have not been removed.
After that, when the photoresist is dissolved in an organic solvent, the gold on the resist is removed together with the resist, leaving only the gold deposited directly on the steel plate surface, forming a gold lattice (marker) on the steel plate surface. .

Next, the steel plate on which the grid markers are formed is held at 300° C. for 2 hours. This releases local strain, which can be observed from the surface as deformation of the lattice. Note that, in order to facilitate the following measurements, this heat treatment is preferably performed in a non-oxidizing atmosphere (for example, an Ar gas atmosphere).
A lattice in which all four lattice points are in the F phase is observed by SEM, and the distances (side lengths) L1 to L4 between the lattice points are measured for each lattice (see FIG. 2). Then, the value ΔLmax/500 obtained by dividing the maximum absolute value ΔLmax (unit: nm) of the difference between them and 500 nm by 500 nm is applied to the lattice portion before performing the heat treatment described above at 300°C for 2 hours. It is defined as the distortion that existed.
In this way, the strain was determined for at least 500 (pieces) of the lattices, and the average value was calculated to be used as the absolute value of the local strain in the F phase.

On the other hand, an experiment was conducted in which the cooling conditions after rolling were varied, and a plurality of steel plates with different average absolute values of local strain in the F phase were produced. The yield strength of these steel sheets was compared with that of a steel sheet made of F phase and P structure made from steel slabs with the same chemical composition, and it was found that the average absolute value of local strain in F phase was 0.050 or more. In some cases, the results were found to be the same or lower. This is why the average absolute value of local strain in the F phase is set to 0.050 or more in the present invention (see FIG. 3, which will be described later).

<Mechanical properties>
According to the hot rolled steel sheet of the present invention, high tensile strength, specifically, a tensile strength of 440 MPa or more can be achieved. The tensile strength is preferably 600 MPa or more, more preferably 800 MPa or more. The upper limit of the tensile strength is not particularly limited, but may generally be 1000 MPa or less.

<Usage form>
The hot-rolled steel sheet of the present invention configured as described above can be formed into a product sheet by surface treatment as necessary, so that it can be formed with a low load, and it can be heat-treated after forming to form a high-strength member. As a hot-rolled steel sheet that can be easily rolled, it can be used in a variety of applications such as architecture, automobiles, home appliances, and industrial machinery.
In addition, by using the steel billet as the original sheet for cold rolled steel sheets in the process of hot rolling and cold rolling into thin steel sheets, it can be cold rolled with a low rolling load, and it can be heat treated after cold rolling to increase its strength. It can be made into a hot-rolled sheet that is easy to roll.

[Production method]
A method for manufacturing a steel plate according to the present invention will be explained.
A slab having a chemical composition that satisfies the above conditions is manufactured. In manufacturing, continuous casting is desirable from the viewpoint of productivity.
Hot rolling is performed directly after casting or after reheating to 1300°C or less. In hot rolling, finish rolling is completed at _three or more Ar points. If the reheating temperature in the case of reheating before rolling exceeds 1300° C., a decrease in yield due to oxidation cannot be overlooked. On the other hand, the lower limit of the reheating temperature may be any temperature as long as finish rolling can be completed at ₃ or more Ar points, and can be set according to the specifications of the rolling equipment.

The reason why the finish rolling is completed at _three or more Ar points is to make the microstructure after cooling to be composed of a bainite structure and an F phase. The reason why the cumulative reduction ratio in finish rolling is set to 50% or more is to set the average grain size of ferrite to 30 μm or less. In the present invention, the cumulative rolling reduction ratio is the sum of the nominal rolling reduction ratio (rolling amount/inlet side plate thickness) for each pass when rolling is performed in multiple passes.

After hot rolling, a first stage of cooling is performed to 400 to 550°C at an average cooling rate of 15 to 35°C/second to effect bainite transformation. If the cooling rate at this time is less than 15° C./sec, ferrite transformation becomes predominant, resulting in a microstructure consisting of an F phase and a P structure. On the other hand, if it exceeds 35° C./sec, the steel sheet will consist mainly of M phase, which is not desirable.
Further, the reason why the temperature at which cooling is performed at an average cooling rate of 15 to 35° C./sec is set to 400 to 550° C. is to increase the occupancy of the B structure and to suppress the formation of the M phase as much as possible.

The second stage of cooling is at 15°C/s or less for 100 seconds from (the time when the bainite structure becomes 50% or more and the austenite phase becomes 35%) to (the time when the bainite structure becomes 50% or more and the austenite phase becomes 20%). Cool to some point between In order to do this, the cooling conditions after rolling may be determined based on a CCT curve predetermined according to the component.
The reason why the cooling rate in the second stage was set to 15° C./sec or less is that if the cooling rate exceeds 15° C., the temperature drops below the Ms point during the progress of the B transformation, and the M phase will predominately be produced.

After achieving such a phase ratio, the third stage of cooling is started, and the cooling rate from the start of cooling to 300° C. is set to 50° C./sec or more to cool down to 300° C. or less, and the material is wound into a coil. Thereby, a fine stress distribution can be maintained and frozen, and strain can remain in the F phase.
There is no need to set an upper limit on this cooling rate for the purpose of suppressing the release of the strain introduced in the F phase, but if it is too fast, the load on the equipment will be large, so the upper limit should be 100°C/sec. .

Here, the reason for starting the third stage of cooling is after the time when the amount of B structure reaches 50% or more is because if the B structure is not at least 50% or more at the start of cooling, the yield ratio will decrease. This is because the strain required to make the martenite phase This is because more sites tend to appear and the amount of ferrite, which has the effect of lowering the yield point, decreases. In addition, the reason why we set the time point 100 seconds after the austenite content is 20% is because if we start cooling after this time, the strain introduced during bainite transformation will be released and the desired effect will not be obtained. It is.

By satisfying the above conditions, a steel plate that satisfies the above conditions can be obtained.

The present invention will be explained by showing examples.
<Example 1>
First, a steel piece having the chemical components listed in Table 1 (the remainder being Fe and unavoidable impurities) was produced. After heating this steel piece to 950°C, it was cooled at several levels of cooling rate between 1 and 60°C/second, and at each cooling rate, the specimen was rapidly cooled at various temperatures during cooling to form an austenite phase. Based on the results, the CCT curves of the steel slabs of each component were determined.

Next, a steel billet having the chemical composition listed in Table 1 was separately produced, and the plate thickness was increased to 2 A hot rolled steel plate with a thickness of .0 to 3.6 mm was used.
In Table 2, the heating temperature of the slab is SRT, the rolling end temperature is FT, the initial cooling rate after rolling is CR1, the slab is cooled to the transformation completion temperature MT (first stage), and the subsequent air cooling time is t _AC The material was air cooled at a cooling rate of CR2 (second stage), and then cooled to the cooling end point temperature (winding temperature) CT at a cooling rate of CR3 (third stage). Note that the austenite fraction at the time of starting cooling in the third stage is ξA, and the time from the time when the austenite fraction reaches 20% until the third stage cooling is started is ty. Here, the austenite fraction was determined by measuring the average value of the plate thickness using a laser ultrasonic method. Further, under condition ii, after the rolling was completed, the sheet was cooled to 600° C. at a rate of 5° C./sec, and the sheet was wound up as it was. Note that the Ar ₃ points of the steels in Table 1 are 750 to 830°C, and the FTs in Table 2 are all Ar ₃ points or higher.

A cross section of the obtained steel plate parallel to the rolling direction was polished and subjected to nital corrosion, and the microstructure composition at a position of 1/4 of the plate thickness was investigated. In addition, a grid marker was made on a steel plate that had been subjected to nital corrosion after chemical polishing, and heat treatment was performed to release the strain. The strain in the F phase was measured, and the average of the absolute values was determined.
Furthermore, a JIS No. 5 tensile test piece was prepared from the obtained steel plate. The test pieces were taken with the tensile direction perpendicular to the rolling direction. A tensile test was conducted on the test piece to determine the yield strength. The yield strength was defined as the upper yield point for those in which the upper yield point was observed, or the 0.2% proof stress for those in which the upper yield point was not observed.
The results are shown in Table 3.

A steel plate (No. 1) obtained by rolling steel a at a cumulative reduction rate of 150% under condition i has a microstructure that satisfies the present invention and has an average absolute value of local strain in the F phase of 0.050 or more. The yield strength was lower than that of the steel plate (No. 2) which was rolled under condition ii at a cumulative reduction rate of 150% to form an F phase and a P structure. In addition, hereinafter, unless otherwise specified, rolling is performed at a cumulative reduction rate of 150%.
Steel plates (No. 4 and 5) obtained by rolling steel b under conditions iii and iv have a microstructure that satisfies the present invention, and have an average absolute value of local strain in the F phase of 0.050 or more. The yield strength was lower than that of the steel plate (No. 3) rolled under condition ii to have an F phase and a P structure.

Steel plates (No. 7, 8, and 9) obtained by rolling steel c under conditions v, vi, and vii have a microstructure that satisfies the present invention, and the average absolute value of local strain in the F phase is 0. .050 or more, and showed a lower yield strength than the steel plate (No. 6) rolled under condition ii to have an F phase and a P structure.
In addition, the steel plates (No. 10 and 11) obtained by rolling steel c under conditions viii and ix have a yield strength higher than that of No. 6, with the average absolute value of local strain in the F phase being less than 0.050. showed that.

A steel plate (No. 13) obtained by rolling steel d under condition x has a microstructure that satisfies the present invention, has an average absolute value of local strain in the F phase of 0.050 or more, and is rolled under condition ii. The yield strength was lower than that of the steel plate (No. 12) with F phase and P structure.

A steel plate (No. 15) obtained by rolling steel e under condition xi has a microstructure that satisfies the present invention, has an average absolute value of local strain in the F phase of 0.050 or more, and is rolled under condition ii. The yield strength was lower than that of the steel plate (No. 14) with F phase and P structure.
In addition, steel plates (No. 16 and 17) obtained by rolling the same steel under conditions xii and .050 or less, showing a higher yield strength than No. 14.
FIG. 3 graphically shows the relationship between the average absolute value of local strain in the F phase and the yield strength for e-steel. When the average absolute value of local strain in the F phase is 0.050 or more, there is a clear decrease in yield strength relative to the steel plate (No. 14) with the F phase and P structure.

A steel plate (No. 19) obtained by rolling steel f under condition xiv has a microstructure composition that satisfies the present invention, an average absolute value of local strain in the F phase is 0.050 or more, and is rolled under condition ii. The yield strength was lower than that of the steel plate (No. 18) with F phase and P structure.
A steel plate (No. 21) obtained by rolling steel g under conditions The yield strength was lower than that of the steel plate (No. 20) with F phase and P structure.
Furthermore, since no F phase was generated in the steel plate (No. 22) obtained by rolling the same steel under condition xvi, no further evaluation was performed.

A steel plate (No. 24) obtained by rolling steel h under condition xvii has a microstructure that satisfies the present invention, has an average absolute value of local strain in the F phase of 0.050 or more, and is rolled under condition ii. The yield strength was lower than that of the steel plate (No. 23) with F phase and P structure.
In addition, a steel plate (No. 25) obtained by rolling the same steel under condition xviii has a microstructure composition outside the range of the present invention, and the average absolute value of local strain in the F phase is 0.050 or less, It showed higher yield strength than No. 23.

Here, in order to investigate the influence of the cumulative reduction rate, a steel plate was prepared by rolling steel h under condition xvii with a cumulative reduction rate of 50%. This steel plate has an average absolute value of local strain in the F phase of 0.050 or more, and has a lower yield strength of 650 MPa than the steel plate (No. 23) rolled under condition ii to form the F phase and P structure. At the same time, the average crystal grain size of the ferrite phase was 30 μm, and a good tensile strength of 902 MPa was exhibited.
On the other hand, a steel plate was also produced by rolling the same steel under condition xvii with a cumulative reduction rate of 40%. This steel sheet had an average absolute value of local strain in the F phase of 0.050 or more and exhibited a low yield strength of 632 MPa, but the average grain size of the ferrite phase was 35 μm and the tensile strength was also low at 430 MPa. It became the value.

<Example 2>
Steel pieces having the chemical components listed in Table 4 (the remainder being Fe and unavoidable impurities) were produced, and hot-rolled steel plates with a thickness of 2.4 to 3.0 mm were prepared under the conditions listed in Table 5. However, under condition 2-i, after the rolling was completed, the sheet was cooled to 600° C. at a rate of 5° C./sec and then wound up as it was.

A cross section of the obtained hot rolled steel sheet parallel to the rolling direction was polished and subjected to nital corrosion, and the microstructure composition at a position of 1/4 of the sheet thickness was investigated. In addition, a grid marker was made on a steel plate that had been subjected to nital corrosion after chemical polishing, and heat treatment was performed to release the strain. The strain in the F phase was measured, and the average of the absolute values was determined.
Next, the obtained hot rolled steel sheet was pickled and then subjected to cold rolling. Rolling was performed to reduce the plate thickness to 1/2 in 7 passes. The line load of each pass was measured and the total value was calculated. Note that the nozzle angle and spray amount of the lubricant were controlled so that the lubrication conditions were uniform for all conditions and all passes. A synthetic ester lubricant (FR-160 manufactured by Nippon Parkerizing Co., Ltd.) was used as the lubricant.
The results are shown in Table 6.

The steel plate (No. 2-1) obtained by rolling Steel 2-a under Condition 2-i had almost no local strain in the F phase, and also had a very high P fraction, so the rolling wire load was relatively low. The value was high.
A steel plate (No. 2-2) obtained by rolling Steel 2-a under Condition 2-ii had an average absolute value of local strain in the F phase of 0.050 or more. It is clear that it can be cold rolled with a slightly lower wire load than the steel plate (No. 2-1) rolled under condition 2-i to have F phase and P structure, and has excellent cold rollability. became.
On the other hand, although the steel plate (No. 2-3) rolled under condition 2-iii exhibits a microstructure mainly consisting of the B structure, the average absolute value of local strain in the F phase does not reach 0.050. First, the rolling wire load was significantly higher than that of No. 2-1 and No. 2-2, and it was found that the cold rolling properties were inferior. One of the reasons is considered to be that CR2 is outside the scope of the present invention.

A steel plate (No. 2-5) obtained by rolling Steel 2-b under Condition 2-iv had an average absolute value of local strain in the F phase of 0.050 or more. It became clear that the steel sheet could be rolled with a slightly lower wire load than the steel plate (No. 2-4) rolled under condition 2-i to have F phase and P structure, and had excellent cold rollability. .
On the other hand, although the steel plate (No. 2-6) rolled under condition 2-v exhibits a microstructure mainly consisting of the B structure, the average absolute value of local strain in the F phase does not reach 0.050. First, the rolling wire load was significantly higher than that of No. 2-4 and No. 2-5, and it was found that the cold rolling properties were inferior. One of the reasons is considered to be that CR2 is outside the scope of the present invention.

A steel plate (No. 2-8) obtained by rolling Steel 2-c under Condition 2-vi had an average absolute value of local strain in the F phase of 0.050 or more. It became clear that the steel sheet could be rolled with a slightly lower wire load than the steel plate (No. 2-7) rolled under condition 2-i to have F phase and P structure, and had excellent cold rollability. .
On the other hand, although the steel plate (No. 2-9) rolled under condition 2-vii has a microstructure mainly composed of the B structure, the average absolute value of local strain in the F phase does not reach 0.050. First, the rolling wire load was significantly higher than that of No. 2-7 and No. 2-8, and it was found that the cold rolling properties were inferior. One of the reasons is considered to be that CR2 is outside the scope of the present invention.

According to the present invention, it is possible to lower only the yield point without changing the tensile strength of the steel plate with high productivity, and further, by lowering the winding temperature, it can be used as a high-strength steel plate as well as a steel plate with a low yield ratio. can do. Therefore, the present invention has high industrial applicability.

Claims (4)

In mass%,
C: 0.05-0.50%,
Si: 0.01-2.0%,
Mn: 0.5-3.0%
Contains
P: 0.03% or less,
S: 0.02% or less,
N: 0.05% or less,
Al: has chemical components each limited to 0.05% or less, with the remainder consisting of Fe and unavoidable impurities,
The microstructure contains a bainite structure of 50% or more and a ferrite phase of 15 to 30% in terms of area ratio, and the average grain size of the ferrite phase is 30 μm or less,
The average absolute value of local strain in the ferrite phase determined by the fine lattice marker method is 0.050 or more, the tensile strength is 440 MPa or more , and the average absolute value of the local strain is 300 ° C. for 2 hours. A high-strength hot-rolled steel sheet, characterized in that the strain is determined based on a numerical value of strain determined by comparing changes in fine lattice markers before and after heating. In addition, the steel plate is in mass%,
Ti: 0.1% or less,
Nb: 0.1% or less,
B: 0.01% or less,
Cr: 1.5% or less,
Cu: 1.0% or less,
The high-strength hot-rolled steel sheet according to claim 1, containing one or more of Ni: 1.0% or less. In addition, the steel plate is in mass%,
Mo: 0.01-1.0%,
W: 0.01-0.5%,
V:0.01~0.5%
The high-strength hot-rolled steel sheet according to claim 1 or 2, characterized in that it contains one or more of the following. A method for producing a high-strength hot-rolled steel sheet according to any one of claims 1 to 3, comprising: directly casting the steel having the chemical composition according to any one of claims 1 to 3; Alternatively, hot rolling is performed by reheating to 1300°C or lower, and in the hot rolling, finish rolling is performed with a cumulative reduction rate of 50% or more that is completed at ₃ or more Ar points, and after hot rolling, the first stage of cooling is performed. Then, in the second stage of cooling, at an average cooling rate of 15°C/s or less, (the bainite structure is 50% or more and the austenite phase is formed). 35%) to (100 seconds after the point when the bainite structure is 50% or more and the austenite phase is 20%), and then as a third stage of cooling. A method for producing a high-strength hot-rolled steel sheet, the method comprising: cooling to 300°C or less at an average cooling rate of 50°C/second or more until the temperature reaches 300°C, and then winding the sheet.

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