WO2016157896A1

WO2016157896A1 - Hot-rolled steel sheet and method for producing same

Info

Publication number: WO2016157896A1
Application number: PCT/JP2016/001834
Authority: WO
Inventors: 田中　孝明; 太郎木津; 俊介豊田
Original assignee: Ｊｆｅスチール株式会社
Priority date: 2015-04-01
Filing date: 2016-03-30
Publication date: 2016-10-06
Also published as: CN107429362A; JP6075517B1; US20180119240A1; MX2017012493A; EP3279353B1; EP3279353A1; KR20170128555A; EP3279353A4; JPWO2016157896A1; KR101989262B1; CN107429362B

Abstract

The present invention has a prescribed component composition and is fashioned into a structure in which: the total surface area of a tempered bainite phase and a tempered martensite phase is 70% or greater; the total surface area of a coarse pearlite phase, a martensite phase, and a residual austenite phase is 10% or less; the tempered bainite phase and the tempered martensite phase have lath averaging 1.0 μm or less in width as a subsidiary structure; the proportion of material having an aspect ratio of 5 or less in Fe carbide precipitated inside the lath and at the lath boundaries is 80% or greater; and MC-type carbide having an average particle diameter of 20 nm or less is dispersed and precipitated inside the lath and at the lath boundaries. Furthermore, the average dislocation density is 1.0 × 10¹⁴m^-2 to 5.0 × 10¹⁵m^-2, inclusive.

Description

Hot-rolled steel sheet and manufacturing method thereof

The present invention has a high tensile strength (TS) of 780 MPa or more, excellent stretch flangeability and punching, suitable for structural steel materials such as automobile parts and other transportation machinery and construction steel materials. Further, the present invention relates to a hot-rolled steel sheet having excellent properties and also excellent in production stability and a method for producing the same.

In order to reduce CO ₂ emissions from the viewpoint of global environmental conservation, maintaining the strength of the car body while reducing its weight and improving the fuel efficiency of the car has always been an important issue in the automobile industry. . In order to reduce the weight of the vehicle body while maintaining the strength of the automobile body, it is effective to reduce the thickness of the steel sheet by increasing the strength of the steel sheet used as a material for automobile parts. For example, in an automobile undercarriage part that often uses a thick steel plate, a significant reduction in weight can be expected by reducing the thickness by increasing the strength.

Generally, automobile underbody parts such as a lower arm are formed by burring, so that a steel plate having excellent stretch flangeability is required. Many researches and developments have been made on hot-rolled steel sheets that have both strength and workability, and various technologies have been proposed. For example, it is known that both a high tensile strength and excellent stretch flangeability can be achieved by making the metal structure substantially a ferrite single phase and precipitating fine carbides in the grains of the ferrite phase.

As such a technique, Patent Document 1 discloses that a steel sheet structure is a ferrite single-phase structure having a low dislocation density and excellent workability, and further, fine carbides are dispersed and precipitated in ferrite to strengthen precipitation, thereby hot rolling. A steel sheet having improved strength while maintaining stretch flangeability of the steel sheet is disclosed.

On the other hand, the normal burring process is performed using a steel sheet punched into a predetermined shape. During actual mass production of parts, the punching clearance usually changes due to die temperature rise or die wear due to continuous pressing. May occur. For these reasons, there is a demand for a steel sheet that always has excellent punchability against fluctuations in punching conditions.

As such a steel sheet, for example, Patent Document 2 discloses that Fe-type carbides having a bainite phase exceeding 92% by volume and an average interval of bainite laths of 0.60 μm or less and precipitated in grains among all Fe-type carbides. A high-strength hot-rolled steel sheet with improved mass production punchability by making the number ratio of 10% or more is disclosed.

JP 2012-26034 A Japanese Unexamined Patent Publication No. 2014-205888

The steel sheet described in Patent Document 1 achieves both high strength and excellent stretch flangeability. However, since the steel sheet structure is substantially a ferrite single phase, there is almost no inclusion that becomes a starting point of voids when the steel sheet is punched. For this reason, in the steel plate described in Patent Document 1, when the conditions such as the clearance and the plate pressing change, the punched end surface may be roughened.

Further, the steel sheet described in Patent Document 2 has a steel sheet structure mainly composed of a predetermined bainite by controlling hot rolling conditions, thereby obtaining excellent punchability. However, such a bainite structure is characterized in that mechanical properties such as tensile strength tend to fluctuate with respect to fluctuations in the coiling temperature. In general, it is not easy to make the steel plate temperature uniform over the entire length of the coil in cooling after hot rolling. For this reason, in the steel plate described in Patent Document 2, the variation in mechanical properties of the steel plate becomes large, and the production stability may be lowered.

The present invention has been developed in view of the above situation, has a tensile strength (TS): high strength of 780 MPa or more, has excellent stretch flangeability and punchability, and further has manufacturing stability. An object of the present invention is to provide a hot-rolled steel sheet that is excellent in combination with its advantageous production method.

The present inventors diligently studied a method for increasing the strength of a steel sheet while maintaining workability, particularly stretch flangeability, and having excellent punchability and suppressing variations in mechanical properties with respect to variations in manufacturing conditions.
As described above, in order to improve the stretch flangeability of the steel sheet, it is effective to make the strength in the metal structure uniform. As such a method, a method of increasing strength by solid solution strengthening or precipitation strengthening as a ferrite single phase structure, or a method of increasing strength by strengthening structure as a bainite single phase structure can be considered. However, in steel sheets with a ferrite single-phase structure, there is almost no inclusion that becomes the starting point of voids when punching the steel sheet, and therefore there is a risk that the punched end face will be rough when conditions such as clearance and sheet pressing change. is there.

On the other hand, a steel sheet having a bainite single-phase structure has excellent stretch flangeability. In addition, a steel sheet having a bainite single-phase structure has a large number of Fe-based carbides in the bainite structure, which serves as a starting point for voids at the time of punching, and thus has excellent punchability. However, since the mechanical properties such as strength greatly vary depending on the transformation temperature in the bainite structure, there is a concern that the variation in mechanical properties with respect to variations in hot rolling conditions such as the coiling temperature will increase.

Therefore, the present inventors considered to reduce the influence of fluctuations in hot rolling conditions by adding a tempering treatment to a structure mainly composed of bainite and martensite.
In general, by applying a tempering treatment to a bainite or martensite structure, variations in mechanical properties due to changes in hot rolling conditions are greatly reduced, but at the same time the steel sheet strength is greatly reduced. Moreover, since the form of the Fe-based carbide in the tempered bainite or tempered martensite phase varies depending on the annealing conditions, the steel sheet does not necessarily have excellent punchability depending on the annealing conditions.

For this reason, the present inventors suppressed the reduction of the steel sheet strength as described above when adding a tempering treatment to the structure mainly composed of bainite and further martensite, and excellent stretch flangeability and punchability. We studied earnestly about the method of combining the two.
As a result, MC type carbides such as TiC are dispersed and precipitated inside the lath and within the lath boundary, which suppresses coarsening of the lath during annealing and further disappearance of the lath due to recovery, and maintains high steel sheet strength even after annealing. I found out that I can do it. In addition, it has been found that excellent punchability can be obtained by ensuring a certain ratio of Fe-based carbides precipitated in the lath and at the lath boundary with an aspect ratio of 5 or less.
Further, the inventors have further studied, in particular, adding 0.03% or more of Ti, and knowing that the steel sheet structure can be stably made the above structure by optimizing the thermal history in the annealing process. It came to.
In addition, MC type carbide is TiC, NbC, VC, (Ti where the atomic ratio of M element (M element includes Ti, Nb, V, Mo, etc.) and C is approximately 1: 1. , Mo) C and other carbides. Here, the M element does not have to be a single type, and may be a composite carbide containing a plurality of metal elements. Also, N-containing carbonitrides and composite carbonitrides may be used.

In addition, the present inventors have further studied earnestly, by appropriately controlling the thermal history when cooling from the highest heating temperature to room temperature in the annealing process, the remaining structure other than the tempered martensite phase and the tempered bainite phase, In particular, the inventors have found that the formation of martensite phase, coarse pearlite phase and retained austenite phase is suppressed, and that, in addition to high strength and excellent punchability, it can also have excellent stretch flangeability.
The present invention was completed after further studies based on the above findings.

That is, the gist configuration of the present invention is as follows.
1. % By mass
C: 0.03% or more and 0.20% or less, Si: 0.4% or less,
Mn: 0.5% to 2.0%, P: 0.03% or less,
S: 0.03% or less, Al: 0.1% or less,
N: 0.01% or less and Ti: 0.03% or more and 0.15% or less, with the balance consisting of Fe and inevitable impurities,
The total area ratio of the tempered bainite phase and the tempered martensite phase is 70% or more, and the total area ratio of the coarse pearlite phase, the martensite phase and the retained austenite phase is 10% or less,
The tempered bainite phase and the tempered martensite phase have lath having an average width of 1.0 μm or less as a substructure, and among Fe-based carbides precipitated in the lath and at the lath boundary, the proportion of those having an aspect ratio of 5 or less 80% or more, and MC type carbide having an average particle size of 20 nm or less dispersed and precipitated inside the lath and at the lath boundary,
A hot-rolled steel sheet having an average dislocation density of 1.0 × 10 ¹⁴ m ⁻² or more and 5.0 × 10 ¹⁵ m ⁻² or less.

2. As the composition, further in mass%,
V: 0.01% to 0.3%,
2. The hot rolled steel sheet according to 1 above, containing at least one or more of Nb: 0.01% to 0.1% and Mo: 0.01% to 0.3%.

3. As the composition, further in mass%,
B: 0.0002% to 0.010%,
The hot-rolled steel sheet according to 1 or 2, comprising:

4). As the composition, further, by mass%, one or more of REM, Zr, As, Cu, Ni, Sn, Pb, Ta, W, Cr, Sb, Mg, Ca, Co, Se, Zn and Cs The hot-rolled steel sheet according to any one of 1 to 3 above, containing 1.0% or less in total.

5. The hot-rolled steel sheet according to any one of 1 to 4, which has a plating layer on the surface thereof.

6). The steel material having the composition according to any one of 1 to 4 above is heated to an austenite single phase region, subjected to hot rolling consisting of rough rolling and finish rolling, and after the finish rolling is finished, A hot rolling process for cooling and winding,
After the hot rolling step, pickling the steel sheet, and then having a continuous annealing step for continuous annealing,
In the hot rolling step, the finish rolling temperature is 850 ° C. or more and 1000 ° C. or less, after the finish rolling is finished, the average cooling rate to 500 ° C. is 30 ° C./s or more, and the winding temperature is 500 ° C. or less,
In the continuous annealing step,
The maximum heating temperature of the steel sheet is 700 ° C. or more (A ₃ points + A ₁ point) / 2 or less,
The time when the steel plate temperature when heating the steel plate to the maximum heating temperature is 600 ° C. or more and 700 ° C. or less is 20 seconds or more and 1000 seconds or less,
The time when the steel sheet temperature is over 700 ° C. is set to 200 seconds or less,
The average cooling rate up to 530 ° C when cooling the steel sheet from the maximum heating temperature is 8 ° C / s or more and 25 ° C / s or less, and after the cooling is stopped, the holding time in the temperature range of 470 ° C or more and 530 ° C or less is set. A method for producing a hot-rolled steel sheet for 10 seconds or more.

7). The method for producing a hot-rolled steel sheet according to 6, further comprising a step of performing a plating treatment after the continuous annealing step.

According to the present invention, tensile strength (TS): high strength of 780 MPa or more, excellent stretch flangeability, suitable for structural steel materials such as automobile machinery and other transportation machinery parts and construction steel materials. Thus, a hot-rolled steel sheet that has both punchability and suppresses variations in mechanical properties with respect to fluctuations in manufacturing conditions can be obtained. Thereby, the further use expansion | deployment of a hot-rolled steel plate is attained, and there exists a remarkable effect on industry.

FIG. 2 is a schematic diagram showing an example of a structure in which a tempered bainite phase and a tempered martensite phase have lath as a substructure, and Fe-based carbides are precipitated and lath MC-type carbides are dispersed and precipitated in the lath inside and lath boundaries.

Hereinafter, the present invention will be specifically described.
First, the component composition in the hot-rolled steel sheet of the present invention will be described. The unit of element content in the component composition is “mass%”, but hereinafter, it is simply indicated by “%” unless otherwise specified.

C: 0.03% to 0.20% C improves the strength of the steel and promotes the formation of bainite and martensite during hot rolling. Therefore, in the present invention, the C content needs to be 0.03% or more. On the other hand, when the C content exceeds 0.20%, the carbon equivalent becomes excessively large and the weldability of the steel sheet is deteriorated. Therefore, the C content is set to 0.03% or more and 0.20% or less. Preferably, it is 0.04% or more and 0.18% or less, more preferably more than 0.05% and 0.15% or less.

Si: 0.4% or less Si is usually positively contained in a high-strength steel sheet as an effective element for improving the steel sheet strength without reducing ductility (elongation). However, if the Si content exceeds 0.4%, an oxide is formed on the surface of the steel sheet during heat treatment, which causes deterioration of plating adhesion. Therefore, the Si content is 0.4% or less. Preferably it is 0.3% or less, More preferably, it is 0.2% or less. Note that Si may be added to an impurity level, or may be O%.

Mn: 0.5% or more and 2.0% or less Mn is an element that contributes to increasing the strength of steel by solid solution. Mn is an element that promotes the formation of bainite and martensite during hot rolling by improving hardenability. In order to obtain such an effect, the Mn content needs to be 0.5% or more. On the other hand, if the Mn content exceeds 2.0%, the austenite becomes excessively stable, and the steel sheet has a structure that excessively contains martensite and residual austenite. Therefore, stretch flangeability deteriorates. Therefore, the Mn content is 0.5% or more and 2.0% or less. In addition, Preferably they are 0.8% or more and 1.8% or less, More preferably, they are 1.0% or more and 1.7% or less.

P: 0.03% or less P is a harmful element that segregates at grain boundaries to reduce elongation, induces cracking during processing, and further degrades impact resistance. Therefore, the P content is 0.03% or less. However, excessive P removal leads to an increase in refining time and cost, so the P content is preferably 0.002% or more.

S: 0.03% or less S is present in steel as MnS or TiS, and promotes the generation of voids during the punching of hot-rolled steel sheets. Moreover, since S becomes a starting point of generation of voids during processing, it reduces stretch flangeability. Therefore, it is preferable to reduce the S content as much as possible, and the content is 0.03% or less. Preferably it is 0.01% or less. However, since excessive desulfurization leads to an increase in refining time and cost, the S content is preferably 0.0002% or more.

Al: 0.1% or less Al is an element that acts as a deoxidizer. In order to obtain such an effect, it is preferable to contain 0.01% or more of Al. However, if Al exceeds 0.1%, it remains as an Al oxide in the steel sheet, and the Al oxide tends to aggregate and become coarse and deteriorate stretch flangeability. Therefore, the Al content is 0.1% or less.

N: 0.01% or less N is present as coarse TiN in the steel, and promotes the generation of coarse voids during the punching of hot-rolled steel sheets. Moreover, since N becomes a starting point of generation of coarse voids during processing, it reduces stretch flangeability. For this reason, it is preferable to reduce N content as much as possible, and set it as 0.01% or less. Preferably it is 0.006% or less. However, excessive N removal leads to an increase in refining time and cost, so the N content is preferably 0.0005% or more.

Ti: 0.03% or more and 0.15% or less Ti is an indispensable element for increasing the strength of a steel sheet by forming MC-type carbides and suppressing lath coarsening in the annealing process. Moreover, MC type carbide increases the strength of the steel sheet by precipitation strengthening. Here, when the Ti content is less than 0.03%, these effects cannot be sufficiently obtained. For this reason, the strength of the steel sheet decreases due to the coarsening of the lath and the decrease in the amount of precipitation, making it difficult to obtain the desired steel sheet strength (tensile strength: 780 MPa or more). On the other hand, if the Ti content exceeds 0.15%, the central segregation becomes prominent and causes punching deterioration. Therefore, the Ti content is set to 0.03% or more and 0.15% or less. Preferably they are 0.04% or more and 0.14% or less, More preferably, they are 0.05% or more and 0.13% or less.

The basic components have been described above. In the hot-rolled steel sheet of the present invention, V: 0.01% or more and 0.3% or less, Nb: 0.01% or more and 0.1% or less, and Mo: 0.01 as necessary for the purpose of further increasing the strength. % Or more and 0.3% or less can be contained.

V: 0.01% or more and 0.3% or less V forms MC-type carbides and, like Ti, contributes to increasing the strength of the steel sheet by suppressing coarsening of the lath and precipitation strengthening in the annealing process. In order to acquire such an effect, it is necessary to contain V 0.01% or more. On the other hand, when the V content exceeds 0.3%, the central segregation becomes prominent and causes punching deterioration. Therefore, the V content is preferably 0.01% or more and 0.3% or less. In addition, More preferably, it is 0.01% or more and 0.2% or less, More preferably, it is 0.01% or more and 0.15% or less. Note that V may form MC-type carbides alone, or may form composite carbides with Ti, Nb, and Mo. These carbide compositions have no influence on the effects of the invention.

Nb: 0.01% or more and 0.1% or less Nb forms MC-type carbides and, like Ti, contributes to increasing the strength of the steel sheet by suppressing lath coarsening and precipitation strengthening in the annealing process. In order to obtain such an effect, it is necessary to contain 0.01% or more of Nb. On the other hand, even if Nb is added excessively in excess of 0.1%, the effect is saturated because it does not form a solid solution in the heating furnace at the time of hot rolling, and this causes unnecessarily high alloy costs. Therefore, the Nb content is preferably 0.01% or more and 0.1% or less. In addition, More preferably, it is 0.01% or more and 0.08% or less, More preferably, it is 0.01% or more and 0.06% or less. Nb may form MC-type carbides alone, or may form composite carbides with Ti, V, and Mo. These carbide compositions have no influence on the effects of the invention.

Mo: 0.01% or more and 0.3% or less Mo forms MC-type composite carbide by compound addition with Ti, and, like Ti, increases the strength of the steel sheet by suppressing lath coarsening and precipitation strengthening in the annealing process. Contribute. In order to obtain such an effect, it is necessary to contain 0.01% or more of Mo. On the other hand, if the Mo content exceeds 0.3%, the central segregation becomes prominent and causes punching deterioration. Therefore, the Mo content is preferably 0.01% or more and 0.3% or less. Mo may form a composite carbide with Nb or V, but these carbide compositions do not affect the effects of the invention.

Further, the hot-rolled steel sheet of the present invention may contain B: 0.0002% to 0.010% as necessary for the purpose of improving the hardenability during hot rolling.

B: 0.0002% or more and 0.010% or less B is an element that segregates at austenite grain boundaries and suppresses the formation and growth of ferrite to improve hardenability and promote the formation of bainite and martensite. In order to obtain these effects, the B content is preferably 0.0002% or more. However, if the B content exceeds 0.010%, a hard iron boride is formed, which causes stretch flangeability deterioration. Therefore, when it contains B, it is preferable to make the content into 0.0002% or more and 0.010% or less. Further, it is more preferably 0.0002% or more and 0.0050% or less, and further preferably 0.0004% or more and 0.0030% or less.

Furthermore, the hot-rolled steel sheet of the present invention further includes REM, Zr, As, Cu, Ni, Sn, Pb, Ta, W, Cr, Sb, Mg, Ca, Co, Se, Zn, and the above composition. One or more of Cs may be contained in a total of 1.0% or less.
Components other than the above are Fe and inevitable impurities.

Next, the reason for limiting the structure in the hot-rolled steel sheet of the present invention will be described.
Total area ratio of tempered bainite phase and tempered martensite phase: 70% or more The hot-rolled steel sheet of the present invention has a structure mainly composed of tempered bainite and tempered martensite, which has both high strength and excellent punchability. Here, when the sum of the area ratios of the tempered bainite phase and the tempered martensite phase is less than 70%, a hot rolled steel sheet having desired high strength and punchability cannot be obtained. The reason why the tempered bainite phase and the tempered martensite fraction are not individually defined is that the tempered bainite and tempered martensite after annealing have a structure that cannot be distinguished. This is also a major factor that can suppress variations in mechanical properties after annealing even when the manufacturing conditions during hot rolling vary. Accordingly, the total area ratio of the tempered bainite phase and the tempered martensite phase is 70% or more. Preferably it is 75% or more, more preferably 80% or more. Further, the total area ratio of the tempered bainite phase and the tempered martensite phase may be 100%.

Total area ratio of coarse pearlite phase, martensite phase and residual austenite phase: 10% or less As described above, the hot-rolled steel sheet of the present invention has a structure mainly composed of tempered bainite and tempered martensite. Examples of the remaining structure other than martensite include Fe carbide, coarse pearlite, fine pearlite, pseudo pearlite, bainite, martensite, and retained austenite. Among these structures, especially when coarse pearlite, martensite, and retained austenite are present in the metal structure, stretch flangeability is significantly deteriorated. Therefore, the sum of the area ratios of the coarse pearlite phase, martensite phase and residual austenite phase is set to 10% or less. In addition, Preferably it is 8% or less, More preferably, it is 5% or less. Further, it may be 0%.
Here, it is assumed that coarse pearlite has a lamella spacing of 0.2 μm or more, fine pearlite has a lamella spacing of less than 0.2 μm, and pseudo pearlite does not clearly observe a pearlite lamella. The lamella spacing can be obtained by observing the structure with a scanning electron microscope.

Further, examples of the remaining structure other than the tempered bainite phase, the tempered martensite phase, the coarse pearlite phase, the martensite phase, and the retained austenite phase include a ferrite phase, a pseudo pearlite phase, and a fine pearlite phase. Such a remaining structure is acceptable if the total area ratio is 30% or less.

The average width of the lath that the tempered bainite phase and the tempered martensite phase have as a substructure: 1.0 μm or less In order to increase the strength by the tempered bainite phase and the tempered martensite phase, the average width of these substructures: 1.0 μm or less It is important to have a fine lath. FIG. 1 is a schematic diagram showing an example of a structure in which a tempered bainite phase and a tempered martensite phase have lath as a substructure, and Fe-based carbides precipitate and MC-type carbides disperse and precipitate in the lath inside and lath boundaries. The figure is shown. Here, when the lath disappears due to recovery or the average width of the lath exceeds 1.0 μm, a predetermined high strength cannot be achieved. Therefore, the average width of the laths that the tempered bainite phase and the tempered martensite phase have as a substructure is 1.0 μm or less. Preferably it is 0.8 micrometer or less, More preferably, it is 0.6 micrometer or less. The lower limit is not particularly limited, but is usually about 0.1 μm.

Percentage of Fe-based carbides precipitated in the lath and lath boundary with an aspect ratio of 5 or less: 80% or more As shown in Fig. 1, Fe-based carbides deposited in the lath and lath boundary are the origin of voids during punching This contributes to improved punchability. In particular, an Fe-based carbide having an aspect ratio of 5 or less has a large effect, and an excellent punchability can be exhibited by setting the ratio to 80% or more. Therefore, the proportion of Fe-based carbides precipitated in the lath and at the lath boundary is set to 80% or more with an aspect ratio of 5 or less. Preferably it is 85% or more. The upper limit is not particularly limited and may be 100%.
Here, the Fe-based carbide is θ carbide (cementite) or ε carbide. The alloy element may be dissolved in the carbide. The aspect ratio is the ratio of the length of the major axis to the minor axis of the Fe-based carbide precipitated in the lath and at the lath boundary.

MC type carbide dispersed and precipitated in the lath and lath boundary: 20 nm or less As shown in Fig. 1, MC type carbide finely dispersed and precipitated in the lath and lath boundary has a pinning effect during annealing of the steel sheet. By suppressing the coarsening of the lath and further suppressing the disappearance of the lath due to the recovery, it contributes to an increase in strength. However, when the average particle diameter of MC type carbide exceeds 20 nm, the number of MC type carbide particles contributing to pinning is insufficient, the pinning effect is insufficient, and the steel sheet strength is reduced. On the other hand, when the average particle size of the MC type carbide is 20 nm or less, a sufficient number of MC type carbides exhibit a pinning effect and suppress a decrease in steel sheet strength. Therefore, the average particle diameter of MC type carbides dispersed and precipitated in the lath and lath boundaries of the tempered bainite phase and the tempered martensite phase is set to 20 nm or less. Preferably it is 15 nm or less. The lower limit is not particularly limited, but is usually about 1 nm. However, the proportion of MC type carbide having a particle size exceeding 50 nm is preferably 10% or less.

In the hot rolled steel sheet of the present invention, it is important that the average dislocation density is in the following range.
Average dislocation density: 1.0 × 10 ¹⁴ m ⁻² or more and 5.0 × 10 ¹⁵ m ⁻² or less In the hot-rolled steel sheet of the present invention, the steel sheet having a bainite and martensite structure is tempered to vary the variation in hot rolling conditions. Is reduced. Here, when the average dislocation density of the steel sheet after annealing exceeds 5.0 × 10 ¹⁵ m ⁻² , the tempering of the steel sheet is insufficient, and the influence of fluctuations in hot rolling conditions cannot be sufficiently mitigated. On the other hand, even when fully tempered, the average dislocation density is usually 1.0 × 10 ¹⁴ m ⁻² or more. Therefore, the average dislocation density is 1.0 × 10 ¹⁴ m ⁻² or more and 5.0 × 10 ¹⁵ m ⁻² or less. Preferably, it is 1.0 × 10 ¹⁴ m ⁻² or more and 2.0 × 10 ¹⁵ m ⁻² or less.

Next, the manufacturing method of the hot rolled steel sheet of the present invention will be described.
The method for producing a hot-rolled steel sheet of the present invention was obtained after heating the steel material having the above-described composition to the austenite single-phase region, subjecting it to hot rolling consisting of rough rolling and finish rolling, and finishing the finish rolling. A hot rolling process for cooling and winding the steel sheet;
After the hot rolling step, pickling the steel sheet, and then having a continuous annealing step for continuous annealing,
In the hot rolling step, the finish rolling temperature is 850 ° C. or more and 1000 ° C. or less, after the finish rolling is finished, the average cooling rate to 500 ° C. is 30 ° C./s or more, and the winding temperature is 500 ° C. or less,
In the continuous annealing step,
The maximum heating temperature of the steel sheet is 700 ° C. or more (A ₃ points + A ₁ point) / 2 or less,
The time when the steel plate temperature when heating the steel plate to the maximum heating temperature is 600 ° C. or more and 700 ° C. or less is 20 seconds or more and 1000 seconds or less,
The time when the steel sheet temperature is over 700 ° C. is set to 200 seconds or less,
The average cooling rate up to 530 ° C when cooling the steel sheet from the maximum heating temperature is 8 ° C / s or more and 25 ° C / s or less, and after the cooling is stopped, the holding time in the temperature range of 470 ° C or more and 530 ° C or less is set. 10 seconds or more. Moreover, you may further provide the process of giving a plating process after the said continuous annealing process.

In addition, the melting method of a steel raw material is not specifically limited, Well-known melting methods, such as a converter and an electric furnace, are employable. Moreover, after melting, it is preferable to use a continuous casting method to form a slab (steel material) from the viewpoint of productivity and the like, but as a slab by a known casting method such as an ingot-bundling rolling method or a thin slab continuous casting method. Also good.

Further, the steel material obtained as described above is subjected to hot rolling consisting of rough rolling and finish rolling, and the steel material is heated in the austenite single phase region prior to rough rolling. If the steel material before rough rolling is not heated in the austenite single-phase region, remelting of Ti carbide, etc. present in the steel material will not proceed, and MC type carbide will precipitate finely during annealing after hot rolling. I will not. Therefore, prior to rough rolling, the steel material is heated to an austenite single phase region, preferably 1150 ° C. or higher. The upper limit of the heating temperature is not particularly specified, but when the heating temperature becomes higher than necessary, the yield decreases due to oxidation of the slab surface, so the heating temperature is usually 1350 ° C. or lower. In addition, when hot rolling the steel material, if the steel material (slab) after casting has a temperature in the austenite single-phase region, the steel material is not heated or directly heated after being heated for a short time. You may roll.

Next, the reason for limiting the manufacturing conditions in the hot rolling process will be described.
Finish rolling temperature: 850 ° C. or more and 1000 ° C. or less When the finish rolling temperature is lowered, ferrite transformation is promoted in cooling after rolling, so that the bainite and martensite fractions of the hot-rolled steel sheet obtained after hot rolling are lowered. As a result, the predetermined tempered bainite and tempered martensite fraction cannot be obtained after annealing. Therefore, the temperature on the finish rolling delivery side needs to be 850 ° C. or higher, preferably 880 ° C. or higher. Further, when the finish rolling temperature exceeds 1000 ° C., the surface properties of the steel sheet deteriorate. For this reason, the upper limit of finish rolling temperature shall be 1000 degrees C or less. Preferably it is 970 degrees C or less. Each temperature such as the coiling temperature including the above finish rolling temperature is the temperature of the steel sheet surface.

After finishing rolling, cooling rate to 500 ° C: 30 ° C / s or more When cooling the steel plate after finishing rolling, if the cooling rate is insufficient, ferrite cannot be suppressed sufficiently, and after hot rolling The bainite and martensite fractions of the resulting hot-rolled steel sheet are reduced. As a result, the predetermined tempered bainite and tempered martensite fraction cannot be obtained after annealing. Therefore, the cooling rate to 500 ° C. after finishing rolling needs to be 30 ° C./s or more. Preferably, it is 50 ° C./s or more. The upper limit of the cooling rate is not particularly limited, but is usually about 300 ° C./s.

Winding temperature: 500 ° C. or less Optimization of the winding temperature is important for controlling the steel sheet structure after hot rolling. When the coiling temperature exceeds 500 ° C., the lath width of the bainite becomes large, so that the lath width of the tempered bainite after annealing cannot be set to a predetermined value. On the other hand, the lower limit of the coiling temperature is not particularly limited, but if the coiling temperature is excessively lowered, the cooling cost is only unnecessarily expensive. For this reason, the winding temperature is preferably 0 ° C. or higher. More preferably, it is 200 ° C. or higher.

After the hot rolling step described above, the hot rolled steel sheet is pickled and subjected to a continuous annealing step for continuous annealing. Hereinafter, the reason for limiting the manufacturing conditions in this continuous annealing process will be described.
Maximum heating temperature of steel sheet: 700 ° C or higher (A ₃ points + A ₁ point) / 2 or lower Optimization of the maximum heating temperature of steel sheets in the continuous annealing process is sufficient for the effects of manufacturing conditions fluctuations during hot rolling due to annealing It is important to reduce and achieve the desired high strength. If the maximum heating temperature of the steel sheet is less than 700 ° C, it is difficult to control the dislocation density in bainite and martensite within an appropriate range. For this reason, the influence of fluctuations in manufacturing conditions during hot rolling is sufficiently reduced. Can not. In addition, when the heating temperature of the steel sheet is less than 700 ° C, the aspect ratio of Fe-based carbides inside and between the laths tends to be large, and the proportion of Fe-based carbides with an aspect ratio of 5 or less should be in the desired range Is difficult. On the other hand, when the maximum heating temperature of the steel sheet exceeds (A ₃ points + A ₁ point) / 2, the MC type carbides become prominently coarsened, so that the lath coarsening in bainite and martensite can be sufficiently suppressed. Disappear. Further, by promoting austenitization, the bainite and martensite fractions are lowered, and the desired tempered bainite and tempered martensite fractions cannot be obtained. Therefore, the maximum heating temperature of the steel sheet in the continuous annealing process is 700 ° C. or more (A ₃ points + A ₁ point) / 2 or less. The temperature is preferably 700 ° C. or higher and {(A ₃ points + A ₁ point) / 2} −10 ° C. or lower.
The points A ₁ and A ₃ can be calculated by the following formula.
A ₁ point = 751-26.6 x [% C] + 17.6 x [% Si]-11.6 x [% Mn] + 22.5 x [% Mo] +
233 × [% Nb] −39.7 × [% V] −57 × [% Ti] −895 × [% B] −169 × [% Al]
A ₃ points = 937-476.5 x [% C] + 56 x [% Si]-19.7 x [% Mn] + 38.1 x [% Mo] +
124.8 × [% V] + 136.3 × [% Ti] −19 × [% Nb] + 3315 × [% B]
Here, [% X] means the content of X element in steel (mass%).

Time when the steel plate temperature is 600 ° C or higher and 700 ° C or lower when the steel plate is heated to the maximum heating temperature: 20 seconds or more and 1000 seconds or less It is desirable to appropriately control the heating heat history when heating the steel plate to the maximum heating temperature This is important for imparting high strength and excellent punchability to the steel sheet. As described above, the pinning effect by MC type carbide is used to suppress the coarsening of the lath. In order to exhibit this pinning effect, it is necessary to sufficiently disperse MC type carbides in bainite and martensite before the lath starts coarsening. According to the study by the present inventors, precipitation of MC type carbide begins to occur remarkably at 600 ° C. or higher. On the other hand, the coarsening and disappearance of lath occurs remarkably when the temperature exceeds 700 ° C. Therefore, the coarsening and disappearance of the lath can be suppressed by maintaining the steel plate temperature in the temperature range of 600 ° C. or more and 700 ° C. or less for a certain period of time and sufficiently depositing the MC type carbide. Here, in order to sufficiently precipitate the MC-type carbide, it is necessary to hold for 20 seconds or more in this temperature range. In addition, when the holding time in this temperature range is insufficient, since the coarsening of the lath is started before the MC type carbide is sufficiently precipitated, the pinning effect is not sufficiently exhibited and the lath becomes coarse. . Preferably it is 35 seconds or more, more preferably 50 seconds or more.
On the other hand, when the holding time in the temperature range of 600 ° C. or more and 700 ° C. or less exceeds 1000 seconds, the Fe-based carbides precipitated inside and between the laths re-dissolve, and the prior austenite grain boundaries and packet grain boundaries In other words, there is no Fe-based carbide in the lath and between the laths, which moves to the block grain boundaries and effectively contributes to the improvement of punchability. Therefore, in order to obtain a steel sheet having excellent punchability, the holding time in the temperature range where the steel sheet temperature is 600 ° C. or higher and 700 ° C. or lower needs to be 1000 seconds or shorter. Preferably it is 800 seconds or less, more preferably 500 seconds or less. In addition, the steel plate temperature here is the temperature of the steel plate surface.

Time when the steel plate temperature exceeds 700 ° C .: 200 seconds or less In the temperature range where the steel plate temperature exceeds 700 ° C., lath coarsening occurs remarkably. As described above, in the present invention, the movement of the lath boundary is suppressed and the coarsening of the lath is suppressed by the pinning effect by the MC type carbide finely dispersed and precipitated. And thereby, the steel plate strength is maintained. However, if the holding time in this temperature range becomes excessively long, the coarsening of the lath cannot be suppressed. For this reason, from the viewpoint of preventing the coarsening of the lath, the holding time in the temperature range where the steel plate temperature exceeds 700 ° C. is set to 200 seconds or less. Preferably it is 180 seconds or less, more preferably 150 seconds or less. However, when the time when the steel plate temperature is higher than 700 ° C. is less than 10 seconds, the ductility of the steel plate is somewhat inferior, and therefore it is preferable to set it to 10 seconds or more.

Average cooling rate up to 530 ° C when cooling the steel plate from the maximum heating temperature: 8 ° C / s or more and 25 ° C / s or less Appropriate cooling rate when cooling after heating the steel plate to the maximum heating temperature in the continuous annealing process Controlling is important for obtaining excellent stretch flangeability. In particular, when the average cooling rate up to 530 ° C. is lower than 8 ° C./s, pearlite transformation cannot be suppressed during cooling, and coarse pearlite is produced in a predetermined amount or more. For this reason, stretch flangeability falls. On the other hand, when the average cooling rate is excessively high, it is difficult to maintain for a predetermined time in a temperature range of 470 ° C. to 530 ° C., which will be described later. Therefore, the average cooling rate up to 530 ° C. when cooling the steel plate from the maximum heating temperature is 25 ° C./s or less.

Holding time in the temperature range of 470 ° C or higher and 530 ° C or lower: 10 seconds or longer In the continuous annealing process, the steel plate is controlled and cooled as described above, and then held in the temperature range of 470 ° C or higher and 530 ° C or lower for a certain period of time. It is important to obtain excellent stretch flangeability. Here, when the holding temperature after the cooling is stopped exceeds 530 ° C., coarse pearlite is generated, so that stretch flangeability is deteriorated. On the other hand, when the holding temperature after stopping cooling is lower than 470 ° C., the transformation from austenite to bainite is delayed. As a result, C concentrates in the untransformed austenite region and stabilizes the austenite, so the transformation is not completed. Then, in the subsequent cooling, untransformed austenite is transformed into martensite or remains in the steel sheet structure as retained austenite, so that stretch flangeability is lowered. Further, when the steel plate is held for 10 seconds or more in the temperature range of 470 ° C. or more and 530 ° C. or less, the transformation of most austenite to bainite is completed, and then the martensite fraction generated when cooled is determined to be a predetermined value. Can be reduced to a range. Accordingly, after the controlled cooling is stopped, the holding time in the temperature range from 470 ° C. to 530 ° C. is set to 10 seconds or more. Preferably it is 20 seconds or more, more preferably 30 seconds or more. The upper limit of the holding time is not particularly limited, but is usually 300 seconds or less.
Since the control of the steel sheet structure is completed by holding in the temperature range of 470 ° C. or higher and 530 ° C. or lower, the subsequent cooling conditions are not particularly limited, and may be cooled to room temperature by any cooling method.

In addition, even if the steel sheet is reheated to a temperature of 700 ° C or lower after being held in the temperature range of 470 ° C to 530 ° C, the total holding time in the temperature range of 600 ° C to 700 ° C is 1000 seconds or less. In this case, a desired steel sheet structure can be obtained and there is no problem.
For example, after holding in a temperature range of 470 ° C. or more and 530 ° C. or less, the steel plate may be immersed in a zinc pot to be a hot dip galvanized steel plate, or may be further heat-treated to obtain an alloyed hot dip galvanized steel plate. . In addition to zinc, aluminum or aluminum alloy can be plated for hot dipping.

Further, after the above-described continuous annealing step, temper rolling may be applied to the steel sheet continuously in the annealing line or using another line according to a conventional method.

Further, the hot-rolled steel sheet produced as described above may be separately subjected to electrogalvanizing treatment or hot dip galvanizing. The hot-rolled steel sheet of the present invention is suitable as a steel sheet for automobile undercarriage, and is suitable for press forming performed at normal room temperature, and has excellent heat treatment characteristics. For this reason, the hot-rolled steel sheet produced as described above is also suitable as a warm-formed material steel sheet that is immediately press-formed after heating the steel sheet before pressing from 400 ° C to 700 ° C.

The molten steel having the composition shown in Table 1 was melted and continuously cast by a generally known method to obtain a slab (steel material) having a thickness of 300 mm. These slabs are heated to the temperatures shown in Table 2, roughly rolled, and finish rolling is completed at the temperatures shown in Table 2. After finishing rolling, the slabs are cooled at the average cooling rate shown in Table 2, and the windings shown in Table 2 are used. The coil was wound at the coiling temperature to obtain a hot rolled steel sheet having a thickness of 3.2 mm. Furthermore, these hot-rolled steel sheets were pickled by a generally known technique and annealed under the conditions shown in Table 2 in a continuous annealing line. Moreover, about some steel plates, the hot dip galvanization process and also the alloying process were performed in the continuous annealing line, and it was set as the hot dip galvanized steel plate and the galvannealed steel plate.

Test specimens were collected from the hot-rolled steel sheets thus obtained, and subjected to structure observation, measurement of average dislocation density, tensile test, hole expansion test, punching test, and production stability evaluation. The evaluation results are shown in Table 3. The test method was as follows.

(I) Microstructure observation A test piece was taken from the obtained hot-rolled steel sheet, a cross section (L cross section) parallel to the rolling direction of the test piece was polished and corroded with nital, and then a scanning electron microscope (magnification: 1000, 3000 and 5000 times), and using an image analyzer, the total area ratio of the tempered bainite phase and the tempered martensite phase, the area ratio of the coarse pearlite phase, the martensite phase and the residual austenite phase (MA) The sum of the area ratios and the area ratios of the other phases were determined. Although it is difficult to distinguish the martensite phase from the retained austenite phase with a scanning electron micrograph, the sum of the area ratios of the coarse pearlite phase, martensite phase and retained austenite phase is important here. Therefore, the area ratio of the sum of the martensite phase and the retained austenite phase (MA) was determined without particularly distinguishing between the martensite phase and the retained austenite phase.

In addition, the thin film produced from the hot-rolled steel sheet is observed with a transmission electron microscope, and the lath width of tempered bainite and tempered martensite is measured. The average particle size of the MC type carbides precipitated inside the lath and at the lath boundary was determined.
Here, for the measurement of lath width of tempered bainite and tempered martensite, transmission electron micrographs of 120 mm x 80 mm in size taken at 10 fields of view at a magnification of 30000 times were measured for three or more consecutive laths. Draw 5 straight lines at an interval of 10mm perpendicular to the long axis, measure the length of each line segment that intersects the lath boundary, and calculate the average length of the obtained line segments as the average lath width. It was.
The proportion of Fe-based carbides deposited inside the lath and at the lath boundary with an aspect ratio of 5 or less is a minimum of 100 in the total of 5 fields deposited inside the lath and at the lath boundary using photographs taken at a magnification of 165000 times. By measuring the lengths of the major axis and minor axis for each Fe-based carbide and calculating the aspect ratio, the ratio of those having an aspect ratio of 5 or less was determined.
In addition, the average particle size of MC type carbide was measured using a photograph taken at a magnification of 300,000 times, and the diameter of MC type carbides such as TiC of at least 100 total in 5 fields was measured. The diameter d _def ) was obtained. The lower limit of the measured particle diameter is 2 nm.

(Ii) Measurement of average dislocation density Samples were taken from the obtained hot-rolled steel sheet, the transition density at 1/4 part thickness was measured, and the dislocation density at 1/4 part thickness was the average dislocation of the steel sheet. The measured value was regarded as the average dislocation density. The collected specimen was polished with 0.1 mm of oxalic acid in addition to mechanical grinding, so that the sample was adjusted so that a 1/4 part thickness was exposed on the surface. Here, the polishing with oxalic acid was performed in order to remove the processed layer by grinding.
About the sample adjusted as mentioned above, the distortion | strain of the steel plate was measured with the X-ray-diffraction apparatus. The measurement chart uses an X-ray diffractometer to measure the diffraction intensity of the (110), (211), and (220) planes of 1 / 4-thick α-iron using CoKα rays. Then, the half width of the peak value of the reflection intensity of each crystal plane is obtained, and the local strain ε ′ applied to the steel sheet is determined by the following equations (1) and (2).
βcosθ / λ = 0.9 / D + 2ε'sinθ / λ (1)
here,
β: Half width of peak value (however, the value corrected by equation (2) is used)
θ: diffraction angle λ: wavelength of CoKα ray (0.1790 nm)
D: Crystallite size (dislocation cell, crystal grain size)
ε ′: Local strain β ² = β _m ² −β ₀ ² (2)
here,
β _m : half width of the peak of the sample whose dislocation density is measured β ₀ : half width of the peak of the sample without strain.
Note that βcos θ / λ is plotted against sin θ / λ, and ε ′ and D are obtained from the slope and intercept. The dislocation density ρ is determined from the obtained local strain ε ′ by the following equation (3).
ρ = 14.4ε ' ² / b ² (3)
here,
b: Burgers vector (0.248nm)
It is.

(Iii) Tensile test From the obtained hot-rolled steel sheet, a JIS No. 5 tensile test piece (JIS Z 2001) with the direction perpendicular to the rolling direction (C direction) as the tensile direction was collected and complied with the provisions of JIS Z 2241 The tensile test was performed, and the yield strength (YS), tensile strength (TS), and elongation (El) were measured.

(Iv) Hole expansion test A test piece (size: 100 mm x 100 mm) was taken from the obtained hot-rolled steel sheet, and a hole with an initial diameter d ₀ : 10 mmφ was punched into the test piece (clearance: test piece thickness) 12.5%). Using these test pieces, a hole expansion test was performed. That is, a conical punch having an apex angle of 60 ° is inserted from a punch side at the time of punching into a hole having an initial diameter d _{0 of} 10 mmφ, the hole is expanded, and the diameter of the hole when a crack penetrates a steel plate (test piece). d (mm) was measured, and the hole expansion ratio λ (%) was calculated by the following formula.
Hole expansion ratio λ (%) = {(d−d ₀ ) / d ₀ } × 100
Here, when the tensile strength (TS) × {hole expansion rate (λ)} ^0.5 is 6200 · MPa% ^0.5 or more, it was determined that the stretch flangeability was good.

(V) Punching test A test piece (size: 30 mm x 30 mm) is taken from the obtained hot-rolled steel sheet, and a hole with a diameter d ₀ : 10 mmφ is punched into the test piece (clearance: 20% of the test piece plate thickness) 30%). After punching, the fracture condition of the punched end face was observed with a microscope (50 times magnification) over the entire circumference of the punch hole, and the presence of cracks, chips and brittle fracture surfaces was observed. The punching property was evaluated by making a crack, a chip, and a brittle fracture surface not good (good), and other things not good (fail).

From Table 3, in all the examples of the present invention, a hot-rolled steel sheet having a high tensile strength (TS): 780 MPa or more and having both excellent stretch flangeability and punchability is obtained.

In addition, in order to evaluate the variation in mechanical properties of the steel sheet, JIS No. 5 tensile test piece (JIS Z 2001) with the perpendicular direction (C direction) as the tensile direction is arbitrarily selected from the full length of the hot rolled steel sheet as an example of the present invention. The
100 samples were collected, subjected to a tensile test in accordance with the provisions of JIS Z 2241, measured for tensile strength (TS), and their standard deviation σ was obtained. As a result, in all of the inventive examples, the standard deviation of the tensile strength (TS) was within 10 MPa.
As described above, it can be said that all of the examples of the present invention have small variations in the mechanical properties of the steel sheet such as tensile strength (TS) and are excellent in manufacturing stability.

Claims

% By mass
C: 0.03% or more and 0.20% or less, Si: 0.4% or less,
Mn: 0.5% to 2.0%, P: 0.03% or less,
S: 0.03% or less, Al: 0.1% or less,
N: 0.01% or less and Ti: 0.03% or more and 0.15% or less, with the balance consisting of Fe and inevitable impurities,
The total area ratio of the tempered bainite phase and the tempered martensite phase is 70% or more, and the total area ratio of the coarse pearlite phase, the martensite phase and the retained austenite phase is 10% or less,
The tempered bainite phase and the tempered martensite phase have lath having an average width of 1.0 μm or less as a substructure, and among Fe-based carbides precipitated in the lath and at the lath boundary, the proportion of those having an aspect ratio of 5 or less 80% or more, and MC type carbide having an average particle size of 20 nm or less dispersed and precipitated inside the lath and at the lath boundary,
A hot-rolled steel sheet having an average dislocation density of 1.0 × 10 14 m −2 or more and 5.0 × 10 15 m −2 or less.
As the composition, further in mass%,
V: 0.01% to 0.3%,
The hot-rolled steel sheet according to claim 1, comprising at least one or more of Nb: 0.01% to 0.1% and Mo: 0.01% to 0.3%.
As the composition, further in mass%,
B: 0.0002% to 0.010%,
The hot-rolled steel sheet according to claim 1 or 2, comprising:
As the composition, further, by mass%, one or more of REM, Zr, As, Cu, Ni, Sn, Pb, Ta, W, Cr, Sb, Mg, Ca, Co, Se, Zn and Cs The hot rolled steel sheet according to any one of claims 1 to 3, which contains 1.0% or less in total.
The hot-rolled steel sheet according to any one of claims 1 to 4, which has a plating layer on the surface thereof.
A steel material obtained by heating the steel material having the composition according to any one of claims 1 to 4 to an austenite single phase region, subjecting the steel material to hot rolling consisting of rough rolling and finish rolling, and finishing the finish rolling. A hot rolling step of cooling and winding up,
After the hot rolling step, pickling the steel sheet, and then having a continuous annealing step for continuous annealing,
In the hot rolling step, the finish rolling temperature is 850 ° C. or more and 1000 ° C. or less, after the finish rolling is finished, the average cooling rate to 500 ° C. is 30 ° C./s or more, and the winding temperature is 500 ° C. or less,
In the continuous annealing step,
The maximum heating temperature of the steel sheet is 700 ° C. or more (A 3 points + A 1 point) / 2 or less,
The time when the steel plate temperature when heating the steel plate to the maximum heating temperature is 600 ° C. or more and 700 ° C. or less is 20 seconds or more and 1000 seconds or less,
The time when the steel sheet temperature is over 700 ° C. is set to 200 seconds or less,
The average cooling rate up to 530 ° C when cooling the steel sheet from the maximum heating temperature is 8 ° C / s or more and 25 ° C / s or less, and after the cooling is stopped, the holding time in the temperature range of 470 ° C or more and 530 ° C or less is set. A method for producing a hot-rolled steel sheet for 10 seconds or more.
The method for producing a hot-rolled steel sheet according to claim 6, further comprising a step of performing a plating treatment after the continuous annealing step.