KR101989262B1 - Hot rolled steel sheet and method of manufacturing same - Google Patents

Abstract

The sum of the area percentages of the tempered bainite phase and the tempered martensite is not less than 70%, and the sum of the area percentages of the coarse pearlite phase, the martensite phase and the retained austenite phase is not more than 10% , The tempered bainite phase and the tempered martensite phase have a lath having an average width of 1.0 탆 or less as a substructure and a ratio of an Fe-based carbide precipitated in the lath and the lath boundary to an aspect ratio of 5 or less is 80 %, And an MC type carbide having an average particle diameter of not more than 20 nm dispersed and precipitated in the interior of the glass and the ras boundary, and has an average dislocation density of not less than 1.0 x 10 ¹⁴ m ^{-2 and not} more than 5.0 x 10 ¹⁵ m ^-2 Or less.

Description

TECHNICAL FIELD [0001] The present invention relates to a hot rolled steel sheet,

The present invention relates to a steel sheet having a high strength of 780 MPa or more and a tensile strength (TS) of 780 MPa or more, which is suitable for structural steels such as components of transportation machinery, Steel sheet and a manufacturing method thereof.

In order to reduce CO ₂ emissions from the viewpoint of global environmental conservation, it is always an important task in the automobile industry to reduce the weight of the vehicle body while maintaining the strength of the vehicle body, and to improve the fuel efficiency of the automobile. In order to reduce the weight of the vehicle body while maintaining the strength of the vehicle body, it is effective to make the steel sheet thinner by increasing the strength of the steel sheet as a material for automobile parts. For example, in a suspension component of an automobile, in which a steel plate having a large thickness is frequently used, it is expected to be considerably reduced in weight by thinning due to high strength.

Generally, automobile suspension parts such as low arm are formed by burring, and therefore, a steel sheet having excellent stretch flangeability is required. As for the hot-rolled steel sheet having strength and workability, various researches and developments have been made and various technologies have been proposed. For example, it is known that a high tensile strength and an excellent stretch flangeability can be achieved by making a metal structure substantially ferrite single phase and precipitating fine carbide in the ferrite phase particles.

As such a technique, Patent Document 1 discloses a technique in which a steel sheet structure is formed into a ferrite single phase structure having a low dislocation density and excellent in workability, and precipitation hardening is carried out by dispersing and precipitating fine carbide in ferrite, Is improved.

On the other hand, normal burring is carried out by using a steel sheet punched out in a predetermined shape. In manufacturing actual parts, it is common that the clearance of the part of the component is changed due to the temperature rise of the mold by the continuous press and the wear of the mold. When the clearance fluctuates, cracking and chipping Sometimes a problem arises. For this reason, there is a demand for a steel plate which has always excellent puncture property against fluctuation of punching conditions.

As such a steel sheet, for example, Patent Document 2 discloses a steel sheet having a bainite phase in a volume ratio of more than 92%, an average spacing of bainite laths of 0.60 탆 or less, and Fe A high-strength hot-rolled steel sheet in which the ratio of the number of carbides in the system is set to 10% or more, thereby improving mass-production puncture resistance.

Japanese Laid-Open Patent Publication No. 2012-26034 Japanese Laid-Open Patent Publication No. 2014-205888

In the steel sheet described in Patent Document 1, both high strength and excellent stretch flangeability are satisfied. However, since the steel sheet structure is substantially made into a ferrite single phase, there are almost no inclusions that serve as starting points of the voids when the steel sheet is punched out. Therefore, in the steel sheet described in Patent Document 1, when the conditions such as the clearance and the blank holder are changed, there is a fear that roughness may occur on the end surface of the stamp.

Further, in the steel sheet described in Patent Document 2, a steel sheet structure having a predetermined bainite as a main body is controlled by controlling the hot rolling condition, whereby excellent scratch resistance is obtained. However, such a bainite structure is characterized in that mechanical characteristics such as tensile strength are likely to fluctuate with fluctuation of the coiling temperature. Generally, it is not easy to make the steel sheet temperature uniform over the entire length of the entire coil in the cooling after hot rolling. Therefore, in the steel sheet described in Patent Document 2, the deviation of the mechanical characteristics of the steel sheet becomes large, and there is a fear that the manufacturing stability is lowered.

The present invention has been made in view of the above circumstances and provides a hot-rolled steel sheet having a high tensile strength (TS) of 780 MPa or more, excellent stretch flangeability and scratch resistance, and excellent manufacturing stability, And the like.

The present inventors have extensively studied a method capable of strengthening a steel sheet while retaining workability, particularly stretch flangeability, and also capable of suppressing a deviation in mechanical characteristics due to excellent puncture property and variation 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. With this technique, a ferrite single phase structure may be solid solution strengthened or precipitated strengthened to increase the strength, or a bainite single phase structure may be strengthened by strengthening the structure. However, in the steel sheet of the ferrite single-phase structure, there is almost no inclusion to be a starting point of voids in the steel sheet, so that there is a possibility that tearing may occur on the piercing end face when the conditions such as clearance and blank holder fluctuate.

On the other hand, the steel sheet of the bainite single phase structure has excellent stretch flangeability. Further, the steel sheet of the bainite single-phase structure has a large number of Fe-based carbides in the bainite structure, and since this is the starting point of the voids at the time of punching, the steel sheet also has excellent punchability. However, since the mechanical properties such as strength vary greatly depending on the transformation temperature, the bainite structure is likely to have a large variation in the mechanical characteristics due to the fluctuation of the hot rolling condition such as the coiling temperature.

Therefore, the present inventors considered to mitigate the influence of the hot rolling condition fluctuation by adding a tempering treatment to the bainite and martensite-based structure.

In general, by adding a tempering treatment to the bainite or martensitic structure, the deviation of the mechanical properties due to the change of the hot rolling condition is greatly reduced, but at the same time, the strength of the steel sheet is greatly reduced. The shape of the Fe-based carbide in the tempered bainite or the tempered martensitic phase changes depending on the annealing condition, so that the steel sheet is not necessarily excellent in the peelability depending on the annealing condition.

Thus, the present inventors have found that, in the case of adding a tempering treatment to a structure mainly composed of bainite and martensite, a method of suppressing a decrease in strength of the steel sheet as described above and also having excellent stretch flange formability and scratch resistance The present inventors have repeatedly studied the present invention.

As a result, by dispersing and depositing MC type carbide such as TiC in the lath inner and the lath boundary, the lath annihilation due to coarsening and recovery of lath during annealing is suppressed, and high steel sheet strength can be maintained even after annealing Respectively. Further, it has been found that excellent saturability can be obtained by securing a ratio of the Fe-based carbide precipitated in the lath boundary and the lath boundary to those having an aspect ratio of 5 or less.

The inventors have further found that the steel sheet structure can be stably formed by adding more than 0.03% of Ti and optimizing the thermal history in the annealing step.

The MC type carbide is a carbide of TiC, NbC, VC, (Ti, Mo, and the like) in which the atomic ratio of M element (M element can be Ti, Nb, V, ) C and the like. Here, the M element does not have to be one kind but may be a complex carbide containing a plurality of metal elements. N-containing carbonitrides and composite carbonitrides may also be used.

The inventors of the present invention have further found that by appropriately controlling the thermal history at the time of cooling from the maximum heating temperature to the room temperature in the annealing step, the residual structure other than the tempered martensite phase and the tempered bainite phase, In particular, it has been found that the formation of the martensitic phase, the coarse pearlite phase and the retained austenite phase is suppressed, whereby the high tensile strength and excellent punching performance can be obtained, and excellent stretch flangeability can be obtained.

The present invention has been completed based on the above findings and has been completed.

That is, the structure of the present invention is as follows.

1.% by mass,

C: not less than 0.03% and not more than 0.20%, Si: not more than 0.4%

Mn: not less than 0.5% to not more than 2.0%, P: not more than 0.03%

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

And a balance of Fe and inevitable impurities,

The total area ratio of the tempered bainite phase and the tempered martensite phase is 70% or more, the sum of the area percentages of the coarse pearlite phase, the martensitic phase and the retained austenite phase is 10%

Wherein the tempered bainite phase and the tempered martensite phase have a lath having an average width of not more than 1.0 mu m as a substructure and a ratio of an Fe-based carbide precipitated in the lath boundary and the lath boundary to an aspect ratio of not more than 5 is 80 % Or more and an MC type carbide having an average particle diameter of 20 nm or less dispersed and precipitated in a grain boundary and a lath boundary,

Wherein the average dislocation density is not less than 1.0 × 10 ¹⁴ m ^{-2 and not} more than 5.0 × 10 ¹⁵ m ^-2 .

2. The composition according to claim 1,

V: not less than 0.01% and not more than 0.3%

Nb: not less than 0.01% and not more than 0.1%

Mo: 0.01% or more and 0.3% or less

The hot-rolled steel sheet according to the above-mentioned 1, wherein the hot-rolled steel sheet contains at least one or two or more of them.

3. The composition according to claim 1,

B: not less than 0.0002% and not more than 0.010%

(1) or (2) above.

4. One or two 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 the above items 1 to 3, which contains not less than 1.0% in total of the species.

5. The hot-rolled steel sheet according to any one of 1 to 4 above, wherein the hot-rolled steel sheet has a plated layer on its surface.

6. A steel material having a composition as described in any one of 1 to 4 above, which is heated to austenite single phase and subjected to hot rolling comprising rough rolling and finish rolling, and after completion of the finish rolling, Cooling and winding the hot rolled steel sheet,

And a continuous annealing step of pickling the steel sheet after the hot rolling step and then continuously annealing the steel sheet,

In the hot rolling step, the finish rolling temperature is 850 DEG C to 1000 DEG C, the average cooling rate to 500 DEG C is 30 DEG C / s or more after the completion of the finish rolling, the coiling temperature is 500 DEG C or less,

In the continuous annealing step,

The maximum heating temperature of the steel sheet is 700 ° C or higher (A ₃ point + A ₁ point) / 2 or lower,

The time at which the steel sheet temperature at the time of heating the steel sheet 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 at which the steel sheet temperature exceeds 700 DEG C is set to 200 seconds or less,

The average cooling rate up to 530 캜 at the time of cooling the steel sheet from the maximum heating temperature is 8 캜 / s to 25 캜 / s or less, and the holding time at a temperature range of 470 캜 to 530 캜 10 seconds or more.

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

INDUSTRIAL APPLICABILITY According to the present invention, it is possible to provide a steel having high tensile strength (TS) of 780 MPa or more, excellent stretch flangeability and scratch resistance, A deviation in the mechanical properties of the hot-rolled steel sheet can be suppressed. As a result, the hot-rolled steel sheet can be further improved in its application and exhibits remarkable effects in industry.

Fig. 1 is a schematic view showing an example of a structure in which a Fe-based carbide is precipitated in the lath and a lath boundary, and a MC-type carbide is dispersed and precipitated, in which a tempered bainite phase and a tempered martensite phase have a lath as a lower structure to be.

Hereinafter, the present invention will be described in detail.

First, the composition of the hot-rolled steel sheet of the present invention will be described. The unit of the content of elements in the composition of the components is "% by mass ", which is hereinafter simply expressed as "% "

C: not less than 0.03% and not more than 0.20%

C improves the strength of the steel and promotes the production 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, if the C content exceeds 0.20%, the carbon equivalent becomes excessively large, deteriorating the weldability of the steel sheet. Therefore, the C content is set to 0.03% or more and 0.20% or less. , Preferably not less than 0.04% and not more than 0.18%, and more preferably not less than 0.05% and not more than 0.15%.

Si: not more than 0.4%

Si is an effective element that improves steel sheet strength without causing a ductility (elongation) deterioration, and is usually contained positively in a high-strength steel sheet. However, when the Si content exceeds 0.4%, an oxide is formed on the surface of the steel sheet at the time of heat treatment, which causes deterioration of the plating adhesion. Therefore, the Si content is set to 0.4% or less. , Preferably not more than 0.3%, more preferably not more than 0.2%. In addition, Si may reduce the addition amount to the impurity level or may be 0%.

Mn: not less than 0.5% and not more than 2.0%

Mn is an element which is solved and contributes to increase the strength of steel. Mn is an element promoting the formation of bainite and martensite during hot rolling due to improvement in castability. In order to obtain such an effect, it is necessary to set the Mn content to 0.5% or more. On the other hand, if the Mn content exceeds 2.0%, the austenite is excessively stabilized, resulting in a structure in which martensite or retained austenite is excessively contained in the steel sheet. Thus, the stretch flangeability deteriorates. Therefore, the Mn content is set to 0.5% or more and 2.0% or less. Further, it is preferably 0.8% or more and 1.8% or less, and more preferably 1.0% or more and 1.7% or less.

P: not more than 0.03%

P is a harmful element that is segregated at the grain boundaries to lower the stretching, cause cracking at the time of processing, and deteriorate impact resistance. Therefore, the P content is set to 0.03% or less. However, excessive P removal results in an increase in refining time and an increase in cost, so that the P content is preferably 0.002% or more.

S: not more than 0.03%

S exists as MnS or TiS in the steel, thereby promoting the generation of voids during punching of the hot-rolled steel sheet. Also, S is a starting point of generation of voids during processing, and therefore, the stretch flangeability is deteriorated. Therefore, it is preferable to reduce the S content as much as possible, and it is made 0.03% or less. It is preferably not more than 0.01%. However, since excess sintering causes an increase in refining time and an increase in cost, the S content is preferably 0.0002% or more.

Al: 0.1% or less

Al is an element serving as a de-oxidizing agent. In order to obtain such an effect, it is preferable to contain Al in an amount of 0.01% or more. However, if the content of Al exceeds 0.1%, Al oxide remains in the steel sheet and the Al oxide coagulates and becomes coarse, thereby deteriorating the stretch flangeability. Therefore, the Al content is set to 0.1% or less.

N: not more than 0.01%

N is present as coarse TiN in the steel and promotes the generation of coarse voids during punching of the hot-rolled steel sheet. Further, N is a starting point of generation of coarse voids during processing, and therefore, the stretch flangeability is deteriorated. Therefore, it is preferable to reduce the N content as much as possible, and it is set to 0.01% or less. It is preferably not more than 0.006%. However, since excessive denitrification causes an increase in refining time and an increase in cost, the N content is preferably 0.0005% or more.

Ti: not less than 0.03% and not more than 0.15%

Ti is an indispensable element for forming a MC type carbide to suppress Raschogonite in an annealing process and to increase the strength of a steel sheet. Further, the MC type carbide increases the strength of the steel sheet by precipitation strengthening. If the Ti content is less than 0.03%, these effects can not be sufficiently obtained. Therefore, the strength of the steel sheet is lowered due to the coarsening of the lath and the decrease in the deposition amount, and it becomes 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 center segregation becomes remarkable, which causes deterioration deterioration. Therefore, the Ti content is set to 0.03% or more and 0.15% or less. , Preferably not less than 0.04% and not more than 0.14%, and more preferably not less than 0.05% and not more than 0.13%.

In the hot-rolled steel sheet of the present invention, the content of V is not less than 0.01% and not more than 0.3%, Nb is not less than 0.01% and not more than 0.1%, and Mo is not less than 0.01% 0.3% or less.

V: not less than 0.01% and not more than 0.3%

V forms an MC-type carbide and contributes to strengthening of the steel sheet by suppressing Raschogonite and strengthening precipitation in an annealing process like Ti. In order to obtain such an effect, it is necessary to contain V of 0.01% or more. On the other hand, if the V content exceeds 0.3%, the center segregation becomes remarkable, which causes deterioration of detonation. Therefore, the V content is preferably 0.01% or more and 0.3% or less. More preferably, it is not less than 0.01% and not more than 0.2%, and more preferably not less than 0.01% and not more than 0.15%. In addition, V may form MC type carbide alone or may form a complex carbide with Ti, Nb and Mo. These carbide compositions do not affect the effect of the invention at all.

Nb: 0.01% or more and 0.1% or less

Nb forms an MC-type carbide and contributes to strengthening of the steel sheet by suppressing Rasco's conversions and strengthening precipitation in an annealing process like Ti. In order to obtain such an effect, it is necessary to contain Nb in an amount of 0.01% or more. On the other hand, if Nb is added in an excess amount exceeding 0.1%, the effect is saturated because it is not dissolved in the heating furnace at the time of hot rolling, which causes the cost of the alloy to increase unnecessarily. Therefore, it is preferable that the Nb content is 0.01% or more and 0.1% or less. It is more preferably not less than 0.01% and not more than 0.08%, and still more preferably not less than 0.01% and not more than 0.06%. Nb may form MC type carbide alone or may form a complex carbide with Ti, V, or Mo. These carbide compositions do not affect the effect of the invention at all.

Mo: 0.01% or more and 0.3% or less

Mo combined with Ti forms a MC type complex carbide and contributes to the strengthening of the steel sheet by the suppression of rascoaction and precipitation strengthening in the annealing step like Ti. In order to obtain such an effect, it is necessary to contain Mo in an amount of 0.01% or more. On the other hand, if the Mo content exceeds 0.3%, the center segregation becomes remarkable, which causes deterioration deterioration. Therefore, the Mo content is preferably 0.01% or more and 0.3% or less. In some cases, Mo forms a complex carbide with Nb or V, but these carbide compositions do not affect the effect of the invention at all.

In the hot-rolled steel sheet of the present invention, B may be contained in an amount of not less than 0.0002% and not more than 0.010%, as required, for the purpose of improving uniformity in hot rolling.

B: not less than 0.0002% and not more than 0.010%

B is segregated at the austenitic grain boundaries and is an element which improves the quenching by inhibiting the generation and growth of ferrite and promotes the production of bainite and martensite. In order to obtain these effects, the B content is preferably 0.0002% or more. However, when the B content exceeds 0.010%, rigid ferro-borides are formed to cause deterioration of the stretch flangeability. Therefore, when B is contained, the content thereof is preferably 0.0002% or more and 0.010% or less. Further, it is more preferably 0.0002% or more and 0.0050% or less, and still more preferably 0.0004% or more and 0.0030% or less.

In addition, the hot-rolled steel sheet of the present invention may further contain one or more of REM, Zr, As, Cu, Ni, Sn, Pb, Ta, W, Cr, Sb, Mg, Ca, Co, Se, Zn and Cs May contain one or more species in a total amount of 1.0% or less.

The components other than the above are Fe and inevitable impurities.

Next, reasons for limiting the structure of the hot-rolled steel sheet of the present invention will be described.

Total area ratio of tempered bainite phase and tempered martensite: not less than 70%

In the hot-rolled steel sheet of the present invention, a structure mainly composed of tempered bainite and tempered martensite, which have high strength and excellent brittleness, is used. If the sum of the area ratios of the tempered bainite phase and the tempered martensite is less than 70%, a hot-rolled steel sheet having desired high strength and scratch resistance can not be obtained. Also, the reason for not individually defining the fraction of the tempered bainite phase and the tempered martensite is that the tempered bainite after annealing and the tempered martensite become indistinguishable structures. This is also a large factor that can suppress the deviation of the mechanical properties after annealing even when the manufacturing conditions at the time of hot rolling are changed. Therefore, the sum of the area ratios of the tempered bainite phase and the tempered martensite is not less than 70%. , Preferably 75% or more, and more preferably 80% or more. In addition, 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 retained austenite phase: 10% or less

As described above, in the hot-rolled steel sheet of the present invention, the structure is mainly composed of tempered bainite and tempered martensite, and the remaining structure other than tempered bainite and tempered martensite includes Fe-based carbide , Coarse pearlite, fine pearlite, pseudo pearlite, bainite, martensite, retained austenite, and the like. Among these structures, in particular, when coarse pearlite, martensite and retained austenite are present in the metal structure, the stretch flangeability is remarkably deteriorated. Therefore, the total area ratio of the coarse pearlite phase, the martensite phase and the retained austenite phase is 10% or less. It is preferably at most 8%, more preferably at most 5%. It may be 0%.

Here, it is assumed that coarse pearlite has a lamellar spacing of 0.2 탆 or more, that a fine pearlite has a lamellar spacing of less than 0.2 탆, and that pearlite lamellar is not clearly observed. The lamellar spacing can be obtained by observation of the structure by a scanning electron microscope.

Examples of the residual structure other than the tempered bainite phase, the tempered martensitic phase, the coarse pearlite phase, the martensitic phase and the retained austenite phase include ferrite phase, pseudo-pearlite phase and fine perlite phase. Such residual structure can be allowed if the total area ratio is 30% or less.

Average width of lasess having a tempered bainite phase and a tempered martensite phase as a substructure: 1.0 m or less

In order to increase the strength by the tempered bainite phase and the tempered martensitic phase, it is important to have fine lathes having an average width of not more than 1.0 mu m as the underlying structure. Fig. 1 shows an example of a structure in which a Fe-based carbide is precipitated in the lath boundary and a lattice boundary in which the tempered bainite phase and the tempered martensite phase have a lath as the underlying structure and an MC type carbide is dispersed and precipitated It shows a schematic diagram. Here, when the lath is lost by the recovery or the average width of the lath exceeds 1.0 탆, the predetermined high strength can not be attained. Therefore, the average width of the lashed structure having the tempered bainite phase and the tempered martensite phase as the substructure is set to 1.0 μm or less. Preferably 0.8 μm or less, and more preferably 0.6 μm or less. The lower limit is not particularly limited, but is usually about 0.1 mu m.

The proportion of the Fe-based carbides precipitated in the lath and the lath boundary having an aspect ratio of 5 or less: 80% or more

As shown in Fig. 1, the Fe-based carbides precipitated in the lath boundary and the lath boundary contribute to the improvement of the saturability by becoming the origin of the voids at the time of punching. Particularly, the Fe-based carbide having an aspect ratio of 5 or less has such a large effect, and when the ratio is 80% or more, superior saturability can be exhibited. Therefore, the proportion of the Fe-based carbides precipitated in the lath boundary and the lath boundary with an aspect ratio of 5 or less is 80% or more. It is preferably at least 85%. The upper limit is not particularly limited, and may be 100%.

Here, the Fe-based carbide is? Carbide (cementite) or? Carbide. The carbide may contain an alloy element. The aspect ratio is defined as the ratio of the lengths of the long diameter and the short diameter of the Fe-based carbides precipitated in the lath and the rath boundary.

Average particle diameter of MC type carbide dispersed and precipitated in the lath boundary and the lath boundary: 20 nm or less

As shown in Fig. 1, the MC type carbide finely dispersed and precipitated in the lath boundary and the lath boundary suppresses coarsening of the lass by the pin fixing effect at the time of annealing of the steel sheet, and suppresses disappearance of lass due to recovery , Contributing to the enhancement of strength. However, when the average particle diameter of the MC-type carbide exceeds 20 nm, the number of particles of the MC-type carbide contributing to pinning is insufficient, the pin-fixing effect becomes insufficient and the strength of the steel sheet is lowered. On the other hand, when the average particle diameter of the MC-type carbide is 20 nm or less, a sufficient number of MC-type carbides exhibit the pin-fixing effect and suppress the decrease of the strength of the steel sheet. Therefore, the average particle diameter of the MC-type carbide dispersed and precipitated on the inner and lath boundaries of the tempered bainite phase and the tempered martensite is set to 20 nm or less. And preferably 15 nm or less. The lower limit is not particularly limited, but is usually about 1 nm. However, it is preferable that the ratio of MC type carbide having a particle diameter exceeding 50 nm is 10% or less.

In the hot-rolled steel sheet of the present invention, it is important to set the average dislocation density to the following range.

Average Dislocation Density: 1.0 x 10 ¹⁴ m ^-2 or more and 5.0 x 10 ¹⁵ m ^-2 or less

In the hot-rolled steel sheet of the present invention, the variation in the hot rolling condition is reduced by tempering the steel sheet having bainite and martensite structure. If the average dislocation density of the steel sheet after annealing exceeds 5.0 x 10 < ¹⁵ > m < ^{2 &} gt ^; , the tempering of the steel sheet is insufficient and the influence of the fluctuation of the hot rolling condition can not be sufficiently mitigated. On the other hand, even when sufficiently tempered, the average dislocation density is usually 1.0 x 10 ¹⁴ m ^-2 or more. Therefore, the average dislocation density should be 1.0 x 10 ¹⁴ m ^-2 to 5.0 x 10 ¹⁵ m ^-2 . And preferably not less than 1.0 × 10 ¹⁴ m ^{-2 and not} more than 2.0 × 10 ¹⁵ m ^-2 .

Next, a method of manufacturing the hot-rolled steel sheet of the present invention will be described.

In the method of manufacturing a hot-rolled steel sheet of the present invention, the steel material having the above-mentioned composition is heated to austenite single phase, hot rolling comprising rough rolling and finish rolling is performed, and after completion of the finish rolling, , A hot rolling step in which hot rolling is carried out,

And a continuous annealing step of pickling the steel sheet after the hot rolling step and then continuously annealing the steel sheet,

In the hot rolling step, the finish rolling temperature is 850 DEG C to 1000 DEG C, the average cooling rate to 500 DEG C is 30 DEG C / s or more after the completion of the finish rolling, the coiling temperature is 500 DEG C or less,

In the continuous annealing step,

The maximum heating temperature of the steel sheet is 700 ° C or higher (A ₃ point + A ₁ point) / 2 or lower,

The time at which the steel sheet temperature at the time of heating the steel sheet 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 at which the steel sheet temperature exceeds 700 DEG C is set to 200 seconds or less,

The average cooling rate up to 530 캜 at the time of cooling the steel sheet from the maximum heating temperature is 8 캜 / s to 25 캜 / s or less, and the holding time at a temperature range of 470 캜 to 530 캜 10 seconds or more. Further, a step of performing a plating process after the continuous annealing process may be further provided.

The method of the solvent for the steel material is not particularly limited, and a known solvent method such as a converter, an electric furnace or the like may be employed. In addition, after the solvent, it is preferable to make the slab (steel material) by the continuous casting method because of productivity and the like, and the slab may be formed by a known casting method such as a roughing-slab rolling method and a thin slab dredging method.

The steel material obtained as described above is subjected to hot rolling consisting of rough rolling and finish rolling, but the steel material is heated in the austenite single phase before rough rolling. If the steel material before rough rolling is not heated in a single phase of austenite, redissolving of the Ti carbide or the like present in the steel material does not proceed, and microcrystallization of MC type carbide is not performed in annealing after hot rolling. Therefore, prior to rough rolling, the steel material is heated to austenite single phase, preferably at least 1150 占 폚. The upper limit of the heating temperature is not specifically defined. However, if the heating temperature is higher than necessary, the reduction in yield due to oxidation of the slab surface becomes remarkable. Therefore, the heating temperature is usually 1350 占 폚 or lower. In the case where the steel material (slab) after casting has a temperature of the austenite single phase in the hot rolling of the steel material, the steel material may be directly rolled without heating or after heating for a short time.

Next, reasons for limiting the production conditions in the hot rolling step will be described.

Finishing rolling temperature: 850 ℃ or more and 1000 ℃ or less

When the finish rolling temperature is lowered, ferrite transformation is promoted in cooling after rolling, so that the bainite and martensite fraction of the hot-rolled steel sheet obtained after hot rolling is lowered. As a result, a predetermined tempered bainite and a tempered martensitic fraction can not be obtained after annealing. Therefore, the temperature at the finishing rolling outlet side needs to be 850 DEG C or higher, preferably 880 DEG C or higher. When the finishing rolling temperature exceeds 1000 캜, the surface properties of the steel sheet deteriorate. Therefore, the upper limit of the finishing rolling temperature is set to 1000 占 폚 or less. Preferably 970 占 폚 or less. Incidentally, each temperature such as the winding temperature including the finish rolling temperature is the temperature of the surface of the steel sheet.

After completion of finish rolling, cooling rate to 500 캜: 30 캜 / s or more

When the cooling rate is insufficient in cooling the steel sheet after finishing rolling, the ferrite can not be sufficiently suppressed, and the bainite and martensite fraction of the hot-rolled steel sheet obtained after hot rolling is lowered. As a result, a predetermined tempered bainite and a tempered martensitic fraction can not be obtained after annealing. Therefore, after completion of the finish rolling, the cooling rate to 500 占 폚 needs to be 30 占 폚 / s or more. Preferably 50 DEG C / s or more. The upper limit of the cooling rate is not particularly limited, but is usually about 300 ° C / s.

Coiling temperature: 500 占 폚 or less

Optimization of the coiling temperature is important in controlling the steel sheet structure after hot rolling. If the coiling temperature exceeds 500 캜, the las width of the bainite becomes large, and therefore the las width of the tempered bainite after annealing can not be a predetermined value. On the other hand, although the lower limit of the coiling temperature is not particularly limited, the cooling cost is unnecessarily expensive when the coiling temperature is excessively reduced. Therefore, the coiling temperature is preferably 0 DEG C or higher. More preferably, it is 200 DEG C or more.

After the above hot rolling step, the hot-rolled steel sheet is pickled and subjected to a continuous annealing process for continuous annealing. Hereinafter, the reason for limiting the manufacturing conditions in this continuous annealing step will be described.

Maximum heating temperature of steel plate: 700 ° C or more (A ₃ points + A ₁ point) / 2 or less

Optimization of the maximum heating temperature of the steel sheet in the continuous annealing step is important in sufficiently reducing the influence of fluctuations in the manufacturing conditions during hot rolling by annealing and achieving a desired high strength. If the maximum heating temperature of the steel sheet is less than 700 deg. C, it is difficult to control the dislocation density in the bainite and martensite to an appropriate range, so that the influence of fluctuation of the manufacturing conditions during hot rolling can not be sufficiently reduced. In addition, when the heating temperature of the steel sheet is less than 700 占 폚, the aspect ratio of the Fe-based carbide in the lath and the lath easily increases, and it is difficult to set the ratio of the Fe-based carbide having the aspect ratio of 5 or less to a desired range. On the other hand, when the maximum heating temperature of the steel sheet exceeds (A ₃ point + A ₁ point) / 2, coarsening of the MC type carbide remarkably occurs, so that coarsening of lath in bainite and martensite can be sufficiently suppressed I will not. In addition, the austenitization is promoted, whereby the bainite and martensite fractions are lowered, and the desired tempered bainite and tempered martensitic fractions can not be obtained. Therefore, the maximum heating temperature of the steel sheet in the continuous annealing step is 700 ° C or higher (A ₃ point + A ₁ point) / 2 or lower. Further, it is preferably 700 ° C or more {(A ₃ point + A ₁ point) / 2} - 10 ° C or less.

The points A ₁ and A ₃ can be calculated by the following equations.

% Mo] + 233 占 [% Nb] - 39.7 占 [% V] - 57 占 폚 - ₁ point = 751 - 26.6 x [% C] + 17.6 x [% Si] [% Ti] - 895 x [% B] - 169 x [% Al]

A ₃ point = 937 - 476.5 x [% C] + 56 x [% Si] - 19.7 x [% Mn] + 38.1 x [% Mo] + 124.8 x [% V] + 136.3 x [ [% Nb] + 3315 x [% B]

Here, [% X] means the content (mass%) in the steel of the X element.

Time when the steel sheet temperature is 600 ° C or more and 700 ° C or less when heating the steel sheet to the maximum heating temperature: 20 seconds or more and 1000 seconds or less

In heating the steel sheet to the maximum heating temperature, it is important to suitably control the heating heat history in imparting desired high strength and excellent punching ability to the steel sheet. As described above, the pin fixing effect by the 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 the MC type carbide in the bainite and the martensite before the lath starts coarsening. According to a study made by the present inventors, precipitation of the MC type carbide starts to occur remarkably at a temperature of 600 ° C or higher. On the other hand, the coarsening and disappearance of lathes remarkably occurs by exceeding 700 캜. Therefore, coarsening and disappearance of lath can be suppressed by keeping the steel sheet temperature at a temperature range of 600 DEG C or more and 700 DEG C or less for a certain period of time and sufficiently depositing MC type carbide. Here, in order to sufficiently precipitate the MC-type carbide, it is necessary to hold it for 20 seconds or more at this temperature range. When the holding time at this temperature range is insufficient, coarsening of the lath is started before the MC type carbide sufficiently precipitates, so that the pin fixing effect is not sufficiently exhibited and the lath becomes coarse. Preferably 35 seconds or more, and more preferably 50 seconds or more.

On the other hand, if the holding time at the temperature range of the steel sheet temperature of 600 ° C or more and 700 ° C or less exceeds 1000 seconds, the Fe-based carbides precipitated in the lath and between the lath are reused and the old austenite grains, And there is no Fe-based carbide between the lath and the lath which effectively contributes to the improvement of the scratch resistance. Therefore, in order to obtain a steel sheet having excellent puncture property, it is necessary to set the holding time at a temperature range of 600 ° C or more and 700 ° C or less to 1000 seconds or less. Preferably 800 seconds or less, and more preferably 500 seconds or less. The steel sheet temperature referred to herein is the temperature of the steel sheet surface.

Time when steel plate temperature exceeds 700 ℃: 200 seconds or less

In the temperature range where the steel sheet temperature exceeds 700 ° C, rough coarsening of the lath is remarkable. As described above, according to the present invention, the movement of the ras boundary is suppressed by the pin fixing effect by the finely dispersed and precipitated MC type carbide, thereby suppressing coarsening of the ras. Thus, the strength of the steel sheet is maintained. However, if the holding time at this temperature range becomes excessively long, the coarsening of lath can not be suppressed at all. Therefore, from the viewpoint of preventing the coarsening of the lath, the holding time at the temperature range in which the steel sheet temperature is higher than 700 deg. C is 200 seconds or less. Preferably 180 seconds or less, more preferably 150 seconds or less. However, if the steel sheet temperature exceeds 700 ° C for less than 10 seconds, the ductility of the steel sheet becomes somewhat inferior, so it is preferable to set it to 10 seconds or more.

Average cooling rate up to 530 ℃ when cooling the steel sheet from the maximum heating temperature: 8 ℃ / s to 25 ℃ / s or less

In the continuous annealing process, it is important to appropriately control the cooling rate in cooling the steel sheet after it is heated up to the maximum heating temperature, which is important in obtaining excellent stretch flangeability. Particularly, when the average cooling rate up to 530 占 폚 is less than 8 占 폚 / s, pearlite transformation can not be suppressed during cooling, and coarse pearlite is produced in a predetermined amount or more. Thus, the stretch flangeability is deteriorated. On the other hand, when the average cooling rate is excessively high, it becomes difficult to maintain the temperature at a temperature range of 470 DEG C to 530 DEG C, which will be described later, for a predetermined time. Therefore, when cooling the steel sheet from the maximum heating temperature, the average cooling rate up to 530 占 폚 is 25 占 폚 / s or less.

Holding time at a temperature range of 470 ° C to 530 ° C: 10 seconds or more

In the continuous annealing process, it is important to maintain the steel sheet for a predetermined period of time in the temperature range of 470 DEG C to 530 DEG C after controlling and cooling the steel sheet as described above in order to obtain excellent stretch flangeability. Here, if the holding temperature after the cooling stop exceeds 530 占 폚, coarse pearlite is produced and the stretch flangeability is deteriorated. On the other hand, if the holding temperature after cooling is stopped below 470 캜, the transformation from austenite to bainite is delayed. As a result, C is concentrated in the untransformed austenite region to stabilize the austenite, so that the transformation is not completed. In the subsequent cooling, the untransformed austenite is transformed into martensite or remains in the steel sheet structure as the retained austenite, so that the stretch flangeability is lowered. Further, when the steel sheet is held for 10 seconds or more at a temperature range of 470 DEG C to 530 DEG C, most of the transformation of the austenite into bainite is completed and the martensite fraction generated upon cooling is then decreased to a predetermined range Can be reduced. Therefore, after the control cooling is stopped, the holding time at a temperature range of 470 DEG C to 530 DEG C is 10 seconds or more. Preferably 20 seconds or more, and 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 the holding at the temperature range of 470 DEG C or higher and 530 DEG C or lower, the cooling condition thereafter is not particularly limited and may be cooled to room temperature by an arbitrary cooling method.

Further, even if the steel sheet is reheated to a temperature of 700 DEG C or lower after holding at a temperature range of 470 DEG C or higher and 530 DEG C or lower, a desired steel sheet structure can be obtained if the cumulative holding time at a temperature range of 600 DEG C to 700 DEG C is 1000 seconds or less There is no problem at all.

For example, the steel sheet may be immersed in a zinc port to be a hot-dip galvanized steel sheet after holding at a temperature range of 470 DEG C to 530 DEG C, and then subjected to a heat treatment to form an alloyed hot-dip galvanized steel sheet. In addition to zinc, aluminum or aluminum alloy may be plated for the hot-dip plating.

After the continuous annealing step, the steel sheet may be subjected to temper rolling in an annealing line continuously or in a separate line according to a conventional method.

Further, the hot-rolled steel sheet produced as described above may be subjected to electro-galvanizing treatment, or may be subjected to hot-dip galvanizing. The hot-rolled steel sheet of the present invention is suitable for press forming carried out at ordinary room temperature in addition to being preferable as a steel sheet for an automobile suspension, and has excellent heat-resistant treatment characteristics. Thus, the hot-rolled steel sheet produced as described above is also preferably used as a hot-rolled material steel sheet in which the steel sheet before pressing is heated from 400 ° C to 700 ° C and then press-formed.

Example

Molten steel having the composition shown in Table 1 was formed into a slab (steel material) having a thickness of 300 mm by solvent and continuous casting by a generally known method. These slabs were heated to the temperatures shown in Table 2, subjected to rough rolling, finished at the temperature shown in Table 2, and cooled at the average cooling rate shown in Table 2 after finishing rolling, To obtain a hot-rolled steel sheet having a thickness of 3.2 mm. These hot-rolled steel sheets were pickled by a commonly known method and subjected to an annealing treatment under the conditions shown in Table 2 on a continuous annealing line. For some steel sheets, hot dip galvanizing treatment and alloying treatment were carried out in a continuous annealing line to obtain hot-dip galvanized steel sheets and galvannealed hot-dip galvanized steel sheets.

A test piece was taken from the hot-rolled steel sheet thus obtained to evaluate the structure, the average dislocation density, the tensile test, the hole expansion test, the punch test, and the manufacturing stability. The evaluation results are shown in Table 3. The test method was as follows.

(i) Tissue observation

The test piece was taken from the obtained hot-rolled steel sheet, and the cross-section (L section) parallel to the rolling direction of the test piece was polished, and the steel sheet was taken by a scanning electron microscope (magnification: 1000, Using the photographs, the sum of the area ratios of the tempered bainite phase and the tempered martensite, the area ratio of the coarse pearlite phase, the sum of the area ratios of the martensite phase and the retained austenite phase (MA) The area ratios of the other images were obtained. It is difficult to distinguish the martensitic phase from the residual austenite phase by a scanning electron micrograph. Here, since the sum of the area ratios of the coarse pearlite phase, the martensitic phase and the retained austenite phase is important, The area ratio of the sum of the martensite phase and the residual austenite phase (MA) was determined without distinguishing between the asphalt phase and the retained austenite phase.

Further, the thin film produced from the hot-rolled steel sheet was observed by a transmission electron microscope, and the las width of tempered bainite and tempered martensite was measured. In addition, the Fe-based carbide precipitated in the lath boundary and the lath boundary, The ratio of the ratio of 5 or less, and the average particle size of the MC type carbide precipitated in the lath boundary and the lath boundary were determined.

Here, for the measurement of the lasso width of the tempered bainite and the tempered martensite, a transmission electron microscope photograph of 120 mm x 80 mm size photographed at 10:00 with a magnification of 30000 times, Five straight lines were drawn at an interval of 10 mm perpendicular to the long axis of the test piece, and the lengths of the line segments intersecting the ras boundary were measured. The average value of the lengths of the obtained line segments was taken as the average width of the ras.

The ratio of the Fe-based carbides precipitated in the lath boundary and the lath boundary to those having an aspect ratio of 5 or less was obtained by using photographs taken at a magnification of 165000 times, The lengths of the long diameter and the short diameter of the Fe-based carbide were measured and the aspect ratios were calculated to obtain the ratios of those having the aspect ratio of 5 or less.

Further, the average particle size of the MC-type carbides, using the pictures taken with 300,000 times magnification, and measuring the diameter with respect to MC-type carbides, such as minimum 100 TiC in 5 visual field totals, arithmetic mean (average particle diameter d _def ). The lower limit of the measured particle size is 2 nm.

(ii) Measurement of average dislocation density

A test piece was taken from the obtained hot-rolled steel sheet to measure the dislocation density at 1/4 of the plate thickness. The dislocation density at 1/4 of the plate thickness was considered to represent the average dislocation density of the steel sheet. Respectively. The sampled specimens were polished with 0.1 mm oxalic acid in addition to mechanical grinding to prepare samples so that 1/4 of the plate thickness was exposed to the surface. Here, polishing with oxalic acid is performed in order to remove the machining layer by grinding.

The deformation of the steel sheet was measured by the X-ray diffractometer with respect to the sample prepared as described above. The diffraction intensity of (110) plane, (211) plane and (220) plane of? Iron of 1/4 plate thickness was measured using an X-ray diffraction apparatus using a X-ray diffraction apparatus, The half width of the peak value of the reflection intensity on the crystal face is determined, and the local strain? 'Given to the steel sheet is determined by the following equations (1) and (2).

?? cos ?? /? = 0.9 / D + 2 ?? sin ?? /? (1)

here,

beta: half width of the peak value (however, the value corrected by the formula (2) is used)

θ: diffraction angle

?: wavelength of CoK? ray (0.1790 nm)

D: crystallite size (dislocation cell, grain size)

ε ': Local strain

? ² =? _m ² -? ₀ ² (2)

here,

beta _m : half width of the peak of the sample for measuring the dislocation density

β ₀ : half width of the peak of the sample without deformation

to be.

Further,? Cos? /? Is plotted against sin? /?, And? 'And D are obtained from the slope and the slice. The dislocation density ρ is determined from the obtained local strain ε 'by the following equation (3).

ρ = 14.4 ε ' ² / b ² (3)

here,

b: Burgers vector (0.248 nm)

to be.

(iii) tensile test

JIS No. 5 tensile test specimen (JIS Z 2001) was taken from the obtained hot-rolled steel sheet in the direction perpendicular to the rolling direction (C direction) in the tensile direction, and subjected to a tensile test according to JIS Z 2241, YS), tensile strength (TS) and elongation (El).

(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 having an initial diameter d ₀ : 10 mmφ was formed in the test piece (clearance: 12.5% of the thickness of the test piece plate). Using these test pieces, a hole expansion test was carried out. That is, a cone punch having an apex angle of 60 ° is inserted into a hole having an initial diameter d _{0 of} 10 mmφ from the punch side at the time of punching, and the hole is pushed to expand the diameter d of the hole when the crack passes through the steel plate (mm) was measured, and the hole expanding ratio? (%) was calculated by the following equation.

Hole expansion factor? (%) = {(D - d ₀ ) / d ₀ } × 100

In this case, it was judged that the stretch flangeability was good when the tensile strength (TS) x (hole expanding rate (?)) ^0.5 was 6200 占 MPa% ^0.5 or more.

(v) Punch test

A test piece (size: 30 mm x 30 mm) was taken from the obtained hot-rolled steel sheet, and a hole having a diameter d ₀ : 10 mmφ was formed in the test piece by punching (clearance: 20% of the thickness of the test piece plate, 30%). After punching, the presence of cracking, chipping, and brittle fracture was observed by observing the fracture surface condition of the punch hole over the entire circumference of the punch hole with a microscope (magnification: 50 times). Cracking, chipping, and brittle fracture were evaluated as " OK ", and the others were evaluated as " x "

In Table 3, the hot-rolled steel sheet of the present invention has a tensile strength (TS) of 780 MPa or more and excellent tensile strength and tensile strength.

Further, in order to evaluate the deviation of the mechanical properties of the steel sheet, a tensile test specimen (JIS Z 2001) of JIS No. 5 (JIS Z 2001) in which the tensile direction was arbitrarily orthogonal to the entire length of the hot-rolled steel sheet And tensile tests were carried out in accordance with JIS Z 2241, tensile strengths (TS) were measured, and their standard deviations? Were determined. As a result, in the present invention, the standard deviation of tensile strength (TS) was all within 10 MPa.

As described above, in the present invention, all of the deviations of the mechanical properties of the steel sheet such as the tensile strength (TS) are small, and the manufacturing stability is also excellent.

Claims (7)

In terms of% by mass,
C: not less than 0.03% and not more than 0.20%, Si: not more than 0.4%
Mn: not less than 0.5% to not more than 2.0%, P: not more than 0.03%
S: 0.03% or less, Al: more than 0% to 0.1%
N: 0.01% or less and Ti: 0.03% or more and 0.15% or less
And a balance of Fe and inevitable impurities,
The total area ratio of the tempered bainite phase and the tempered martensite phase is 70% or more, the sum of the area percentages of the coarse pearlite phase, the martensitic phase and the retained austenite phase is 10%
Wherein the tempered bainite phase and the tempered martensite phase have a lath having an average width of not more than 1.0 mu m as a substructure and a ratio of an Fe-based carbide precipitated in the lath boundary and the lath boundary to an aspect ratio of not more than 5 is 80 % Or more and an MC type carbide having an average particle diameter of 20 nm or less dispersed and precipitated in a grain boundary and a lath boundary,
Wherein the average dislocation density is not less than 1.0 × 10 ¹⁴ m ^{-2 and not} more than 5.0 × 10 ¹⁵ m ^-2 . The method according to claim 1,
Wherein the composition further comprises at least one group selected from the following (A) to (C).
At least one or more of V: at least 0.01% and not more than 0.3%, Nb: at least 0.01% and not more than 0.1%, and Mo: not less than 0.01% and not more than 0.3%
(B) in mass%, B: not less than 0.0002% and not more than 0.010%;
(C) at least one of REM, Zr, As, Cu, Ni, Sn, Pb, Ta, W, Cr, Sb, Mg, Ca, Co, Se, Zn, Not more than 1.0%. The hot-rolled steel sheet according to claim 1 or 2, wherein the hot-rolled steel sheet has a plated layer on its surface. A steel material having a composition according to any one of claims 1 to 3 is heated to austenite single phase and subjected to hot rolling consisting of rough rolling and finish rolling to cool the obtained steel sheet after finishing rolling, A rolling process,
And a continuous annealing step of pickling the steel sheet after the hot rolling step and then continuously annealing the steel sheet,
In the hot rolling step, the finish rolling temperature is 850 DEG C to 1000 DEG C, the average cooling rate to 500 DEG C is 30 DEG C / s or more after the completion of the finish rolling, the coiling temperature is 500 DEG C or less,
In the continuous annealing step,
The maximum heating temperature of the steel sheet is 700 ° C or higher (A ₃ point + A ₁ point) / 2 or lower,
The time at which the steel sheet temperature at the time of heating the steel sheet 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 at which the steel sheet temperature exceeds 700 DEG C is set to 200 seconds or less,
The average cooling rate up to 530 캜 at the time of cooling the steel sheet from the maximum heating temperature is 8 캜 / s to 25 캜 / s or less, and the holding time at a temperature range of 470 캜 to 530 캜 10 seconds or more. 5. The method of claim 4,
Further comprising a step of performing a plating treatment after the continuous annealing step. delete delete

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