JP6699307B2 - Hot-rolled steel sheet and its manufacturing method - Google Patents

Description

  TECHNICAL FIELD The present invention relates to a high-strength hot-rolled steel sheet excellent in ductility, fatigue characteristics, and surface machinability, and a method for manufacturing the same.

  2. Description of the Related Art In recent years, weight reduction of various parts constituting an automobile has been promoted for the purpose of improving fuel efficiency of the automobile. The weight reduction means differs depending on the required performance of each part. For example, in the case of frame parts, the strength of the steel sheet is increased to make it thinner, and in panel parts, the steel sheet is replaced with a light metal such as an Al alloy. However, as compared with steel, light metals such as Al alloys are expensive, so that they are currently applied mainly to luxury cars.

  On the other hand, automobile demand is shifting from developed countries to emerging countries, and it is expected that both weight reduction and price reduction will be required in the future. Regardless of the parts, it is necessary to achieve higher strength and thinner steel to reduce weight.

  Aluminum wheels and forged products were advantageous for passenger car wheels from the viewpoint of design. Recently, however, products with the same design characteristics as aluminum wheels have emerged for steel pressed products by devising the materials and construction methods.

  Up to now, as properties required for steel plates for wheel discs, overhanging workability, drawing workability and fatigue durability have been particularly emphasized. This is because the processing of the hat portion is strict in the wheel disc forming process, and the fatigue durability is controlled by the strictest standard in the member characteristics of the wheel. Fatigue durability is classified into two types, that is, the fatigue characteristics of a smooth material without notches and the notch fatigue characteristics under stress concentration, and it is desirable that both of the fatigue characteristics be good in a wheel disk.

  Further, in the case of a wheel disk that is required to have particularly high designability, machinability of the surface layer may be required in addition to the above-mentioned workability and fatigue durability. In the case of a high design wheel in which a disc is welded to the rim, the surface of the disc is cut by about 10 to 50 μm before welding the disc and the rim, and the gap between the contact surface with the rim is eliminated to perform welding. At this time, if the machinability of the disk is good, the manufacturing cost is improved by improving the yield and extending the life of the cutting tool.

  Currently, as a high-strength hot-rolled steel sheet for wheel discs, a 590 MPa class ferrite-martensite composite structure steel sheet (so-called Dual Phase steel), which is excellent in fatigue characteristics, is used with emphasis on fatigue durability of members. .. However, the strength level required for these steel sheets is increasing from 590 MPa to 780 MPa.

  Generally, it is said that even if the strength is the same, uniform elongation can be improved by making the microstructure of a steel plate into a composite structure such as Dual Phase steel (hereinafter referred to as DP steel) composed of ferrite and martensite. ..

  On the other hand, DP steel is known to have a low local deformability represented by bending, hole expanding, and burring. This is because the strength difference between the ferrite and martensite is large, so that large strain concentration and stress concentration occur in the ferrite near the martensite due to molding, and cracks occur.

  Based on this knowledge, a high-strength steel sheet having a high hole expansion ratio by reducing the difference in strength between structures has been developed. Patent Document 1 proposes a steel sheet in which bainite or bainitic ferrite is used as a main phase to secure strength and greatly improve hole expandability. By using a single structure steel, the above-mentioned strain concentration and stress concentration do not occur, and a high hole expansion ratio can be obtained.

  However, the single structure steel of bainite or bainitic ferrite deteriorates the elongation greatly, and it is not possible to achieve both elongation and hole expandability.

  Further, in recent years, a high-strength steel sheet has been proposed in which ferrite having excellent elongation is used as the structure of a single-structure steel and the strength is enhanced by using carbides such as Ti and Mo (for example, Patent Document 2, Patent Documents 3 and 4).

  However, the steel sheet proposed in Patent Document 2 contains a large amount of Mo. The steel sheet proposed in Patent Document 3 contains a large amount of V. Furthermore, the steel sheet proposed in Patent Document 4 requires cooling during rolling in order to refine the crystal grains. Therefore, there is a problem that alloy cost and manufacturing cost increase. Also, in these steel sheets, the elongation has deteriorated due to the large strength of ferrite itself. Although the elongations of bainite and bainitic ferrite single-structure steels are higher, the elongation-hole expandability balance was not always sufficient.

  Patent Document 5 proposes a steel sheet having a composite structure in which the martensite in the DP steel is bainite and the strength difference between the structures of the steel and ferrite is reduced to improve the hole expandability.

  However, as a result of increasing the area fraction of the bainite structure in order to secure the strength, the elongation deteriorates and the elongation-hole expandability balance is not sufficient.

  Further, Patent Document 7, Patent Document 8 and Patent Document 9 also propose a steel sheet in which the strength difference from the hard structure is reduced by precipitation strengthening the ferrite of DP steel.

  However, in this technique, Mo is an essential element, and there is a problem that the manufacturing cost becomes high. Even if ferrite is precipitation-strengthened, the strength difference with martensite, which is a hard structure, is large, and a high hole expandability improving effect is not obtained.

  On the other hand, in order to improve the fatigue characteristics of a material having no notch, it is effective to make the structure fine. For example, in Patent Document 10 and Patent Document 11, strength-ductility balance and fatigue limit ratio (ratio between fatigue strength and tensile strength (TS) (fatigue strength/TS)) in which the average grain size of ferrite is 2 μm or less. Of good hot rolled steel sheet. Further, since fatigue cracks are generated from the vicinity of the surface of the material, it is described in Patent Document 12 that the structure in the vicinity of the surface is made finer, and in Patent Document 13 that further improvement of the fatigue property by making the martensite structure finer. ing. However, although grain refinement improves the fatigue properties of materials without notches and does not impair ductility, the strength increase is small compared to strengthening mechanisms such as solid solution strengthening and precipitation strengthening. However, there was a limit to ensuring the fatigue characteristics by only fine graining.

  Solid solution strengthening and precipitation strengthening generally give a large increase in strength and are applied to high strength steel sheets. In Non-Patent Document 1, the effect of solid solution strengthening, precipitation strengthening, and fine grain strengthening on tensile strength and fatigue properties was investigated, and the amount of increase in fatigue strength with respect to the amount of increase in tensile strength was solid solution strengthening>precipitation strengthening. > It is reported that the order of fine grain strengthening is affected. Non-Patent Document 2 reports that when the state of Cu in steel is changed from solid solution to precipitation, the fatigue properties deteriorate. Although precipitation strengthening has a larger increase in tensile strength per alloy addition than solid solution strengthening, it has been reported that the fatigue property increase is small compared to solid solution strengthening as described above, so fatigue properties are important. Steel plates with solid solution strengthening have been used preferentially for these parts.

  For the improvement of notch fatigue properties, it has been reported that it is effective to reduce the crack propagation speed by forming a composite texture. In Patent Document 14, hard bainite or martensite is dispersed in a structure having fine ferrite as a main phase, thereby achieving both the fatigue characteristics and the notch fatigue characteristics of a material having no notches. However, no method for improving the press formability is described, and no particular attention is paid to the hardness and shape of bainite and martensite, and it is considered that the press formability is not good.

  Patent Documents 15 and 16 report that the crack propagation speed can be reduced by increasing the aspect ratio of martensite in the composite structure. However, all of them are technologies applied to thick plates, and do not have workability such as ductility and hole expandability required for press molding. Therefore, it is difficult to use the steel plates described in Patent Documents 15 and 16 as automobile steel plates. As described above, the fatigue property of a material without cracks in precipitation strengthened steel is an issue, and regarding the compatibility of notch fatigue property, fatigue property of smooth material without notch and press formability, an effective solution means is proposed. It has not been.

When it comes to improving machinability, generally, "excellent machinability" means
・The tool used during cutting has little wear and has a long tool life,
・Chips ejected during cutting are finely divided and have excellent chip disposability,
・Low cutting resistance, which is the force that acts on the tool during machining,
・Good finish on the cutting and grinding surfaces,
And so on.

  The improvement of machinability is often dealt with mainly by optimizing the cutting method (cutting conditions, tool material, tool shape, etc.). However, the machinability may be governed by the characteristics of the steel material itself, and in particular, in order to stably secure a long tool life, improvement of the machinability of the steel material is an important issue.

  It is generally well known that the machinability of steel materials is improved by the addition of Pb. However, the addition of Pb not only causes an increase in the steel material price, but also may cause environmental pollution. Therefore, research on a technique for improving the machinability of a steel material without adding Pb has been advanced. A typical example thereof is a machinability improving technique by utilizing MnS that is a sulfide-based inclusion, and many studies have been made and some have been put into practical use.

  However, with respect to automobile steel sheets, MnS often becomes a starting point of cracking and causes a decrease in hole expandability and elongation. Therefore, inclusion of a large amount of S is often avoided. That is, regarding automobile steel sheets, it is strongly desired to improve machinability without positively utilizing inclusions.

  For example, in Patent Document 17, machinability is improved by utilizing spheroidized cementite to lower the yield stress, but since tool steel is in mind, a long heat treatment is required when manufacturing, and It is difficult to apply it to steel sheets for use. Further, in Patent Document 18, machinability is improved by controlling the ferrite fraction in the ferrite-pearlite structure, but since the entire steel sheet is softened, a carburization-heat treatment step is required to secure fatigue durability. Therefore, it is difficult to apply it to steel sheets for automobiles.

JP, 2003-193190, A JP, 2003-089848, A JP, 2007-063668, A JP, 2004-143518, A JP, 2004-204326, A JP, 2007-302918, A JP, 2003-321737, A JP, 2003-321738, A JP, 2003-321739, A JP-A-11-92859 JP, 11-152544, A JP 2004-212199 A JP, 2010-70789, A Japanese Patent Laid-Open No. 04-337026 JP, 2005-320619, A JP, 07-90478, A JP 2012-126953 A JP, 2011-89189, A

Takashi Abe et al.: Iron and Steel, 70th Year (1984) No. 10, p. 145 T. Yokoi et al.: Journal of Materials science, 36th (2001) page 5757. Masai Mizui et al.: CAMP ISIJ, 5th year (1992), page 1867.

  The present invention has been devised in view of the above-mentioned problems, and an object thereof is to provide a high-strength hot-rolled steel sheet excellent in ductility and fatigue characteristics and surface layer machinability, and a manufacturing method thereof. Is.

  The ductility defined here is the uniform elongation required for press formability. If the uniform elongation is small, the plate thickness is likely to decrease due to necking during press forming, which causes press cracking. Further, the uniform elongation means the plastic elongation (%) at the maximum test force of JIS Z 2241 2011, and may be referred to as u-El (uniform Elongation).

  The uniform elongation (u-El) may satisfy tensile strength (TS) ≧540 MPa and product of tensile strength and uniform elongation: TS×u-El ≧8000 MPa% in order to secure press formability. is necessary. In addition, the hole expansion ratio (λ) measured based on JIS Z 2256 2010 is set so that tensile strength and hole strength are set so as to prevent cracking due to insufficient stretch flangeability when the underbody parts are pressed. Product of expansion ratio: It is necessary to satisfy TS×λ≧36000 MPa%. However, the tensile strength (TS) represents the tensile strength measured based on JIS Z 2241 2011.

Further, the surface layer machinability defined here is the average cutting resistance (P _ave ) and the difference (ΔP) between the maximum value and the minimum value of the cutting resistance when the surface layer cutting test is performed. When P _ave and ΔP are large, the wear rate of the cutting tool is increased and chipping of the tip of the cutting tool due to galling is induced, so that the life of the cutting tool is shortened. In order to maintain a cutting tool life comparable to that of mild steel, it is necessary to satisfy the ratio of average cutting resistance and yield stress (YP): P _ave /YP≦0.60 (N/MPa) and ΔP≦300N. However, the yield stress (YP) is Rr _0.2 , which is the proof stress (total elongation method) measured based on JIS Z 2241 2011. The surface cutting test will be described later.

By optimizing the composition and manufacturing conditions of the high-strength hot-rolled steel sheet and controlling the structure of the steel sheet, we succeeded in manufacturing a high-strength hot-rolled steel sheet with excellent ductility, fatigue properties and surface machinability. The summary is as follows.
(1) The chemical composition is mass%,
C: 0.030 to 0.200%,
Mn: 0.10 to 3.00%,
Si: 2.00% or less,
P: 0.100% or less,
S: 0.0300% or less,
Al: 2.00% or less,
N: 0.0100% or less (0 is not included),
O: 0.0100% or less,
Ti: 0.010 to 0.380%,
Si+5Al: 0.50 to 12.00%,
A steel sheet whose balance is Fe and unavoidable impurities,
Tief represented by the formula (a) is 0.01 to 0.30%,
In the cross section of the steel sheet in the rolling direction viewed from the width direction, the case where the orientation difference between adjacent crystals is 15° or more is defined as a grain boundary, which is a region surrounded by the grain boundaries , and the equivalent circle diameter of the region is 0. .3 μm or more is used as the crystal grains, the crystal grains having an average orientation difference of 0.5° or less in the crystal grains have an area distribution at a depth of ¼ of the plate thickness from the steel plate surface. The ratio is 50.0% or more, and the area fraction is 95.0% or more at a depth of 10 μm from the steel plate surface.
Further, the total of martensite, tempered martensite, and retained austenite is 2.0% or more and 10.0% or less in area fraction at a depth of 1/4 of the plate thickness from the steel plate surface, and 10 μm depth from the steel plate surface. Then, the area fraction is 0.5% or more and 10.0% or less,
Ti mass in Ti carbide, relative Tief of formula (a), not less than 40% by mass ratio, the mass of the circle-equivalent particle diameter of 7nm more 20nm following are among the Ti carbides, the 50% or more of the mass of the Ti carbide having a circle-equivalent particle diameter of 1 nm or more and 100 nm or less , and Vickers hardness HV1 in the ferrite grains at a depth of 10 μm from the steel plate surface and 100 μm depth from the steel plate surface. Hot rolling characterized in that Vickers hardness HV2 in the ferrite grain and Vickers hardness HV3 in the ferrite grain at a depth of ¼ of the plate thickness from the surface of the steel plate satisfy the formulas (b) and (c). steel sheet.
Tief=[Ti]−48/14×[N]−48/32×[S] (a)
However, [Ti][N][S] represent mass% of Ti, N, and S, respectively.
HV1/HV3≦0.80 (b)
HV2/HV3≧0.80...(c)
(2) Further, instead of part of the Fe, in mass%,
Nb: 0.010 to 0.100%,
V: 0.010 to 0.300%,
Cu: 0.01 to 1.20%,
Ni: 0.01-0.60%,
Cr: 0.01 to 2.00%,
Mo: 0.01 to 1.00%,
The hot-rolled steel sheet according to (1), containing one or more of the above.
(3) Further, in place of part of the Fe, in mass%,
Mg: 0.0005 to 0.0100%,
Ca: 0.0005 to 0.0100%,
REM: 0.0005 to 0.1000%,
The hot-rolled steel sheet according to any one of (1) and (2), which contains one or more of the above.
(4) Further, in place of part of the Fe, in mass%,
B: 0.0002 to 0.0020%,
The hot-rolled steel sheet according to any one of (1) to (3), characterized by containing.
(5) A method for manufacturing a hot-rolled steel sheet according to any one of (1) to (4), comprising a slab having the component composition according to any one of (1) to (4), When heating to a temperature of T1 or higher defined by the formula (d) and hot rolling the heated slab including finish rolling consisting of continuous rolling of a plurality of stages,
Of the stages of finish rolling in which the time (t1) from the central temperature of the steel sheet to 1000° C. to the Ar3 temperature determined by the formula (f) is set to 9.0 seconds or less, and the reduction ratio in the finish rolling is 10% or more. The second stage from the rearmost stage is the finish rolling stage so that the rolling roll temperature is 450°C or higher, and the rearmost stage is the finishing rolling stage, and the rolling roll temperature is 300°C or lower. And the rolling temperature of the steel plate after rolling in the rearmost stage is (T2-20)° C. or higher and (T2+100)° C. or lower with respect to the temperature T2 defined by the formula (e), and then The obtained hot-rolled steel sheet is cooled at an average cooling rate of 20° C./second or more to a temperature of 730° C. or more and 830° C. or less, and then held in a temperature range of 730° C. or more and 830° C. or less for 3 seconds or more, and then an average cooling rate. A method for producing a hot-rolled steel sheet, which comprises cooling at 40° C./sec or more and then winding at a temperature of 300° C. or less.
T1 (° C.)=7000/{2.75-log([Ti]×[C])}-273...Equation (d)
T2 (° C.)=870+10×([C]+[N])×[Mn]+350×[Nb]+250×[Ti]+40×[B]+10×[Cr]+100×[Mo]+100×[V] ...Equation (e)
Ar3(°C)=868-396×[C]-68.1×[Mn]+24.6×[Si]-36.1×[Ni]-24.8×[Cr]-20.7×[Cu ]+250×[Al]... Formula (f)
However, when the Ar3 temperature calculated by the formula (f) exceeds 900°C, Ar3=900°C. In addition, the [elemental symbol] in formula (a), formula (d), formula (e), and formula (f) indicates the mass% of the element in parentheses in the steel sheet (hereinafter the same in the present specification. ).

  The hot-rolled steel sheet according to the present invention has an effect of being excellent in ductility, fatigue characteristics, and surface machinability. If this steel sheet is used, it is possible to extend the press formability and fatigue life of the material applied to the underbody parts of the automobile material at low cost, and it is considered that the weight reduction of the automobile body is promoted, and it is industrially possible. The contribution of is remarkable.

It is a figure explaining a repeated stress strain curve and a repeated yield stress. It is a figure which shows the example of the low cycle fatigue test piece without a notch. It is a figure which shows the example of the fatigue test piece with a notch. It is a figure which shows the outline of a surface layer machinability evaluation test. It is a figure which shows the definition of cutting resistance.

  The contents of the present invention will be described in detail below.

[Chemical composition of steel sheet]
First, the reasons for limiting the chemical components of the hot-rolled steel sheet of the present invention will be described. In addition,% of the content regarding a component is mass %. Further, unless otherwise specified, the notation of [elemental symbol] indicates the mass% of the element symbol contained. For example, [Si] indicates mass% of contained Si.

(C: 0.030% to 0.200%)
C is one of the important elements in the present invention. C forms martensite and stabilizes austenite, and since it forms Ti carbide, it greatly contributes to the improvement of strength of the hot-rolled steel sheet by structural strengthening and precipitation strengthening. If it is less than 0.030%, a strength of 540 МPa cannot be secured. Preferably, since the notch fatigue property tends to improve as the hard phase fraction increases, 0.050% or more is preferably added.
On the other hand, if added in excess of 0.200%, the area fraction of the low temperature transformation product which is the hard second phase increases and the hole expandability deteriorates. Since the hole expandability tends to improve as the added amount decreases, it is preferably 0.180% or less, and more preferably 0.150% or less. Therefore, the content of C is set to 0.030% to 0.200%.

(Si: 2.00% or less)
Si is an element that is a deoxidizing element and at the same time is involved in the formation of ferrite, and expands the ferrite region temperature to a high temperature side as the content thereof increases, thereby expanding the two-phase region temperature region of ferrite and austenite. If added in excess of 2.0%, the ductility will drop significantly, so the Si content should be 2.00% or less. Originally, it is desirable to contain Si in order to obtain the composite structure steel of the present invention, but it is not limited to the case where Al that expands the two-phase temperature range is added as in the case of Si. Therefore, as described later, Si+5Al only needs to satisfy 0.50% to 12.00%, and therefore the lower limit of Si is not particularly limited.
If the Si content is less than 0.05%, spindle-shaped scale defects may increase. Therefore, it is desirable that the Si content be 0.05% or more. On the other hand, in order to avoid a decrease in ductility, it is desirable that the content be 1.5% or less.

(Mn: 0.10 to 3.00% or less)
In addition to solid solution strengthening, Mn is added to enhance hardenability and generate martensite or austenite in the steel sheet structure. If Mn is added in excess of 3.00%, a segregation zone of Мn is generated in the central portion of the steel sheet in the sheet thickness direction, and this segregation zone serves as a starting point of cracking, so that the hole expansion rate decreases. On the other hand, if the Mn content is less than 0.10%, it is difficult to exert the effect of suppressing pearlite, which causes a reduction in the hole expansion rate during cooling. Therefore, the Mn content is set to 0.10% to 3.00%.
From the viewpoint of ensuring sufficient hardenability and ensuring the effect of suppressing pearlite, the Mn content is preferably 0.30% or more, and more preferably 0.50% or more. On the other hand, from the viewpoint of suppressing the decrease in ductility due to an increase in center segregation, the Mn content is preferably 2.50% or less, and more preferably 2.00% or less.

(P: 0.100% or less)
P is an impurity contained in the hot metal, and is an element that segregates at the grain boundaries and decreases the low temperature toughness as the content increases. For this reason, the lower the P content is, the more preferable it is. If the P content exceeds 0.100%, workability and weldability are adversely affected. Therefore, the P content is set to 0.100% or less. Particularly, considering the weldability, the P content is preferably 0.030% or less.

(S: 0.0300% or less)
S is an impurity contained in the hot metal, and when the content is too large, it is an element that not only causes cracks during hot rolling but also produces inclusions such as MnS that deteriorate the hole expandability. Therefore, the content of S should be reduced as much as possible, but if it is 0.0300% or less, it is in an allowable range, so it is set to 0.0300% or less. However, when the hole expandability is required to some extent, the S content is preferably 0.0100% or less, more preferably 0.0050% or less.

(Al: 2.00% or less)
Al is an element effective as a deoxidizing agent for molten steel. In addition, Al is a strong ferrite-forming element and has the effect of increasing the Ar3 temperature, so that the crystal grains having an average orientation difference of 0 to 0.5° are 50.0% in area fraction. It is an effective element for achieving the above. If the Al content exceeds 2.00%, cracking may occur during rolling. Therefore, the upper limit of the Al content is 2.00%.
In the present invention, as will be described later, Si+5Al only needs to satisfy 0.50% to 12.00%, so the lower limit of the Al content is not particularly limited.
In addition, when the amount of Al is small, deoxidation may be insufficient, which may cause flaws on the surface of the slab. Therefore, the Al content is preferably 0.02% or more. On the other hand, if the Al content exceeds 0.50%, the toughness and ductility may start to deteriorate, so the upper limit of the Al content is preferably 0.50%.

(Si+5Al: 0.50% to 12.00%)
Since Si and Al are elements that expand the ferrite region temperature to the high temperature side and expand the two-phase region temperature region of ferrite and austenite, a crystal having an average misorientation in crystal grains of 0 to 0.5° In order to make the grains have an area fraction of 50.0% or more, the lower limit of the content of Si+5Al is 0.5%, preferably 1.0%. Further, if Si+5Al exceeds 12%, ductility and toughness are significantly lowered and cracks may occur during rolling. Therefore, the upper limit of Si+5Al is preferably set to 12%.

(N: 0.0100% or less (0 is not included))
Since N exists as TiN, it contributes to the improvement of the low temperature toughness through the refinement of the crystal grain size at the time of heating the slab, so N may be added. However, since nitride in steel lowers the hole expansion rate, it is necessary to set it to 0.0100% or less. Desirably, it is 0.0050% or less. On the other hand, it is economically undesirable to set the content to less than 0.0005%, so it is preferable to set it to 0.0005% or more.

(O: 0.0100% or less)
O forms an oxide and deteriorates formability, so it is necessary to suppress the addition amount. In particular, when O exceeds 0.0100%, this tendency becomes remarkable, so it is preferable to set it to 0.0100% or less. On the other hand, since it is economically unfavorable to set it to less than 0.0010%, it is desirable to set it to 0.0010% or more.

(Ti: 0.010 to 0.380%)
Ti is added in order to achieve both excellent fatigue strength and high strength due to precipitation strengthening. If Ti is less than 0.010%, the effect of precipitation strengthening cannot be obtained, so 0.010% or more may be added. If over 0.380% is added, the above effect is saturated and the economical efficiency is lowered. Therefore, the Ti content is set to 0.010% to 0.380%.

(Tief: 0.010 to 0.300%)
Tief=[Ti]−48/14×[N]−48/32×[S] (a)
Ti nitride and Ti sulfide are formed at a higher temperature than Ti carbide. For this reason, Ti carbide cannot fully be generated when there are many N and S in steel. Therefore, the index called Tief represented by the formula (a) was used as an index related to the formation of Ti carbide. When Tief is less than 0.010%, the precipitation amount of Ti carbide becomes small, and it becomes impossible to achieve both fatigue strength and high strength due to Ti carbide. Desirably, Tief is 0.025% or more. Further, even if Tief is added in an amount of more than 0.300%, the above effect is saturated and the economical efficiency is lowered. Further, if Tief exceeds 0.150%, the tundish nozzle may be easily clogged during casting, so 0.150% or less is desirable.

  The above is the basic chemical composition of the hot rolled steel sheet of the present invention, and the balance is composed of Fe and unavoidable impurities. Instead of part of Fe, the following components can be further contained.

(Nb: 0.100% or less)
Nb may be added because carbonitride or solid solution Nb delays the grain growth during hot rolling to make the grain size of the hot rolled sheet finer and improve the low temperature toughness. Even if the Nb content exceeds 0.100%, the above effect is saturated and the economical efficiency is lowered. Further, if the Nb content is less than 0.010%, the above effect cannot be sufficiently obtained, so the Nb content is preferably set to 0.010% or more.

(V: 0.300% or less)
V is an element that has the effect of improving the strength of the hot rolled steel sheet by precipitation strengthening or solid solution strengthening, and may be added. Even if the V content exceeds 0.300%, the above effect is saturated and the economical efficiency is lowered. Further, if the V content is less than 0.010%, the above effect cannot be sufficiently obtained, so the V content is preferably 0.010% or more.

(Cu: 2.00% or less)
Cu is an element that has the effect of improving the strength of the hot rolled steel sheet by precipitation strengthening or solid solution strengthening, and may be added. Even if the Cu content exceeds 2.00%, the above effect is saturated and the economical efficiency is deteriorated. Therefore, it is preferable to contain 2.00% or less. If the Cu content exceeds 1.20%, scratches due to scale may occur on the surface of the steel sheet. Therefore, the Cu content is preferably set to 1.20% or less. On the other hand, if the content of Cu is less than 0.01%, the above effect cannot be sufficiently obtained, so the content of Cu is preferably 0.01% or more.

(Ni: 2.00% or less)
Ni is an element that has the effect of improving the strength of the hot-rolled steel sheet by precipitation strengthening or solid solution strengthening, and may be added, but even if the Ni content exceeds 2.00%, the above effect is not obtained. Since it saturates and the economic efficiency decreases, it is preferable to contain 2.00% or less. If the Ni content exceeds 0.60%, the ductility may start to deteriorate. Therefore, the Ni content is preferably 0.60% or less. Further, if the Ni content is less than 0.01%, the above effect cannot be sufficiently obtained, so the Ni content is preferably 0.01% or more.

(Cr: 2.00% or less)
Cr is an element that has the effect of improving the strength of hot-rolled steel sheet by precipitation strengthening or solid solution strengthening, and may be added, or the above effect is saturated even if the Cr content exceeds 2.00%. Therefore, the economical efficiency is lowered, so it is preferable to contain 2.00% or less. Further, if the content of Cr is less than 0.01%, the above effect cannot be sufficiently obtained, so the content is preferably 0.01% or more.

(Mo: 1.00% or less)
Mo is an element that has the effect of improving the strength of the hot-rolled steel sheet by precipitation strengthening or solid solution strengthening, and may be added. Even if the Mo content exceeds 1.00%, the above effect is saturated. Therefore, the economical efficiency is lowered, so that it is preferable to contain 1.00% or less. Further, if the Mo content is less than 0.01%, the above effect cannot be sufficiently obtained, so the Mo content is preferably 0.01% or more.

(Mg: 0.0100% or less)
Mg is an element that acts as a starting point of fracture and controls the morphology of non-metallic inclusions that cause workability to deteriorate, and improves workability. Therefore, Mg may be added, and the content of Mg is 0. Even if added over 0100%, the above effect is saturated and the economical efficiency is lowered, so 0.0100% or less is preferably contained. Moreover, since the effect becomes remarkable when the content of Mg is 0.0005% or more, the content of Mg is preferably 0.0005% or more.

(Ca: 0.0100% or less)
Since Ca is an element that acts as a starting point of fracture and controls the morphology of non-metallic inclusions that cause deterioration of workability and improves workability, it may be added, and the content of Ca is 0. Even if added over 0100%, the above effect is saturated and the economical efficiency is lowered, so 0.0100% or less is preferably contained. Further, since the effect becomes remarkable when the content of Ca is 0.0005% or more, the content of Ca is preferably 0.0005% or more.

(REM: 0.1000% or less)
REM (rare earth element) is an element that acts as a starting point of fracture and controls the morphology of non-metallic inclusions that cause workability to deteriorate, and improves workability, so it may be added, and REM is included. Even if added in an amount of more than 0.1000%, the above effect is saturated and the economical efficiency is deteriorated. The content of REM becomes 0.0005% or more, and the effect becomes remarkable. Therefore, it is preferable that the content of REM be 0.0005% or more.

(B: 0.0100% or less)
B segregates at the grain boundaries and improves the low temperature toughness by increasing the grain boundary strength. From this, it may be added, and if the added amount of B exceeds 0.0100%, the effect is saturated and the economy is poor, so 0.0100% or less is preferably contained. Further, this effect becomes remarkable when the amount of B added to the steel sheet is 0.0002% or more. Further, B is a strong quenching element, and when more than 0.0020% is added, the area fraction of the crystal grains such that the average orientation difference in the crystal grains is 0 to 0.5° is reduced. Therefore, the B content is preferably 0.0002% or more.

  It has been confirmed that the effects of the present invention are not impaired even if the total content of Sn, Zr, Co, Zn, and W of other elements is 1% or less. Among these elements, Sn is preferably contained in 0.05% or less because it may cause flaws during hot rolling.

[Microstructure of steel sheet]
The microstructure of the steel sheet will be described.
The microstructure of the steel sheet is subjected to EBSD analysis at a measurement interval of 1 μm or less using the EBSD method (electron beam backscattering diffraction pattern analysis method) often used for crystal orientation analysis. Regarding the crystal orientations measured at the measurement intervals of 1 μm or less obtained by EBSD analysis, the case where the crystal orientation difference between adjacent crystals is 15° or more is defined as a grain boundary, which is a region surrounded by the grain boundaries. A grain having a circle-equivalent grain size of 0.3 μm or more is defined as a crystal grain.

Then, with respect to the orientations of all the measurement points inside the crystal grains, the orientation difference between the adjacent measurement points is obtained, and the average value of these orientation differences is set as the average of the orientation differences within the crystal grains.
In the steel sheet of the present invention, the crystal grains having an average orientation difference of 0.5° or less in the crystal grains are located at a position at a depth of ¼ of the plate thickness from the surface of the steel sheet (“a thickness of ¼ depth”). In some cases, it is referred to as “position”), and the area fraction is 50.0% or more with respect to all the crystal grains, and the position is 10 μm deep from the surface of the steel sheet (may be referred to as “position of 10 μm depth”. In the above, it is preferable that the area ratio of all the crystal grains is 95.0% or more. Since such crystal grains have high ductility and are precipitation strengthened by Ti carbide, by securing such crystal grains at a certain ratio or more, while maintaining the tensile strength (TS) at 540 MPa or more, The ductility can be improved.

The ratio of crystal grains having an average orientation difference within the crystal grains of 0.5° or less can be measured, for example, by the following method.
In the width direction of the steel plate, at a position of 1/4 W (width) or 3/4 W (width) of the plate width from the plate end, the cross section (width direction cross section) of the steel plate in the rolling direction is the observation surface. Samples are collected so that the area of 20,000 μm ² or more of the steel sheet is EBSD-analyzed at a measurement interval of 0.2 μm at positions of ¼ depth and 10 μm depth from the steel sheet surface.

Here, the EBSD analysis uses a device composed of a thermal field emission scanning electron microscope (for example, JEOL JSM-7001F) and an EBSD detector (for example, TSL HIKARI detector) at 50 to 300 points/sec. Perform at analysis speed. With respect to the obtained crystal orientation information, crystal grains having an orientation difference of 15° or more in the crystal grains and a circle equivalent grain size (diameter) of 0.3 μm or more are extracted, and the average orientation difference in the crystal grains is calculated. Then, the area ratio of the crystal grains having an average orientation difference within the crystal grains of 0.5° or less is obtained. The definition of the crystal grains and the calculation of the average orientation difference within the crystal grains can be obtained using, for example, the software “OIM Analysis ^™ ” attached to the EBSD analysis device.

  When the average grain orientation difference within the crystal grains is 0.5° or less at a position of 1/4 depth of thickness and the area fraction is less than 50.0%, the ductility deteriorates and the tensile strength increases. Strength (TS)×uniform elongation (u-El)≧8000 MPa% cannot be satisfied. The higher the area fraction of the crystal grains having an average orientation difference of 0.5° or less, the more the ductility is improved. Therefore, the area fraction is preferably 60.0% or more, and more preferably the area fraction. It is better to set it to 80.0% or more. Area fraction

When the area of the crystal grains having an average orientation difference of 0.5° or less at the depth of 10 μm is less than 95.0% of all the crystal grains, the cutting of the surface layer is performed. The resistance increases, and P _ave /YP≦0.60 cannot be satisfied, and the machinability of the surface layer deteriorates.

  The crystal grains in which the average orientation difference in the crystal grains is 0.5° or less in the present embodiment are not directly related to the ferrite defined by the observation result of the optical microscope. In other words, for example, even if there is a hot-rolled steel sheet having a ferrite area fraction of 50.0% or more, the proportion of crystal grains in which the average orientation difference within the crystal grains is 0.5° or less is 50.0%. This is not always the case. Therefore, the characteristics equivalent to the hot-rolled steel sheet according to the present embodiment cannot be obtained only by controlling the ferrite area fraction.

[Martensite + tempered martensite + retained austenite area fraction]
In the steel sheet of the present invention, the total area fraction of martensite, tempered martensite, and retained austenite is 2.0% or more and 10.0% or less in the area fraction at the position of 1/4 depth of the plate thickness. There is preferably 0.5% or more and 10.0% or less at the position of 10 μm depth. The hard phase composed of martensite or tempered martensite or retained austenite is an obstacle to fatigue crack propagation in the soft phase and has the effect of reducing the fatigue crack propagation rate, thus improving notch fatigue properties. Contribute. From this, the total of the hard phases of martensite, tempered martensite, and retained austenite is 2.0% or more in area fraction at the 1/4 depth position of the plate thickness. When the fraction of the hard phase is less than 2.0%, the ratio of the fatigue limit (c-FL) and the tensile strength (TS) in the notch fatigue test, which is a guideline for the high-strength steel sheet having excellent notch fatigue properties, is: Since the c-FL/TS is less than 0.25, the hard phase fraction is preferably 2.0% or more. More preferably, it is 5% or more. The method for obtaining the fatigue limit (c-FL) in the notch fatigue test will be described later.

  On the other hand, when the area fraction of these structures exceeds 10.0%, the hole expansion ratio (λ), which is one of the press formability, decreases, so that martensite, tempered martensite, and retained austenite are combined. The area fraction must be 10.0% or less.

  The presence of martensite, tempered martensite, and retained austenite at a depth of 10 μm from the surface of the steel sheet causes the surface layer machinability, more specifically, the difference (ΔP) between the maximum and minimum values of cutting resistance. It has a great influence. When the size of the chips generated at the tip of the cutting tool during cutting becomes maximum, the cutting resistance becomes maximum. On the other hand, since the cutting resistance is minimized when the chips are separated from the work material, it is desirable to separate the chips from the work material before the chips grow large in order to reduce ΔP. Martensite, tempered martensite, and retained austenite at a depth of 10 μm from the surface of the steel sheet serve as void (defect) generation sites during surface layer cutting, which promotes chip separation, resulting in smaller ΔP and cutting. Has the effect of extending the tool life. This effect is exhibited when the area fraction of martensite, tempered martensite, and retained austenite at a position 10 μm deep from the steel plate surface is 0.5% or more. On the other hand, if the area fraction of martensite, tempered martensite, and retained austenite at a depth of 10 μm from the surface of the steel plate exceeds 10.0%, the average cutting resistance increases, so that 10 μm from the surface of the steel plate. The area fraction of the martensite at the depth position, the tempered martensite, and the retained austenite are preferably 0.5% or more and 10.0% or less.

  The area fractions of martensite and tempered martensite that compose the steel sheet structure of the present invention are measured using FE-SEM (field emission scanning electron microscope). The area fraction of retained austenite is measured using the EBSD method (electron beam backscattering diffraction pattern analysis method). Specifically, the area fractions of martensite and tempered martensite that compose the steel sheet structure of the present invention are measured from one end of the steel sheet to a width W at a position of 1/4 W or 3/4 W in the width direction. A sample is taken so that the cross section becomes the observation surface (hereinafter, the collected material is referred to as a “texture measurement sample”), the observation surface is polished, and nital-etched, and a quarter of the plate thickness, and It was determined by observing a portion having a depth of 10 μm from the surface with an FE-SEM (field emission scanning electron microscope). A martensite having a lath-like (thin and long plate-like) structure in which no carbide was precipitated was observed as observed by FE-SEM. A structure having a lath-like structure in which carbide was precipitated in a multivariant manner (cementite was precipitated in various directions) was designated as tempered martensite.

It should be noted that, when the carbide was precipitated in a single variant (the cementite was uniformly precipitated in one direction), it was determined to be bainite. Using FE-SEM, a region of 120,000 μm ² or more was measured at a magnification of 1000 times for 10 fields of view of a quarter of the plate thickness and a portion 10 μm deep from the surface. The area fractions of martensite and tempered martensite were obtained for each field of view, and the average value thereof was used as a typical area fraction of martensite and tempered martensite.

The area fraction of retained austenite was measured using the EBSD method (electron beam backscattering diffraction pattern analysis method) after removing the processed layer by electrolytic polishing of the structure measurement sample. The crystal grains for which backscattering characteristic of FCC metal was obtained were defined as retained austenite. Using the EBSD method, a region having a thickness of 1/4 of the plate thickness and a portion having a depth of 10 μm from the surface was observed at a region of 120,000 μm ² or more. The area fraction of retained austenite was obtained for each range, and the average value of them was used as a typical area fraction of retained austenite.

[Vickers hardness in ferrite grains]
Next, the Vickers hardness in the ferrite grains will be described. Incidentally, the ferrite grains referred to here, when the steel plate is mirror-polished after the mirror polishing and observed with an optical microscope, for the crystal grains partitioned by the grain boundary portion to be corroded in black, the corrosion unevenness in the crystal grains is small, A crystal grain that appears to be a single color. When measuring the Vickers hardness of ferrite grains, it is necessary to polish the steel sheet in advance and then perform nital corrosion.

The steel sheet according to the present invention has a Vickers hardness of HV1 in a ferrite grain having a depth of 10 μm from the steel sheet surface, a Vickers hardness of HV2 in a ferrite grain having a depth of 100 μm from the steel sheet surface, and a sheet thickness of 1/4 depth from the steel sheet surface. When the Vickers hardness in the ferrite grains is HV3, it is necessary to satisfy the expressions (b) and (c).
HV1/HV3≦0.80 (b)
HV2/HV3≧0.80...(c)
If Expression (b) is satisfied, the value of P _ave /YP will be 0.60 (N/MPa) or less. That is, if the hardness (HV1) of the portion to be cut during the surface layer cutting is smaller than 80% of the hardness at the plate thickness 1/4 depth position that is representative of the hardness of the entire steel plate, the average cutting of the surface layer is performed. Resistance is low with respect to yield stress and cutting tool life is extended. On the other hand, when the expression (b) is not satisfied, the value of P _ave /YP exceeds 0.60, and it cannot be said that the material has good surface layer machinability.

If Expression (c) is satisfied, the value of c-YP/YP becomes 0.90 or more. In the material that does not satisfy the formula (c), the structure with low hardness spreads from the surface of the steel plate to a depth of more than 100 μm, so P _ave /YP≦0.60 is satisfied, but it is repeated compared with the yield stress (YP). The yield stress (c-YP) becomes low, and the ratio of cyclic yield stress (c-YP) to yield stress (YP): c-YP/YP becomes less than 0.90. That is, the hardness at the position immediately below the surface layer to be cut (the position at a depth of 100 μm from the steel plate surface) is about the same as the hardness at the center of the steel plate (1/4 depth position of the plate thickness) (80% in Hv value). In the above case, only the hardness of the surface portion (surface layer) of the steel sheet is small. It is necessary to reduce only the ferrite hardness of the surface portion to be cut in order to achieve both surface machinability and fatigue characteristics.

[Ti carbide]
Next, the state and amount of Ti carbide in the structure will be described. It has been conventionally reported that precipitation strengthening by Ti carbide is inferior in fatigue properties to solid solution strengthening by Si or the like. However, as a result of intensive studies by the inventors, Ti of 40% or more of Tief (obtained by the formula (a)) calculated from Ti, N, and S in steel is precipitated as Ti carbide, Among the precipitated Ti carbides, those having a circle-equivalent particle diameter (diameter) of the Ti carbide (the particle diameter in this specification simply means the circle-equivalent particle diameter) of 7 nm to 20 nm are Ti carbides. It has been found that, when the equivalent circle diameter is 1 nm or more and 100 nm or less, when the mass fraction is 50% or more, fatigue properties equivalent to or higher than those of solid solution strengthening can be obtained.

As described in Non-Patent Document 3, a composite microstructure steel having a high cyclic yield stress (c-YP) has good low cycle fatigue properties and high cycle fatigue properties. Cyclic yield stress (c-YP) is the resistance of a material to deformation after being subjected to repeated deformation, and solid solution strengthened steel that does not lose the strengthening function of the steel sheet even after repeated deformation, Among the precipitation strengthening, those that utilize precipitates having a large grain size will be large. It is known that a material having a large ratio of cyclic yield stress (c-YP) to yield stress (YP) measured by a tensile test has good low cycle fatigue characteristics and high cycle fatigue characteristics, and A steel plate satisfying the ratio (c-YP/YP) of 0.90 or more has a good yielding property (YP) despite its low yield stress (YP), and has an excellent balance between productivity and fatigue properties during press forming. Precipitation strengthening when the particle size of Ti carbide is small and the area fraction of Ti carbide having a particle size of 7 nm to 20 nm is less than 50% of the Ti carbide having a circle equivalent particle size of 1 nm or more and 100 nm or less. As a result, the yield stress (YP) is greatly increased, but the increase amount of the repeated yield stress (c-YP) is small, and the ratio (c-YP/YP) thereof is less than 0.90. This is because, as a result of the dislocation motion due to repeated deformation, the Ti carbide having a small grain size undergoes shear fracture, and the effect of suppressing the dislocation motion by the Ti carbide decreases. The method for obtaining the repeated yield stress (c-YP) will be described later.

On the other hand, when the Ti carbide having a grain size of 7 nm to 20 nm is sufficiently present, it is considered that the dislocation moves around the Ti carbide. Since a circular dislocation called an Orowan loop remains around the bypassed Ti carbide, the Orowan loop grows due to the dislocation motion, dislocation density increases and dislocation strengthening occurs, and the yield stress increases more than before repeated deformation. Then, as a result, the ratio of the repeated yield stress to the yield stress (c-YP/YP)≧0.90. In addition, the particle size is too large Ti carbide, if the particle size 7nm weight fraction of 20 nm Ti carbides of a circle equivalent diameter of the T i carbide is less than 50% of the weight of 1nm or 100nm following ones In particular, the effect of suppressing the dislocation motion by Ti carbide becomes small, and the effect of improving these fatigue characteristics becomes small.

  Similarly, even when Ti corresponding to less than 40% by mass of Tief calculated from Ti, N, and S in steel is precipitated as Ti carbide, the fatigue property is small because the suppression of dislocation motion by Ti carbide is small. The improvement effect becomes smaller. Therefore, the above-mentioned ratio (c-YP/YP) of repeated yield stress and yield stress cannot satisfy 0.90. Therefore, Ti corresponding to a mass of 40% or more of Tief calculated from Ti, N, and S in steel needs to be precipitated as Ti carbide. From the viewpoint of securing the effect of improving the fatigue characteristics, the Ti precipitated as Ti carbide is preferably 45% or more of Tief.

  The Ti carbide precipitation amount can be measured by electrolyzing the steel sheet and identifying the weight of Ti in the residual residue by chemical analysis or the like. Specifically, the total weight of precipitated Ti is obtained from the weight of Ti in the residue. In addition, the weight of Ti deposited as Ti nitride (TiN) can be obtained from the weight of nitrogen contained in the steel. By subtracting the weight of Ti deposited as TiN from the total weight of Ti deposited, the weight of Ti in the Ti carbide can be obtained. Thereby, the mass% of Ti contained in the Ti carbide can be obtained.

  Further, the means for identifying the Ti carbide and measuring the particle size is not particularly specified, but by using, for example, 3D-AP (three-dimensional atom probe), the Ti carbide can be measured by measuring the existing positions of Ti and C in the steel sheet. The particle size of Ti carbide having a small particle size can be measured with high accuracy. Specifically, at least 20 fields of view having a width of 10 μm×10 μm are observed at a ¼ depth position of the plate thickness, and the magnification is increased according to the particle size of Ti carbide, and the particle size of 1 nm to 100 nm is used. The particle size distribution of the Ti carbide in the range may be determined, and the ratio of the particle size of 7 nm to 20 nm in the range may be determined by the weight ratio. If the Ti carbide has a particle size of 5 nm or more, an electron diffraction pattern is obtained using Fe-TEM, and the Ti carbide can be identified and measured by measuring the azimuth relationship with the matrix Fe and the lattice spacing. It is possible.

  The cyclic yield stress (c-YP) can be measured by executing a low cycle fatigue test under the condition that the strain amplitude is different. In this study, the low cycle fatigue test was performed at the strain amplitudes of 0.2%, 0.3%, 0.5%, 0.8% and 1.0% to determine the number of fatigue tests at the time of fatigue fracture. The maximum stress of each strain amplitude at the time when half the number of fatigue tests was performed was connected to create a repeated stress-strain curve. As shown in FIG. 1, a straight line having a Young's modulus slope is inserted into the cyclic stress-strain curve as shown in FIG. 1, and the intersection point with the cyclic stress-strain curve is defined as the cyclic yield stress (c-YP). Defined.

  The high-strength hot-rolled steel sheet of the present invention having the structure and composition as described above is provided with a hot-dip galvanized layer by hot dip galvanizing treatment on the surface, and further, an alloyed zinc-plated layer after alloying treatment after plating. The surface machinability can be improved by using such a material. The plating layer is not limited to pure zinc, and elements such as Si, Mg, Al, Fe, Mn, Ca, and Zr may be added to further improve the machinability of the surface layer. The provision of such a plating layer does not impair the excellent ductility and fatigue characteristics and surface machinability of the present invention. Further, the effect of the present invention can be obtained regardless of whether the surface treatment layer is formed by organic film formation, film lamination, organic salt/inorganic salt treatment, non-color treatment or the like.

[Steel plate manufacturing method]
The manufacturing method prior to hot rolling is not particularly limited. That is, various secondary smelting processes are performed after smelting in a blast furnace, an electric furnace or the like to adjust the composition to the above-described composition, and then casting may be performed by a method such as normal continuous casting or thin slab casting. At that time, scrap may be used as a raw material as long as it can be controlled within the component range of the present invention.

  The casting slab is heated to a predetermined temperature before starting hot rolling. In the case of continuous casting, the material may be once cooled to a low temperature and then reheated and then hot-rolled, or without being particularly cooled, continuous casting may be followed by heating and hot-rolling.

The slab heating time in hot rolling needs to be T1° C. or more represented by the formula (d). When ordinary slab casting is performed, the slab temperature drops to the Ar3 temperature or lower, so that Ti carbide precipitates in the structure. In order to control the grain size of Ti carbide, it is first necessary to solutionize the Ti carbide precipitated in the slab. If the slab heating temperature is lower than T1°C, the Ti carbide precipitated in the slab is not sufficiently solutionized, and the grain size of the Ti carbide cannot be controlled. Therefore, the temperature of the heating furnace is set to T1°C or higher.
The upper limit of the slab heating temperature is not specified. However, if the heating temperature is excessively high, the slab surface is oxidized and becomes scale, which is not economically preferable. From this, it is desirable that the upper limit of the slab heating temperature be 1300°C.
T1 (° C.)=7000/{2.75-log([Ti]×[C])}-273...Equation (d)

After heating the slab, the slab extracted from the heating furnace is subjected to a rough rolling step of hot rolling and a finishing rolling step thereafter to obtain a hot rolled steel sheet. During this hot rolling, the time (t1) from when the central temperature of the steel sheet falls to 1000° C. to when the Ar3 temperature (transformation point temperature) represented by the formula (f) is reached is preferably 9.0 seconds or less. ..
Ar3(°C)=868-396×[C]-68.1×[Mn]+24.6×[Si]-36.1×[Ni]-24.8×[Cr]-20.7×[Cu ]+250×[Al]... Formula (f)
However, when the Ar3 temperature calculated by the formula (f) exceeds 900° C., Ar3=900° C. and the time from when the steel plate temperature drops to 1000° C. to 900° C. is t1. The center temperature of the steel sheet described here can be obtained by thermal analysis from the temperature and heat history of the steel sheet surface.

In order to obtain the effect of the present invention, it is necessary that 50% or more by mass fraction of the precipitated Ti carbide having a circle equivalent particle size of 1 nm to 100 nm has a circle equivalent particle size of 7 nm to 20 nm. Is. The Ti carbide that precipitates in the austenite region has a thickness of 20 nm or more and does not contribute to the improvement of fatigue characteristics. Therefore, it is necessary to suppress the precipitation of Ti carbide in the austenite region. Precipitation of Ti carbide in the austenite region is remarkable when the center temperature of the steel sheet is 1000° C. or lower, and therefore it is effective to shorten the time during which the center temperature of the steel sheet is maintained at the Ar3 temperature or higher and 1000° C. or lower. According to the study by the present inventors, when this time is longer than 9.0 seconds, coarse Ti carbide is precipitated, and the area fraction of Ti carbide having a grain size of 7 nm to 20 nm is reduced to a desired mass fraction. 50% cannot be obtained. There is a possibility that the central temperature of the steel sheet may rise from the Ar3 temperature to 1000°C in the rough rolling rear stage, finish rolling, and the front stage part of the cooling device, and the temperature history of this process is controlled to maintain the temperature in this temperature range. It must be within 9.0 seconds.

The finish rolling is usually performed by multi-stage (for example, 6 or 7 stages) continuous rolling. In the finish rolling performed by this multi-stage continuous rolling, the rolling reduction is higher on the front side (upstream side) than on the rear side (downstream side), and the rolling rate is reduced on the rear side (downstream side). Sometimes. In the present invention, in the finish rolling performed in the multi-stage continuous rolling, among the finish rolling stages having a reduction ratio of 10% or more, the temperature (surface of the second rolling roll from the rearmost stage (downstream side)) Temperature) is set to 450° C. or higher, the temperature (surface temperature) of the rolling roll on the rearmost side is set to 300° C. or lower, and the rolling temperature for finish rolling on the rearmost side is the temperature T2 defined by the formula (e). Therefore, the temperature is preferably (T2-20)° C. or higher and (T2+100)° C. or lower. Desirably, the temperature is set to (T2-20)°C or higher and (T2+30)°C or lower.
T2 (° C.)=870+10×([C]+[N])×[Mn]+350×[Nb]+250×[Ti]+40×[B]+10×[Cr]+100×[Mo]+100×[V] ...Equation (e)

That is, it is not always necessary to control the rolling roll temperature and the rolling temperature of the rearmost finish rolling stage (downstream side), but the rearmost rolling stage (downstream side) of the finishing rolling stages having a rolling reduction of 10% or more. And, the rolling roll temperature and rolling temperature of the second stage finish rolling are controlled.
The rolling reduction is calculated by the following formula at each stage.
Reduction ratio (%) = (plate thickness on inlet side of rolling mill-plate thickness on outlet side of rolling mill)/(plate thickness on inlet side of rolling mill) x 100%

  Among finish rolling stages having a rolling reduction of 10% or more, the temperature of the rolling roll on the most rear stage side is set to 300° C. or less so that the hardness of the surface layer 10 μm portion of the steel sheet is reduced to satisfy the formula (b). This is because By lowering the temperature of the rolling rolls by utilizing the roll high-cooling equipment, the temperature of the steel sheet surface can be temporarily lowered and rolling can be performed in the completely unrecrystallized region. Since the austenite grains rolled in the completely unrecrystallized region have a high dislocation density, they are transformed at a high temperature, and it is possible to increase the proportion of crystal grains having an average orientation difference of 0.5° or less in the crystal grains, Ferrite hardness can be reduced. On the other hand, when the temperature of the rolling roll on the most rear side is more than 300° C., the formula (b) cannot be satisfied, and the machinability of the surface layer deteriorates.

  Further, as a result of intensive studies by the inventors, it was found that controlling the temperature of the rolling roll at the second stage from the rearmost stage is also important. If the temperature of the rolling roll is less than 450° C., the temperature of the steel sheet surface layer is significantly reduced, and the ferrite hardness of the portion 100 μm from the surface layer is also reduced by the mechanism described above, so that the formula (c) cannot be satisfied. As a result, the fatigue characteristics deteriorate, so the temperature of the second rolling roll from the rearmost stage is preferably 450° C. or higher.

  When the rolling temperature of the finishing rolling stage on the most rear side among the finishing rolling stages having a rolling reduction of 10% or more is less than (T2-20)° C., rolling is performed with austenite recrystallization suppressed. By doing so, the aspect ratio of austenite is increased. Since the shape is inherited by the hard phase such as martensite, tempered martensite, and retained austenite, the aspect ratio of the hard phase particularly in the central part of the plate thickness increases and the hole expansion ratio decreases.

  When the rolling temperature of the finishing rolling stage on the most rear side among the finishing rolling stages with a rolling reduction of 10% or more is higher than (T2+100)° C., the average orientation difference within the crystal grains is 0.5° or less. The area fraction of the crystal grains is less than 50.0%, and the ductility decreases. It is speculated that this is because the austenite coarsened after recrystallization and the transformation temperature decreased. Further, of the finishing rolling stages having a reduction ratio of 10% or more, the finishing rolling stage on the most rear side is (T2-20)° C. or more (T2+100) with respect to the temperature T2 defined by the formula (e). The reason why the rolling temperature is not higher than 0° C. is that if the rolling reduction is less than 10%, the grain refining effect is not obtained, and the rolling reduction is required to be 10% or more.

  Rolling with a rolling reduction of 40% or more imposes a heavy load on the rolling mill, and thus the rolling reduction is preferably less than 40%.

  After finishing rolling of the finishing rolling with a rolling reduction of 10% or more, the obtained steel sheet is cooled to a temperature of 730° C. or more and 830° C. or less (primary cooling). In this primary cooling, it is preferable to cool at an average cooling rate of 20° C./s or more. If the average cooling rate of primary cooling is less than 20° C./s, the area fraction of the crystal grains having an average orientation difference within the crystal grains of 0.5° or less is less than 50.0%, and the ductility decreases.

  It is considered that this is because austenite coarsens due to grain growth after rolling, and thus the austenite grain boundary interface integration rate, which is a nucleation site for ferrite transformation, decreases and the transformation temperature decreases. If the average cooling rate of the primary cooling is high, the transformation temperature rises and the ductility improves, so the upper limit is not particularly limited. However, if the average cooling rate exceeds 200° C./s, it is difficult to control the cooling stop temperature, so it is desirable to set it to 200° C./s or less.

  Following the primary cooling, the temperature is maintained in the temperature range of 730°C to 830°C for 3 seconds or more. As a result, the area fraction of crystal grains having an average orientation difference of 0.5° or less can be 50.0% or more. The holding time is preferably 5 seconds or longer in a temperature range of 750° C. or higher and 830° C. or lower, and more preferably 5 seconds or longer in the temperature range of 780° C. or higher and 830° C. or lower. The holding method and the heat pattern during the holding are not particularly limited as long as the material is retained in the holding temperature range for the holding time. For example, you may cool (air cooling etc.) so that it may stay in the said holding temperature range for 3 seconds or more.

  By increasing the holding temperature, the area fraction of the crystal grains having an average orientation difference of 0.5° or less increases. Therefore, the average orientation difference of the crystal grains is 0.5° or less. There is a possibility that the crystal grain of is a ferrite that has been transformed at a higher temperature among the ferrites that have been defined by conventional optical microscope observation. The upper limit of the holding time is not particularly specified, but if it is 15 seconds or more, the effect of increasing the area fraction of the crystal grains having an average orientation difference of 0.5° or less is saturated and the productivity decreases. Therefore, the holding time is preferably less than 15 seconds.

  If the holding temperature exceeds 830° C., only Ti having a mass corresponding to less than 40% of Tief represented by the formula (a) will precipitate as Ti carbide in the structure, and the fatigue characteristics will deteriorate. Further, if the holding temperature is lower than 730° C., the area fraction of the crystal grains having an average orientation difference within the crystal grains of 0.5° or less becomes less than 50.0%, and the ductility decreases.

  Then, secondary cooling is performed. In the secondary cooling, cooling is performed at an average cooling rate of 40° C./s or more, and the steel sheet temperature is wound up at 300° C. or less. If the average cooling rate of the secondary cooling is less than 40° C./s, the area fraction of the total structure of martensite, tempered martensite, and retained austenite is less than 2.0%, and the notch fatigue properties deteriorate.

  As the average cooling rate of the secondary cooling is higher, hard martensite, tempered martensite, and retained austenite are obtained, and the notch fatigue property becomes superior. For this reason, the upper limit of the average cooling rate is not limited, but large-scale equipment investment is required to perform cooling at 500° C./s or more, so it is desirable to set it to less than 500° C./s.

  In the winding step, the winding temperature is set to 300°C or lower. When the coiling temperature is higher than 300° C., the bainite fraction in the structure increases, it is difficult to secure the fraction of the hard phase, and the notch fatigue characteristics deteriorate. When notch fatigue characteristics are emphasized, a desirable winding temperature is 150°C or lower. This is because the notch fatigue characteristics improve as the hardness of the hard phase increases.

  For the purpose of improving ductility by correcting the shape of the steel sheet and introducing movable dislocations, skin pass rolling with a rolling reduction of 0.1% or more and 2% or less may be performed after the completion of all steps. After the completion of all the steps, the hot-rolled steel sheet obtained may be pickled, if necessary, for the purpose of removing scale attached to the surface of the obtained hot-rolled steel sheet. Further, after pickling, the obtained hot-rolled steel sheet may be subjected to in-line or off-line skin pass with a rolling reduction of 10% or less or cold rolling.

  In addition to the hot rolling step described above, even if a part of the continuous casting, pickling, and the like that accompanies the manufacturing is carried out, the excellent ductility and fatigue characteristics that are the effects of the present invention can be secured. Further, even if the hot rolled steel sheet is once produced and then heat-treated at a temperature range of 100 to 600° C. online or offline for the purpose of improving the ductility, the effect of the present invention is excellent in the rolling direction. Fatigue characteristics and workability can be secured.

The test results for confirming the action and effect of the present invention will be described.
Table 1 shows the components of the steel used in the test.
Table 2-1, Table 2-2, and Table 2-3 (in the present specification, these are collectively referred to as Table 2) show the steel types of the test pieces used in the test and the manufacturing conditions thereof.
Each test piece is shown in Table 3-1, Table 3-2, Table 3-3, Table 3-4, Table 3-5, and Table 3-6 (in the present specification, these are collectively referred to as Table 3). The evaluation results of are shown.

Among the mechanical properties, tensile strength characteristics (yield stress, tensile strength, uniform elongation) were measured according to JIS Z 2241 2011, which was taken at the position of 1/4 W or 3/4 W of the plate width with the width direction as the longitudinal direction. The No. 5 test piece of No. 5 was used for evaluation in accordance with JIS Z 2241 2011. For yield stress, 0.2% proof stress may be used, and for uniform elongation, plastic elongation (%) at maximum test force was used. The hole expansion ratio was measured based on JIS Z 2256 2010 using a hole expansion test piece collected from the same position as the tensile test piece collection position.
Further, the steel sheet according to the present invention has a tensile strength (TS)≧540 MPa, a tensile strength (TS)×uniform elongation (u-El)≧8000 MPa%, and a tensile strength (TS)×a hole expansion ratio (λ). It was confirmed that the steel plate satisfies ≧36000 MPa%.

  As a test piece for evaluating fatigue characteristics, a fatigue test piece having a shape shown in FIGS. 2 and 3 was sampled from the same position as the tensile test piece sampling position so that the width direction was the long side, and the test piece was subjected to the fatigue test. .. The fatigue test piece shown in FIG. 2 is a test piece for obtaining repeated yield stress, which is an index of the fatigue characteristics of a material having no notch, and the fatigue test piece shown in FIG. 3 is for obtaining fatigue strength of the notch material. It is a notch test piece produced in. The fatigue test piece was ground to a depth of about 0.05 mm from the outermost layer.

  A low cycle fatigue test is performed by using the test piece of FIG. 2 to perform strain control of both swings at a strain rate of 0.4%/s, a strain amplitude of 0.2%, 0.3%, 0.5% and 1.0%. Was executed. The maximum stress of each strain amplitude at the time of performing the fatigue test number of half of the fatigue test number at the time of fatigue fracture was connected, and the repeated stress-strain curve was created. As shown in FIG. 1, a straight line having a Young's modulus slope is inserted into this cyclic stress-strain curve at a point of 0.2% strain, and the intersection point with the stress-strain curve is defined as cyclic yield stress (c-YP). did.

  Using the test piece of FIG. 3, a stress control shaft fatigue test was performed at a stress ratio R=0.1 and a frequency of 5 Hz to evaluate notch fatigue characteristics. The stress that does not break after 10 million cycles was defined as the notch fatigue limit (c-FL).

The outline of the surface layer machinability evaluation test is shown in FIG. Two-dimensional cutting was performed under the cutting conditions of a cutting speed of 1.0 m/min, a cutting depth of 0.02 mm, a cutting width of 3.5 mm, and a cutting length of 30 mm, and the cutting resistance during processing was measured. As the tool, P20 type cemented carbide was used according to JIS B 4053:2013 (classification of use of cutting super hard tool material and how to assign a reference symbol). For the measurement of cutting resistance, a four-component dynamometer “9272” manufactured by KISTLER was used, and a cutting test was performed with the test material fixed on the dynamometer. FIG. 5 shows an example of test data obtained by the cutting test. The cutting resistance on the vertical axis is the composition of the vectors of the stress components obtained from the four-component dynamometer, and cutting resistance = √ {(stress component in the cutting direction) ² + (stress component in the direction perpendicular to the cutting direction) ) ² }. The average value P _ave of the cutting resistance and the difference ΔP between the maximum value and the minimum value of the cutting resistance are calculated from the data of FIG. 5, and P _ave /YP≦0.60 (N/MPa) and ΔP≦300 (N) are satisfied. The steel plate was a steel plate having excellent surface machinability.

  As shown in Table 3, it was confirmed that the hot-rolled steel sheet according to the example of the present invention had excellent ductility and fatigue characteristics and surface machinability.

On the other hand, Steel No. 2, 20 is the heating temperature is less than or equal to T1 ° C. defined by the formula (d), those Ti carbide does not condition of equivalent circle diameter of the T i carbides 100nm of less than 1nm The mass of the Ti carbide having a particle size of 7 to 20 nm was less than 50% with respect to the mass of , and the repeated yield stress (c-YP)/yield stress (YP)≧0.90 could not be satisfied.

  For steel Nos. 3 and 21, the rolling temperature of the finishing rolling stage on the most rear side among the finishing rolling stages having a rolling reduction of 10% or more was (T2-20 )° C., the tensile strength (TS)×hole expansion ratio (λ)≧36000 MPa% could not be satisfied. This is because rolling was performed in a state where austenite recrystallization was suppressed, the aspect ratio of austenite was increased, and its shape was inherited by hard phases such as martensite and tempered martensite and retained austenite, It is considered that this is because the aspect ratio of the hard phase in the thick center portion increased.

  Steel Nos. 6 and 24 are (T2+100)° C. with respect to the temperature T2 defined by the formula (e) for the rolling temperature of the finishing rolling stage on the most rear side among the finishing rolling stages having a rolling reduction of 10% or more. As described above, the area fraction of crystal grains having an average orientation difference of 0.5° or less was less than 50.0%, and tensile strength (TS)×uniform elongation (u-El)≧ 8000 MPa% could not be satisfied.

For Steel No. 7, 25 is the time until the 1000 ℃ ~Ar3 temperature was at least 9.0 seconds, the particle size relative to the mass of those circle-equivalent particle diameter is 1nm or more 100nm or less of the T i carbides 7~20nm The mass of Ti carbide was less than 50%, and c-YP/YP≧0.90 could not be satisfied.

  Steel Nos. 8 and 26 were the primary coolings, and the average cooling rate was less than 20° C./s, so the area fraction of the crystal grains having an average orientation difference of 0.5° or less was 50.0%. It was less than less than TS×u-El≧8000 MPa%.

  Since the holding temperatures of Steel Nos. 10 and 28 were less than 730° C., the area fraction of the crystal grains having an average orientation difference within the crystal grains of 0.5° or less was less than 50.0%, and TS×u- El≧8000 MPa% could not be satisfied.

  Since the holding start temperatures of Steel Nos. 13 and 31 were 830° C. or higher, less than 40% by weight of Tief represented by the formula (a), Ti, was not precipitated as Ti carbide in the structure, and c-YP /YP≧0.90 could not be satisfied.

  Steel Nos. 14 and 32 had a holding time of less than 3 seconds, so the average area difference of crystal grains in the crystal grains at a depth of 1/4 from the surface layer of the plate thickness was 0.5° or less. Was less than 50.0%, and TS×u-El≧8000 MPa% could not be satisfied.

  Steel Nos. 16 and 34 had an average cooling rate of secondary cooling of less than 40° C./s, so that the total structure of martensite, tempered martensite, and retained austenite at a depth of 1/4 from the surface layer of the plate thickness ( The area ratio of the (hard phase) was less than 2.0%, and c-FL/TS≧0.25 could not be satisfied.

  Steel Nos. 17 and 35 had a winding temperature of over 300° C., so the total structure (hard phase) of martensite, tempered martensite, and retained austenite at a depth of 1/4 from the surface layer of the plate thickness The ratio was less than 2.0%, and c-FL/TS≧0.25 could not be satisfied.

  Since the addition amount of C was less than 0.030% in Steel No. 37, the total structure (hard phase) of martensite, tempered martensite, and retained austenite at a depth of 1/4 from the surface layer of the plate thickness is the area. The fraction was less than 2.0% and the strength of 540 МPa could not be secured.

  Steel No. 40 contained more than 0.200% of C, so the total structure (hard phase) of martensite, tempered martensite, and retained austenite at a depth of 1/4 from the surface of the plate thickness is the area. The fraction was 10.0% or more, and TS×λ≧36000 MPa% could not be satisfied.

  Steel No. 46 could not satisfy TS×λ≧36000 MPa% because the added amount of Mn was more than 3.00%.

  In Steel No. 47, since the amount of Mn added was less than 0.10%, pearlite was generated during cooling and TS×λ≧36000 MPa% could not be satisfied.

  In Steel No. 48, since the amount of P added was more than 0.100%, the workability was deteriorated, and TS×u-El≧8000 MPa% and TS×λ≧36000 MPa% could not be satisfied.

  In Steel No. 49, the amount of S added was more than 0.0300%, so the hole expansion ratio decreased, and TS×λ≧36000 MPa% could not be satisfied.

  In Steel No. 50, the amount of Al added was more than 2.00%, so cracking occurred during rolling.

  In Steel No. 55, Si+5Al was less than 0.50%, so the average area difference of crystal grains in the crystal grains at a depth of 1/4 from the surface layer of the plate thickness is 0.5° or less It was less than 50.0% and TS×u-El≧8000 MPa% could not be satisfied.

  In Steel No. 56, the amount of N added was more than 0.0100%, so that the hole expansion ratio decreased and TS×λ≧36000 MPa% could not be satisfied.

  Steel No. 57 had a Tief of more than 0.300%, so the tundish nozzle was clogged during casting.

  Steel No. 59 had a Tief of less than 0.010%, so only Ti having a weight of less than 40% of Tief represented by the formula (a) was precipitated as Ti carbide in the structure, and c-YP/YP ≧0.90 could not be satisfied.

  Steel No. 68 had cracks during rolling because the amount of Si added was more than 2.00%.

  Steel Nos. 69 and 71 had a rolling reduction of 10% or more, and the rolling roll temperature of the second finishing rolling stage from the rearmost side was less than 450° C., so that the depth from the surface layer Since the tissue at a position of 100 μm in the direction softened, it was not possible to satisfy c-YP/TS≧0.90.

  Steel Nos. 70 and 72 were able to soften the surface layer structure because the rolling roll temperature of the finishing rolling stage on the most downstream side was more than 300° C. among the finishing rolling stages having a rolling reduction of 10% or more. In other words, Pave/YP≦0.60 could not be satisfied.

  The hot rolled steel sheet according to the present invention is excellent in ductility, fatigue characteristics, and surface machinability. Therefore, the hot-rolled steel sheet according to the present invention can be used for machine products requiring designability and the like. In particular, it can be applied to wheels for automobiles.

Claims (5)

The chemical composition is% by mass,
C: 0.030 to 0.200%,
Mn: 0.10 to 3.00%,
Si: 2.00% or less,
P: 0.100% or less,
S: 0.0300% or less,
Al: 2.00% or less,
N: 0.0100% or less (0 is not included),
O: 0.0100% or less,
Ti: 0.010 to 0.380%,
Si+5Al: 0.50 to 12.00%,
A steel sheet with the balance being Fe and inevitable impurities,
Tief represented by the formula (a) is 0.01 to 0.30%,
In a cross section of the steel sheet in the rolling direction viewed from the width direction, the case where the adjacent crystal has a misorientation of 15° or more is a grain boundary, which is a region surrounded by the grain boundaries , and the circle-equivalent grain size of the region is 0. .3 μm or more, when the crystal grains have an average misorientation of 0.5° or less, the crystal grains have an area distribution at a depth of ¼ of the plate thickness from the steel plate surface. The ratio is 50.0% or more, and the area fraction is 95.0% or more at a depth of 10 μm from the steel plate surface.
Further, the total of martensite, tempered martensite, and retained austenite is 2.0% or more and 10.0% or less in area fraction at a depth of 1/4 of the plate thickness from the steel plate surface, and 10 μm depth from the steel plate surface. Then, the area fraction is 0.5% or more and 10.0% or less,
Ti mass in Ti carbide, relative Tief of formula (a), not less than 40% by mass ratio, the mass of the circle-equivalent particle diameter of 7nm more 20nm following are among the Ti carbides, the 50% or more of the mass of the Ti carbide having a circle-equivalent particle diameter of 1 nm or more and 100 nm or less ,
Further, the Vickers hardness HV1 in the ferrite grains at a depth of 10 μm from the steel plate surface, the Vickers hardness HV2 in the ferrite grains at a depth of 100 μm from the steel plate surface, and the ferrite at a depth of ¼ of the plate thickness from the steel plate surface A hot-rolled steel sheet having a Vickers hardness HV3 in the grain satisfying the formulas (b) and (c).
Tief=[Ti]−48/14×[N]−48/32×[S] Equation (a)
However, [Ti][N][S] represent mass% of Ti, N, and S, respectively.
HV1/HV3≦0.80...Equation (b)
HV2/HV3≧0.80...Equation (c) Further, instead of part of the Fe, in mass%,
Nb: 0.010 to 0.100%,
V: 0.010 to 0.300%,
Cu: 0.01 to 1.20%,
Ni: 0.01-0.60%,
Cr: 0.01 to 2.00%,
Mo: 0.01 to 1.00%,
The hot-rolled steel sheet according to claim 1, containing one or more of the above. Further, instead of part of the Fe, in mass%,
Mg: 0.0005 to 0.0100%,
Ca: 0.0005 to 0.0100%,
REM: 0.0005 to 0.1000%,
1 or 2 or more types are contained, The hot-rolled steel plate of Claim 1 or Claim 2 characterized by the above-mentioned. Further, instead of part of the Fe, in mass%,
B: 0.0002 to 0.0020%,
The hot-rolled steel sheet according to claim 1, wherein the hot-rolled steel sheet comprises: A method for producing a high-strength hot-rolled steel sheet excellent in ductility, fatigue characteristics, and surface machinability according to any one of claims 1 to 4, wherein the method is any one of claims 1 to 4. When the slab having the component composition described in the item is heated to a temperature of T1 or higher defined by the formula (d), and the heated slab is subjected to hot rolling including finish rolling consisting of continuous rolling of a plurality of stages,
Of the stages of finish rolling in which the time (t1) from the central temperature of the steel sheet to 1000° C. to the Ar3 temperature determined by the formula (f) is set to 9.0 seconds or less, and the reduction ratio in the finish rolling is 10% or more. The second stage from the rearmost stage is the finish rolling stage so that the rolling roll temperature is 450°C or higher, and the rearmost stage is the finishing rolling stage, and the rolling roll temperature is 300°C or lower. At a rolling temperature of (T2-20)° C. or higher and (T2+100)° C. or lower with respect to the temperature T2 defined by the formula (e). The obtained hot-rolled steel sheet is then cooled at an average cooling rate of 20° C./sec or higher to a temperature of 730° C. or higher and 830° C. or lower, and then held in a temperature range of 730° C. or higher and 830° C. or lower for 3 seconds or longer. A method for producing a hot-rolled steel sheet, comprising cooling at an average cooling rate of 40° C./sec or more, and then winding at a temperature of 300° C. or less.
T1 (° C.)=7000/{2.75-log([Ti]×[C])}-273...Equation (d)
T2 (° C.)=870+10×([C]+[N])×[Mn]+350×[Nb]+250×[Ti]+40×[B]+10×[Cr]+100×[Mo]+100×[V] ...Equation (e)
Ar3(°C)=868-396×[C]-68.1×[Mn]+24.6×[Si]-36.1×[Ni]-24.8×[Cr]-20.7×[Cu ]+250×[Al]... Formula (f)
However, when the Ar3 temperature calculated by the formula (f) exceeds 900°C, Ar3=900°C. In addition, the [elemental symbol] in formula (a), formula (d), formula (e), and formula (f) indicates the mass% of the element in the brackets in the steel sheet.

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