Patent Application
20140238555
This disclosure relates to a high strength hot rolled steel sheet suitable for parts of automobiles and other transportation machines and structural steels, e.g., construction steels, and which has high strength of tensile strength (TS): 780 MPa or more and excellent formability in combination, and a method of manufacturing the same.
In the automobile industry, it is always an important issue to reduce the weight of an automotive body while maintaining the strength thereof and enhance the automotive fuel economy to reduce the amount of CO2 output from the viewpoint of global environmental conservation. It is effective to decrease the thickness of a steel sheet by enhancing the strength of the steel sheet serving as a raw material for automotive parts for the purpose of weight reduction of automotive bodies while maintaining their strength. On the other hand, most automotive parts using the steel sheet as a raw material are formed by press forming, burring, or the like and, therefore, the steel sheet for automotive parts is required to have excellent ductility and stretch flange formability. Consequently, the strength of, for example, tensile strength: 780 MPa or more and, in addition formability are regarded as important for the steel sheet for automotive parts so that a high strength steel sheet with excellent formability, e.g., stretch flange formability, is desired.
However, in general, formability of the iron and steel material is degraded along with enhancement in strength so that it is not easy to improve formability, e.g., stretch flange formability, to the high strength steel sheet without impairing the strength. For example, a technology is known, wherein a steel sheet microstructure is a complex microstructure in which a hard, low temperature transformed phase, e.g., martensite, is dispersed in mild ferrite to prepare a high strength steel sheet with excellent ductility. Such a technology is aimed at ensuring the compatibility between high strength and high ductility by optimizing the amount of martensite dispersed in ferrite. However, there is a problem that when the steel sheet having the above-described complex microstructure is subjected to expansion of a punched portion, so-called “stretch flange” forming, cracking occurs from the interface between soft ferrite and a hard low temperature transformed phase, e.g., martensite, to easily cause fractures. That is, sufficient stretch flange formability is not obtained by the complex microstructure high strength steel sheet composed of soft ferrite and a hard low temperature transformed phase, e.g., martensite.
Meanwhile, Japanese Unexamined Patent Application Publication No. 6-172924 proposes to improve stretch flange formability of a high strength hot rolled steel sheet with tensile strength: 500 N/mm2 (500 MPa) or more by containing C: 0.03% to 0.20%, Si: 0.2% to 2.0%, Mn: 2.5% or less, P: 0.08% or less, and S: 0.005% on a percent by weight basis and employing a microstructure primarily containing bainitic.ferrite or a microstructure containing ferrite and bainitic.ferrite as the steel sheet microstructure. That technology discloses that high stretch flange formability can be given to a high strength material by generating a bainitic.ferrite microstructure which has a lath microstructure and in which no carbide is generated and the dislocation density is high, in the steel. Also, it is disclosed that, when a ferritic microstructure including reduced dislocation and having high ductility and good stretch flange formability is generated in addition to the bainitic.ferrite microstructure, both the strength and stretch flange formability become good.
On the other hand, although not only the stretch flange formability is noted, Japanese Unexamined Patent Application Publication No. 2000-328186 proposes improving fatigue strength and stretch flange formability of a high strength hot rolled steel sheet with tensile strength (TS): 490 MPa or more by specifying the composition to contain C: 0.01% to 0.10%, Si: 1.5% or less, Mn: more than 1.0% to 2.5%, P: 0.15% or less, S: 0.008% or less, Al: 0.01% to 0.08%, B: 0.0005% to 0.0030%, and one of Ti and Nb or total of the two: 0.10% to 0.60% on a percent by weight basis and employing a microstructure in which the amount of ferrite is 95% or more on an area fraction basis, the average grain size of ferrite is 2.0 to 10.0 μm, and martensite and retained austenite are not contained.
Also, Japanese Unexamined Patent Application Publication No. 8-73985 proposes to ensure bendability and weldability of a hot rolled steel sheet and allowing the tensile strength (TS) thereof to become 950 N/mm2 (950 MPa) or more by specifying the composition to contain C: 0.05% to 0.15%, Si: 1.5% or less, Mn: 0.70% to 2.50%, Ni: 0.25% to 1.5%, Ti: 0.12% to 0.30%, B: 0.0005% to 0.0030%, P: 0.020% or less, S: 0,010% or less, sol. Al: 0.010% to 0.10%, and N: 0.0050% or less on a weight ratio basis, specifying the ferrite grain size to be 10.0 μm or less, and precipitating TiC having a size of 10 nm or less and iron carbide having a size of 10 μm or less. That technology discloses that the steel sheet strength is enhanced and, in addition, bendability is improved by allowing ferrite grains and TiC to become finer and specifying the Mn content to be 0.70% or more.
Also, Japanese Unexamined Patent Application Publication No. 6-200351 proposes to prepare a hot rolled steel sheet having excellent stretch flange formability and, in addition, having a tensile strength (TS) of 70 kgf/mm2 (686 MPa) or more by specifying the composition to contain C: 0.02% to 0.10%, Si≦2.0%, Mn: 0.5% to 2.0%, P≦0.08%, S≦0.006%, N≦0.005%, Al: 0.01% to 0.1%, and Ti: 0.06% to 0.3% on a percent by weight basis, where the amount of Ti satisfies 0.50<(Ti−3.43N−1.55)/4C, and employing the microstructure in which the area ratio of the low temperature transformed product and pearlite is 15% or less and TiC is dispersed in polygonal ferrite. That technology discloses that most of the steel sheet microstructure is polygonal ferrite containing a small amount of solid solution C, and the tensile strength (TS) is enhanced and, in addition, excellent stretch flange formability is obtained by TiC precipitation hardening and solid solution hardening due to Mn (content: 0.5% or more) and P.
Japanese Unexamined Patent Application Publication No. 2002-322539 proposes a thin steel sheet which is substantially composed of a matrix having a ferrite single phase and fine precipitates having a grain size of less than 10 nm and dispersing in the matrix and which has a tensile strength of 550 MPa or more and excellent press formability. The above-described technology discloses that, preferably, the composition contains C<0.10%, Ti: 0.03% to 0.10%, and Mo: 0.05% to 0.6% on a percent by weight basis, where Fe is a primary component, and thereby, a thin steel sheet exhibiting good hole expanding ratio and good total elongation in spite of high strength is prepared. In addition, an example containing Si: 0.04% to 0.08% and Mn: 1.59% to 1.67% is shown.
Japanese Unexamined Patent Application Publication No. 2007-302992 proposes to allow a hot rolled steel sheet to have a tensile strength of 690 to 850 MPa and, in addition, allow the hole expanding ratio to become 40% or more by specifying the composition to contain C: 0.015% to 0.06%, Si: less than 0.5%, Mn: 0.1% to 2.5%, P≦0.10%, S≦0.01%, Al: 0.005% to 0.3%, N≦0.01%, Ti: 0.01% to 0.30%, and B: 2 to 50 ppm on a percent by mass basis, where a component balance between C, Ti, N, and S and Mn, Si, and B is regulated, and further employing the microstructure in which the area fraction of ferrite and bainitic ferrite is 90% or more in total and the area fraction of cementite is 5% or less.
Japanese Unexamined Patent Application Publication No. 2005-298924 proposes to allow a hot rolled steel sheet to have a tensile strength of 690 MPa or more and, in addition, improve the punching ability and hole expanding property by specifying the composition to contain C: 0.01% to 0.07%, Si: 0.01% to 2%, Mn: 0.05% to 3%, Al: 0.005% to 0.5%, N≦0.005%, S≦0.005%, and Ti: 0.03% to 0.2% on a percent by mass basis, where furthermore the P content is reduced to 0.01% or less, and employing the microstructure in which a ferrite or bainitic ferrite microstructure is a phase having a maximum area fraction and the area fraction of hard second phase and cementite is 3% or less.
However, as for Japanese Unexamined Patent Application Publication No. 6-172924, if the ferrite content increases, further enhancement of strength cannot be expected. Meanwhile, when a complex microstructure is employed by adding a hard second phase instead of ferrite for the purpose of enhancing the strength, there is a problem that, in stretch flange forming, cracking occurs from the interface between bainitic.ferrite and a hard second phase to easily cause fractures and stretch flange formability is degraded as with the above-described ferrite-martensite complex microstructure steel sheet.
Also, in Japanese Unexamined Patent Application Publication No. 2000-328186 stretch flange formability of the steel sheet is improved by allowing crystal grains to become finer. However, the tensile strength (TS) of the resulting steel sheet is about 680 MPa at most (refer to the example in Japanese Unexamined Patent Application Publication No. 2000-328186). Therefore, there is a problem that further enhancement of strength cannot be expected. In addition, in Japanese Unexamined Patent Application Publication No. 2000-328186, it is necessary that more than 1% of Mn be contained. Consequently, fractures resulting from segregation of Mn occurs easily during forming, and it is difficult to stably ensure excellent stretch flange formability.
Also, in Japanese Unexamined Patent Application Publication No. 8-73985, bendability of the steel sheet has been studied, but stretch flange formability of the steel sheet has not been studied. Bending and hole expanding (stretch flange forming) are different in the forming property sort, and bendability and stretch flange formability are different in the properties required of the steel sheet. Therefore, there is a problem that a high strength steel sheet with excellent bendability does not always have good stretch flange formability.
As for Japanese Unexamined Patent Application Publication No. 6-200351, large amounts of Mn and, furthermore, Si are contained to enhance strength. Therefore, hardenability of the steel is enhanced and it is difficult to obtain a microstructure primarily containing polygonal ferrite stably. Meanwhile, significant segregation occurs during casting because of these elements, so that fractures along the segregation easily occurs during forming and stretch flange formability tends to be degraded. In addition, as is shown in the example thereof, the tensile strength of 780 MPa or more is not stably obtained in spite of the fact that addition of 1% or more of Mn is necessary.
Also, in Japanese Unexamined Patent Application Publication No. 2002-322539, the examples containing 1.59% to 1.67% of Mn are shown. Therefore, fractures due to segregation of Mn easily occur during forming and there is a problem that it is difficult to stably ensure excellent stretch flange formability even by the above-described technology.
In Japanese Unexamined Patent Application Publication No. 2007-302992, as is shown in the examples thereof, it is necessary that 1% or more of Mn be added to increase the tensile strength of the steel sheet to 780 MPa or more, and if the Mn content is reduced to about 0.5%, the resulting tensile strength is as little as less than 750 MPa. That is, even Japanese Unexamined Patent Application Publication No. 2007-302992 cannot increase the tensile strength of the steel sheet to 780 MPa or more while the amount of Mn is reduced and excellent stretch flange formability is ensured.
In Japanese Unexamined Patent Application Publication No. 2005-298924, as is shown in the examples thereof, the strength cannot be obtained unless at least about 1% of Mn is added, and it is difficult to stably obtain stretch flange formability because of segregation due to this addition of Mn. Meanwhile, Japanese Unexamined Patent Application Publication No. 2005-298924 discloses an example in which Ti, V, Nb, and Mo are added to C=0.066%, Si=0.06%, and Mn=0.31%. In that example, it is necessary that a bainitic ferrite microstructure be formed by performing coiling at a low temperature of 540° C. in order to avoid pearlite generation and, therefore, stable stretch flange formability is not obtained. Furthermore, Japanese Unexamined Patent Application Publication No. 2005-298924 also discloses an example in which the Mn content is 0.24% and, in addition, the tensile strength is 810 MPa. That example contains a large amount, 1.25%, of easy-to-segregate Si in compensation for the strength and, likewise, stable stretch flange formability is not obtained.
As described above, it is not preferable that the steel sheet microstructure is a complex microstructure from the viewpoint of stretch flange formability. In this regard, stretch flange formability is improved by specifying the steel sheet microstructure to be a ferritic single phase, although it is difficult to ensure high strength of the ferritic single phase steel sheet in the related art while excellent stretch flange formability is maintained.
It could therefore be helpful to provide a high strength hot rolled steel sheet having tensile strength (TS): 780 MPa or more and excellent stretch flange formability and a method of manufacturing the same.
We thus provide:
[1] A high strength hot rolled steel sheet having a composition containing
C: more than 0.035% and 0.07% or less, Si: 0.3% or less,
Mn: more than 0.35% and 0.7% or less, P: 0.03% or less,
S: 0.03% or less, Al: 0.1% or less,
N: 0.01% or less, Ti: 0.135% or more and 0.235% or less,
and the remainder composed of Fe and incidental impurities, on a percent by mass basis, in such a way that C, S, N, and Ti satisfy the formula (1) described below, a microstructure in which a matrix includes more than 95% of ferritic phase on an area fraction basis and fine Ti carbides having an average grain size of less than 10 nm are precipitated in the grains of the above-described ferritic phase, and a tensile strength of 780 MPa or more.
((Ti−(48/14)N−(48/32)S)/48)/(C/12)<1.0 (1)
(C, S, N, and Ti: content of the respective elements (percent by mass))
[2] The high strength hot rolled steel sheet according to the above-described item [1], wherein the above-described composition further contains B: 0.0025% or less on a percent by mass basis.
[3] The high strength hot rolled steel sheet according to the above-described item [1], wherein the above-described composition further contains 1.0% or less of at least one of REM, Zr, Nb, V, As, Cu, Ni, Sn, Pb, Ta, W, Mo, Cr, Sb, Mg, Ca, Co, Se, Zn, and Cs in total on a percent by mass basis.
[4] The high strength hot rolled steel sheet according to the above-described item [2], wherein the above-described composition further contains 1.0% or less of at least one of REM, Zr, Nb, V, As, Cu, Ni, Sn, Pb, Ta, W, Mo, Cr, Sb, Mg, Ca, Co, Se, Zn, and Cs in total on a percent by mass basis.
[5] The high strength hot rolled steel sheet according to any one of the above-described items [1] to [4], wherein a coating layer is included on the steel sheet surface.
[6] A method of manufacturing a high strength hot rolled steel sheet including the steps of heating a semi-manufactured steel to an austenitic single phase region, performing hot rolling composed of rough rolling and finish rolling, and performing cooling and coiling after completion of the finish rolling to produce a hot rolled steel sheet,
wherein the above-described semi-manufactured steel has a composition containing
C: more than 0.035% and 0.07% or less, Si: 0.3% or less,
Mn: more than 0.35% and 0.7% or less, P: 0.03% or less,
S: 0.03% or less, Al: 0.1% or less,
N: 0.01% or less, Ti: 0.135% or more and 0.235% or less,
and the remainder composed of Fe and incidental impurities, on a percent by mass basis, in such a way that C, S, N, and Ti satisfy the formula (1) described below,
the finishing temperature of the above-described finish rolling is specified to be 900° C. or higher, the average cooling rate in the above-described cooling from 900° C. to 750° C. is specified to be 10° C./sec. or more, and the coiling temperature in the above-described coiling is specified to be 580° C. or higher and 750° C. or lower.
((Ti−(48/14)N−(48/32)S)/48)/(C/12)<1.0 (1)
(C, S, N, and Ti: content of the respective elements (percent by mass))
[7] The method of manufacturing a high strength hot rolled steel sheet according to the above-described item [6], wherein the above-described composition further contains B: 0.0025% or less on a percent by mass basis.
[8] The method of manufacturing a high strength hot rolled steel sheet according to the above-described item [6], wherein the above-described composition further contains 1.0% or less of at least one of REM, Zr, Nb, V, As, Cu, Ni, Sn, Pb, Ta, W, Mo, Cr, Sb, Mg, Ca, Co, Se, Zn, and Cs in total on a percent by mass basis.
[9] The method of manufacturing a high strength hot rolled steel sheet according to the above-described item [7], wherein the above-described composition further contains 1.0% or less of at least one of REM, Zr, Nb, V, As, Cu, Ni, Sn, Pb, Ta, W, Mo, Cr, Sb, Mg, Ca, Co, Se, Zn, and Cs in total on a percent by mass basis.
[10] The method of manufacturing a high strength hot rolled steel sheet according to any one of the above-described items [6] to [9], wherein the above-described hot rolled steel sheet is subjected to a coating treatment.
[11] The method of manufacturing a high strength hot rolled steel sheet according to the above-described item [10], wherein the above-described hot rolled steel sheet is subjected to an alloying treatment following the above-described coating treatment.
A high strength hot rolled steel sheet suitable for parts of automobiles and other transportation machines and structural steels, e.g., construction steels, and which has high strength of tensile strength (TS): 780 MPa or more and excellent stretch flange formability in combination, is obtained. Therefore, further uses of the high strength hot rolled steel sheet can be developed and a remarkable industrial effect is exerted.
We studied various factors exerting an influence on enhancement of strength and stretch flange formability of hot rolled steel sheets. As a result, we found that Mn and Si, which had been previously considered to be very useful as solid solution hardening elements in enhancing the strength of the steel sheet and which had been positively contained in the high strength hot rolled steel sheet, adversely affect the stretch flange formability of the steel sheet.
We then observed the microstructure of a hot rolled steel sheet containing a large amount, about 1%, of Mn and Si. Consequently, we ascertained that segregation of Mn and Si unavoidably occurred in the sheet thickness center portion thereof and we found that a change in the shape of the microstructure resulting from segregation and unevenness in dislocation density and the like adversely affected stretch flange formability. Then, with respect to the composition of the hot rolled steel sheet, we found that the above-described influence of the segregation microstructure can be suppressed by reducing the Mn content and the Si content to a predetermined amount or less, specifically reducing both the Mn content and the Si content to contents further smaller than 1%.
On the other hand, reduction in the steel sheet strength along with reduction in the contents of Mn and Si serving as the solid solution hardening elements is not avoided. Therefore, we attempted to apply precipitation hardening due to Ti carbides as a hardening mechanism alternative to solid solution hardening due to Mn and Si. An effect of improving the steel sheet strength to a great extent can be expected by precipitating fine Ti carbides in the steel sheet. However, the Ti carbides easily become coarse. Consequently, as is shown by examples in the above-described patent literatures, it was difficult to precipitate Ti carbides in a fine state in the steel sheet and maintain the fine state by only simply containing Ti serving as a carbide-forming element into the steel sheet composition, and a sufficient strength enhancing effect was not obtained.
We performed further studies to search for a way to precipitate Ti carbides in a fine state in the steel sheet and suppress coarsening thereof. As a result, we found that coarsening of Ti carbides was able to be suppressed and Ti carbides were able to be made fine by adjusting the steel sheet composition to control the concentration ratio of the amount of Ti, which does not bond to N and S, but bonds to C, to the amount of C.
In this regard, the reasons the microstructure caused by Mn segregation present in the vicinity of the sheet thickness center portion of the steel sheet adversely affects the stretch flange formability are not certain. However, we believe that when a hole is punched and stretch flange forming to further expand the hole is performed, if a microstructure caused by segregation (microstructure having a flat shape or a high dislocation density) is present in the center portion, early-stage cracking is easily formed around it, fracturing proceeds in the sheet thickness direction because of forming (hole expanding forming) thereafter and, thereby, the hole expanding ratio is reduced.
Our steel sheets and methods will be described below in detail. Our hot rolled steel sheets are characterized in that substantially a ferritic single phase is employed and, in addition, improvement in the stretch flange formability is aimed by reduction and rendering harmless of Mn segregation and, in addition, Si segregation in the sheet thickness center portion through decreases in the Mn content and, in addition, the Si content in the steel sheet. Also, the hot rolled steel sheets are characterized in that enhancement of the strength of the steel sheet is achieved by precipitating fine Ti carbides, controlling the amount of C to be bonded to the amount of Ti in the steel composition in such a way as to become larger than the amount of Ti and generate no pearlite, and suppressing growth and coarsening of fine Ti carbides through reduction of the amount of solid solution Ti.
To begin, reasons for the limitation of the microstructure and carbides of the steel sheet will be described.
The hot rolled steel sheet has a microstructure in which a matrix includes more than 95% of ferritic phase on an area fraction basis and fine Ti carbides having an average grain size of less than 10 nm are precipitated in the grains of the above-described ferritic phase.
Ferritic Phase: More than 95% in the Matrix on an Area Fraction Basis
Formation of the ferritic phase is indispensable to ensure stretch flange formability of the hot rolled steel sheet. It is effective in improving ductility and stretch flange formability of the hot rolled steel sheet for the matrix microstructure of the hot rolled steel sheet to be a ferritic phase having a small dislocation density and excellent ductility. In particular, it is preferable in improving stretch flange formability that the matrix microstructure of the hot rolled steel sheet is a ferritic single phase. Even when not a perfect ferritic single phase, a substantially ferritic single phase, that is, a ferrite phase constituting more than 95% of the whole matrix microstructure on an area fraction basis exerts the above-described effect sufficiently. Therefore, the area fraction of the ferritic phase is more than 95%, and preferably 97% or more.
Meanwhile, examples of microstructures which may be contained in the matrix other than the ferritic phase include cementite, pearlite, a bainitic phase, a martensitic phase, and a retained austenitic phase. If these microstructures are present in the matrix, stretch flange formability is degraded. However, these microstructures are permitted when a total area fraction relative to the whole matrix microstructure is less than about 5%, and preferably about 3% or less.
Ti Carbide
As described above, an increase in the steel sheet strength due to solid solution hardening cannot be expected because the contents of Mn and Si serving as solid solution hardening elements are reduced for the purpose of suppressing Mn segregation and, in addition, Si segregation in the sheet thickness center portion, which adversely affect stretch flange formability. Consequently, precipitation of fine Ti carbides in grains of the ferritic phase is indispensable to ensure the strength. In this regard, the Ti carbides are carbides precipitated at interfaces of phases at the same time with the transformation from austenite to ferrite during cooling after completion of finish rolling in a hot rolled steel sheet production process or aging precipitation carbides precipitated in ferrite after ferrite transformation.
Average Grain Size of Ti Carbides: Less than 10 nm
The average grain size of Ti carbides is very important in achieving predetermined strength (tensile strength: 780 MPa or more) for the hot rolled steel sheet. The average grain size of Ti carbides is less than 10 nm. When fine Ti carbides are precipitated in grains of the above-described ferritic phase, Ti carbides act as resistance against movement of dislocations generated when deformation is applied to the steel sheet and, thereby, the hot rolled steel sheet is strengthened. However, Ti carbides become sparse along with coarsening of Ti carbides, and the distance of stopping of dislocation increases, so that precipitation hardening ability is degraded. Then, if the average grain size of Ti carbides becomes 10 nm or more, the steel sheet strengthening ability sufficient to compensate for the amount of reduction in the steel sheet strength resulting from reduction in the contents of Mn and Si serving as solid solution hardening elements is not obtained. Therefore, the average grain size of Ti carbides is less than 10 nm, and more preferably 6 nm or less.
In this regard, we ascertained that the shape of the Ti carbide was nearly the shape of a disc (shape of a circular plate) as schematically shown in
Meanwhile, although the effect is not specifically limited, the form of precipitation of fine Ti carbides may be observed in the shape of a row. However, even in this case, precipitation is at random in the plane containing the row of the individual row-shaped precipitates and, in many cases, precipitates are not observed in the shape of a row on the basis of actual observation with a transmission electron microscope.
Next, reasons for the limitation of the chemical composition of the hot rolled steel sheet will be described. In this regard, “%” expressing the chemical composition hereafter refers to “percent by mass” unless otherwise specified.
C: More than 0.035% and 0.07% or Less
Carbon is an element indispensable in strengthening the hot rolled steel sheet by forming Ti carbides in the steel sheet. If the C content is 0.035% or less, Ti carbides to bring the tensile strength to 780 MPa or more cannot be ensured and the tensile strength of 780 MPa or more is not obtained. On the other hand, if the C content is more than 0.07%, pearlite is easily generated and stretch flange formability is degraded. Therefore, the C content is more than 0.035% and 0.07% or less, and more preferably 0.04% or more and 0.06% or less.
Si: 0.3% or Less
Silicon is an element effective in enhancing steel sheet strength without causing degradation in ductility (elongation) and is usually positively contained in a high strength steel sheet. However, Si facilitates Mn segregation in the sheet thickness center portion, which should be avoided in the hot rolled steel sheet and, in addition, Si in itself is an element which segregates. Therefore, the Si content is limited to 0.3% or less for the purpose of suppressing the above-described Mn segregation and suppressing Si segregation. The Si content is more preferably 0.1% or less, and further preferably 0.05% or less.
Mn: More than 0.35% and 0.7% or Less
Manganese is a solid solution hardening element and, as with Si, is positively contained in a common high strength steel sheet. However, if Mn is positively contained in the steel sheet, Mn segregation in the sheet thickness center portion is not avoided, and degradation in stretch flange formability of the steel sheet is brought about. Therefore, the Mn content is limited to 0.7% or less for the purpose of suppressing the above-described Mn segregation. The Mn content is more preferably 0.6% or less, and further preferably 0.5% or less. On the other hand, if the Mn content is 0.35% or less, the austenite-ferrite transformation temperature increases and, thereby, it is difficult to make Ti carbides finer. As described above, Ti carbides are precipitated at the same time with the transformation from austenite to ferrite during cooling after completion of finish rolling in the hot rolled steel sheet production process or are aging-precipitated in ferrite. At this time, if the austenite-ferrite transformation temperature becomes a high temperature, Ti carbides become coarse because of precipitation in a high temperature range. Therefore, the lower limit of the Mn content is more than 0.35%.
P: 0.03% or Less
Phosphorus is a harmful element which segregates at grain boundaries to reduce elongation and induce fractures during forming. Therefore, the P content is 0.03% or less, more preferably 0.020% or less, and further preferably 0.010% or less.
S: 0.03% or Less
Sulfur is present as MnS or TiS in the steel to facilitate generation of voids during punching of the hot rolled steel sheet and, furthermore, serve as a starting point of voids during forming so that stretch flange formability is degraded. Therefore, the S content is preferably minimized and is 0.03% or less. The S content is more preferably 0.010% or less, and further preferably 0.0030% or less.
Al: 0.1% or Less
Aluminum is an element serving as a deoxidizing agent. It is desirable that the content be 0.01% or more to obtain such an effect. However, if Al is more than 0.1%, Al oxides remain in the steel sheet, the Al oxides tend to aggregate and become coarse easily and cause degradation in stretch flange formability. Therefore, the Al content is 0.1% or less, and more preferably 0.065% or less.
N: 0.01% or Less
N is a harmful element and preferably minimized. Nitrogen is bonded to Ti to form TiN. If the N content is more than 0.01%, stretch flange formability is degraded because of an increase in the amount of TiN formed. Therefore, the N content is 0.01% or less, and more preferably 0.006% or less.
Ti: 0.135% or More and 0.235% or Less
Titanium is an element indispensable in forming Ti carbides and enhancing the strength of the steel sheet. It becomes difficult to ensure a predetermined hot rolled steel sheet strength (tensile strength: 780 MPa or more) if the Ti content is less than 0.135%. On the other hand, if the Ti content is more than 0.235%, Ti carbides tend to become coarse, and it becomes difficult to ensure a predetermined hot rolled steel sheet strength (tensile strength: 780 MPa or more). Therefore, the Ti content is 0.135% or more and 0.235% or less, and more preferably 0.15% or more and 0.20% or less.
The hot rolled steel sheet contains C, S, N, and Ti in the above-described ranges such that formula (1) is satisfied.
((Ti−(48/14)N−(48/32)S)/48)/(C/12)<1.0 (1)
(C, S, N, and Ti: content of the respective elements (percent by mass))
The above-described formula (1) is a requirement to be satisfied for the average grain size of Ti carbides to be less than 10 nm and is a very important indicator.
As described above, a predetermined steel sheet strength is ensured by precipitating fine Ti carbides in the hot rolled steel sheet. Here, the Ti carbides tend to become fine carbides having a very small average grain size. However, if the atomic concentration of Ti contained in the steel becomes larger than or equal to the atomic concentration of C, the Ti carbides become coarse easily. Then, along with coarsening of carbides, it becomes difficult to ensure a predetermined hot rolled steel sheet strength (tensile strength: 780 MPa or more). It is necessary that the atomic percent of C ((percent by mass of C)/12) contained in the steel is larger than the atomic percent of Ti ((percent by mass of Ti)/48) contributable to carbide generation. In this regard, when the steel composition is controlled as described above, the number of Ti atoms in the Ti carbides becomes smaller than the number of C atoms so that the effect of suppressing Ti carbides from becoming coarse is enhanced.
Meanwhile, as described later, a predetermined amount of Ti is added to a steel, carbides in the steel is melted by heating before hot rolling, and Ti carbides are precipitated mainly during coiling after hot rolling. However, the whole amount of Ti added to the steel does not contribute to carbide generation and part of Ti added to the steel is consumed to form nitrides and sulfides. In a temperature region higher than the coiling temperature, Ti forms nitrides and sulfides easily as compared to carbides because Ti forms nitrides and sulfides before the coiling step in production of the hot rolled steel sheet. Therefore, the amount of Ti, which contributes to carbide generation, in Ti added to the steel can be represented by “Ti−(48/14)N−(48/32)S”.
For the reasons described above the individual elements, C, S, N, and Ti, are contained in such a way as to satisfy the above-described formula (1), that is, ((Ti−(48/14)N−(48/32)S)/48)/(C/12)<1.0, for the purpose of allowing the atomic percent of C(C/12) to become larger than the atomic percent of Ti ((Ti−(48/14)N−(48132)S)148) contributable to carbide generation. In the case where the above-described formula (1) is not satisfied, Ti carbides generated in ferrite grains cannot be maintained in a fine state (average grain size of less than 10 nm), and it becomes difficult to obtain a predetermined steel sheet strength (tensile strength: 780 MPa or more).
In this regard, for the purpose of allowing Ti carbides to become finer, the value of the left side of the above-described formula (1), that is, the value of ((Ti−(48/14)N−(48/32)S)/48)/(C/12), is preferably 0.5 or more and 0.95 or less, and more preferably 0.6 or more and 0.9 or less.
Meanwhile, carbides in the steel are melted by heating of the steel before hot rolling and, usually, the Ti carbides are precipitated at interfaces of phases at the same time with the transformation from austenite to ferrite during cooling after hot rolling or are aging-precipitated in ferrite. Here, if the temperature of transformation from austenite to ferrite of the steel is high, Ti carbides are precipitated in a high temperature region, in which the diffusion rate of Ti is large, after the hot rolling so that Ti carbides become coarse easily. However, Ti carbides are suppressed from becoming coarse effectively by lowering the temperature of transformation from austenite to ferrite (Ar3 transformation temperature) to the coiling temperature range (that is, the temperature region in which the diffusion rate of Ti is small).
Then, B: 0.0025% or less can be further contained in addition to the above-described composition for the purpose of retarding transformation from austenite to ferrite of the steel and stably lowering the precipitation temperature (Ar3 transformation temperature) of Ti carbides to the coiling temperature range described later.
B: 0.0025% or Less
Boron is an element that retards the start of austenite-ferrite transformation of the steel and lowers the precipitation temperature of Ti carbides by suppressing the austenite-ferrite transformation to contribute to making the carbides finer. In particular, when Mn is reduced to a great extent for the purpose of avoiding segregation, lowering the Ar3 transformation point due to Mn cannot be expected and, therefore, it is preferable that the austenite-ferrite transformation be retarded by containing B. Consequently, when the Mn content is reduced to a great extent (for example, Mn: 0.5% or less), Ti carbides can stably be made finer. On the other hand, if the B content is more than 0.0025%, a bainite transformation effect due to B is enhanced, and it becomes difficult to convert to a ferrite microstructure. Therefore, the B content is 0.0025% or less. On the other hand, addition of more than 0.0010% may reduce elongation because solid solution B inhibits the movement of dislocation so that the B content is more preferably 0.0002% or more and 0.0010% or less, and further preferably 0.0002% or more and 0.0007% or less.
The hot rolled steel sheet may further contain 1.0% or less of at least one of REM, Zr, Nb, V, As, Cu, Ni, Sn, Pb, Ta, W, Mo, Cr, Sb, Mg, Ca, Co, Se, Zn, and Cs in total in addition to the above-described composition. In this regard, the components other than those described above are Fe and incidental impurities.
Meanwhile, in the case where a coating layer is disposed on the surface of the hot rolled steel sheet for the purpose of giving the corrosion resistance to the steel sheet, the above-described effects are not impaired. In this regard, the type of the coating layer disposed on the steel sheet surface is not specifically limited, and any type, e.g., electroplated coating and hot dip coating, may be employed. Also, examples of hot dip coating include hot dip galvanized coating. Furthermore, hot dip galvannealed coating may be employed, where an alloying treatment is performed after coating.
Next, a method of manufacturing our hot rolled steel sheets will be described.
A semi-manufactured steel having the above-described composition is heated to an austenitic single phase region, hot rolling composed of rough rolling and finish rolling is performed, and cooling and coiling are performed after completion of the finish rolling to produce a hot rolled steel sheet. At this time, the finishing temperature of the finish rolling is 900° C. or higher, the average cooling rate in the cooling from 900° C. to 750° C. is 10° C./sec. or more, and the coiling temperature is 580° C. or higher and 750° C. or lower.
The method of melt-refining the steel is not specifically limited, and a known melt-refining method, e.g., a converter or an electric furnace, can be adopted. In this regard, after the melt-refining, it is preferable that a slab (semi-manufactured steel) be prepared by a continuous casting method from the viewpoint of productivity and the like. However, the slab may be prepared by a known casting method, e.g., an ingot making-roughing method or a thin slab continuous casting method. Meanwhile, the Mn content and the Si content causing segregation are reduced for the purpose of improving formability (stretch flange formability and the like). Consequently, when the continuous casting method advantageous to suppress segregation is adopted, the effects become still more considerable.
The semi-manufactured steel obtained as described above is subjected to rough rolling and finish rolling. The semi-manufactured steel is heated to an austenitic single phase region prior to rough rolling. If the semi-manufactured steel before rough rolling is not heated to the austenitic single phase region, remelting of Ti carbides present in the semi-manufactured steel does not proceed and precipitation of fine Ti carbides after rolling is not achieved. Therefore, the semi-manufactured steel is heated to the austenitic single phase region, preferably 1,200° C. or higher, prior to the rough rolling. However, if the heating temperature of the semi-manufactured steel is too high, the surface is excessively oxidized, Ti is consumed because of generation of TiO2, and reduction in hardness occurs easily in the vicinity of the surface of the resulting steel sheet. Consequently, the above-described heating temperature is more preferably 1,350° C. or lower. In this regard, when the semi-manufactured steel is subjected to hot rolling, when the temperature of the semi-manufactured steel (slab) after casting is in the austenitic single phase region, the semi-manufactured steel may be directly rolled without heating the semi-manufactured steel or after short-time heating. Meanwhile, it is not necessary that the rough rolling condition is specifically limited.
Finish Rolling Temperature: 900° C. or Higher
Controlling the finish rolling temperature is important to ensure stretch flange formability of the hot rolled steel sheet. If the finish rolling temperature is lower than 900° C., a band-shaped microstructure is formed easily at the position, at which Mn has segregated, in the sheet thickness center portion of a finally obtained hot rolled steel sheet, and stretch flange formability is easily degraded. Therefore, the finish rolling temperature is 900° C. or higher, and more preferably 920° C. or higher. Also, the finish rolling temperature is more preferably 1,050° C. or lower from the viewpoint of prevention of a flaw and roughing due to secondary scales of the surface.
Average Cooling Rate: 10° C./sec. or More
As described above, it is necessary that fine Ti carbides be precipitated. For this purpose, the temperature of the austenite-ferrite transformation is lowered to facilitate precipitation of fine Ti carbides, suppress coarsening, and ensure a predetermined average grain size (less than 10 nm). Here, Ti carbides are precipitated on the basis of ferrite transformation of the steel microstructure from austenite after completion of the above-described finish rolling. If the austenite-ferrite transformation point (Ar3 transformation temperature) is higher than 750° C., large Ti carbides grow easily.
Then, to establish the austenite-ferrite transformation point (Ar3 transformation temperature) at 750° C. or lower, after the finish rolling is completed, the average cooling rate in the cooling from 900° C. to 750° C. is 10° C./sec. or more, and more preferably 30° C./sec. or more.
Consequently, the average cooling rate is increased and, thereby, the austenite-ferrite transformation point (Ar3 transformation temperature) is 750° C. or lower, that is, in the temperature region of the coiling temperature described later so that disc-shaped Ti carbides are maintained in a fine state. However, if the above-described average cooling rate increases excessively, there may be a problem that a hardened microstructure is generated only in the surface layer easily. Therefore, the average cooling rate in the cooling from 900° C. to 750° C. after the finish rolling is completed is more preferably 600° C./sec. or less.
Coiling Temperature: 580° C. or Higher and 750° C. or Lower
Controlling the coiling temperature is important to establish the above-described austenite-ferrite transformation point (Ar3 transformation temperature) to be 750° C. or lower, and allowing the hot rolled steel sheet to have a predetermined matrix microstructure (area fraction of ferritic phase: more than 95%). If the coiling temperature is lower than 580° C., martensite and bainite are easily generated and it becomes difficult for the matrix to be substantially a ferritic single phase. On the other hand, if the coiling temperature is higher than 750° C., pearlite is easily generated and stretch flange formability is degraded. Also, if the coiling temperature is higher than 750° C., the austenite-ferrite transformation temperature cannot be made 750° C. or lower, and coarsening of Ti carbides is induced. Therefore, the coiling temperature is 580° C. or higher and 750° C. or lower, and more preferably 610° C. or higher and 690° C. or lower.
As described above, after cooling following finish rolling, austenite-ferrite transformation is allowed to occur in an temperature region of 750° C. or lower. Consequently, the austenite-ferrite transformation occurs easily in the vicinity of the coiling temperature, and the coiling temperature tends to substantially agree with the austenite-ferrite transformation temperature.
Meanwhile, it is further preferable that the coil after coiling be held at 580° C. to 750° C. for 60 sec. or more because a uniform microstructure is obtained easily.
In addition, a coating layer may be formed on the steel sheet surface by subjecting the hot rolled steel sheet produced as described above to a coating treatment. The coating treatment may be any one of electroplating and hot dipping. For example, a galvanized coating layer can be formed by applying a galvanizing treatment as the coating treatment. Alternatively, galvannealed coating layer may be formed by further applying an alloying treatment after the above-described galvanizing treatment. Also, the hot dipped coating can be coated with aluminum, an aluminum alloy, or the like other than zinc. The high strength hot rolled steel sheet is suitable for common press forming performed at ambient temperature and, in addition, is suitable for warm forming in which a steel sheet before pressing is heated from 400° C. to 750° C. and is immediately subjected to forming.
Molten steels having compositions shown in Table 1 (Table 1-1 and Table 1-2) were melt-refined and continuously cast by a usually known technique to prepare slabs (semi-manufactured steels) having a thickness of 300 mm. These slabs were heated to temperatures shown in Table 2 and were roughly rolled. Finish rolling at a finish rolling temperature shown in Table 2 was applied and after finish rolling was completed, the temperature region of from 900° C. to 750° C. was cooled at an average cooling rate shown in Table 2 and coiling was performed at a coiling temperature shown in Table 2, so that a hot rolled steel sheet having a sheet thickness: 2.3 mm was produced. In this regard, it was separately ascertained that the transformation from austenite to ferrite did not occur during the cooling up to the coiling except Steel No. 22.
Subsequently, the hot rolled steel sheets obtained as described above were pickled to remove surface layer scales. Thereafter, part of the hot rolled steel sheets (Steel Nos. 6 and 7) were dipped into a galvanizing bath (0.1% Al—Zn) at 480° C. to form galvanized coating layers on both surfaces of the steel sheet, where the amount of deposit per one surface was 45 g/m2 so that galvanized steel sheets were produced. Also, other part of the hot rolled steel sheets (Steel Nos. 8, 9, and 10) were provided with galvanized coating layers in the same manner as that described above and were subjected to the alloying treatment at 520° C. so that galvannealed steel sheets were produced.
0.031
1.132
0.090
12
0.088
15
0.299
1.724
20
1.02
21
0.035
0.130
0.0030
22
0.029
1.25
0.24
1.117
23
0.030
1.02
1.49
0.110
41
0.76
1.53
0.250
42
0.68
1.59
1.025
12
15
20
21
22
890
550
23
870
41
880
500
42
880
500
850
540
760
Test pieces were taken from the hot rolled steel sheets (hot rolled steel sheet, galvanized steel sheet, and galvannealed steel sheet) obtained as described above. A microstructure observation, a tensile test, and a hole expanding test were performed and, thereby, the area fraction of ferritic phase, the types an the area ratios of microstructures other than the ferritic phase, the average grain size of Ti carbides, the tensile strength, the elongation, and the hole expanding ratio (stretch flange formability) were determined. The test methods were as described below.
(i) Microstructure Observation
A test piece was taken from the resulting hot rolled steel sheet, a cross-section (L cross-section) parallel to a rolling direction of the test piece was polished, and corrosion with nital was performed. Thereafter, microstructure photographs taken with an optical microscope (magnification: 400 times) and a scanning electron microscope (magnification: 5,000 times) were used, and the types of ferritic phases and microstructures other than the ferritic phase and the area fractions thereof were determined.
In addition, a thin film produced from the hot rolled steel sheet was observed with a transmission electron microscope and the average grain size of Ti carbides was determined. A photograph taken with the transmission electron microscope (magnification: 340,000 times) was used, the maximum diameter d (diameter of a largest portion of the upper and lower surfaces of the disc) of 100 Ti carbides in total of five fields of view and the thickness t of the disc-shaped precipitate in a direction orthogonal to the upper and lower surfaces of the disc were measured, and the average grain size of Ti carbides was determined as the above-described arithmetic average value (average grain size ddef).
(ii) Tensile Test
A JIS No. 5 tensile test piece (JIS Z 2201) was taken from the resulting hot rolled steel sheet, where the tensile direction was the direction at a right angle to the rolling direction. A tensile test was performed in conformity with the specification of JIS Z 2241 and the tensile strength (TS) and the elongation (EL) were measured.
(iii) Hole Expanding Test
A test piece (size: 130 mm×130 mm) was taken from the resulting hot rolled steel sheet. A hole having initial diameter d0: 10 mmφ was formed in the test piece by punching (clearance: 12.5% of test piece sheet thickness). A hole expanding test was performed using the resulting test pieces. That is, a cone punch having an apex angle: 60° was pushed into the hole from the punch side in the punching to expand the hole, and a hole diameter d1 when a crack penetrated the steel sheet (test piece) was measured. The hole expanding ratio λ(%) was calculated on the basis of the following formula.
hole expanding ratio λ(%)={(d1−d0)/d0}×100
The results are shown in Table 3.
712
70 (remainder pearlite)
11.2
721
12
85 (remainder pearlite)
664
15
13.0
644
20
87 (remainder pearlite)
21
89 (remainder bainitic ferrite)
721
65 (remainder bainite)
12.3
778
23
745
41
85 (remainder bainitic ferrite)
12.0
775
42
90 (remainder bainitic ferrite)
13.0
765
16.2
758
88 (remainder bainitic ferrite +
11.2
770
martensite)
12.6
756
All our examples are hot rolled steel sheets having high strength of tensile strength TS: 780 MPa and excellent formability of elongation EL: 20% or more and hole expanding ratio λ:100% or more in combination. On the other hand, the comparative examples are unable to ensure predetermined high strength or are unable to ensure a sufficient hole expanding ratio.
| Number | Date | Country | Kind |
|---|---|---|---|
| 2011-242299 | Nov 2011 | JP | national |
| Filing Document | Filing Date | Country | Kind | 371c Date |
|---|---|---|---|---|
| PCT/JP2012/007021 | 11/1/2012 | WO | 00 | 4/25/2014 |
