Classifications

• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
• • C—CHEMISTRY; METALLURGY
  • C21—METALLURGY OF IRON
  • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
  • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
  • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
  • C21D8/04—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for deep-drawing
  • C21D8/0421—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for deep-drawing characterised by the working steps
  • C21D8/0426—Hot rolling
• • C—CHEMISTRY; METALLURGY
  • C21—METALLURGY OF IRON
  • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
  • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
  • C21D9/46—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals
  • C21D9/48—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals deep-drawing sheets
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/12—Ferrous alloys, e.g. steel alloys containing tungsten, tantalum, molybdenum, vanadium, or niobium
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/14—Ferrous alloys, e.g. steel alloys containing titanium or zirconium
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
  • C22C38/24—Ferrous alloys, e.g. steel alloys containing chromium with vanadium
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
  • C22C38/28—Ferrous alloys, e.g. steel alloys containing chromium with titanium or zirconium
• • C—CHEMISTRY; METALLURGY
  • C21—METALLURGY OF IRON
  • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
  • C21D2211/00—Microstructure comprising significant phases
  • C21D2211/004—Dispersions; Precipitations
• • C—CHEMISTRY; METALLURGY
  • C21—METALLURGY OF IRON
  • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
  • C21D2211/00—Microstructure comprising significant phases
  • C21D2211/005—Ferrite

Definitions

• the present invention relates to a high-strength steel sheet having high stretch flangeability after working and corrosion resistance after painting.
• Automobile parts such as chassis and truck frames, require formability (mainly elongation and stretch flangeability), and steel having a tensile strength on the order of 590 MPa has been used for such applications.
• formability mainly elongation and stretch flangeability
• steel having a tensile strength on the order of 590 MPa has been used for such applications.
• use of higher-strength automotive steel sheets has been promoted in recent years, and use of steel having a tensile strength on the order of 780 MPa is being investigated.
• Patent Documents 1 to 6 describe techniques for improving elongation and stretch flangeability.
• Patent Document 1 discloses a technique relating to high-workability high-strength steel sheet having a tensile strength of 590 MPa or more, wherein the steel sheet has a substantially ferritic single phase in which carbide containing Ti and Mo having an average particle size of less than 10 nm is dispersedly precipitated.
• Patent Document 2 discloses a technique relating to a high-strength hot-rolled steel sheet having a strength of 880 MPa or more and a yield ratio of 0.80 or more.
• the steel sheet has a steel structure that contains, on the basis of mass, C: 0.08% to 0.20%, Si: 0.001% or more but less than 0.2%, Mn: more than 1.0% but not more than 3.0%, Al: 0.001% to 0.5%, V: more than 0.1% but not more than 0.5%, Ti: 0.05% or more but less than 0.2%, and Nb: 0.005% to 0.5%, provided that the following three formulae are satisfied, the remainder being Fe and incidental impurities, and that contains 70% by volume or more ferrite having an average particle size of 5 ⁇ m or less and a hardness of 250 Hv or more.
• Patent Document 3 discloses a technique relating to a hot-rolled steel sheet that contains, on the basis of mass, C: 0.05% to 0.2%, Si: 0.001% to 3.0%, Mn: 0.5 to 3.0, P: 0.001% to 0.2%, Al: 0.001% to 3%, and V: more than 0.1% but not more than 1.5%, the remainder being Fe and impurities, and has a structure mainly composed of ferrite phase having an average particle size in the range of 1 to 5 ⁇ m, the ferrite particles containing carbonitride of V having an average particle size of 50 nm or less.
• Patent Document 4 discloses a thermally stable high-strength thin steel sheet that contains precipitated carbide in the steel structure.
• carbide has a NaCl-type crystal structure represented by MC wherein M denotes a metallic element composed of at least two metals, and the at least two metals are regularly spaced in a crystal lattice, forming a superlattice.
• Patent Document 5 discloses the following hot-rolled steel sheet.
• the steel sheet has a composition of C: 0.0002% to 0.25%, Si: 0.003% to 3.0%, Mn: 0.003% to 3.0%, and Al: 0.002% to 2.0% on the basis of mass percent, the remainder being Fe and incidental impurities, the impurities containing 0.15% or less P, 0.05% or less S, and 0.01% or less N.
• a ferrite phase accounts for 70% by area or more of the metal structure and has an average grain size of 20 ⁇ m or less and an aspect ratio of 3 or less. Seventy percent or more of ferrite grain boundaries are high-angle grain boundaries.
• the area percentage of precipitates having a maximum diameter of 30 ⁇ m or less and a minimum diameter of 5 nm or more is 2% or less of the metal structure.
• Second phases having the largest area percentage among phases other than the ferrite phases and the precipitates have an average grain size of 20 ⁇ m or less.
• High-angle grain boundaries of ferrite phases are disposed between the nearest second phases.
• Patent Document 6 discloses a drawable high-strength thin steel sheet that has excellent shape fixability and burring characteristics, wherein the thin steel sheet contains, on the basis of mass percent, C: 0.01% to 0.1%, S ⁇ 0.03%, N ⁇ 0.005%, and Ti: 0.05% to 0.5%, the Ti content satisfying Ti-48/12C-48/14N-48/32S ⁇ 0%, the remainder being Fe and incidental impurities, at least the mean values of X-ray random intensity ratios in a plane at half the thickness of the steel sheet are 3 or more for ⁇ 100 ⁇ 011> to ⁇ 223 ⁇ 110> orientations and 3.5 or less for three orientations of ⁇ 554 ⁇ 225>, ⁇ 111 ⁇ 112>, and ⁇ 111 ⁇ 110>, the arithmetical mean roughness Ra of at least one of the surfaces of the steel sheet ranges from 1 to 3.5, and the steel sheet is coated with a lubricating composition.
• Patent Document 5 discloses single-phase ferritic steel sheets having a tensile strength TS of 422 MPa or less (for example, test numbers 1 to 5 in Table 6 and test number 45 in Table 8 in Examples) and multiphase steel sheets composed of a ferrite phase and a second phase and having a tensile strength TS of 780 MPa or more (for example, test numbers 33 to 36 in Table 6 and test number 49 in Table 8 in Examples).
• These steel sheets described in Patent Document 5 mainly take advantage of solid-solution strengthening due to Si or Mn and transformation hardening utilizing a hard second phase.
• These steel sheets must therefore be cooled to a temperature in the range of 600°C to 800°C at an average cooling rate of 30°C/s or more within two seconds after finish rolling, air-cooled for 3 to 15 seconds, and then water-cooled at an average cooling rate of 30°C/s or more before coiling.
• This promotes two-phase separation during ferrite transformation, allowing the steel sheets to have a mixed structure of the ferrite phase and the second phase.
• the finish-rolling temperature ranges from (Ae3 point + 100°C) to Ae3 point, which is lower than the temperature range suitable for manufacture according to the present invention described below.
• the finish-rolling temperature for multiphase steel sheets having a tensile strength TS of 780 MPa or more ranged from 871°C to 800°C.
• a low finish-rolling temperature results in a decrease in the solubility limit of a carbide-forming element, such as Ti, in an austenite phase.
• a carbide-forming element such as Ti
• strain-induced precipitation increases the amount of precipitates having a size of 20 nm or more.
• Patent Document 5 also discloses a technique in which a ferritic single phase can be manufactured by greatly decreasing the C content and decreasing the amount of austenite forming element, Mn, in a steel composition (see steel numbers AA to AE in Table 2 in Examples).
• Mn austenite forming element
• a decrease in the amount of Mn, which is also a solid-solution strengthening element lowers the solid-solution strengthening level.
• a decrease in C content results in a decrease in the amount of precipitated carbide, for example, of Ti or Nb, which has precipitation hardening effects, thereby lowering the precipitation hardening level.
• a single-phase ferritic steel sheet cannot have a strength of 780 MPa or more (see test numbers 1 to 5 in Table 6 and test number 45 in Table 8 in Examples).
• an object of the present invention that is, a steel sheet that has a substantially ferritic single phase, a tensile strength of 780 MPa or more, and other characteristics cannot be manufactured by the technique described in Patent Document 5.
• Patent Document 6 discloses steel sheets having a tensile strength ⁇ B of 780 MPa or more (for example, steel symbols A-4, A-8, A-10, C, E, and H in Table 2 in Examples).
• the YRs of these steel sheets (YR represents ⁇ Y / ⁇ B x 100 (%)) are as low as 69% to 74%, indicating that these steel sheets contain a hard second phase, such as a martensite phase.
• Patent Document 5 the possible basic ideas behind the design of a steel sheet having a strength of 780 MPa or more according to Patent Document 6 mainly take advantage of solid-solution strengthening due to Si or Mn and transformation hardening utilizing a hard second phase. As described in Patent Document 5, therefore, rolling at a total reduction of 25% or more must be performed at a finish-rolling temperature (Ar3 point + 100°C or less) lower than the temperature range suitable for manufacture according to the present invention described below.
• the finish-rolling temperature for a steel sheet having a tensile strength ⁇ B of 780 MPa or more ranged from 800°C to 890°C.
• High-strength steel sheets according to the present invention have a tensile strength (hereinafter also referred to as TS) of 780 MPa or more and include hot-rolled steel sheets and surface-treated steel sheets, which are high-strength steel sheets subjected to surface treatment, such as plating.
• TS tensile strength
• Target characteristics of the present invention include a stretch flangeability ( ⁇ 10 ) of 60% or more after rolling at an elongation percentage of 10% and a one-side maximum peel width of 3.0 mm or less after a tape peel test in a warm salt water immersion test (SDT) described below.
• the present invention provides a high-strength hot-rolled steel sheet that has high stretch flangeability after working, corrosion resistance after painting, and a TS of 780 MPa or more.
• the present invention has these advantages without the addition of Mo and can therefore reduce costs.
• a high-strength hot-rolled steel sheet according to the present invention in automobile chassis and truck frames should allow thickness reduction, reduce the effects of automobiles on the environment, and markedly improve crashworthiness of automobiles.
• C can be precipitated in ferrite as carbide with Ti or V, thereby contributing to high strength of a steel sheet.
• 0.02% or more C is required to achieve a TS of 780 MPa or more.
• more than 0.20% C results in coarsening of precipitates and the formation of a second phase, lowering stretch flangeability after working.
• the C content ranges from 0.02% to 0.20%, preferably 0.03% to 0.15%.
• Si can contribute to solid-solution strengthening, the addition of more than 0.3% Si results in the formation of cementite at grain boundaries, lowering stretch flangeability after working.
• the Si content is 0.3% or less, preferably 0.001% to 0.2%.
• Mn can contribute to solid-solution strengthening.
• the TS is less than 780 MPa at a Mn content of less than 0.5%.
• the addition of more than 2.5% Mn markedly lowers weldability.
• the Mn content ranges from 0.5% to 2.5%, preferably 0.6% to 2.0%.
• the P can segregate at prior austenite grain boundaries, lowering workability and low-temperature toughness.
• the P content is preferably minimized and is 0.06% or less, preferably in the range of 0.001% to 0.055%.
• S can segregate at prior austenite grain boundaries or can be precipitated as MnS.
• the segregation or a large amount of MnS lowers low-temperature toughness.
• S also markedly lowers stretch flangeability, regardless of the presence or absence of working.
• the S content is preferably minimized and is 0.01% or less, preferably in the range of 0.0001% to 0.005%.
• Al can be added to steel as a deoxidizer and effectively improves the cleanliness of the steel.
• 0.001% or more Al is added to steel to produce this effect.
• more than 0.1% Al results in the generation of a large number of inclusions, causing flaws in a steel sheet.
• the Al content is 0.1% or less, preferably 0.01% to 0.04%.
• Ti is very important for the precipitation hardening of ferrite and is an important factor for the advantages of the present invention.
• a required strength is difficult to achieve at a Ti content of less than 0.05%.
• the effects of Ti become saturated at a Ti content of more than 0.25%, and more than 0.25% Ti only increases costs.
• the Ti content ranges from 0.05% to 0.25%, preferably 0.08% to 0.20%.
• V 0.05% to 0.25%
• V can contribute to an improvement in strength by precipitation hardening or solid-solution strengthening. Like Ti, V is therefore an important factor for the advantages of the present invention. A proper amount of V, together with Ti, tends to be precipitated as fine Ti-V carbide having a particle size (hereinafter also referred to as "size") of less than 20 nm. Unlike Mo, V does not lower corrosion resistance after painting. Less than 0.05% V is insufficient for the effects described above. However, the effects of V become saturated at a V content of more than 0.25%, and more than 0.25% V only increases costs. Thus, the V content ranges from 0.05% to 0.25%, preferably 0.06% to 0.20%.
• steel according to the present invention can have target characteristics.
• any one or two or more of Cr: 0.01% to 0.5%, W: 0.005% to 0.2%, and Zr: 0.0005% to 0.05% may be added for the following reasons.
• Cr, W, and Zr can strengthen ferrite as a precipitate or solid solution.
• Less than 0.01% Cr, less than 0.005% W, or less than 0.0005% Zr makes a negligible contribution to high strength of steel.
• more than 0.5% Cr, more than 0.2% W, or more than 0.05% Zr lowers workability.
• Cr, W, and Zr when any one or two or more of Cr, W, and Zr are added, their amounts are Cr: 0.01% to 0.5%, W: 0.005% to 0.2%, and Zr: 0.0005% to 0.05%, preferably Cr: 0.03% to 0.3%, W: 0.01% to 0.18%, and Zr: 0.001% to 0.04%.
• the remainder consists of Fe and incidental impurities.
• incidental impurity for example, O forms a non-metallic inclusion and has adverse effects on the quality of steel. O is therefore desirably decreased to 0.003% or less.
• 0.1% or less Cu, Ni, Sn, and/or Sb may be contained as a trace element without compromising the operational advantages of the present invention.
• a substantially ferritic single phase refers to allowance for a minute amount of another phase or precipitate other than carbide of the present invention, and the volume percentage of ferrite is preferably 95% or more.
• a substantially ferritic single phase may contain up to 5% by volume of cementite, pearlite, and/or bainite without affecting the characteristics of the present invention.
• the volume percentage of ferrite can be determined by exposing a microstructure in the vertical cross-section parallel to the rolling direction using 3% nital, observing the microstructure at a quarter thickness in the depth direction with a scanning electron microscope (SEM) at a magnification of 1500, and determining the ferrite area ratio, for example, using an image-processing software "Ryusi Kaiseki (particle analysis) II" from Sumitomo Metal Technology, Inc.
• a precipitate having a size of 20 nm or more has a small effect in preventing dislocation movement and cannot sufficiently increase the hardness of ferrite, sometimes resulting in low strength.
• a precipitate preferably has a size of less than 20 nm.
• a fine precipitate having a size of less than 20 nm can be formed by the addition of both Ti and V.
• V forms a complex carbide mainly with Ti.
• the Ti content and the V content of precipitates having a size of less than 20 nm are less than 200 ppm and less than 150 ppm, respectively, the number density of the precipitates is small, and the distance between precipitates increases. The precipitates therefore have a small effect in preventing dislocation movement. Thus, the precipitates cannot sufficiently increase the hardness of ferrite, and therefore the TS cannot be 780 MPa or more.
• the precipitates When the Ti content and the V content of precipitates having a size of less than 20 nm are 200 ppm or more and less than 150 ppm, respectively, the precipitates have a tendency to become coarse, and therefore the TS may be less than 780 MPa.
• the Ti content and the V content of precipitates having a size of less than 20 nm are less than 200 ppm and 150 ppm or more, respectively, the precipitation efficiency of V decreases, and therefore the TS may be less than 780 MPa.
• the Ti content or the V content of precipitates having a size of less than 20 nm is more than 1750 ppm, the corrosion resistance after painting decreases, and therefore the target characteristics cannot be achieved. This is probably because a large number of fine precipitates prevent the formation or growth of crystals on the surface of a steel sheet during chemical conversion.
• the amounts of precipitated Ti and V in precipitates having a size of less than 20 nm must be satisfactorily controlled.
• the TS can be 785 MPa or more, thus achieving more suitable conditions. Although there is no clear reason, optimization of the ratio of Ti to V should improve heat stability.
• the Ti content and the V content of precipitates having a size of less than 20 nm range from 200 to 1750 ppm and 150 to 1750 ppm, respectively. Furthermore, the ratio of the Ti content to the V content of precipitates having a size of less than 20 nm preferably satisfies 0.4 ⁇ (Ti/48)/(V/51) ⁇ 2.5.
• a precipitate and/or an inclusion is hereinafter also collectively referred to as a precipitate or the like.
• the Ti content and the V content can be controlled by the coiling temperature.
• the coiling temperature preferably ranges from 500°C to 700°C. At a coiling temperature above 700°C, precipitates become coarse, and the amounts of precipitated Ti and V in precipitates having a size of less than 20 nm are less than 200 ppm and less than 150 ppm, respectively, and the TS cannot be 780 MPa or more. At a coiling temperature below 500°C, the amounts of precipitated Ti and V in precipitates having a size of less than 20 nm are also less than 200 ppm and less than 150 ppm, respectively. Such a low coiling temperature should result in insufficient diffusion of Ti and V.
• the Ti content and the V content of precipitates having a size of less than 20 nm can be determined by the following method.
• the sample After a predetermined amount of sample is electrolyzed in an electrolyte, the sample is removed from the electrolyte and is immersed in a dispersive solution. Precipitates in the solution is filtered with a filter having a pore size of 20 nm. Precipitates in filtrate passing through the filter having a pore size of 20 nm have a size of less than 20 nm.
• the filtrate after filtration is appropriately analyzed by inductively coupled plasma (ICP) emission spectroscopic analysis, ICP mass spectrometry, atomic absorption spectrometry, or the like to determine the Ti content and the V content of precipitates having a size of less than 20 nm.
• ICP inductively coupled plasma
• V in solid solution is the most important factor. Solid solution of V is important in improving stretch flangeability after working. Less than 200 ppm V in solid solution has an insufficient effect, and 200 ppm or more V in solid solution is required to produce the effect described above. 1750 ppm or more V in solid solution exhibits a saturated effect and is considered as an upper limit.
• the amount of V in solid solution is 200 ppm or more but less than 1750 ppm.
• the workability of steel according to the present invention slightly deteriorates with increasing strength, when the Ti content and the V content of precipitates having a size of less than 20 nm are both 1750 ppm or less, 200 ppm or more V in solid solution can sufficiently ensure target stretch flangeability after working.
• 200 ppm or more but less than 1750 ppm V in solid solution can be measured, for example, by the following method.
• the analysis method may be inductively coupled plasma (ICP) emission spectroscopic analysis, ICP mass spectrometry, or atomic absorption spectrometry.
• ICP inductively coupled plasma
• a high-strength steel sheet according to the present invention can be manufactured by heating a steel slab adjusted within the chemical component ranges described above at a temperature in the range of 1150°C to 1350°C, hot-rolling the steel slab at a finish-rolling temperature in the range of 850°C to 1100°C, and coiling the rolled steel at a temperature in the range of 500°C to 700°C. Conditions suitable for these processes will be described in detail below.
• a carbide-forming element such as Ti or V
• Ti or V is mostly present as a precipitate in a steel slab.
• a precipitate in the form of carbide must be temporarily dissolved before hot rolling.
• a precipitate must therefore be heated at 1150°C or more.
• carbide having a size of 20 nm or more which does not contribute to precipitation hardening or corrosion resistance after painting, remains. This reduces the amount of Ti and V involved in the formation of fine precipitates having a size of less than 20 nm required for the advantages of the present invention. A target amount of precipitates having a size of less than 20 nm cannot therefore be obtained in coiling described below.
• carbide containing Ti or V remains dissolved during slab heating and finish rolling, and is precipitated as fine carbide containing Ti or V during coiling after finish rolling.
• the heating temperature is therefore more preferably 1170°C or more so that carbide can be dissolved almost completely.
• heating at a temperature above 1350°C excessively increases the crystal grain size, lowering stretch flangeability and elongation after working. Taking subsequent heat-treatment conditions into consideration, an increase in crystal grain size can be almost completely prevented at a heating temperature of 1300°C or less.
• the slab heating temperature preferably ranges from 1150°C to 1350°C, more preferably 1170°C to 1300°C.
• finish-rolling temperature is important in ensuring the Ti content and the V content of precipitates having a size of less than 20 nm according to the present invention.
• a steel slab after working is hot-rolled at a finish-rolling temperature in the range of 850°C to 1100°C, which is the final temperature of hot rolling.
• a finish-rolling temperature below 850°C a steel slab is rolled in a ferrite + austenite region and has an elongated ferrite phase. This may lower stretch flangeability or elongation after working.
• a finish-rolling temperature above 1100°C may result in coarsening of ferrite particles and a TS below 780 MPa.
• the finish-rolling temperature is more preferably 990°C or less to prevent coarsening of ferrite particles.
• the finish-rolling temperature preferably ranges from 850°C to 1100°C, more preferably 935°C to 990°C.
• the control of coiling temperature is important in ensuring the Ti content and the V content of precipitates having a size of less than 20 nm in the present invention. As described above, this is because, in the most desirable manufacturing form, this coiling process yields a large number of precipitation sites from which carbide is precipitated, thus preventing carbide grains from growing to 20 nm or more.
• the coiling temperature preferably ranges from 500°C to 700°C so that steel has a substantially ferritic single phase and the characteristics of the present invention can be achieved.
• a coiling temperature below 500°C may result in an insufficient amount of precipitated carbide containing Ti and/or V and reduced strength.
• a bainite phase may be formed in place of a ferritic single phase.
• the coiling temperature is preferably 500°C or more, more preferably 550°C or more.
• a coiling temperature above 700°C may result in coarsening of precipitated carbide and reduced strength.
• a coiling temperature above 700°C may also promote the formation of a pearlite phase, lowering stretch flangeability after working.
• the coiling temperature is more preferably 650°C or less to prevent coarsening of precipitated carbide without fail.
• the coiling temperature preferably ranges from 500°C to 700°C, more preferably 550°C to 650°C.
• Steel sheets according to the present invention include surface-treated steel sheets and surface-coated steel sheets.
• a steel sheet according to the present invention may be subjected to hot-dip galvanizing to form a galvanized steel sheet, and the present invention can be suitably applied to such a galvanized steel sheet. Because a steel sheet according to the present invention has excellent workability, such a galvanized steel sheet can also have excellent workability.
• Hot-dip galvanizing is zinc and zinc-based (approximately 90% or more) hot dipping and includes hot dipping including an alloying element, such as Al or Cr, as well as zinc. Hot-dip galvanizing may be performed alone or followed by alloying.
• a steel melting method is not particularly limited, and any known melting method may be suitable.
• a suitable melting method involves melting in a converter or an electric furnace and secondary refining in a vacuum degassing furnace.
• a casting method is preferably continuous casting in terms of productivity and quality. After casting, hot direct rolling may be performed immediately or after concurrent heating, without compromising the advantages of the present invention.
• a hot-rolled material may be heated after rough rolling and before finish rolling, continuous hot rolling in which rolled materials are joined may be performed after rough rolling, or heating and continuous rolling of a heating material of a rolled material may be performed simultaneously. These do not compromise the advantages of the present invention.
• the microstructure of the hot-rolled steel sheet was analyzed by the following method to determine the Ti content and the V content of precipitates having a size of less than 20 nm and the amount of V in solid solution.
• the tensile strength TS, the stretch flangeability after working ⁇ 10 , and the corrosion resistance after painting (SDT one-side maximum peel width) were measured.
• the hot-rolled steel sheet thus formed was cut into an appropriate size. Approximately 0.2 g of hot-rolled steel sheet was subjected to constant-current electrolysis at an electric current density of 20 mA/cm 2 in 10% AA electrolyte (10% by volume acetylacetone-1% by mass tetramethylammonium chloride-methanol).
• a test piece on which a precipitate was deposited was removed from the electrolyte and was immersed in aqueous sodium hexametaphosphate (500 mg/l) (hereinafter referred to as aqueous SHMP). Ultrasonic vibration was applied to the test piece to detach and extract the precipitate from the test piece in aqueous SHMP.
• the aqueous SHMP containing the precipitate was then passed through a filter having a pore size of 20 nm. The filtrate was analyzed with an ICP spectrometer to measure the absolute amounts of Ti and V in the filtrate.
• the absolute amounts of Ti and V were divided by the weight of the electrolyzed sample to calculate the Ti content and the V content of precipitates having a size of less than 20 nm.
• the weight of electrolyzed sample was calculated by subtracting the sample weight after the detachment of the precipitate from the sample weight before electrolysis.
• the concentrations of V and a comparative element Fe in the electrolyte were measured by ICP mass spectrometry. On the basis of the concentrations thus measured, the ratio of the concentration of V to the concentration of Fe was calculated. The ratio was multiplied by the Fe content of the sample to calculate the amount of V in solid solution. The Fe content of the sample can be calculated by subtracting the summation of compositions other than Fe from 100%.
• a tensile test according to JIS Z 2241 was performed with a JIS No. 5 specimen in the tensile direction parallel to the rolling direction to measure TS.
• a chemical conversion treatment was performed under more adverse temperature and concentration conditions than the standard conditions using a degreasing agent, Surfcleaner EC090, a surface conditioner, Surffine 5N-10, and a chemical conversion treatment agent, Surfdine SD2800, all manufactured by Nippon Paint Co., Ltd.
• a degreasing process included a concentration of 16 g/l, a treatment temperature in the range of 42°C to 44°C, a treatment time of 120 s, and spray degreasing, and a surface conditioning process included a total alkalinity in the range of 1.5 to 2.5 points, a free acidity in the range of 0.7 to 0.9 points, an accelerator concentration in the range of 2.8 to 3.5 points, a treatment temperature of 44°C, and a treatment time of 120 s.
• a treatment temperature in a chemical conversion treatment process was decreased to 38°C.
• electrodeposition coating was performed using an electrodeposition paint, V-50, manufactured by Nippon Paint Co., Ltd.
• the target amount of deposited chemical conversion film ranged from 2 to 2.5 g/m 2 , and the target film thickness in electrodeposition coating was 25 ⁇ m.
• Corrosion resistance after painting was determined in a warm salt water immersion test (SDT).
• SDT warm salt water immersion test
• a crosscut was formed with a cutter in a sample subjected to chemical conversion treatment and electrodeposition coating. The sample was immersed in warm salt water (5% NaCl at 55°C) for 10 days, was then washed with water, and was dried. Tape peeling on the crosscut was performed to measure the maximum peel width on the left and right sides of the crosscut. A one-side maximum peel width of 3.0 mm or less was considered as high corrosion resistance after painting.
• Table 2 shows the results, together with manufacturing conditions.
• Table 2 No Type of steel Slab heating temperature (°C) Finish-rolling temperature (°C) Coiling temperature (°C) TS (MPa) Elongation after prestraining (%) Stretch flangeability after working: ⁇ 10 (%) Precipitated Ti content for ⁇ 20 nm (mass ppm) Precipitated V content for ⁇ 20 nm (mass ppm) Amount of V in solid solution (mass ppm) One-side maximum peel width (mm) Phase Note 1 A 1250 920 630 812 20 87 752 818 340 1.7 Ferrite: 100% Example 2 B 1300 926 632 952 18 79 1580 765 231 2.5 Ferrite: 100% Example 3 C 1270 911 650 966 17 81 703 1700 380 2.2 Ferrite: 100% Example 4 C 1270 900 580 865 17 95 635 657 1350 1.2 Ferrite: 99%, Remainder: Cementite 1% Example 5 D 1270
• Table 2 shows that the working examples had a TS of 780 MPa or more, ⁇ 10 of 60% or more, and an SDT one-side maximum peel width of 3.0 mm or less, indicating that the hot-rolled steel sheets had high stretch flangeability after working and corrosion resistance after painting.
• the comparative examples had a low TS (strength), small ⁇ 10 (stretch flangeability after working), and/or a large SDT one-side maximum peel width (corrosion resistance after painting).
• Example 1 the microstructure of the hot-rolled steel sheet thus formed was analyzed to determine the Ti content and the V content of precipitates having a size of less than 20 nm and the amount of V in solid solution.
• the tensile strength TS, the stretch flangeability after working ⁇ 10 , and the corrosion resistance after painting (SDT one-side maximum peel width) were measured.
• Table 4 shows the results. Table 4 No Type of steel Slab heating temperature (°C) Finish-rolling temperature (°C) Coiling temperature (°C) TS (MPa) Elongation after prestraining (%) Stretch flangeability after working: ⁇ 10 (%) Precipitated Ti content for ⁇ 20 nm (mass ppm) Precipitated V content for ⁇ 20 nm (mass ppm) Amount of V in solid solution (mass ppm) One-side maximum peel width (mm) Phase Note 25 T 1250 921 625 832 17 99 750 815 250 2.5 Ferrite: 100% Example 26 U 1250 918 620 830 18 90 753 760 252 2.2 Ferrite: 100% Example 27 V 1250 920 621 829 17 93 753 770 250 2.0 Ferrite: 100% Example 28 W 1250 921 620 842 18 98 760 823 251 2.6 Ferrite: 100% Example 35 T 1250 940 600 835 18 92 780 820 240 2.2 Ferrite: 100% Example 36
• Table 4 shows that the working examples had a TS of 780 MPa or more, ⁇ 10 of 60% or more, and an SDT one-side maximum peel width of 3.0 mm or less, indicating that the hot-rolled steel sheets had high stretch flangeability after working and corrosion resistance after painting.
• a steel sheet according to the present invention had high strength, high stretch flangeability after working, and high corrosion resistance after painting, and is therefore most suitable for, for example, automobile and truck frames, and components that require elongation and stretch flangeability.

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