EP2058411B1 - Wärmebehandelter stahldraht für hochfeste feder - Google Patents

Wärmebehandelter stahldraht für hochfeste feder Download PDF

Info

Publication number
EP2058411B1
EP2058411B1 EP06823432.7A EP06823432A EP2058411B1 EP 2058411 B1 EP2058411 B1 EP 2058411B1 EP 06823432 A EP06823432 A EP 06823432A EP 2058411 B1 EP2058411 B1 EP 2058411B1
Authority
EP
European Patent Office
Prior art keywords
carbides
amount
less
strength
spring
Prior art date
Legal status (The legal status is an assumption and is not a legal conclusion. Google has not performed a legal analysis and makes no representation as to the accuracy of the status listed.)
Active
Application number
EP06823432.7A
Other languages
English (en)
French (fr)
Other versions
EP2058411A4 (de
EP2058411A1 (de
Inventor
Masayuki Hashimura
Hiroshi Hagiwara
Takayuki Kisu
Kouichi Yamazaki
Tatsuro Ochi
Takashi Fujita
Current Assignee (The listed assignees may be inaccurate. Google has not performed a legal analysis and makes no representation or warranty as to the accuracy of the list.)
Nippon Steel Corp
Original Assignee
Nippon Steel and Sumitomo Metal Corp
Priority date (The priority date is an assumption and is not a legal conclusion. Google has not performed a legal analysis and makes no representation as to the accuracy of the date listed.)
Filing date
Publication date
Application filed by Nippon Steel and Sumitomo Metal Corp filed Critical Nippon Steel and Sumitomo Metal Corp
Publication of EP2058411A1 publication Critical patent/EP2058411A1/de
Publication of EP2058411A4 publication Critical patent/EP2058411A4/de
Application granted granted Critical
Publication of EP2058411B1 publication Critical patent/EP2058411B1/de
Active legal-status Critical Current
Anticipated expiration legal-status Critical

Links

Images

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
    • C22C38/18—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
    • C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
    • 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
    • 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
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/34—Ferrous alloys, e.g. steel alloys containing chromium with more than 1.5% by weight of silicon

Definitions

  • the present invention relates to a heat treated steel wire for a spring.
  • the method of production of a coil spring made of steel includes hot coiling heating the steel to the austenite region for coiling, then quenching and tempering it and cold coiling a high strength steel wire comprised of steel quenched and tempered beforehand.
  • cold coiling oil tempering, high frequency treatment, etc. enabling rapid heating and rapid cooling when producing the steel wire can be used, so it is possible to reduce the grain size of the prior austenite of the spring material and as a result produce a spring excellent in breakage property.
  • this method has the advantage that the heating furnace and other equipment on the spring manufacturing line can be simplified, so leads to a reduction of the equipment cost for spring makers. Cold coiling of springs is therefore being shifted to in recent years. In suspension springs as well, while a large diameter of steel wire is used compared with valve springs, cold coiling is introduced due to the above advantage.
  • WO 2006/059784 US 2003/0201036 A1 and US 6,338,763 B1 disclose a high-strength spring steel and a spring steel wire in which Nb is added with a high amount as an optional element.
  • the present invention has as its task to provide a heat treated steel wire for a spring with a tensile strength of 2000 MPa or more which is coiled in a cold state and can achieve both sufficient strength in the atmosphere and coilability and spring steel used for that steel wire.
  • the inventors discovered that by controlling the N, which was not focused on until now, even if adding alloy elements, it is possible to suppress the formation of undissolved carbides and possible to secure toughness and workability and thereby developed a heat treated steel wire for a spring achieving both a high strength and coilability.
  • the present inventors set the chemical ingredients to obtain a high strength and controlled the shape of the carbides in the steel by heat treatment so as to secure a coiling property sufficient for producing a spring in a steel wire and thereby reached the present invention.
  • C is an element which greatly affects the basic strength of a steel material and is set to 0.5 to 0.9% so as to obtain a strength more sufficient that the past. If less than 0.5%, a sufficient strength cannot be obtained. In particular, even when omitting the nitriding for improving spring performance, 0.5% or more of C is required to secure a sufficient spring strength. If over 0.9%, a substantial hypereutectoid appears, and a large amount of coarse cementite precipitates, therefore the toughness is remarkably lowered. This simultaneously lowers the coiling property. Further, the relationship with the microstructure is also close.
  • the number of carbides is small, so the regions where carbide distribution is locally smaller than other parts (hereafter described as "carbide poor regions") easily increase and sufficient strength and toughness or coilability (ductility) are hard to obtain. Therefore, preferably it is 0.55% or more, from the viewpoint of the balance of strength-coiling, more preferably 0.6% or more.
  • the alloy-based and cementite-base carbides tend to be hard to dissolve by the heat during the quenching.
  • the heating temperature in the heat treatment is high or when the heating time is short, the strength and coilability are often insufficient.
  • the undissolved carbides also affect the carbide poor regions. If the C in steel forms undissolved carbides, the de facto C in the matrix is decreased, so as previously explained, the area ratio of the carbide poor regions sometimes increase. Further, if the amount of C increases, the form of the martensite during tempering becomes the general lath martensite in medium carbon steel, while when the amount of C is great, it is known that the form changes to lenticular martensite.
  • the carbide distribution of the tempered martensite structure formed by tempering the lenticular martensite is lower in carbide density compared with the case of tempering the lath martensite. Consequently, by increasing the amount of C, the increase of the lenticular martensite and undissolved carbides sometimes causes the carbide poor regions to increase. For this reason, it is preferably 0.7% or less. More preferably, by making it 0.65% or less, it is possible to relatively easily reduce the carbide poor regions.
  • Si is an element necessary for securing strength, hardness, and settling resistance of a spring. If the amount is small, the required strength and settling resistance are insufficient, therefore 1.0% is made the lower limit. Further, Si has the effect of spheroidizing and refining the carbide precipitates of the grain boundary. By actively adding it, there is the effect of reducing the grain boundary area percentage of grain boundary precipitates. However, when adding too large an amount, the material not only hardens, but also embrittles. Therefore, 3.0% is set as an upper limit to prevent embrittlement after quenching and tempering. Further, Si is an element contributing to tempering softening resistance. To prepare a high strength wire rod, it is preferable to add a large amount to a certain extent. Specifically, it is preferable to add 2% or more. On the other hand, to obtain a stable coilability, it is preferable to make it 2.6%.
  • Mn is used for deoxidation and for fixing S in the steel as MnS, to raise quenching, and to sufficiently obtain hardness after heat treatment.
  • 0.1% is set as the lower limit in order to secure this stability.
  • the upper limit is set at 2.0% in order to prevent embrittlement caused by Mn.
  • it is preferably 0.3 to 1%. Further, when giving priority to coiling, making it 1.0% or less is effective.
  • Cr is an effective element to improve quenching and softening resistance in tempering. Further, it is an effective element not only for securing tempering hardness, but also for increasing the surface layer hardness after nitridation and the depth of the hardened layer in nitridation such as seen in recent high strength valve springs. However, if the added amount is large, not only is an increase in cost incurred, but also the cementite seen after quenching and tempering coarsens. Further, it has the effect of stabilizing and coarsening the alloy-based carbides. As a result, the wire rod becomes brittle, so there is also the negative effect that the rod easily breaks during coiling. Consequently, when adding Cr, if not 1.0% or more, the effect is not clear.
  • the carbides are finely controlled by setting the N, so a large amount of Cr can be added, therefore the amount added was made one easily giving a high strength.
  • the addition of Cr enables the hardened layer obtained by the nitridation to be made deeper. Because of this, addition of 1.1% or more is preferable. Further, to make the rod suitable for nitridation for an unprecedented high strength spring, addition of 1.2% or more is preferable.
  • the added amount is made 2.0% or less. More preferably it is made 1.7% or less.
  • V can be utilized for the hardening of the steel wire at the tempering temperature and the hardening of the surface layer during nitriding due to the secondary precipitation and hardening for precipitating and hardening the carbides during tempering. Further it is effective for suppressing coarsening of the austenite grain size due to the formation of nitrides, carbides, and carbonitrides. Addition is therefore preferable. However, until now, because the nitrides, carbides, and carbonitrides of V are formed at even the austenitizing temperature A 3 point of steel, when insufficiently dissolved, they easily remain as undissolved carbides (nitrides). The undissolved carbides not only become the cause of breakage during spring coiling, but also "wastefully consume the V".
  • controlling the amount of N enables the formation of V-based nitrides, carbides, and carbonitrides at the austenitizing temperature A 3 point or more to be suppressed, so it is possible to add a larger amount of V by that amount.
  • the added amount of V was therefore made over 0.15% to 1.0%.
  • the added amount is 0.15% or less, there is little effect of adding V such as the improvement of hardness of the nitrided layer and increase of depth of the nitrided layer and a sufficient fatigue limit (durability) of conventional steel or more cannot be secured. Further, if the added amount is over 1.0%, coarse undissolved inclusions are formed and the toughness is reduced. In the same way as Mo, an overcooled structure is easily formed and cracks or breakage during drawing are easily caused. For those reasons, 1.0%, where industrially stable handling is easy, was made the upper limit.
  • Nitrides, carbides, and carbonitrides of V are formed even at the austenitizing temperature A 3 point of steel or more, so when dissolution is insufficient, they easily remain as undissolved carbides (nitrides). Therefore, if considering the current ability to control the amount of nitrogen industrially, making it industrially 0.5% or less is preferable and making it 0.4% or less is more preferable.
  • the rod is heated to a maximum of a temperature of 300°C or more, so to suppress hardening of the top surface layer and softening of the inner portion hardness by nitridation, it is necessary to add over 0.15%. Preferably, addition of 0.2% or more is preferable.
  • Al is a deoxidation element and influences the formation of oxides.
  • hard oxides such as Al 2 O 3 easily become the starting points of breakage, so it is necessary to avoid this. For this reason, it is important to strictly control the amount of Al.
  • the tensile strength as a heat treated steel wire is over 2100 MPa, strict control of the oxide-forming elements is essential to reduce fluctuations in the fatigue strength.
  • Al was set to 0.005% or less. This is because if over 0.005%, Al 2 O 3 -based oxides are easily formed, so breakage caused by the oxides occurs and a sufficient fatigue strength and quality stability cannot be secured. Further, when requiring high strength fatigue, it is preferable made 0.003% or less.
  • N is an important point.
  • a strict limit value of N ⁇ 0.007% is set. This is because in the present invention, the role of N is newly focused on.
  • the effects of N control and the reasons for the provisions in the present invention will be explained below.
  • the effects of N are as follows: 1) N is present in ferrite as dissolved N which suppresses the movement of the dislocations in the ferrite and thereby causes the ferrite to harden. 2) nitrides are formed with Ti, Nb, V, Al, B, and other alloy elements and affect the performance of the steel material. The mechanism and the like will be explained later. 3) N affects precipitation behavior of cementite and other iron-based carbides and affects the performance of the steel material performance.
  • These chemical compositions are mainly nitrides at a high temperature.
  • the form changes to carbonitrides and carbides along with cooling. Consequently, the nitrides formed at a high temperatures easily become nuclei for the precipitation of V carbides.
  • the tempering temperature is made 300 to 500°C.
  • Nb and optionally Ti are added.
  • the N amount can be suppressed to 0.003% or less, good performance is obtained without adding one or both of Ti and Nb, but industrially stably making the amount 0.003% or less becomes disadvantageous in the point of manufacturing cost. Therefore, small amounts of Nb and optionally Ti are added. If adding Ti and Nb, these elements form nitrides at a high temperature, so substantially reduce the dissolved nitrogen. Therefore, the same effect as with reducing the amount of N added can be obtained. Because of that, the upper limit of the added amount of N may be increased. However, if the amount of N exceeds 0.007%, the amount of V, Nb, or Ti nitrides becomes larger. As a result, the undissolved carbides become greater and the TiN and other hard inclusions increase, so the toughness falls and the fatigue limit characteristics and coiling characteristics fall. Therefore, the upper limit of the amount of N was limited to 0.007%.
  • the upper limit of the amount of N is preferably made 0.005% or less, more preferably 0.004% or less.
  • Nb forms nitrides, carbides, and carbonitrides.
  • the nitrides are produced at a higher temperature than with V. Due to this, formation of Nb nitrides during cooling consumes the N in the steel and can suppress the formation of V-based nitrides. As a result, the formation of V-based undissolved carbides can be suppressed, so temper softening resistance, workability and coilability can be secured.
  • Nb-based carbonitrides suppressing the coarsening of the austenite grain size, they can be utilized for hardening the steel wire at the tempering temperature and hardening the surface layer during nitriding.
  • the added amount is too great, undissolved carbides with Nb-based nitride nuclei easily remain, so addition of a large amount should be avoided.
  • the Nb added amount is less than 0.001%, almost no effect of addition is recognized.
  • 0.01% or more the large amount of addition forms coarse undissolved inclusions, lowers the toughness, and, like with Mo, easily forms an overcooled structure easily causing cracks and breakage during drawing. Therefore, the amount is made the 0.01% where industrially stable handling is easy.
  • FIG. 1 is a view showing the results of measurement of the impact values of materials of the chemical compositions shown in Table 1, that is, the results of measurement of the impact values of the samples A and B heat treated by the method of the examples described below.
  • Table 1 Steel compositions (mass%) C Si Mn P S Cr Mo V: W Nb N S-Al Sample A 0.61 2.20 0.53 0.002 0.004 1.21 0.13 0.20 0.16 - 0.0049 0.002 Sample B 0.61 2.21 0.54 0.002 0.004 1.19 0.13 0.20 0.16 0.009 0.0050 0.002
  • this added amount when adding Ti, this added amount is 0.001% to less than 0.005%.
  • Ti is a deoxidizing element and an element forming nitrides and sulfides, so has an effect on the formation of oxides, nitrides, and sulfides. Consequently, addition of a large amount facilitates formation of hard oxides and nitrides, so if adding this carelessly, it forms hard carbides and lowers the fatigue limit. Like with Al, in particular in high strength springs, it lowers the stability of fluctuation of the fatigue strength more than the fatigue limit itself of the spring. If the amount of Ti is great, the breakage rate due to inclusions becomes greater, so it is necessary to control this amount. The amount is made less than 0.005%.
  • Ti forms TiN in the molten steel at a high temperature, so acts to reduce the sol.N in the molten steel.
  • limiting the N to suppress the formation of the V-based nitrides and further suppress the growth of the V-based undissolved carbides is the point of the technology. For this reason, if consuming the N at a temperature of the V-based nitride formation temperature or more, it is possible to suppress the growth of V-based nitrides and V-based carbonitrides growing using these as nuclei during cooling. That is, adding Ti substantially reduces the amount of N bonding with V, so reduces the temperature of formation of the V-based nitrides and further suppresses V-based undissolved carbides.
  • the added amount is 0.001% or more. If less than 0.001%, there is no effect of N consumption, there is no effect of suppressing V-based undissolved carbides, and the effect of improvement of the workability (coilability) cannot be seen. However, the amount of addition of Ti is preferably 0.003% or less.
  • the steel of the present invention has the above stated chemical compositions as basic compositions and further may have added to it compositions in order to improve the properties of the steel. That is, further, one or both of W and Mo are added when strengthening the temper softening resistance. W not only improves quenching, but also acts to form carbides in the steel and raise the strength and is effective for conferring temper softening resistance. Therefore, adding as much as possible is preferable. W forms carbides at a lower temperature compared to Ti, Nb, and the like, so it does not easily form undissolved carbides. Further, it can confer temper softening resistance by precipitation hardening. That is, in nitriding and strain relief annealing as well, the inner hardness will not be greatly decreased.
  • the added amount is 0.05% or less, the effect is not seen, while if 0.5% or more, coarse carbides are formed and conversely the ductility and other mechanical properties are liable to be impaired, so the added amount of W was set to 0.05 to 0.5%. Further, if considering the ease of heat treatment, 0.1 to 0.4% is preferable. In particular, to avoid an overcooled structure right after rolling and other negative effects and obtain the maximum extent of temper softening resistance, adding 0.15% or more is further preferable.
  • Mo improves hardenability and precipitates as carbides at a temperature of about the tempering and nitriding temperature, so can confer temper softening resistance. Therefore, even after high temperature tempering, strain relief annealing or nitriding in the process, or other heat treatment, the steel does not soften and can exhibit a high strength. This suppresses the decrease of the spring internal hardness after nitriding and facilitates hot setting and strain relief annealing, so improves the fatigue characteristics of the final spring. That is, it is possible to make the tempering temperature higher when controlling the strength. Making the tempering temperature higher is advantageous in decreasing the grain boundary area percentage of the grain boundary carbides.
  • the Mo forms Mo-based carbides separate from cementite in the steel.
  • V etc. it has a lower precipitation temperature, so has the effect of suppressing the coarsening of the carbides.
  • the added amount is 0.05% or less, no effect is recognized.
  • the added amount is great, an overcooled structure easily is formed in the softening heat treatment before rolling or drawing and easily causes cracks and breakage at the time of drawing. That is, when drawing, it is preferable to first patent the steel material to convert it to a ferrite-pearlite structure.
  • Mo is an element which confers large hardenability, so when the added amount becomes large, the time until the end of the pearlite transformation becomes longer, an overcooled structure is easily formed in the cooling after rolling or in the patenting process and becomes a cause of breakage when drawing, or, when not breaking but having internal cracks, the characteristics of the final product are greatly degraded. If Mo exceeds 0.5%, the hardenability becomes great and industrially making a ferrite-pearlite structure becomes difficult, so this was made the upper limit. To suppress the formation of a martensite structure causing a drop in the production ability in the rolling, drawing, or other production process and facilitate industrially stable rolling and drawing, 0.4% or less is preferable and 0.2% or so is more preferable.
  • V, Nb, and Ti similarly having an effect of strengthening the temper softening resistance
  • V, Nb, and Ti form nitrides as explained above and facilitate the growth of carbides with these as nuclei
  • W and Mo do not form nitrides much at all, so are free of the effects of the amount of N and can strengthen the softening resistance if added.
  • strengthening of the softening resistance is possible even with V, Nb, and T, but the amounts added end up being self restricted for addition for strengthening the softening resistance while avoiding undissolved carbides.
  • Ni, Cu, Co, and B may be added to secure strength by strengthening the matrix when the optimal balance of the softening resistance and workability by control of the carbides cannot be obtained in achieving both strength and workability.
  • Ni improves the hardenability and enables stable increase of strength by heat treatment. Further, it improves the ductility of the matrix and improves the coilability. However, quenching and tempering increase the retained austenite, so the settling and uniformity of the material are inferior after forming the spring. If the added amount is 0.05% or less, an effect in increasing the strength and improving the ductility cannot be recognized. On the other hand, addition of a large amount of Ni is not preferable. At 3.0% or more, the negative effect of the greater retained austenite becoming greater becomes remarkable and the effect of improving the hardenability and ductility become saturated, which is disadvantageous from the point of cost and the like.
  • adding Cu can prevent decarburization.
  • a decarburized layer decreases fatigue life after processing the spring, so effort is made to reduce it as much as possible. Further, when the decarburized layer becomes deep, the surface layer is removed by the process called peeling. Further, in the same way as Ni, there is an effect of improving the corrosion resistance. By suppressing a decarburized layer, the improvement of the fatigue life and peeling of the spring can be omitted.
  • the effect of suppression of decarburization and effect of improvement of the corrosion resistance by Cu can be exhibited at 0.05% or more. As explained later, even if adding Ni, even over 0.5%, embrittlement easily causes rolling marks. Therefore, the lower limit was made 0.05% and the upper limit was made 0.5%.
  • the addition of Cu does not harm the mechanical properties much at all, but when adding Cu 0.3%, the hot rollability is degraded, so cracks sometimes are formed at the billet surface at the time of rolling. Because of that, it is preferable to change the amount of Ni added to prevent cracking during rolling in accordance with the amount of Cu added, that is, [Cu%] ⁇ [Ni%]. In the range of Cu of 0.3% or less, no rolling marks will be caused, so there is no need to define the amount of Ni added for the purpose of preventing rolling marks.
  • Co decreases hardenability in some cases, but can improve the high temperature strength. Further, it inhibits the formation of carbides, so acts to suppress the formation of coarse carbides at issue in the present invention. Consequently, it can suppress the coarsening of the cementite and other carbides. Therefore, addition is preferable. When added, if 0.05% or less, the effect is small. However, if a large amount is added, the hardness of the ferrite phase increases and the ductility is lowered, so the upper limit was made 3.0%. Industrially, stable performance is obtainable at 0.5% or less.
  • B is an element improving the hardenability and effective for cleaning the austenite grain boundaries.
  • the P, S, and other elements segregating at the grain boundaries and lowering the toughness are rendered harmless and the breakage characteristics are improved by addition of B.
  • the added amount was therefore made 0.0005% where the effect becomes clear as the lower limit and 0.0060% where the effect is saturated as the upper limit.
  • the amount is 0.003 or less. More preferably, it is effective to use the Ti, Nb, and other nitride-forming elements to fix the free N and make B 0.0010 to 0.0020%.
  • Ni, Cu, Co, and B are mainly effective for strengthening the ferrite phase of the matrix. These are elements effective when securing strength by strengthening the matrix when an optimal balance of softening resistance and workability cannot be obtained by control of carbides in order to achieve both strength and workability.
  • Te, Sb, Mg, Zr, Ca, and Hf are added as elements to control the form of the oxides and sulfides when further higher performance and stabler performance are sought.
  • Te has the effect of making the MnS spheroidal. If less than 0.0002%, that effect is not clear, while if over 0.01%, the negative effects of decreasing the toughness of the matrix, causing hot breakage, and decreasing the fatigue durability become remarkable, so 0.01% is set as the upper limit.
  • Sb has the effect of spheroidizing MnS. If less than 0.0002%, that effect is not clear, while if over 0.01%, the negative effects of decreasing the toughness of the matrix, causing hot breakage, and decreasing the fatigue limit become remarkable, so 0.01% is set as the upper limit.
  • Mg forms oxides in molten steel higher than the MnS formation temperature and is already present in the molten steel at the time of MnS formation. Therefore, it can be used as nuclei for precipitation of MnS and thereby can control the distribution of MnS. Further, in numerical distribution, Mg-based oxides disperse in the molten steel finer than the Si- and Al-based oxides often seen in conventional steels, so the MnS with Mg-based oxides as nuclei finely disperse in the steel. Therefore, even with the same S content, the MnS distribution differs depending upon the presence of Mg. By adding this, the MnS grain size becomes further finer. This effect can be sufficiently obtained even in small amounts.
  • the MnS becomes finer. However, if over 0.0005%, hard oxides easily form. Further, MgS and other sulfides also start to be formed and a drop in fatigue strength and a drop in coilability are incurred. Therefore, the amount of Mg added is made 0.0001 to 0.0005%. If used for a high strength spring, 0.0003% or less is preferable. These elements are used in very small amounts, but by using large amounts of Mg-based refractories, about 0.0001% can be added. Further, by carefully selecting the auxiliary materials and using auxiliary materials having a low Mg content, the amount of Mg added can be controlled.
  • Zr is an oxide- and sulfide-forming element.
  • spring steel it finely disperses the oxides, so like with Mg forms the nuclei for precipitation of MnS. Due to this, the fatigue limit is improved and the ductility is increased, so the coilability is improved. If less than 0.0001%, this effect is not seen, while even if added in an amount over 0.0005%, formation of hard oxides is promoted, so even if the sulfides are finely dispersed, trouble due to the oxides easily occurs.
  • the amount was made 0.0005% or less. Further, when used for a high strength spring, this added amount is preferably made 0.0003% or less. These elements are very small in amounts, but can be controlled by carefully selecting the auxiliary materials and precisely controlling the refractories and the like.
  • Ca is an oxide- and sulfide-forming element.
  • spring steel it makes MnS spheroidal and thereby can suppress the length of the MnS acting as the starting point of fatigue and other breakage and render it harmless. This effect is not clear if less than 0.0002%, while even if added in an amount over 0.01%, not only does the yield become poor, but also oxides and CaS and other sulfides are formed and manufacturing trouble and a decrease in fatigue limit characteristics are caused, so the amount was made 0.01% or less.
  • the added amount is preferably 0.001% or less.
  • Hf is an oxide-forming element and becomes the nuclei for precipitation of MnS. Due to this, by fine dispersal, Zr is an oxide- and sulfide-forming element. In spring steel, the oxides finely disperse, so like Mg, these become nuclei for precipitation. Due to this, the fatigue limit is improved and the ductility is increased, so the coilability is improved. The effect is not clear if less than 0.0002%, while even if added in an amount of over 0.01%, not only does the yield become poor, but also oxides and ZrN, ZrS, and other nitrides and sulfides are formed and manufacturing trouble and a decrease in fatigue limit characteristics are caused, so the amount was made 0.01% of less. The added amount is preferably 0.003% or less.
  • P makes the steel harden, but further causes segregation and makes the material become brittle.
  • the P segregated at the austenite grain boundaries causes a drop in impact value and delayed breakage due to penetration of hydrogen and the like. Therefore, less is better.
  • the amount is preferably made P: 0.015% or less where this embrittlement tendency becomes remarkable. Further, in the case of a high strength where the tensile strength of the heat treated steel wire exceeds 2150 MPa, the amount is preferably made less than 0.01%.
  • S causes embrittlement of the steel when present in the steel. Its influence is made much smaller by Mn, but MnS also takes the form of inclusions, so the breakage characteristics decline. In particular, in high strength steel, breakage sometimes occurs due to a very fine amount of MnS, so the S is preferably reduced as much as possible. Making the amount 0.015% or less where these negative effects become remarkable is preferable. Further, in the case of a high strength where the tensile strength of the heat treated steel wire exceeds 2150 MPa, the amount is preferably made less than 0.01%.
  • t-O is made 0.0002 to 0.01%.
  • Steel contains oxides formed through the deoxidation process and dissolved O.
  • t-O total amount of oxygen
  • the size of the oxide-based inclusions is small, they will not affect spring performance, but if there is a large amount of large oxides present, they will have a great effect on spring performance.
  • the amount of oxygen is over 0.01%, the spring performance is remarkably reduced, so the upper limit is preferably made 0.01%. The smaller the amount of oxygen the better, but even if less than 0.0002%, the effect is saturated, so this is preferably made the lower limit. If considering ease in the actual deoxidation process and the like, adjustment to 0.0005 to 0.005% is preferable.
  • the tensile strength is preferably made 2000 MPa or more. If the tensile strength is high, the fatigue characteristics of the spring tend to be improved. Further, even when applying nitridation or other surface hardening treatment, if the basic strength of the steel wire is high, high fatigue characteristics and settling characteristics can be obtained. On the other hand, if the strength is high, the coilability declines and spring production becomes difficult. Because of this, it is important to not only improve the strength, but also to impart ductility enabling coiling.
  • TS ⁇ 2000 MPa is made the lower limit. Further, when applied to a high strength spring, further higher strength is preferable.
  • the amount is preferably 2200 MPa or more and further, for application to a high strength spring, increase of the strength to 2250 or 2300 MPa or more in a range not impairing the coilability is preferable.
  • the undissolved carbides are generally spheroidal and include ones mainly made of alloy elements and ones mainly made of cementite.
  • FIG. 2 shows a typical example of observation.
  • (a) shows an example of observation of undissolved carbides by a scan type electron microscope
  • (b) shows an example of elemental analysis by X-rays of alloy-based undissolved carbides X
  • (c) shows an example of elemental analysis by X-rays of cementite-based undissolved carbides Y.
  • two types of structures are recognized in the steel: needle-shaped structures and spheroidal structures of the matrix.
  • the "undissolved carbides” referred to here include not only so-called alloy-based spheroidal carbides (X) where the above alloys form nitrides, carbides, and carbonitrides, but also cementite-based spheroidal carbides (Y) mainly comprising Fe carbides (cementite).
  • FIG. 2(b) and (c) show examples of analysis by EDX attached to an SEM.
  • Conventional inventions focus on only the V, Nb, and other alloy element-based carbides.
  • One example is FIG. 2(b) .
  • This is characterized in that the Fe peak is relatively small and the alloy peak (in this example V) is large in the carbides.
  • the alloy-based carbides (X) strictly speaking are mostly composite carbides with nitrides (so-called carbonitrides), so here these alloy-based carbides and nitrides and their composite alloy-based spheroidal precipitates will be collectively referred to as "alloy-based spheroidal carbides”.
  • spheroidal carbides are believed to be carbides which do not sufficiently dissolve in the quenching and tempering by oil tempering and high frequency treatment and become spheroidal and grow or shrink in the quenching and tempering process.
  • the carbides of these dimensions not only do not contribute at all to the strength and toughness due to quenching and tempering, but conversely degrade them. That is, they fix the C in the steel and consume the C added to become the source of strength and further coarsen the same so become a source of stress concentration as well, so the mechanical properties of the steel wire are reduced.
  • the area percentage of the carbides with a circle equivalent diameter of 0.2 ⁇ m or more is 7% or less.
  • the density of carbides with a circle equivalent diameter of 0.2 ⁇ m or more is 1 carbide/ ⁇ m 2 or less
  • the undissolved spheroidal carbides affect the coiling characteristics, that is, the bending characteristics up to breaking.
  • the general practice was to add large amounts of not only C, but also Cr, V, and other alloy elements, but there were the negative effects that the strength became too high, the deformation ability became insufficient, and the coiling characteristics were degraded. It is believed that the cause was the coarse carbides precipitating in the steel.
  • alloy-based and cementite-based carbides in the steel can be observed by etching a mirror polished sample by picral, electrolytic etching, or the like, but for detailed observation and evaluation of the dimensions and the like, a scan type electron microscope must be used for observation at a high magnification of 3000X or more.
  • the alloy-based spheroidal carbides and the cementite-based spheroidal carbides covered here have circle equivalent diameters of 0.2 ⁇ m or more.
  • carbides are essential for securing the strength of the steel and temper softening resistance, but if the effective particle size is 0.1 ⁇ m or less or conversely over 1 ⁇ m, rather there is no contribution to the strength or increased fineness of the austenite particle size and the deformation characteristics are just degraded.
  • the prior art only focuses on V, Nb, and other alloy-based carbides.
  • alloy-based and cementite-based carbides are observed by electrolytically etching a mirror polished sample and using a scan type electron microscope to observe it at 10000X observing in 10 fields or more. If the area percentage of the spheroidal carbides exceeds 7%, the workability is extremely inferior, so this was set as the upper limit.
  • the reason for making the prior austenite grain size number #10 or larger is that in steel wire having basically a tempered martensite structure, the prior austenite grain size has a great effect on the basic properties of the steel wire along with the carbides. That is, a smaller prior austenite grain size means superior fatigue characteristics and coilability. However, no matter how small the austenite grain size, if the carbides are contained in over the prescribed amount, the effect is small. Generally, to reduce the austenite grain size, it is effective to lower the heating temperature during quenching, but this conversely increases the undissolved spheroidal carbides. Therefore, it is important to finish the steel wire to one balanced in the amount of carbides and prior austenite grain size.
  • the prior austenite grain size number is less than #10, sufficient fatigue characteristics and coilability cannot be obtained, so the prior austenite grain size number was made #10 or larger.
  • finer grains are preferable.
  • the reason for making the retained austenite 15 mass% or less is that retained austenite often remains at the segregated parts or prior austenite grain boundaries or near regions surrounded by subgrains.
  • the retained austenite becomes martensite by work-induced transformation. If transformation is induced during spring formation, locally high hardness parts are formed in the material and, rather, the coiling characteristics as a spring are reduced. Further, recent springs are strengthened at their surfaces by shot-peening, setting, and other plastic deformation, but if the production process includes a plurality of steps of applying such plastic deformation, the work-induced martensite formed at an early stage will lower the fracture strain and lower the workability and the breakage characteristics of the spring during use.
  • the wire will easily break during coiling. Further, by gradually breaking down in nitriding, strain relief annealing, and other heat treatment, the mechanical properties are changed, the strength reduced, the coilability reduced, and other negative effects are caused. Therefore, the retained austenite is reduced as much as possible and formation of work-induced martensite is suppressed so as to improve the workability. Specifically, if the amount of retained austenite exceeds 15% (mass%), the sensitivity to strike marks etc. becomes higher and the wire easily breaks during coiling or other handling, so 15% or less was restricted to.
  • the amount of retained austenite changes depending on the amount of C, Mn, and other alloy elements added and the heat treatment conditions. Therefore, improvement of not only the design of the compositions, but the heat treatment conditions is important.
  • the martensite formation temperature (starting temperature Ms point, finishing temperature Mf point) becomes a low temperature, martensite will not be formed unless the temperature is made considerably low during quenching. Retained austenite will easily remain.
  • control becomes necessary to maintain the cooling refrigerant at a low temperature, maintain a low temperature as much as possible after cooling, secure a long transformation time to martensite, and the like.
  • the material is processed by a continuous line, so the temperature of the cooling refrigerant easily rises to near 100°C, but it is preferably maintained at 60°C or less. A low temperature of 40°C or less is more preferable. Further, to sufficiently promote martensite transformation, the material must be held in the cooling medium for at least 1 second. Securing a holding time after cooling is also important.
  • the tensile strength, hardness after annealing, impact value, and reduction in area as measured by a tensile test are shown as evaluation items.
  • the tensile strength is directly linked with the fatigue limit of the spring. The higher the strength, the higher the fatigue limit shown.
  • the reduction in area measured simultaneously with the measurement of the tensile strength shows the plastic deformation behavior of the material and is an evaluation indicator of workability into a spring (coiling characteristic).
  • the larger the reduction in area the easier workability shown, but in general the higher the strength, the smaller the reduction in area.
  • the prepared test piece is obtained by quenching and tempering a material of ⁇ 13 mm to exceed 2200 MPa, then preparing a No. 9 test piece of JIS Z 2201. This is tested based on JIS Z 2241. The tensile strength was calculated from this breaking load.
  • springs are often being made higher in strength by hardening by nitridation of the surface layers.
  • the nitridation is performed by heating the spring at 400 to 500°C in a nitriding atmospheric gas and holding it there for several minutes to 1 hour so as to harden the surface layer. At this time, the inside where the nitrogen does not penetrate is heated, so is annealed and softened. It is important to suppress this softening, so the hardness after annealing simulating nitriding was used as an item for evaluation of the softening resistance.
  • the Charpy impact value was made an evaluation item. Generally, it is believed that a material which has an excellent impact value is also good in breakage resistance including fatigue characteristics. Further, a brittle material is also inferior in workability, so a material with a high toughness is considered to be excellent in workability as well.
  • the Charpy impact value of a material heat treated in the same way as one measured for tensile strength after quenching and tempering was measured. The Charpy impact value is influenced by the austenite grain size, so the austenite grain size of the same material was also measured. Note that the Charpy impact test piece is comprised of a so-called half size (5x10 mm cross-section) material obtained from a ⁇ 13 mm heat treated material and formed with a U-notch.
  • the heat treatment is ended in a relatively short time. Because of this, it is known that undissolved carbides easily remain and the workability is decreased. Consequently, in the invention examples as well, the material was patented and drawn to ⁇ 4 mm and the drawn wire was heated treated to measure the distribution of carbides and austenite grain size. Generally, if the heating temperature is low and the time is short, the austenite grain size becomes small, but the undissolved carbides tend to increase. A balance of the two should be used for overall evaluation. The results appear in the tensile strength and the elongation, so these two were evaluated. With a fine diameter material of ⁇ 5 mm or less, since the cross-sectional area is small, in the plastic deformation behavior, a clearer difference appears in the elongation rather than the reduction in area.
  • a tensile test was conducted based on the JIS by preparing a test piece with a parallel portion of ⁇ 6 mm and measuring the tensile strength and elongation.
  • the amount of retained austenite was determined by mirror polishing after quenching and tempering and measurement by X-rays.
  • the hardness after annealing was determined by mirror polishing after heat treatment and measurement of the Vickers hardness at the depth of 1/2 from the surface of the radius at three points. The average value was used as the hardness after annealing.
  • Invention Example 16 of the present invention produced the material by a 2t vacuum melting furnace, then rolled this into a billet. At that time, in the invention examples, the high temperature of 1200°C or more was held for a certain time. Next, in each case, the billet was rolled to ⁇ 13 mm.
  • the material was melted in a 16 kg vacuum melting furnace, then forged by forging to ⁇ 13 mm x 600 mm, then heat treated. At this time as well, in the same way, the material was held at a 1200°C or more high temperature for a certain time, then heat treated to become a predetermined high strength.
  • the material was held at 1200°C ⁇ 15 min ⁇ air-cooled, then heated at 950°C for 10 minutes, then charged into a lead bath heated to 650°C, further heated at 950°C x 10 min, charged into a 60°C oil bath for quenching, then, in the invention examples, adjusted in tempering temperature so that the tensile strength exceeded 2200 MPa.
  • the tensile strength, drawability, and Charpy impact value with this heat treatment were measured.
  • This tempering temperature differs depending on the chemical compositions, but in regards to the present invention, the materials are heat treated in accordance with the chemical compositions so that the tensile strength becomes 2200 MPa or more. On the other hand, in regards to the comparative examples, the materials are heat treated just to match the tensile strength. Further, the materials were annealed at 400°C x 20 min simulating nitriding and measured for hardness so as to evaluate the softening resistance.
  • the rods were held at 1200°C x 15 min ⁇ air-cooled, then cut to ⁇ 10 mm, heated at 950°C for 10 minutes, then charged to a lead bath heated to 650°C. Further, this was drawn to reduce it in diameter to ⁇ 4 mm, heated at 950°C x 5 min, then charged into a 60°C oil bath for quenching, then adjusted in tempering temperature to give a tensile strength exceeding 2200 MPa. Further, the stress able to give a number of load cycles exceeding 10 7 in a Nakamura type rotary bending test was deemed the fatigue strength.
  • Tables 2 to 9 show the chemical compositions of the present invention and the comparative steels when treated at ⁇ 4 mm, the cementite-based carbide poor region area ratio, the area percentage of the alloy-based/cementite-based spheroidal carbides, the density of cementite-based spheroidal carbides of a circle equivalent diameter of 0.2 to 3 ⁇ m, the density of cementite-based spheroidal carbides of a circle equivalent diameter of over 3 ⁇ m, the maximum oxide diameter, the prior austenite grain size number, the amount of retained austenite (mass%), and the resultant obtained tensile strength, hardness after annealing, impact value, and reduction in area as measured in the tensile test.
  • Tables 2 and 3 show the chemical compositions of Invention Example Nos. 1 to 17, while Tables 4 and 5 show the chemical compositions of Invention Example Nos. 18 to 41.
  • Table 6 shows the chemical compositions of Comparative Example Nos. 42 to 67.
  • Table 7 shows the characteristics of Invention Example Nos. 1 to 17 and Table 8 shows them for Invention Example Nos. 18 to 41 respectively with drawing and without drawing.
  • Table 9 shows the characteristics of Comparative Example Nos. 42 to 67 with drawing and without drawing.
  • Examples 42 and 43 are cases where neither Ti nor Nb is included. A large amount of V and Cr is added, so undissolved carbides with nitrides as nuclei are formed, so the reduction in area in the tensile test or elongation after drawing is low and the workability is lowered.
  • Examples 46 to 49 Ti is added to fix N as TiN, but the amount of Ti added is excessive and there are negative effects by TiN. Because of this, the distribution of inclusions becomes greater and as a result the reduction in area in the tensile test or elongation after drawing is low and the workability is lowered.
  • Example 47 is the case where the heating temperature during quenching is reduced and thereby a large number of undissolved carbides are formed.
  • Examples 50 to 52 are examples in which Nb is added, but the added amount is excessive, so a large number of undissolved carbides are observed, the reduction in area in the tensile test or elongation after drawing is low, and the workability is lowered.
  • the Al is excessive, so the oxides become larger and the fatigue characteristics decline.
  • Examples 55 and 56 are cases where the added amount of V is excessive.
  • the hardness after annealing simulating nitridation is low, the prior austenite grain size tends to become coarse, and the fatigue characteristics decline.
  • the surface layer hardness is lower, the nitriding depth is shallower even with the same nitriding time, and other differences occur in performance after nitriding.
  • Examples 59 to 61 are cases where the cooling temperature at the time of quenching is high and the cooling time is short. The amount of retained austenite becomes great. Because of this, the hardness after annealing is insufficient, and, in terms of practical application, the areas around the slight handling marks become brittle due to stress induced transformation, so the workability declines.
  • Examples 62 and 63 are examples when the heating temperature during quenching is made too high. The prior austenite grain size becomes larger, the impact value becomes lower, and the fatigue characteristics decline.
  • Examples 64 to 67 are cases when C or Si are smaller than defined. The tensile strength after annealing decreases, so the fatigue strength cannot be secured
  • the present invention steels, because the area percentage and density of the cementite-based and alloy-based spheroidal carbides in the steel wire for cold-coiling springs, the austenite grain size, and the amount of retained austenite are made small, are increased in strength to 2000 MPa or more, are given coilability, and enable production of springs high in strength and excellent in breakage characteristics.

Landscapes

  • Chemical & Material Sciences (AREA)
  • Engineering & Computer Science (AREA)
  • Materials Engineering (AREA)
  • Mechanical Engineering (AREA)
  • Metallurgy (AREA)
  • Organic Chemistry (AREA)
  • Heat Treatment Of Steel (AREA)
  • Heat Treatment Of Strip Materials And Filament Materials (AREA)
  • Heat Treatment Of Articles (AREA)

Claims (1)

  1. Hochfester wärmebehandelter Stahldraht für eine Feder, dadurch gekennzeichnet, dass er eine Stahlzusammensetzung hat, die in Masse-% aufweist:
    0,5 bis 0,9 % C,
    1,0 bis 3,0 % Si,
    0,1 bis 1,5 % Mn,
    1,0 bis 2,5 % Cr,
    über 0,15 bis 1,0 % V und
    höchstens 0,005 % Al,
    höchstens 0,007 % N,
    0,001 bis weniger 0,01 % Nb, optional ein oder mehrere Elemente, die ausgewählt sind aus
    0,001 bis weniger 0,005 % Ti,
    0,05 bis 0,5 % W,
    0,05 bis 0,5 % Mo,
    0,05 bis 3,0 % Ni,
    0,05 bis 0,5 % Cu,
    0,05 bis 3,0 % Co,
    0,0005 bis 0,006 % B,
    0,0002 bis 0,01 % Te,
    0,0002 bis 0,01 % Sb,
    0,0001 bis 0,0005 % Mg,
    0,0001 bis 0,0005 % Zr,
    0,0002 bis 0,01 % Ca und
    0,0002 bis 0,01 % Hf sowie als Rest Eisen und unvermeidliche Verunreinigungen,
    eine Zugfestigkeit von mindestens 2000 MPa hat,
    cementitbasierte kugelige Carbide und legierungsbasierte kugelige Carbide in einem mikroskopischen Gesichtsfeld hat, das erfüllt:
    einen Flächenprozentsatz von Carbiden mit einem Kreisäquivalenzdurchmesser von 0,2 µm bis 3 µm von höchstens 7 % und
    eine Dichte von Carbiden mit einem Kreisäquivalenzdurchmesser von 0,2 µm bis 3 µm von höchstens 1 Korn/µm2,
    eine Voraustenit-Komgrößennummer von mindestens 10 hat und
    höchstens 15 Masse-% Restaustenit hat.
EP06823432.7A 2006-11-09 2006-11-09 Wärmebehandelter stahldraht für hochfeste feder Active EP2058411B1 (de)

Applications Claiming Priority (1)

Application Number Priority Date Filing Date Title
PCT/JP2006/322784 WO2008056428A1 (fr) 2006-11-09 2006-11-09 Acier pour ressorts à haute résistance et fil d'acier traité thermiquement pour ressorts à haute résistance

Publications (3)

Publication Number Publication Date
EP2058411A1 EP2058411A1 (de) 2009-05-13
EP2058411A4 EP2058411A4 (de) 2010-01-13
EP2058411B1 true EP2058411B1 (de) 2014-02-19

Family

ID=39364257

Family Applications (1)

Application Number Title Priority Date Filing Date
EP06823432.7A Active EP2058411B1 (de) 2006-11-09 2006-11-09 Wärmebehandelter stahldraht für hochfeste feder

Country Status (5)

Country Link
EP (1) EP2058411B1 (de)
KR (1) KR100968938B1 (de)
CN (1) CN101287850B (de)
BR (1) BRPI0607042B1 (de)
WO (1) WO2008056428A1 (de)

Families Citing this family (21)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
EP2453033B1 (de) * 2009-07-09 2015-09-09 Nippon Steel & Sumitomo Metal Corporation Stahldraht für eine hochfeste feder
KR101253813B1 (ko) * 2009-09-08 2013-04-12 주식회사 포스코 구상화 열처리 생략형 강재 및 그 제조방법
WO2013024876A1 (ja) * 2011-08-18 2013-02-21 新日鐵住金株式会社 ばね鋼およびばね
CN104060175B (zh) * 2014-05-27 2017-01-11 安徽红桥金属制造有限公司 一种弹簧及其生产方法
RU2578276C1 (ru) * 2014-12-22 2016-03-27 Юлия Алексеевна Щепочкина Сталь
CN104630650A (zh) * 2015-02-06 2015-05-20 铜陵百荣新型材料铸件有限公司 一种耐低温高强度弹簧钢及其制备方法
CN105316591A (zh) * 2015-03-14 2016-02-10 洛阳辰祥机械科技有限公司 一种高性能弹簧的制备方法
JP6453138B2 (ja) * 2015-03-31 2019-01-16 株式会社神戸製鋼所 曲げ加工性に優れた熱処理鋼線
CN104745953B (zh) * 2015-03-31 2017-01-11 马鞍山市兴隆铸造有限公司 一种船用侧板低碳铬合金材料及其制备方法
CN105088090A (zh) * 2015-08-28 2015-11-25 宝山钢铁股份有限公司 一种抗拉强度2000MPa级的防弹钢板及其制造方法
US10844920B2 (en) 2015-09-04 2020-11-24 Nippon Steel Corporation Spring steel wire and spring
KR101745192B1 (ko) 2015-12-04 2017-06-09 현대자동차주식회사 초고강도 스프링강
KR101745196B1 (ko) 2015-12-07 2017-06-09 현대자동차주식회사 초고강도 스프링강
KR101776490B1 (ko) 2016-04-15 2017-09-08 현대자동차주식회사 내식성이 우수한 고강도 스프링강
KR101776491B1 (ko) * 2016-04-15 2017-09-20 현대자동차주식회사 내식성이 우수한 고강도 스프링강
CN105886930B (zh) * 2016-04-26 2017-12-26 宝山钢铁股份有限公司 一种高强度耐蚀弹簧钢及其制造方法
CN114555849B (zh) * 2019-10-16 2022-11-01 日本制铁株式会社 钢线
MX2022010291A (es) * 2020-02-21 2022-10-13 Nippon Steel Corp Alambre de acero.
EP4217518A1 (de) * 2020-09-23 2023-08-02 ArcelorMittal Stahl für blattfedern von kraftfahrzeugen und verfahren zur herstellung eines flügels davon
CN113265590A (zh) * 2021-06-08 2021-08-17 江苏通工金属科技有限公司 一种高强度防锈合金弹簧钢丝及其成型工艺
CN113930680B (zh) * 2021-09-29 2023-03-17 武汉钢铁有限公司 一种耐低温高强度弹簧扁钢及其生产方法

Family Cites Families (11)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
CA2057190C (en) * 1991-02-22 1996-04-16 Tsuyoshi Abe High strength spring steel
JP3403913B2 (ja) * 1997-03-12 2003-05-06 新日本製鐵株式会社 高強度ばね用鋼
JP3577411B2 (ja) * 1997-05-12 2004-10-13 新日本製鐵株式会社 高靭性ばね鋼
JP3595901B2 (ja) * 1998-10-01 2004-12-02 鈴木金属工業株式会社 高強度ばね用鋼線およびその製造方法
JP4464524B2 (ja) * 2000-04-05 2010-05-19 新日本製鐵株式会社 耐水素疲労特性の優れたばね用鋼、およびその製造方法
KR100682150B1 (ko) * 2000-12-20 2007-02-12 가부시키가이샤 고베 세이코쇼 경인발스프링용 강선재, 경인발스프링용 신선재와 경인발스프링 및 경인발스프링의 제조방법
US7789974B2 (en) * 2000-12-20 2010-09-07 Nippon Steel Corporation High-strength spring steel wire
JP3851095B2 (ja) * 2001-02-07 2006-11-29 新日本製鐵株式会社 高強度ばね用熱処理鋼線
KR100851083B1 (ko) * 2004-11-30 2008-08-08 신닛뽄세이테쯔 카부시키카이샤 고강도 스프링용 강 및 강선
JP4476863B2 (ja) * 2005-04-11 2010-06-09 株式会社神戸製鋼所 耐食性に優れた冷間成形ばね用鋼線
JP4478072B2 (ja) * 2005-06-09 2010-06-09 新日本製鐵株式会社 高強度ばね用鋼

Also Published As

Publication number Publication date
CN101287850B (zh) 2011-04-27
KR100968938B1 (ko) 2010-07-14
EP2058411A4 (de) 2010-01-13
CN101287850A (zh) 2008-10-15
WO2008056428A1 (fr) 2008-05-15
EP2058411A1 (de) 2009-05-13
BRPI0607042A2 (pt) 2009-08-04
KR20080057205A (ko) 2008-06-24
BRPI0607042B1 (pt) 2014-08-19

Similar Documents

Publication Publication Date Title
EP2058411B1 (de) Wärmebehandelter stahldraht für hochfeste feder
EP2003223B1 (de) Gehärteter und vergüteter stahl für den einsatz als feder
EP2003222B1 (de) Ein gehärteter und vergüteter stahl für den einsatz als feder
US20100028196A1 (en) High Strength Spring Steel and High Strength Heat Treated Steel Wire for Spring
EP2465963B1 (de) Hochfester Federstahl und Stahldraht
JP4478072B2 (ja) 高強度ばね用鋼
KR100514120B1 (ko) 고강도 스프링강 및 스프링강선
EP2453033B1 (de) Stahldraht für eine hochfeste feder
JP4423254B2 (ja) コイリング性と耐水素脆化特性に優れた高強度ばね鋼線
WO2012005373A1 (ja) 高強度ばね用伸線熱処理鋼線および高強度ばね用伸線前鋼線
JP3851095B2 (ja) 高強度ばね用熱処理鋼線
JP4559959B2 (ja) 高強度ばね用鋼
KR20100077250A (ko) 고강도 스프링강 및 스프링강선
JPH06299296A (ja) 耐脱炭性に優れた高強度ばね用鋼
JP2002180196A (ja) 高強度ばね鋼
JP2003166032A (ja) 高強度ばね鋼
JP2002180197A (ja) 高強度ばね鋼

Legal Events

Date Code Title Description
PUAI Public reference made under article 153(3) epc to a published international application that has entered the european phase

Free format text: ORIGINAL CODE: 0009012

17P Request for examination filed

Effective date: 20070725

AK Designated contracting states

Kind code of ref document: A1

Designated state(s): AT BE BG CH CY CZ DE DK EE ES FI FR GB GR HU IE IS IT LI LT LU LV MC NL PL PT RO SE SI SK TR

AX Request for extension of the european patent

Extension state: AL BA HR MK RS

RBV Designated contracting states (corrected)

Designated state(s): DE FR IT SE

RBV Designated contracting states (corrected)

Designated state(s): DE FR IT SE

A4 Supplementary search report drawn up and despatched

Effective date: 20091215

17Q First examination report despatched

Effective date: 20100408

DAX Request for extension of the european patent (deleted)
RAP1 Party data changed (applicant data changed or rights of an application transferred)

Owner name: NIPPON STEEL & SUMITOMO METAL CORPORATION

GRAP Despatch of communication of intention to grant a patent

Free format text: ORIGINAL CODE: EPIDOSNIGR1

INTG Intention to grant announced

Effective date: 20130726

GRAS Grant fee paid

Free format text: ORIGINAL CODE: EPIDOSNIGR3

GRAA (expected) grant

Free format text: ORIGINAL CODE: 0009210

RIN1 Information on inventor provided before grant (corrected)

Inventor name: KISU, TAKAYUKI

Inventor name: HASHIMURA, MASAYUKI

Inventor name: HAGIWARA, HIROSHI

Inventor name: FUJITA, TAKASHI

Inventor name: OCHI, TATSURO

Inventor name: YAMAZAKI, KOUICHI

AK Designated contracting states

Kind code of ref document: B1

Designated state(s): DE FR IT SE

RBV Designated contracting states (corrected)

Designated state(s): DE FR IT SE

REG Reference to a national code

Ref country code: SE

Ref legal event code: TRGR

REG Reference to a national code

Ref country code: DE

Ref legal event code: R096

Ref document number: 602006040344

Country of ref document: DE

Effective date: 20140403

REG Reference to a national code

Ref country code: DE

Ref legal event code: R097

Ref document number: 602006040344

Country of ref document: DE

PLBE No opposition filed within time limit

Free format text: ORIGINAL CODE: 0009261

STAA Information on the status of an ep patent application or granted ep patent

Free format text: STATUS: NO OPPOSITION FILED WITHIN TIME LIMIT

26N No opposition filed

Effective date: 20141120

REG Reference to a national code

Ref country code: DE

Ref legal event code: R097

Ref document number: 602006040344

Country of ref document: DE

Effective date: 20141120

REG Reference to a national code

Ref country code: FR

Ref legal event code: PLFP

Year of fee payment: 10

REG Reference to a national code

Ref country code: FR

Ref legal event code: PLFP

Year of fee payment: 11

REG Reference to a national code

Ref country code: FR

Ref legal event code: PLFP

Year of fee payment: 12

REG Reference to a national code

Ref country code: FR

Ref legal event code: PLFP

Year of fee payment: 13

REG Reference to a national code

Ref country code: DE

Ref legal event code: R082

Ref document number: 602006040344

Country of ref document: DE

Representative=s name: VOSSIUS & PARTNER PATENTANWAELTE RECHTSANWAELT, DE

Ref country code: DE

Ref legal event code: R081

Ref document number: 602006040344

Country of ref document: DE

Owner name: NIPPON STEEL CORPORATION, JP

Free format text: FORMER OWNER: NIPPON STEEL & SUMITOMO METAL CORP., TOKYO, JP

PGFP Annual fee paid to national office [announced via postgrant information from national office to epo]

Ref country code: FR

Payment date: 20230929

Year of fee payment: 18

PGFP Annual fee paid to national office [announced via postgrant information from national office to epo]

Ref country code: SE

Payment date: 20231002

Year of fee payment: 18

Ref country code: IT

Payment date: 20231010

Year of fee payment: 18

Ref country code: DE

Payment date: 20230929

Year of fee payment: 18