DRAWN HEAT TREATED STEEL WIRE FOR HIGH STRENGTH SPRING USE AND PRE-DRAWN STEEL WIRE FOR HIGH STRENGTH SPRING USE

Abstract

Drawn heat treated steel wire for high strength spring use is provided containing, by mass %, C: 0.67% to less than 0.9%, Si: 2.0 to 3.5%, Mn: 0.5 to 1.2%, Cr: 1.3 to 2.5%, N: 0.003 to 0.007%, and Al: 0.0005% to 0.003%, having Si and Cr satisfying the following formula: 0.3%≦Si−Cr≦1.2%, and having a balance of iron and unavoidable impurities, having as impurities, P: 0.025% or less and S: 0.025% or less, furthermore having a circle equivalent diameter of undissolved spherical carbides of less than 0.2 μm, further having, as a metal structure, at least residual austenite in a volume rate of over 6% to 15%, having a prior austenite grain size number of #10 or more, and having a circle equivalent diameter of undissolved spherical carbides of less than 0.2 μm.

Description

TECHNICAL FIELD

The present invention relates to drawn heat treated steel wire for high strength spring use which can be used as a material for high strength springs produced by cold coiling and to pre-drawn steel wire.

BACKGROUND ART

The springs which are used for automobile engines, clutches, etc. are being required to offer more advanced performance and higher durability in order to deal with the trend toward lighter weights and higher performances of automobiles. For this reason, their materials, that is, drawn heat treated steel wire for high strength spring use, are also being required to offer high material strength. In general, when producing such small sized, high strength springs, the material of the drawn heat treated steel wire for high strength spring use is quenched and tempered to impart higher material strength in the drawn heat treated steel wire for high strength spring use, then is cold coiled to obtain a coil spring shape. Furthermore, stress-relief annealing or other heat treatment and nitriding are performed to obtain a finished coil spring. For this reason, drawn heat treated steel wire for high strength spring use is required to have not only high strength, but also to have a high enough workability that it will not break at the cold coiling and to suppress softening due to the annealing, nitriding, and other heat treatment performed after coiling, that is, to have temper softening resistance.

A spring is required to have fatigue characteristics, so drawn heat treated steel wire for high strength spring use is used as a material and further nitrided or shot peened to raise the hardness of the surface layer of the spring. The durability of a spring includes fatigue characteristics and a sag property. The fatigue characteristics are affected by the surface layer hardness. The sag property (property of the spring ending up plastically deforming in the load direction during use) is greatly affected by not only the surface layer hardness, but also the hardness of the base material of the spring. For this reason, in steel wire for high strength spring use, the surface layer hardness after nitriding and the temper softening resistance at the inside where nitrogen is not introduced by nitriding are important.

Furthermore, when producing a spring by cold coiling, when producing the material of the drawn heat treated steel wire for high strength spring use, oil tempering, induction hardening treatment, etc. where rapid heating and rapid cooling are possible may be used.

For this reason, the drawn heat treated steel wire for high strength spring use can be reduced in prior austenite grain size, so a spring with excellent fracture characteristics can be obtained. However, if drawn heat treated steel wire for high strength spring use becomes higher in strength, in cold coiling, breakage may occur and the spring shape may not be able to be formed.

To deal with this problem, some of the inventors proposed drawn heat treated steel wire for high strength spring use obtained by controlling the carbides, making the prior austenite finer, and achieving both strength and cold coiling ability (PLT 1). Furthermore, they proposed drawn heat treated steel wire for high strength spring use obtained by controlling the residual austenite and carbides, refining the prior austenite, and achieving both strength and cold coiling ability (PLT 2 to PLT 4). In particular, the starting points of fracture caused by the formation of coarse oxides and carbides are suppressed and the distribution of fine carbides of cementite required for securing strength is made uniform so as to suppress deterioration of the fatigue characteristics and workability of the drawn heat treated steel wire for high strength spring use.

PLT 2 focuses on the fact that the region of sparse spherical carbides with a circle equivalent diameter of 2 μm or more in the region of a sparse distribution of fine spherical carbides (in particular, cementite) affects the dynamic characteristics and defines that region.

PLT 3 and PLT 4 take note of the effect of precipitation of fine carbides due to the addition of the alloy element V and limits the nitrogen (N) content to suppress undissolved spherical carbides. That is, they utilize the effect of precipitation of carbides, nitrides, and carbonitrides of V to enable utilization for hardening the steel wire at the tempering temperature or hardening the surface layer in nitriding. Furthermore, there is also an effect on suppressing coarsening of the austenite grain size due to the formation of precipitates. The effect of addition of V is remarkable. However, undissolved carbides or nitride easily form, so even if suppressing the nitrogen (N), the control of precipitation has to be performed precisely.

Therefore, PLT 4 quantitatively compares the undissolved spherical carbides and the precipitated carbides and defines the amounts so as to obtain as much precipitated V carbides, which are effective for the final spring performance, as possible. Specifically, it proposes to weigh the residue of V carbides in the electrolytic solution at a constant potential and compare this with the amount of V which passes through the filter (amount of precipitated V).

CITATION LIST

Patent Literature

• PLT 1: Japanese Patent Publication (A) No. 2002-180198
• PLT 2: Japanese Patent Publication (A) No. 2006-183137
• PLT 3: Japanese Patent Publication (A) No. 2006-342400
• PLT 4: International Publication WO2007/114491

SUMMARY OF INVENTION

Technical Problem

In recent years, to raise the durability of high strength springs, surface hardening by nitriding has become a general practice. Furthermore, increasing the nitrided depth and shortening the nitriding time by raising the treatment temperature is being studied. For this reason, drawn heat treated steel wire for high strength spring use is being required to be further improved in temper softening resistance.

That is, a further better cold coiling ability than even with conventional drawn heat treated steel wire for high strength spring use, excellent temper softening resistance even after being held at 500° C. for 1 hour, internal softening kept to a minimum, and greater hardness of the surfacemost layer are being sought.

The above conventional drawn heat treated steel wire for high strength spring use secures a certain extent of uniform dispersion of fine carbides for improving the fatigue characteristics and workability. However, to improve the temper softening resistance, further uniform dispersion is necessary. In particular, the addition of V proposed in PLT 3 and PLT 4 does indeed have the effect of hardening the steel wire at the tempering temperature, hardening the surface layer in nitriding, and refining the austenite. However, on the other hand, control of the nitrogen (N) content is not easy. As a result, coarse carbides, nitrides, and carbonitrides are precipitated and cause degradation in the fatigue strength.

PLT 3 adds Nb and Ti with the aim of the effect of trapping excess nitrogen (N). However, even if doing this, control to a suitable amount of N content is still not easy.

PLT 4 samples the residue of undissolved spherical carbides obtained as a result and compares it with the dissolved carbides. Therefore, it does not proactively control uniform dispersion of fine carbides.

Due to the above, the present invention has as its object to keep to a minimum the addition of V and other alloy elements, that is, without precisely controlling the N content, develop drawn heat treated steel wire for high strength spring use which has excellent yield strength and hardness and excellent workability and which has superior surface layer hardness and internal hardness even after nitriding.

Further, as described in PLT 3 and PLT 4, to obtain excellent yield strength and hardness and excellent workability, the size of the undissolved spherical carbides in the steel should be small. The effective size is preferably 0.1 μm or less. If over 1 μm, the contribution to strength and workability is lost and the deformation characteristics are just degraded. For this reason, the density of presence of undissolved spherical carbides with a circle equivalent diameter of 0.2 μm or more becomes an important indicator. Therefore, the present invention has as its object the development of steel wire for high strength spring use not allowing the presence of undissolved spherical carbides with a circle equivalent diameter of 0.2 μm or more.

Solution to Problem

The inventors engaged in intensive research to solve the above problems and as a result obtained the following discoveries:

(a) It was discovered that by strictly controlling the contents of C, Si, Mn, and Cr in the steel wire to suppress the formation of spherical carbides and by utilizing the residual austenite, even without adding alloy elements such as V, the drawn heat treated steel wire for high strength spring use is improved in strength and cold coiling ability compared with the conventional.

(b) It was also discovered that by adding both Cr and Si in the steel wire in suitable amounts, the formation of undissolved spherical carbides and the softening in annealing or nitriding after coiling are suppressed, and, furthermore, greater hardness of the nitrided layer can be achieved.

That is, for increasing the strength in the fatigue characteristics, addition of Cr is effective, but Cr is an element which easily leaves behind undissolved spherical carbides which would have a detrimental effect on the cold coiling ability. For this reason, the amount of addition had to be restricted. The inventors also took note of Si which suppresses the growth of undissolved spherical carbides and the formation of cementite. They discovered that if adding Si and together increasing the amount of addition of Cr, the drawn heat treated steel wire can be increased in strength. Quantitatively, it is sufficient to add large amounts of both Si and Cr and, as the relationship between them, control the difference in amount of addition of Si and the amount of addition of Cr, that is, (Si−Cr) %.

(c) Further, it was discovered that by heating the bloom to 1250° C. or more, it is possible to make Cr and other alloy elements in the steel material uniformly disperse and suppress the formation of coarse undissolved spherical carbides and, furthermore, make fine carbides uniformly disperse.

Undissolved spherical carbides are present in the steel material just after casting and become causes of not only poor coiling ability, but also breakage in rolling and drawing. For this reason, to prevent a detrimental effect in the steps of blooming after casting, wire rod rollingwire rod, patenting, quenching, and drawing, it is effective to raise the heating temperature in each step and constantly suppress undissolved spherical carbides.

(d) Furthermore, it was discovered that the addition of V has a detrimental effect on the mechanical properties and fatigue strength of steel wire for spring use.

That is, from just after casting to being worked into a spring, a steel material is repeatedly heated. Usually, the undissolved spherical carbides are mainly cementite (Fe₃C). However, by repeating the heating, the undissolved spherical carbides often include Cr, V, etc. It is learned that not only are Cr, V, and other alloy elements wastefully consumed, but there is also a possibility of degrading the mechanical characteristics after nitriding (surface hardness, internal hardness, etc.)

Further, as explained above, with the addition of V, control of the nitrogen (N) content is not easy. As a result, coarse carbides, nitrides, and carbonitrides precipitate and become causes of degradation in fatigue strength.

From these facts, the inventors discovered that by not adding V, or an extremely small amount even though added, and further, as explained above, controlling the amount of Cr in balance with the amount of Si, it is possible to suppress coarsening of the undissolved spherical carbides.

Here, “undissolved spherical carbides” means undissolved carbides with a ratio of the maximum size (long size) and minimum size (short size) (aspect ratio) of 2 or less. Actually, “carbides” and “spherical carbides” are also undissolved. Here, in the sense of emphasis, while synonymous, these respectively are also called “undissolved carbides” and “undissolved spherical carbides”.

The present invention was made based on these discoveries. The gist of the invention is as follows:

(1) Pre-drawn steel wire for high strength spring use characterized by containing, by mass %,

C: 0.67% or greater and less than 0.9%,

Si: 2.0 to 3.5%,

Mn: 0.5 to 1.2%,

Cr: 1.3 to 2.5%,

N: 0.003 to 0.007%, and

Al: 0.0005% to 0.003%,

having Si and Cr satisfying the following formula:

0.3%≦Si−Cr≦1.2%,

having a balance of iron and unavoidable impurities,
having P and S as impurities comprising
P: 0.025% or less and
S: 0.025% or less, and, furthermore,
having a circle equivalent diameter of undissolved spherical carbides of less than 0.2 μm.

(2) Pre-drawn steel wire for high strength spring use as set forth in (1) characterized by, further, containing, by mass %, one or more of

V: 0.03 to 0.10%,

Nb: 0.015% or less

Mo: 0.05 to 0.30%,

W: 0.05 to 0.30%

Mg: 0.002% or less,
Ca: 0.002% or less, and
Zr: 0.003% or less,
when containing V
satisfying 1.4%≦Cr+V≦2.6% and 0.70%≦Mn+V≦1.3%, and, when containing Mo and W,
satisfying 0.05%≦Mo+W≦0.5%.
(3) Drawn heat treated steel wire for high strength spring use characterized by containing, by mass %,
C: 0.67% or greater and less than 0.9%,

Si: 2.0 to 3.5%,

Mn: 0.5 to 1.2%,

Cr: 1.3 to 2.5%,

N: 0.003 to 0.007%, and

Al: 0.0005% to 0.003%,

having Si and Cr satisfying the following formula:

0.3%≦Si−Cr≦1.2%, and

having a balance of iron and unavoidable impurities, having P and S as impurities comprising
P: 0.025% or less and
S: 0.025% or less,
furthermore,
having a metal structure comprised of at least residual austenite in a volume rate of over 6% to 15%,
having prior austenite grain size number of #10 or more, and
having a circle equivalent diameter of undissolved spherical carbides of less than 0.2 μm.
(4) Drawn heat treated steel wire for high strength spring use as set forth in (3) characterized by, further, containing, by mass %, one or more of

V: 0.03 to 0.10%,

Nb: 0.015% or less

Mo: 0.05 to 0.30%,

W: 0.05 to 0.30%

Mg: 0.002% or less,
Ca: 0.002% or less, and
Zr: 0.003% or less,
when containing V
satisfying 1.4%≦Cr+V≦2.6% and 0.70%≦Mn+V≦1.3%, and,
when containing Mo and W,
satisfying 0.05%≦Mo+W≦0.5%.
(5) Drawn heat treated steel wire for high strength spring use as set forth in (3) or (4) characterized in that said drawn heat treated steel wire for high strength spring use has a tensile strength of 2100 to 2400 MPa.
(6) Drawn heat treated steel wire for high strength spring use as set forth in any one of (3) to (5) characterized in that said drawn heat treated steel wire for high strength spring use has a yield stress of 1600 to 1980 MPa.
(7) Drawn heat treated steel wire for high strength spring use as set forth in any one of (3) to (6) characterized said drawn heat treated steel wire for high strength spring use has a a surface Vicker's hardness of HV750 or more and an internal Vicker's hardness of HV570 or more aftersoft nitriding of keeping at 500° C. for 1 hour.
(8) A method of production of pre-drawn steel wire for high strength spring use characterized by taking a bloom containing, by mass %,
C: 0.67% or greater and less than 0.9%,

Si: 2.0 to 3.5%,

Mn: 0.5 to 1.2%,

Cr: 1.3 to 2.5%,

N: 0.003 to 0.007%, and

Al: 0.0005% to 0.003%,

having Si and Cr satisfying the following formula:

0.3%≦Si−Cr≦1.2%,

having a balance of iron and unavoidable impurities,
having P and S as impurities comprising
P: 0.025% or less and
S: 0.025% or less, heating the bloom to 1250° C. or more, then hot rolling the bloom to produce a billet and heating the billet to 1200° C. or more, then hot rolling to produce pre-drawn steel wire.
(9) A method of production of pre-drawn steel wire for high strength spring use as set forth in (8) characterized by the bloom further, containing, by mass %, one or more of

V: 0.03 to 0.10%,

Nb: 0.015% or less

Mo: 0.05 to 0.30%,

W: 0.05 to 0.30%

Mg: 0.002% or less,
Ca: 0.002% or less, and
Zr: 0.003% or less,
when containing V
satisfying 1.4%≦Cr+V≦2.6% and 0.70%≦Mn+V≦1.3%, and,
when containing Mo and W,
satisfying 0.05%≦Mo+W≦0.5%.
(10) A method of production of pre-drawn steel wire for high strength spring use characterized by further heating pre-drawn steel wire as set forth in (8) or (9) to 900° C. or more, then patenting it at 600° C. or less.
(11) A method of production of heat treated steel wire for high strength spring use characterized by drawing said pre-drawn steel wire which was produced by the method of production of pre-drawn steel wire as set forth in any one of (8) to (10), heating by a heating rate of 10° C./sec or more up to an A₃ point, holding at a temperature of the A₃ point or more for 1 minute to 5 minutes, then cooling by a cooling rate of 50° C./sec or more down to 100° C. or less.
(12) A method of production of heat treated steel wire for high strength spring use as set forth in (11) characterized by further holding and tempering it at 400 to 500° C. for 15 minutes or less.

Advantageous Effects of Invention

According to the present invention, in particular, due to the excellent cold coiling ability and temper softening resistance, even with nitriding at 500° C. for 1 hour, drawn heat treated steel wire for high strength spring use with a high surface layer hardness and internal hardness and, furthermore, high strength spring excellent in durability can be provided. The contribution in industry is extremely great.

BRIEF DESCRIPTION OF DRAWINGS

FIG. 1 is a micrograph of the metal structure showing one example of spherical carbides in the drawn heat treated steel wire for high strength spring use of the present invention. At the tips of the arrows in the figure, undissolved spherical carbides are observed.

FIG. 2 is a view showing the shape of a punch for providing a notch in a test piece.

FIG. 3 is a view showing the step of providing a notch in a test piece.

FIG. 4 is a view showing an outline of a notch bending test.

FIG. 5 is a view showing a method of measurement of a notch bending angle.

DESCRIPTION OF EMBODIMENTS

In general, a wire rod for a spring is produced as follows: Of course, production of springs is not limited to this here described process. This is described as just one example.

A bloom made of steel containing predetermined chemical compositions is rolled to obtain a billet. Next, the billet is rolled to produce a predetermined diameter of steel wire. The steel wire which is produced at this stage is called the “pre-drawn steel wire”.

The steel wire which is produced after rolling is patented and drawn to obtain further finer steel wire, then the working stress at the surface layer is removed and subsequent cold coiling workability is obtained by heat treatment (quenching and tempering). The steel wire which is produced at this stage is called the “drawn heat treated steel wire”.

Next, the spring is worked by cold coiling and is improved in strength and surface hardness by nitriding. In this way, a “spring” is produced as a final product.

First, the chemical compositions of the drawn heat treated steel wire for high strength spring use of the present invention and its material, that is, pre-drawn steel wire for high strength spring use, will be explained. Here, the “%” in the chemical compositions means mass % except when otherwise indicated.

C: 0.67% to less than 0.9%

C is an important element which has a great effect on the strength of the steel material and contributes to the formation of residual austenite as well. In the present invention, to obtain sufficient strength, the lower limit of the amount of C is made 0.67% or more. To raise the strength, the amount of C is preferably made 0.70% or more, more preferably 0.75% or more.

On the other hand, if the amount of C becomes 0.9% or more, excessive coprecipitation results, a large amount of coarse cementite is precipitated, and the toughness remarkably falls. Further, if the amount of C is excessive, coarse spherical carbides are formed and the coiling ability is impaired. Therefore, the upper limit of the amount of C is made less than 0.9%. From the viewpoint of suppressing the formation of spherical carbides, the upper limit of the amount of C is preferably 0.85%, more preferably 0.80%.

Si: 2.0 to 3.5%

Si is an important element for improving the temper softening resistance of the steel and the sag property of the spring. To obtain these effects, 2.0% or more has to be added. Further, Si is effective for spheroidization and refinement of the cementite. To suppress the formation of coarse spherical carbides, 2.1% or more of Si is preferably added. To raise the internal hardness after nitriding and other treatment for making the surface layer harder, 2.2% or more of Si is more preferably added. Furthermore, from the balance with Cr, Si is more preferably made 2.3% or more. Si is sometimes made 3.0% or more.

On the other hand, if excessively adding Si, the steel wire hardens and becomes brittle, so the upper limit of the amount of Si is made 3.5% or less. From the viewpoint of the prevention of embrittlement, the upper limit is preferably made 3.4%, more preferably 3.3% or less.

Mn: 0.5 to 1.2%

Mn is an element which is important for raising the quenchability and stably securing the amount of residual austenite. In the present invention, to raise the yield strength of the steel wire and secure the residual austenite, Mn has to be added in 0.5% or more, more preferably 0.65% or more, still more preferably 0.70% or more.

On the other hand, if excessively adding Mn, the residual austenite increases. In working, work-induced martensite is formed and the cold coiling ability is impaired. To prevent embrittlement due to excessive addition of Mn, the upper limit of the amount of Mn is made 1.2% or less, preferably 1.1% or less, more preferably 1.0% or less.

Cr: 1.3 to 2.5%

Cr is an element which is effective for improving the quenchability and temper softening resistance. To obtain these effects, 1.3% or more of Cr has to be added. When performing the nitriding, it is possible to make the hardened layer obtained by nitriding deeper by the addition of Cr. Therefore, when imparting hardening by nitriding and softening resistance at the nitriding temperature, over 1.5% of Cr is preferably added. More preferably, 1.7% or more should be added.

On the other hand, if the amount of Cr is excessive, the production cost becomes higher. Not only this, dissolution of the carbides is impaired, undissolved spherical carbides become greater, and the coiling ability is impaired, so the upper limit of the amount of Cr is made 2.5% or less. Further, if the amount of Cr is large, to suppress formation of coarse cementites, the amount of Cr is preferably suppressed to 2% or less. Furthermore, to obtain both strength and coiling ability, the upper limit of the amount of Cr is preferably made 1.8% or less.

N: 0.003 to 0.007%

N is an element, in the present invention, which forms nitrides with Al etc. included as impurities in the steel. To utilize the fine nitrides and refine the prior austenite, 0.003% or more of N has to be included. On the other hand, if the amount of N is excessive, the nitrides coarsen and the cold coiling ability and fatigue characteristics fall. Therefore, the upper limit of the amount of N is made 0.007% or less. Further, if considering the ease of heat treatment etc., the amount of N is preferably 0.005% or less.

P: 0.025% or less

P is an impurity. It causes the steel to harden, forms segregation, and causes embrittlement, so the upper limit of the amount of P is made 0.025% or less. Further, the P which segregates at the prior austenite grain boundaries causes the toughness and delayed fracture resistance etc. to fall, so the upper limit of the amount of P is preferably made 0.015% or less. Furthermore, when the yield strength of the steel wire will exceed 2150 MPa, the amount of P is preferably limited to less than 0.010%.

S: 0.025% or less

S is also an impurity. If present in steel, it causes the steel to embrittle, so the upper limit of the amount of S is made 0.025% or less. To suppress the effect of S, addition of Mn is effective. However, MnS is an inclusion. In particular in high strength steel, MnS sometimes becomes starting points of fracture. Therefore, to suppress the occurrence of fracture, the upper limit of the amount of S is preferably made 0.015% or less. Furthermore, when the yield strength of the drawn heat treated steel wire for high strength spring use will exceed 2150 MPa, the amount of S is preferably limited to less than 0.01%.

Al: 0.0005 to 0.003%

Al is a deoxidizing element. It affects the formation of oxides. If forming hard oxides, the fatigue durability falls. In particular, in high strength springs, if excessively adding Al, the fatigue strength fluctuates and the stability is impaired. In the drawn heat treated steel wire for high strength spring use of the present invention, if the amount of Al exceeds 0.003%, the rate of occurrence of fracture due to inclusions becomes greater, so the amount of Al is limited to 0.003% or less. The upper limit value of the amount of Al is preferably 0.0028%, more preferably 0.0025%.

On the other hand, if the amount of Al becomes less than 0.0005%, silica-based hard oxides are easily formed. For this reason, the amount of Al is made 0.0005% or more. The lower limit of the amount of Al is preferably 0.0007%, more preferably 0.0008%, further preferably 0.001% or more.

Next, the point of the present invention, that is, the relationship between Si and Cr, will be explained. It is already known that Si and Cr are both important for increasing the strength of spring steel. However, excessive addition causes problems.

0.3%≦Si−Cr≦1.2%

If the amount of Si exceeds the prescribed amount, the embrittlement becomes extreme and the workability in coiling is impaired. Not only that, decarburization in the intermediate processes becomes remarkable. For this reason, in the final product of the spring, the surface layer hardness becomes lower and the durability falls. Further, decarburized parts are randomly formed, so the stability of the strength of the spring product is impaired. When the amount of Si is smaller than the prescribed amount, the strength falls. Furthermore, the sag property is insufficient. This appears in the hardness after nitriding as well. Sufficient hardness cannot be secured both at the surface layer and inside.

However, the relationship between Si and Cr through the cementite in the steel is important. That is, Si is an element which destabilizes cementite. If adding a large amount of Cr or other element which stabilizes cementite, in heating, there is the effect of promoting the formation of a solid solution by the cementite. Therefore, regardless of adding a large amount of Cr, if the amount of addition of Si is small, the amount of undissolved spherical carbides becomes greater and the workability is remarkably reduced. The inventors discovered that it is possible to use the difference between the Si content (mass %) and Cr content (mass %) in the steel, that is, the Si−Cr amount, as a yardstick. That is, when the value of Si−Cr is smaller than 0.3%, the amount of Cr becomes relatively large and undissolved spherical carbides easily remain. On the other hand, if over 1.2%, Si becomes relatively excessive and easily causes embrittlement, decarburization, or other problems. Therefore, the value of Si−Cr should be made 0.3 to 1.2%.

From the viewpoint of suppressing the formation of carbides, a larger amount of Si−Cr enables the undissolved carbides to be suppressed, but industrially, if the Si is too great, the depth of the hardened layer formed by the nitriding will easily become shallow. For this reason, if considering the behavior of undissolved spherical carbides and the hardened layer formed by nitriding, preferably Si−Cr≦0.9%, more preferably Si—Cr≦0.75%. Further, from the viewpoint of relatively reducing the amount of Cr and reducing the residual presence of undissolved spherical carbides, the lower limit is preferably 0.35≦Si−Cr, more preferably 0.4≦Si−Cr.

Next, the selectively added chemical compositions will be explained.

V: 0.03 to 0.10%

V is an element which forms nitrides, carbides, and carbonitrides. Fine V nitrides, carbides, and carbonitrides with a circle equivalent diameter of less than 0.2 μm are effective for refinement of the prior austenite. Further, they may also be utilized for hardening the surface layer by nitriding. However, on the other hand, undissolved carbides and nitrides are easily formed, so even if suppressing the nitrogen (N), it is necessary to precisely control the precipitation.

For this reason, in the present invention, V is not deliberately added.

To obtain such an effect of addition of V, a fine amount can be added. To obtain these effects, V should be added in 0.03% or more, preferably 0.035% or more, more preferably 0.04% or more.

On the other hand, if adding over 0.10% of V, coarse spherical carbides are formed and the cold coiling ability and spring fatigue characteristics are impaired. Therefore, the V content should be made 0.1% or less. Further, by the addition of V, before drawing, a supercooled structure causing cracks and breakage in drawing easily is formed. For this reason, the upper limit of the amount of V is preferably made 0.09% or less, more preferably 0.08% or less, most preferably 0.05% or less. In particular, in the case of adding a fine amount of Nb, the amount of addition of V is preferably made 0.05% or less. Further, V is an element which greatly affects the formation of residual austenite in the same way as Mn, so the amount of V has to be precisely controlled together with the amount of Mn.

Nb: 0.015% or less

Nb is an element which forms nitrides, carbides, and carbonitrides in steel. These precipitates are sometimes used for control of the austenite grain size etc. However, simultaneously, excessive addition reduces the ductility when hot and results in easier cracking in rolling or hot forging. For this reason, excessive addition must be avoided.

Nb is added for the purpose of controlling the amount of N. The precipitates are not directly used for controlling the quality. Valve springs and other springs are produced by quenching, tempering, then cold coiling, but at that time, the dissolved nitrogen obstructs color deformation and reduces the limit strain. For this reason, the coiling ability is impaired. Therefore, by adding Nb and forming nitrides at a high temperature, there is the effect that the dissolved nitrogen in the steel in the steel matrix is lowered and the cold workability is improved.

Further, the addition of a fine amount of Nb is also effective for suppressing V and other undissolved spherical carbides mixed in as unavoidable impurities. V is an element which is effective for improving the temper softening resistance in nitriding and the surfacemost layer hardness. However, if the amount of addition becomes greater, even in the patenting, quenching, and other heating for obtaining an austenite phase for producing drawn heat treated steel wire for high strength spring use, V nitrides, V carbides, and V carbonitrides often are not sufficiently dissolved. The undissolved spherical carbides of V grow from cores of the V-based nitrides formed at the time of normal high temperature. As a result, undissolved spherical carbides remain and the coiling ability is impaired. For this reason, when suppressing the undissolved spherical carbides, it is necessary to suppress the amount of addition of V. In the present invention, V was not made an essential element.

As opposed to this, Nb forms nitrides at a higher temperature than V. For this reason, in the steelmaking process, addition of Nb suppresses the formation of V nitrides. That is, Nb forms nitrides in the high temperature region where V dissolves and does not form nitride. Furthermore, at the high temperature where V nitrides are formed, Nb consumes nitrogen, so V nitrides become harder to form even when cooled. For this reason, the addition of a fine amount of Nb is particularly effective for suppressing undissolved spherical carbides and securing coiling ability when adding a large amount of V.

If the amount of addition of Nb is over 0.015%, the hot ductility is impaired and the occurrence of defects and other problems in rolling becomes easy. For this reason, the amount of addition is made 0.015% or less, preferably 0.010% or less, more preferably 0.005% or less, most preferably less than 0.001%.

On the other hand, the effect of Nb in controlling the amount of N in spring steel appears starting from 0.0005%, so when adding Nb, 0.0005% or more is preferably added. Further, when adding V etc., addition of a fine amount of Nb is more effective. A range of 0.003 to 0.012% is preferable. Furthermore, a range of 0.005 to 0.009% is more preferable. The effect is obtained even at 0.005 to 0.001%.

1.4%≦Cr+V≦2.6%

In the present invention, V is not deliberately added. However, as explained above, addition of a fine amount of V has an effect on the refinement of the prior austenite and formation of residual austenite. By precisely controlling the sum of the amounts of addition of Cr and V with respect to V, it is possible to raise the strength to make the surface layer hardness after nitriding and the internal hardness suitable for high strength springs.

Cr and V are both elements which prevent softening upon the heating by the annealing or nitriding etc. performed after the spring coiling, that is, impart so-called temper softening resistance. In particular, nitriding causes nitrides to precipitate at the nitrided part of the surface layer to thereby improve the surface hardness and increase the nitriding effect. Further, even at the inside where nitriding does not spread, decomposition of the carbides is suppressed. Further, there is the effect of suppressing softening by precipitation of carbides. On the other hand, both are elements which facilitate the formation of undissolved spherical carbides. Cr dissolves in the cementite to increase the stability so in the heating steps for dissolving the cementite (heating at time of patenting and heating at time of quenching), suppresses the dissolution of the cementite, often remains as undissolved spherical carbides. Further, V also has a dissolution temperature of the precipitates higher than the A3 point of steel, so easily remains as undissolved spherical carbides.

If the total of the contents of Cr and V, that is, Cr+V, is less than 1.4%, the surface layer hardness of the high strength spring falls below HV750 and the internal hardness falls below HV570. For this reason, Cr+V is preferably 1.4% or more. Furthermore, 1.5% or more is preferable. On the other hand, excessive addition of Cr+V of over 2.6% leaves behind large amounts of undissolved spherical carbides, so the coiling ability is impaired. Therefore, 2.6% is made the upper limit. Further, the Cr+V is preferably 2% or less, more preferably 1.8% or less.

0.7%≦Mn+V≦1.3%

Mn and V are elements which improve the quenchability and also have a large effect on the formation of residual austenite. If Mn is larger than the prescribed amount, a large amount of residual austenite remains. Therefore, the sum of both Mn and the V which is included as an unavoidable impurity has a direct effect on the austenite behavior. If these exceed their prescribed amounts, the amount of residual austenite increases. Not only is the workability affected, but also the yield strength is greatly affected. Sufficiently durability cannot be secured.

For this reason, in the present invention, the total of the contents of Mn and V, that is, Mn+V, is made 0.7 to 1.3%. To secure a volume rate of over 6% of residual austenite, the lower limit of Mn+V has to be made 0.7% or more.

As a result, transformation induced plasticity causes the ductility to be improved and enables the cold coiling ability to be secured. On the other hand, to make the residual austenite a volume rate of 15% or less, the upper limit value of Mn+V has to be made 1.3% or less. Due to this, the formation of work-induced martensite due to strike marks in cold coiling is suppressed and local embrittlement can be prevented.

Mo: 0.05 to 0.30%

Mo is an element which improves the quenchability. Further, it is also extremely effective for improving the temper softening resistance. In the present invention, in particular, to further improve the temper softening resistance, 0.05% or more of Mo can be added. Further, Mo is an element which forms Mo-based carbides in the steel. The temperature at which the Mo-based carbides precipitate is lower than carbides of V etc. For this reason, addition of a suitable amount of No is also effective for suppressing coarsening of carbides. Addition of 0.10% or more of Mo is preferable. On the other hand, if the amount of addition of Mo is over 0.30%, a supercooled structure easily forms in hot rolling, the patenting before drawing, etc. Therefore, to suppress the formation of a supercooled structure causing cracking or wire breakage in drawing, the upper limit of the amount of Mo is made 0.30% or less, preferably 0.25% or less. Further, if the amount of Mo is large, in the patenting, the time until the end of the pearlite transformation becomes longer, so the amount of Mo is preferably made 0.20% or less. Furthermore, to shorten the patenting time and stably end the pearlite transformation, 0.15% or less is preferable.

W: 0.05 to 0.30%

W, like Mo, is an element which is effective for improvement of the quenchability and temper softening resistance and is an element which precipitates in the steel as carbides. In the present invention, in particular, to improve the temper softening resistance, 0.05% or more of W is added.

On the other hand, if excessively adding W, a supercooled structure is formed which causes cracking or wire breakage in drawing, so the amount of W has to be made 0.30% or less.

Furthermore, if considering the ease of heat treatment etc., the amount of W is preferably 0.1 to 0.2%, more preferably 0.13 to 0.18%.

0.05%≦Mo+W≦0.5%

Mo and W are elements which are effective for improvement of the temper softening resistance. If adding the two combined, the growth of carbides is suppressed and the temper softening resistance can be remarkably improved compared with addition of Mo and W alone. In particular, to improve the temper softening resistance in heating to 500° C., Mo+W has to be made 0.05% or more, preferably 0.15% or more.

On the other hand, if Mo+W is over 0.5%, in hot rolling, patenting before drawing, etc., a so-called supercooled structure of martensite, bainite, etc. is formed. To suppress the formation of a supercooled structure causing cracks or wire breakage in drawing, the upper limit of Mo+W is made 0.5% or less, preferably 0.35% or less.

Next, Mg, Ca, and Zr will be explained.

Mg: 0.002% or less

Mg forms oxides in molten steel higher in temperature than the MnS forming temperature. At the time of formation of MnS, it is already present in the molten steel. Therefore, it can be used as a nuclei for precipitation of MnS. Due to this, the distribution of the MnS can be controlled. Further, in number distribution as well, Mg-based oxides are more finely dispersed in the molten steel compared with the Si- and Al-based oxides which are often seen in conventional steel, so MnS formed around cores of Mg-based oxides are finely dispersed in the steel. Therefore, even with the same S content, depending on the presence of Mg, the MnS distribution differs. Adding these makes the MnS grain size finer. By making the MnS finely disperse, it is possible to render the action as a starting point of fatigue of MnS harmless. The effect is sufficiently obtained even in fine amounts. Preferably Mg 0.0002% or more, more preferably 0.0005% or more, should be added.

However, with addition of over 0.001%, it is difficult for the Mg to remain in the molten steel, there is an effect on the oxide composition, and the rate of appearance of oxides as initiation sites of fatigue becomes higher, so 0.002% is made the upper limit. Therefore, the upper limit of the amount of addition of Mg was made 0.002%, preferably 0.0015% or less. Furthermore, in the case of spring steel, compared with other steel for structural use, the amount of addition of S is suppressed, so if considering the yield etc., 0.001% or less is preferable. Further, when used for a high strength valve spring, the inclusion susceptibility is high, so Mg has the effect of improving the corrosion resistance and resistance to delayed fracture preventing rolling cracks due to the effect of the distribution of MnS etc. Addition of as much as possible is preferable, so control of the amount of addition in the extremely narrow range of 0.0002 to 0.001% is preferable.

Ca: 0.002% or less

Ca is an oxide- and sulfide-forming element. In spring steel, it makes the MnS spherical to thereby suppress the length of MnS serving as initiation sites of fatigue and other fracture and render it harmless. The effect is similar to Mg. Addition of 0.0002% or more is preferable. Further, even if over 0.002% is added, not only is the yield poor, but also oxides and CaS and other sulfides are formed and trouble in production and degradation in spring fatigue durability characteristics are caused, so the amount was made 0.002% or less. Regarding the amount of addition, when used for a high strength valve spring, the inclusion susceptibility is high, so the amount is preferably 0.0015% or less, more preferably 0.001% or less.

Zr: 0.003% or less

Zr is an oxide-, sulfide-, and nitride-forming element. In spring steel, the oxides are finely dispersed, so like with Mg, form nuclei for precipitation of MnS and can make the MnS finely disperse. Due to this, it is possible to improve the fatigue durability and, further, increase the ductility to thereby improve the coiling ability. 0.0002% or more is preferably added. Further, even if over 0.003% is added, not only is the yield poor, but oxides and ZrN, ZrS, and other nitrides and sulfides are formed and trouble in production or degradation in the spring fatigue durability characteristics is caused, so the amount is made 0.003% or less. The amount of addition is preferably 0.0025% or less. Furthermore, when used for high strength valve spring, there is also the effect that the coiling ability is improved by the control of sulfides, so addition is preferred, but to minimize the effects on the dimensions of inclusions, suppression to 0.0015% or less is preferable.

Note that, the above optionally added chemical compositions, if contained in fine amounts, do not impair the effects of the steel wire comprised of the basic chemical compositions of the present invention.

Next, the metal structure of the steel wire for high strength spring use of the present invention will be explained.

Undissolved Spherical Carbides

Undissolved spherical carbides perform the important role of securing strength in steel wire for high strength spring use. On the other hand, the presence of undissolved spherical carbides causes the coiling ability to deteriorate. Further, coarse carbides cause the fatigue characteristics to degrade as well. Therefore, suppressing undissolved spherical carbides in coiling and causing uniform dispersion of fine carbides after the final nitriding are essential for solving the problem of the present invention.

The steel wire for high strength spring use of the present invention has a long size of the undissolved spherical carbides of 0.2 μm or less, that is, is suppressed in coarsening. The undissolved spherical carbides are already present after wire rod rolling (that is, the pre-drawn steel wire).

The undissolved spherical carbides are hard to go into solid-solution in the subsequent heat treatment (patenting, generation of working heat in drawing, and quenching and tempering, for instance). Rather, they sometimes grow in these heat treatment steps and coarsen. That is, the undissolved spherical carbides in the pre-drawn steel wire sometimes act as nuclei for coarsening of themselves.

For this reason, to restrict the coarsened undissolved spherical carbides of the steel wire after heat treatment (heat treated steel wire), it is important to reduce as much as possible the undissolved spherical carbides which are present in the pre-drawn steel wire. Due to the above, the definition regarding the “undissolved spherical carbides” has important meaning in not only the pre-drawn steel wire for high strength spring use according to the present invention, but also the drawn heat treated steel wire for high strength spring use.

The steel wire for high strength spring use of the present invention is increased in strength by having C added, having Mn and Cr added and, further having Mo, W, and other so-called alloy elements added. When adding large amounts of C and, in particular, Cr and other alloy elements which form nitrides, carbides, and carbonitrides, spherical cementite carbides and alloy-based carbides easily remain in the steel. Spherical cementite carbides and alloy-based carbides are undissolved spherical carbides which do not dissolve in the steel in heating in the hot rolling.

Note that, in the present invention, spherical alloy-based carbides and spherical cementite carbides will be referred to all together as spherical carbides. In the steel, there are pin-shaped carbides corresponding to the pin-shaped structure of tempered martensite, but these pin-shaped carbides are not included in the spherical carbides of the present invention. The pin-shaped carbides are not present right after quenching and precipitate in the process of tempering. The tempered martensite structure is a structure suitable for achieving both strength and toughness and workability. Being pin-shaped is, in a certain sense, the ideal form in carbides.

Strictly speaking, if carbides with an aspect ratio of 2 or more (pin-shaped carbides) also coarsen, the workability may be impaired. However, in actuality, pin-shaped carbides become coarse when the tempering temperature is high or the holding time in tempering is extremely long. The effect on performance is to make the strength and hardness insufficient. Problems arise in different areas than with undissolved spherical carbides. In the 2100 MPa or so strength steel wire covered by the present invention, coarse pin-shaped carbides are not formed. Therefore, in the present invention, pin-shaped carbides are not covered. As explained above, the normally precipitated carbides are undissolved, but in the present invention, the term “undissolved” added to the top. This just stresses the undissolved nature. In the present invention, “undissolved spherical carbides” and “spherical carbides” are synonymous.

The undissolved spherical carbides can be observed under a scanning electron microscope (SEM) by polishing a sample obtained from pre-drawn steel wire or drawn heat treated steel wire for high strength spring use to a mirror finish and etching it by picral or electrolytically etching it. Further, they can be observed by the replica method under a transmission type electron microscope (TEM).

FIG. 1 shows an example of a structural photograph of a sample after electrolytic etching as observed under an SEM. In the structural photograph of FIG. 1, the steel is observed to have two types of structures of the matrix, that is, pin-shaped structures and spherical structures. Among these, the pin-shaped structures are tempered martensite formed by quenching and tempering. On the other hand, the spherical structures are carbides 1 made spherical by not dissolving into the steel due to the heating of the hot rolling and by being made spherical by quenching and tempering by oil tempering or induction hardening treatment (undissolved spherical carbides). Spherical carbides can be observed at the front end of the arrow in FIG. 1.

Circle Equivalent Diameter of Undissolved Spherical Carbides of Less Than 0.2 μm

In the present invention, the undissolved spherical carbides affect the characteristics of the drawn heat treated steel wire for high strength spring use, so are controlled in size as follows: Note that, in the present invention, compared with the prior art, further finer spherical carbides are defined for achieving both higher performance and workability. Spherical carbides with a circle equivalent diameter of less than 0.2 μm are extremely effective for securing the strength and temper softening resistance of the steel.

On the other hand, spherical carbides with a circle equivalent diameter of 0.2 μm or more do not contribute to improvement of the strength or temper softening resistance and degrade the cold coiling ability. For this reason, the present invention is characterized by not allowing the formation of spherical carbides with a circle equivalent diameter of 0.2 μm or more.

The pre-drawn steel wire and drawn heat treated steel wire of the present invention is characterized in that the undissolved spherical carbides have a circle equivalent diameter of less than 0.2 μm. For this reason, it is possible to secure strength while securing workability as well.

As explained above, the pre-drawn steel wire has to be then patented, drawn and heated, quenched and tempered, or otherwise heat treated, so the undissolved spherical carbides may grow and coarsen. For this reason, the circle equivalent diameter of the undissolved spherical carbides in the pre-drawn steel wire is preferably made smaller than 0.2 μm.

From the results of experiments of the inventors, the circle equivalent diameter of undissolved carbides of the pre-drawn steel wire is confirmed to be able to be reduced to 0.18 μm or less. Further, it is also confirmed that if making the billet heating temperature 1250° C. or more, the diameter can be made 0.15 μm or less.

Here, the methods of measuring the circle equivalent diameter and density of presence of spherical carbides will be explained. A sample which is taken from steel wire for high strength spring use is polished and electrolytically etched. Note that, the observed location is randomly selected near the center of the radius of the heat treated wire rod (steel wire), that is, the so-called “½R part”, so as to enable elimination of special conditions such as decarburization and center segregation. Further, the measurement area is 300 μm² or more. In electrolytic etching, the surface of the sample is corroded by electrolytic action in an electrolytic solution (a mixture of acetyl acetone 10 mass %, tetramethyl ammonium chloride 1 mass %, and a balance of methyl alcohol) using the sample as the anode and platinum as the cathode using a current generator with a lower potential. The potential becomes constant at a potential suitable for the sample in the range of −50 to −200 mV vs SCE. For the steel wire of the present invention, it is preferable that it become constant at −100 mV vs SCE.

The amount of power run can be found by the total surface area of the sample×0.133 [c/cm²]. Note that, when embedding the sample in a resin, not only the polished surface, but also the area of the sample surface in the resin are added to the total surface area of the sample. The power starts to be run, then the sample is held for 10 seconds, then the power is stopped and the sample is cleaned.

After that, the sample is observed under an SEM and a structural photograph of the spherical carbides is taken. Under the SEM, the structures which appear relatively white and which have a ratio (aspect ratio) of the maximum size (long size) and minimum size (short size) of 2 or less are the spherical carbides. The magnification of the photograph taken under the SEM is X1000 or more, while X5000 to X20000 is preferable. For the measurement locations, 10 fields were randomly selected from locations at a depth of about 0.5 to 1 mm from the surface of the wire rod while avoiding the center segregation parts. The thus captured SEM structural photographs were processed by image processing to measure the minimum size (short size) and maximum size (long size) of the spherical carbides seen in the measured fields and calculate the circle equivalent diameter. The circle equivalent diameter is the diameter when calculating the area of an undissolved carbide in a field by image processing and converting it to a circle of the same area. Further, it is also possible to measure the density of presence of spherical carbides with a circle equivalent diameter of 0.2 μm or more seen in the measurement field.

Metal Structure of Pre-Drawn Steel Wire For High Strength Spring Use and Drawn Heat Treated Steel Wire

The metal structure of the drawn heat treated steel wire for high strength spring use according to the present invention is comprised of, by volume rate, over 6% to 15% of residual austenite and a balance of tempered martensite. Fine inclusions are allowed. The “fine inclusion” are oxides and sulfides. The oxides are deoxidation products of Al and Si etc., while the sulfides correspond to MnS, CaS, etc. Further, the balance of the tempered martensite structure also includes undissolved spherical carbides in fine amounts.

The prior austenite grain size number in the structure is #10 or more, while the circle equivalent diameter of the spherical carbides is less than 0.2 μm.

Further, in the metal structure of the pre-drawn steel wire for high strength spring use according to the present invention, the pearlite structure accounts for 90% or more, preferably 95% or more, more preferably 98% or more. A substantially 100% pearlite structure is ideal.

Prior Austenite Grain Size Number: #10 or more

The drawn heat treated steel wire for high strength spring use of the present invention is mainly comprised of tempered martensite in structure. The prior austenite grain size has a great effect on the characteristics. That is, if refining the grain size of the prior austenite, due to the effect of grain refinement, the fatigue characteristics and the coiling ability are improved. In the present invention, to obtain sufficient fatigue characteristics and coiling ability, the prior austenite grain size number is made #10.

Refining the prior austenite is particularly effective for improving the characteristics of the drawn heat treated steel wire for high strength spring use. The prior austenite grain size number is preferably made #11, more preferably #12. To refine the grain size of prior austenite, it is effective to lower the heating temperature of the quenching. Note that, the “prior austenite grain size number” is based on JIS G 0551. If actually performing the quenching by lowering the heating temperature and shortening the time, the prior austenite grain size can be refined, but unreasonable low temperature, short time treatment not only increases the undissolved spherical carbides, but also sometimes results in insufficient austenite transformation itself and two-phase quenching. Conversely, sometimes the coiling ability and the fatigue characteristics are lowered. For this reason usually #13.5 is the upper limit.

Residual Austenite: Over 6% to 15% (volume rate)

The microstructure at the drawn heat treated steel wire for high strength spring use after quenching and tempering is comprised of tempered martensite, residual austenite, and a slight volume fraction of inclusions (here, precipitates also expressed included in inclusions). Residual austenite is effective for improving the cold coiling ability. In the present invention, to secure the cold coiling ability, the volume rate of the residual austenite is made over 6%, preferably 7% or more, more preferably 8% or more.

On the other hand, if the residual austenite exceeds a volume rate of 15%, the martensite which is formed due to the work-induced transformation causes the cold coiling characteristics to drop. Therefore, the volume rate of the residual austenite is made 15% or less, preferably 14% or less, more preferably 12% or less.

The volume rate of the residual austenite can be found by the X-ray diffraction method and the magnetic measurement method. Among these, the magnetic measurement method enables simple measurement of the volume rate of the residual austenite, so is the preferable measurement method. Here, the volume rate is measured, but the obtained figures are the same as the area rate.

Note that, residual austenite is softer than tempered martensite, so reduces the yield strength. Further, the transformation induced plasticity is used to improve the ductility, so this remarkably contributes to improvement of the cold coiling ability. On the other hand, residual austenite often remains at the segregated parts, prior austenite grain boundaries, and near regions sandwiched by the sag grains, so the martensite which is formed by the work-induced transformation (work-induced martensite) becomes starting points of fracture. Further, if the residual austenite increases, the tempered martensite falls relatively.

For this reason, in the past, the drop in the strength and cold coiling ability due to the residual austenite had been considered an issue. However, in high strength steel wire of over 2000 MPa, the amounts of addition of C, Si, Mn, Cr, etc. become greater, so for improvement of the cold coiling ability, utilization of the residual austenite is extremely effective. Further, recently, high precision spring working technology has made it possible to suppress the deterioration of the coiling characteristics even if high hardness parts are locally formed due to the work-induced martensite formed in shaping the spring.

Next, the mechanical properties of the drawn heat treated steel wire for high strength spring use of the present invention will be explained.

To reduce the size and lighten the weight of a spring, it is effective to make it higher in strength. Further, a spring is required to have a superior fatigue strength. In the present invention, a high strength spring is produced by bending the material of the drawn heat treated steel wire for high strength spring use to a desired shape, then nitriding, shot peening, or otherwise hardening the surface. In the nitriding, the spring is heated to 500° C. or so, so the spring is sometimes softened more than the material of the drawn heat treated steel wire for high strength spring use.

Therefore, to raise the strength of the spring and improve the fatigue characteristics, it is necessary to secure the yield strength of the material of the drawn heat treated steel wire for high strength spring use. Further, in order for the drawn heat treated steel wire for high strength spring use to be worked into the desired shape of a spring, cold coiling ability is demanded, so the upper limit of the yield strength has to be limited.

Yield Strength: 2100 to 2400 MPa

If the drawn heat treated steel wire for high strength spring use is high in yield strength, it is possible to improve the fatigue characteristics and sag property of the spring hardened at the surface by nitriding etc. In the present invention, to improve the fatigue characteristics and sag property of the spring, the yield strength of the drawn heat treated steel wire for high strength spring use is made 2100 MPa or more.

Further, the higher the drawn heat treated steel wire for high strength spring use in yield strength, the better the spring in fatigue characteristics, so the drawn heat treated steel wire for high strength spring use has a yield strength of preferably 2200 MPa or more, more preferably 2250 MPa or more.

On the other hand, if the drawn heat treated steel wire for high strength spring use is too high in yield strength, the cold coiling ability falls, so the yield strength is made 2400 MPa or less.

Yield strength (if yield strength cannot be seen, 0.2% proof stress): 1600 to 1980 MPa

In the present invention, the yield strength or yield point of the drawn heat treated steel wire for high strength spring use means the top yield strength when a yield point is seen at the stress-strain curve in a single-axis tensile test and the 0.2% proof stress when no yield point is seen. To secure the strength or sag resistance of the spring, which elastically deformed by repeated stress, raising the yield strength is preferable. To raise the yield strength of the spring, raising the yield strength of the material, that is, the drawn heat treated steel wire for high strength spring use, is preferable.

On the other hand, if the drawn heat treated steel wire for high strength spring use becomes high in yield strength, the cold coiling ability is sometimes impaired. Therefore, the drawn heat treated steel wire for high strength spring use preferably has a yield strength of 1600 MPa or more for securing the strength and sag property of the spring.

To impart further higher durability, 1700 MPa or more is preferable.

On the other hand, if the yield strength exceeds 1980 MPa, the cold coiling ability is sometimes impaired, so the yield strength is preferably made 1980 MPa or less. Note that to raise the yield strength of the drawn heat treated steel wire for high strength spring use of the material having the same yield strength right after short time quenching and tempering, it is preferable to lower the volume the volume rate of the residual austenite.

Vicker's hardness after nitriding by holding at 500° C. for 1 hour: Surface layer hardness HV≧750, internal hardness HV≧570

A high strength spring is improved in surface layer hardness in nitriding, but the inside softens. For example, in gas soft nitriding at 500° C., if the conventional heating temperature becomes 500° C., it was difficult to suppress softening at the inside of the drawn heat treated steel wire for high strength spring use. The drawn heat treated steel wire for high strength spring use of the present invention is excellent in temper softening resistance and enables fatigue characteristics and the sag property of the spring after heating at 500° C. to be secured. In the present invention, the surface layer hardness and the internal hardness after gas soft nitriding are defined.

The surface layer hardness is made a micro Vicker's hardness at the depth of 50 to 100 μm from the surface layer of 750 or more. If less than 750, the surface layer hardness becomes insufficient and the fatigue durability also becomes inferior, so residual stress after shot peening cannot be sufficiently imparted. Preferably, the surface layer hardness is 780 or more.

On the other hand, in internal hardness, the Vicker's hardness is sometimes measured when, in quenching, the temperature of the surface layer of the steel wire is higher than the inside, so measuring this at a position of 500 μm depth from the surface is preferable. To secure the spring fatigue characteristics and sag property, the Vicker's hardness after heat treatment holding the wire at 500° C. for 1 hour should be 570 or more. Furthermore, 575 or more is preferable.

Note that, the upper limit of the Vicker's hardness after holding at 500° C. for 1 hour for heat treatment is not particularly defined, but to ensure that the Vicker's hardness before the heat treatment is not exceeded, usually it is made 783 or less.

Furthermore, when using the drawn heat treated steel wire for high strength spring use of the present invention as the material for production of high strength springs, the surface layer is hardened by shot peening, nitriding, etc. On the other hand, the Vicker's hardness at a position of 500 μm depth from the surface of the high strength spring (internal hardness) is affected by the heating in nitriding. Therefore, when actually producing a spring, the internal hardness will fluctuate depending on the temperature of the nitriding.

Note that, when using the drawn heat treated steel wire for high strength spring use of the present invention as the material for production of high strength springs, it is cold coiled and nitrided. For this reason, the residual austenite at a position of 500 μm depth from the surface of the high strength springs falls somewhat compared with the material of the drawn heat treated steel wire for high strength spring use.

However, the chemical compositions, spherical carbides, and prior austenite crystal grain size are believed to be little affected by the cold coiling and nitriding. Therefore, the chemical compositions, spherical carbides, and prior austenite crystal grain size of the high strength steel made using the drawn heat treated steel wire for high strength spring use of the present invention as a material are the same extent as the chemical compositions, spherical carbides, and prior austenite crystal grain size of the drawn heat treated steel wire for high strength spring use of the present invention.

Next, the method of production of the drawn heat treated steel wire for high strength spring use of the present invention will be explained.

A steel bloom adjusted to predetermined chemical compositions was rolled to produce a steel billet reduced in size. Further, the billet was heated, then hot rolled to obtain pre-drawn steel wire for high strength spring use. This pre-drawn steel wire for high strength spring use was patented, the shaped and, furthermore, was annealed for softening the hard layer. It was then drawn, quenched, and tempered to produce drawn heat treated steel wire for high strength spring use. The “patenting” is heat treatment for making the structure of the steel wire after hot rolling ferrite and pearlite and is performed for softening the steel wire before drawing. After drawing, oil tempering, induction hardening treatment, and other quenching and tempering are performed to adjust the steel wire in structure and characteristics.

In the method of production of pre-drawn steel wire for high strength spring use of the present invention, the process of preventing coarsening of the spherical carbides is important.

In particular, when containing high C and high Cr like in the present invention, it is extremely important to sufficiently heat the bloom or billet before rolling in that state and ease precipitation inside the steel and to dissolve the internal coarse carbides (alloy carbides and cementite) and make the material uniform. To prevent the formation of coarse spherical carbides, the coarse carbides which are formed at the bloom or billet must be made to dissolve in the steel. Furthermore, causing uniform dispersion in the steel is necessary. For this reason, raising the heating temperature is preferable.

Therefore, first, the bloom or billet after casting is made a heating temperature of 1250° C. or more. Due to this, it is possible to make the undissolved spherical carbides sufficiently dissolve. For this reason, in the heating of the subsequent rolling, patenting, and quenching, the heating temperature and the heating time are insufficient, so undissolved spherical carbides easily remain, but to enable sufficient dissolution from the start, the dimensions of the undissolved spherical carbides can be controlled to less than 0.2 μm. The bloom heating temperature should be 1270° C. or more.

Next, the billet which is produced by rolling the bloom is further hot rolled (wire rod is rolled) to produce pre-drawn steel wire for high strength spring use. At this time, the heating temperature of the billet is made 1200° C. or more. Preferably, the heating temperature of the billet should be made 1250° C. or more.

After extracting the steel from the heating furnace, the temperature falls and precipitates grow. For this reason, after extraction from the heating furnace, the hot rolling is preferably completed within 5 minutes. By the above heating of the bloom and billet, the coarse carbides in the steel are uniformly dispersed and dissolved and can uniformly finely precipitate in the later precipitation.

Note that, when rolling a bloom into steel wire without going through a billet, the heating temperature before rolling of the bloom should be made 1250° C. or more, more preferably 1270° C. or more.

In the above way, to suppress coarsening of the undissolved carbides of the steel wire after heat treatment, even if greatly reducing the undissolved carbides which are present before drawing (that is, after wire rod rolling) and if for example undissolved carbides had been present, it is necessary to make the size finer to prevent easy coarsening.

Therefore, in the rolling step of heating before drawing, it is important to make the bloom heating temperature and the billets heating temperature sufficiently high for the carbides to dissolve. Due to this, the size of the undissolved spherical carbides can be kept small. The rolling of the spring steel is completed in several minutes from extraction of the billet from the heating furnace to a size of material before drawing of about φ10 mm. For this reason, it is important to heat to 1200° C. or more where the effect of the billet heating temperature is the largest. 1250° C. or more is more preferable. 1270° C. or more is more preferable.

After rolling, the wire is taken up in a coil and air cooled at that time as general practice. For this reason, usually the microstructure of the pre-drawn steel wire (steel wire after rolling of wire rod) is comprised of ferrite and pearlite or pearlite with a high pearlite structure fraction since the amount of C is high. Undissolved spherical carbides are present in the base material.

The undissolved spherical carbides can be observed by observing a polished and etched detection sample by an SEM. The undissolved carbides can be clearly differentiated from the lamellar cementite contained in the pearlite structure of the base material since they are spherical. Of course, the magnitude may also be measured.

Due to the above step, a pre-drawn steel wire for spring use (rolled wire rod) is obtained.

After hot rolling, the pre-drawn steel wire for spring use is patented. The heating temperature of this patenting may be made 900° C. or more to promote dissolution of the carbides. A high temperature of 930° C. or more is more preferable. Further, 950° C. or more is preferable. After that, the wire may be patented at 600° C. or less. In the pre-drawn steel wire for spring use according to the present invention, the method of patenting and drawing is not limited. If a general patenting and drawing method for steel wire, the same treatment as usual may be performed.

When drawing by the wire diameter and precision required is omitted, the patenting step before the drawing may be omitted. In this case, by making the heating temperature in the later explained quenching high (for example, 970° C. or more), dissolution of the undissolved spherical carbides is promoted.

The quenching after the drawing is performed by heating to temperature of the A₃ point or more. To promote the dissolution of carbides, it is preferable to raise the heating temperature of the quenching. In the quenching, to suppress the growth of carbides, the heating rate is preferably made 10° C./sec or more and the holding time at the temperature of the A₃ point or more is preferably made 1 minute to 5 minutes. To suppress grain growth of the austenite, it is preferable to shorten the holding time. To promote the quenching and martensite transformation, the cooling rate is preferably made 50° C./sec to 100° C.

The coolant in the quenching process is preferably made 100° C. or less, more preferably a low temperature of 80° C. or less, but in the present invention, to precisely control the amount of residual austenite, the coolant temperature is made 40° C. or more. The coolant is not particularly limited so long as being an oil, a water soluble quenching agent, water, or other coolant which enables quenching. Further, the cooling time may be shortened like with oil tempering and induction hardening treatment. It is preferable to avoid extending the holding time at a low temperature for greatly reducing the residual austenite and lowering the coolant temperature to 30° C. or less. That is, the quenching is preferably ended within 5 minutes.

After quenching, tempering is performed. The tempering suppresses the growth of carbides, so it is preferable to make the heating rate 10° C./sec or more and make the holding time 15 minutes or less. The holding time fluctuates due to the chemical compositions and the targeted strength, but the material is usually held at 400 to 500° C.

The pre-drawn steel wire for high strength spring use is cold coiled to work it to the desired spring shape, is relieved of stress, and is nitrided and shot peened to produce the spring.

The cold coiled steel wire is reheated by stress-relieving annealing, nitriding, etc. At this time, the inside is softened, so the performance of the spring falls. In particular, in the present invention, even if performing the nitriding at a high temperature of about 500° C., sufficient hardness is maintained. As a result, if using the pre-drawn steel wire for high strength spring use of the present invention as a material, it is possible to make the micro Vicker's hardness at a depth of 500 μm from the surface layer of high strength springs HV575 or more. Note that, the micro Vicker's hardness is measured at a depth of 500 μm from the surface layer of the spring so as to evaluate the Vicker's hardness of the base material not affected by nitriding and shot peening for hardening.

EXAMPLES

Steels having the chemical compositions shown in Tables 1-1 to 1-4 were smelted in a 10 kg vacuum melting furnace and cast to obtained blooms or billets. These vacuum melted materials were hot forged up to φ8 mm. After that, the materials hot forged up to φ8 mm were heated at 1270° C.×4 hr. Further, part of the samples were refined in a 250 ton converter, continuously cast to prepare blooms, then heated at 1270° C.×4 hr or more, then made into cross-section 160 mm×160 mm billets. Furthermore, these were rolled to obtained φ8 mm rolled wire rods. The heating temperature of the billets before rolling was made 1200° C. or more.

A diameter 8 mm pre-drawn steel wire (rolled wire rod) is preferably made an easily drawn structure by patenting it before drawing. The heating temperature at the patenting is preferably 900° C. or more so that the carbides etc. sufficiently dissolve. The patenting is performed by heating at 930° C., then charging the sample into a 600° C. flowing bed. After patenting, the wire is drawn to obtain a diameter 4 mm drawn wire rod. In this way, by heating the bloom at a high temperature, then making the temperature in the rolling process, patenting, and quenching as high as possible, it is possible to suppress growth of undissolved spherical carbides and keep the dimensions down to 0.2 μm or less.

To adjust the yield strength of the patented and drawn steel wire, the wire was quenched and tempered to produce pre-drawn steel wire for spring use. Note that, a sample which broke in the drawing was not quenched and tempered. The quenching and tempering were performed by heating the drawn steel wire by a 10° C./sec or more heating rate at 950° C. or 1100° C. (temperature of A₃ point or more), holding at the peak heating temperature for 4 minutes to 5 minutes, then placing the steel in a room temperature water tank so that the cooling rate became 50° C./sec or more and cooling down to 100° C. or less.

As the results of evaluation, the state of wire breakage, prior austenite grain size number, residual austenite amount (vol %), circle equivalent diameter and density of presence of carbides, yield strength, 0.2% proof stress, notch bending angle, average fatigue strength, and Vicker's hardness after gas soft nitriding are shown.

The target values to be passed were made as follows with reference to conventional steel wire for high strength spring use.

Prior austenite grain size number: 10 degrees or more

Residual austenite amount (vol %): 20% or less

Circle equivalent diameter of spherical carbides: 0.2 μm or less

Yield strength: 2100 MPa or more

0.2% proof stress: 1800 MPa or more

Yield ratio: 75% to 95%

Notch bending angle: 28 degrees or more

Average fatigue strength (Nakamura type rotating bending strength): 900 MPa or more

Internal hardness by Vicker's hardness after gas nitriding: 590 Hv or more

Nitrided layer hardness by Vicker's hardness after gas nitriding: 750 Hv or more

Note that, in the steel wire according to the present invention, the strength and workability (coiling ability) both have to be achieved, so if the yield ratio is too high, the workability deteriorates. Therefore, the upper limit of the yield ratio is preferably 90%, more preferably 88% or less.

A sample was taken from the obtained drawn heat treated steel wire for spring use, evaluated for prior austenite grain size, volume rate of residual austenite, and carbides, then was subjected to a tensile test, notch bending test, and micro Vicker's hardness test. Note that, the fatigue characteristics were evaluated by treatment simulating production of a spring (below, referred to as “spring production and treatment”) including gas nitriding simulating nitriding performed on the spring after working (500° C., 60 minutes), shot peening (diameter of cut wire 0.6 mm, 20 minutes), and low temperature stress-relieving treatment (180° C., 20 minutes).

The prior austenite grain size number was measured based on JIS G 0551. The circle equivalent diameter and density of presence of the carbides were measured by using an electrolytically etched sample, obtaining a SEM structural photograph, and analyzing the image. Further, the volume rate of the residual austenite was measured by the magnetic measurement method.

The fatigue test is a Nakamura type rotating bending fatigue test (fatigue test bending by two-point supported weight and turning by motor to apply compressive and tensile stress to surface of wire). The maximum load force of 10 samples showing a lifetime of 10⁷ cycles or more by a probability of 50% or more was made the average fatigue strength. The notch bending test is a test for evaluating the cold coiling ability and is performed as follows.

A punch 2 with an angle of the tip shown in FIG. 2 of 120° was used to provide a groove (notch) of a maximum depth of 30 μm in the test piece. Note that, as shown in FIG. 3, the notch 4 was provided at a right angle to the longitudinal direction at the center of the test piece 3 in the longitudinal direction. Next, as shown in FIG. 4, from the opposite side of the notch 4, a pusher 5 was used to apply a load P of a maximum tensile stress through a load-use fixture 6 and the test piece was deformed by three-point bending. Note that, the radius of curvature r of the tip of the load-use fixture 6 was made 4.0 mm, while the difference L between supports was made L=2r+3D. Here, D is the diameter of the test piece.

The bending deformation continued to be applied until the notch part fractured. The bending angle at the time of fracture (notch bending angle) was measured as shown in FIG. 5. Note that, when the test piece was split, the fractured parts were placed together to measure the notch bending angle θ. In the present invention, a sample with a notch bending angle of 28° or more is judged to be excellent in cold coiling ability.

The micro Vicker's hardness after nitriding was evaluated using the depth of 500 μm or more from the surface layer as the internal hardness was defining the micro Vicker's hardness of a depth of 50 μm from the surface layer as the “nitrided layer hardness”. The measurement weight was 10 g.

The results of these tests are shown in Tables 1-5 to 1-8. Note that, in Tables 1-5 to 1-8, the metal structure is comprised of residual austenite (γ) plus tempered martensite and slight inclusions. Further, the balance of the chemical compositions was iron and unavoidable impurities.

The pre-drawn steel wire (steel wire after rolling wire rod) was evaluated only by the circle equivalent diameter of the undissolved spherical carbides. This is because since this is before heat treatment, even if measuring the mechanical properties or the austenite grain size etc., there is not much meaning to the figures.

Examples 1 to 47 of the present invention all have the indicator of the cold coiling ability, that is, the notch bending angle, of a good 28° or more and have an excellent indicator of the spring durability, that is, the Nakamura type rotating bending fatigue strength (hereinafter simply referred to as the “fatigue durability”) and an excellent indicator of the sag property and temper softening resistance, that is, the nitrided layer hardness.

Comparative Examples 48 and 49 are examples where the amount of addition of C is outside the range of the claims. If C is over the prescribed amount (Comparative Example 48), the undissolved spherical carbides become greater and the indicator of the cold coiling ability, the notch bending angle, is low. On the other hand, if C is smaller than the prescribed amount (Comparative Example 49), a sufficient yield strength cannot be secured. In particular, the internal hardness after nitriding becomes lower and the spring fatigue durability (Nakamura type rotating bending fatigue strength) and the sag property (internal hardness after nitriding).

Comparative Examples 50 and 51 are examples where the amount of addition of Si is outside the range of the claims. If Si exceeds the prescribed amount, the matrix is embrittled and the workability is impaired, that is, the notch bending angle is low. On the other hand, if Si is smaller than the prescribed amount, the quenching and tempering characteristics deteriorate, so sufficient strength cannot be secured after heating by nitriding. In particular, the internal hardness after nitriding and the nitrided layer hardness become low.

Comparative Examples 52 and 53 are examples where the amount of addition of Mn is outside the range of the claims. If Mn is over the prescribed range, the residual austenite becomes greater, the yield strength falls, and the fatigue durability (Nakamura type rotating bending fatigue strength) is inferior. On the other hand, when Mn is smaller than the prescribed amount, the residual austenite falls too much and the workability deteriorates, so the notch bending angle falls.

Comparative Examples 54 and 55 are examples where the amount of addition of Cr is outside the range of the claims. If the Cr is over the prescribed range, cementite stabilizes and even in the high temperature heating of the bloom or billet, quenching and tempering, etc., undissolved carbides increase and the spring workability is greatly reduced. For this reason, the notch bending angle falls. On the other hand, if Cr is smaller than the prescribed amount, the steel ends up softening in the heat treatment in the nitriding etc. and the so-called temper softening resistance otherwise becomes insufficiently so the nitrided layer hardness falls.

Comparative Examples 56, 57, and 58 are examples where the amounts of addition of Mo, W, and Mo+W are over the ranges of the claims. If Mo and W exceed the prescribed amounts, in rolling and cooling and after patenting and other heat treatment, a supercooled structure of martensite, bainite, etc. forms, the wire breaks in the conveyance or drawing process, and the measurement test cannot be performed.

Comparative Example 59 is an example of excessive addition of V. V is an element which forms carbides in the steel. Excessive addition causes undissolved carbides to form around the V, the workability to deteriorate, and the notch bending angle to fall.

Comparative Examples 60 and 61 are cases where the amount of N is excessive compared with the range of the claims. This excessive N raises the temperature of formation of nitrides and carbonitrides of V, Nb, etc. and causes coarsening of carbides and other precipitates using these as nuclei. Further, when used for repeated heating such as in the present invention, the nitrides, carbonitride, and carbides are incompletely dissolved and a large amount of coarse undissolved spherical carbides remain. As a result, the workability is impaired. This is an example where the notch bending angle falls.

Comparative Examples 62 and 63 are examples where the amount of addition of Nb is outside the range of the claims. If Nb exceeds the prescribed amount, the hot ductility is remarkably impaired, numerous surface flaws occur at the rolled material, wire breakage occurs during drawing, and a measurement test could not be run.

Comparative Examples 64 is the case where the sum of the amounts of addition of Mn and V is more than the range explained in the present invention. The amount of residual austenite in the steel wire becomes greater than the prescribed value. In the notch bending test, the notch part hardens due to the stress-induced transformation and the workability falls. This is an example where the notch bending angle falls. While repeating ourselves, V is not added in the present invention, but sometimes V is included as an unavoidable impurity, so this is a limitation for rendering the V harmless.

Comparative Examples 65 is the case where the sum of the amounts of addition of Mn and V is lower than the range explained in the present invention. The amount of residual austenite is smaller than the optimum range, so the workability, that is, the notch bending angle, falls.

Comparative Example 66 is the case where the sum of the amounts of addition of Cr and V is greater than the scope explained in the present invention. The undissolved spherical carbides excessively remain and the workability, that is, the notch bending angle, falls.

Comparative Example 67 is the case where the sum of the amounts of addition of Cr and V is less than the range explained in the present invention. The workability is excellent, but the internal hardness after nitriding and the nitrided layer hardness are insufficient and the spring performance is not sufficient.

Comparative Examples 68 to 70 are cases where the difference between the amount of Si and the amount of Cr ([Si %]-[Cr %]) is off from the scope of the claims and the amount of Cr is greater than the amount of Si. If Cr is excessive with respect to the amount of Si, undissolved spherical carbides remain and the workability is degraded, that is, that is, the notch bending angle falls.

Similarly, Comparative Examples 71 and 72 are the case where the difference of the amount of Si and the amount of Cr ([Si %]-[Cr %]) is larger than the upper limit of the range of the claims. Si is very excessive compared with the amount of Cr. In these cases, the surface layer decarburized layer of the rolled material greatly grows and cannot be sufficiently removed by a slight amount of surface layer shaving. For this reason, the fatigue durability (Nakamura type rotating bending fatigue strength) was inferior.

Comparative Examples 73 and 74 are respectively the Invention Example 1 and Invention Example 23 where the steel is rolled at the billet heating temperature 1100° C. At the start of the rolling, undissolved spherical carbides remain. The effects finally remain, so the workability is degraded, that is, the notch bending angle falls.

Invention Examples 101 to 109 are examples of the pre-drawn steel wires of Invention Examples 1 to 5 and 20 to 23. Comparative Examples 110 and 111 are the Invention Examples 101 and 106 where the billet heating temperature is made 1100° C.

The pre-drawn steel wire is evaluated, so only the maximum circle equivalent diameter of the undissolved spherical carbides is evaluated. If the billet heating temperature is high, it is learned that the circle equivalent diameter of the undissolved spherical carbides becomes smaller.

	TABLE 1-1
	Chemical compositions (mass %)
Ex.	C	Si	Mn	P	S	Cr	Al	N	V	Nb
1	Inv. ex.	0.78	2.48	0.68	0.0076	0.0045	1.57	0.0022	0.0031	—	—
2	Inv. ex.	0.77	2.41	0.68	0.0034	0.0047	2.05	0.0011	0.0042	—	—
3	Inv. ex.	0.68	2.38	0.87	0.0045	0.0061	1.53	0.0013	0.0033	—	—
4	Inv. ex.	0.88	2.50	0.87	0.0063	0.0071	1.71	0.0018	0.0035	—	—
5	Inv. ex.	0.78	2.11	0.84	0.0057	0.0039	1.50	0.0010	0.0031	—	—
6	Inv. ex.	0.72	2.62	0.62	0.0054	0.0031	1.96	0.0022	0.0038	—	—
7	Inv. ex.	0.72	2.67	0.58	0.0037	0.0035	1.51	0.0015	0.0032	—	—
8	Inv. ex.	0.76	2.28	1.02	0.0067	0.0069	1.82	0.0028	0.0031	—	—
9	Inv. ex.	0.73	2.23	0.73	0.0054	0.0077	1.53	0.0015	0.0058	—	—
10	Inv. ex.	0.77	2.35	0.80	0.0041	0.0047	1.52	0.0019	0.0032	—	—
11	Inv. ex.	0.77	2.59	0.83	0.0060	0.0074	1.53	0.0012	0.0063	—	0.008
12	Inv. ex.	0.75	2.52	0.68	0.0054	0.0056	1.67	0.0013	0.0039	0.06	—
13	Inv. ex.	0.75	2.33	0.83	0.0066	0.0060	2.00	0.0030	0.0037	0.09	0.001
14	Inv. ex.	0.72	2.41	0.86	0.0040	0.0068	1.77	0.0013	0.0033	—	—
15	Inv. ex.	0.77	2.57	0.75	0.0078	0.0075	1.91	0.0018	0.0043	—	—
16	Inv. ex.	0.72	2.42	0.84	0.0044	0.0053	1.90	0.0019	0.0056	—	—
17	Inv. ex.	0.76	2.53	0.67	0.0061	0.0076	1.65	0.0028	0.0039	—	—
18	Inv. ex.	0.73	2.46	0.66	0.0051	0.0060	1.60	0.0023	0.0035	—	—
19	Inv. ex.	0.73	2.34	0.75	0.0039	0.0071	1.97	0.0028	0.0032	—	—
20	Inv. ex.	0.73	2.35	0.70	0.0046	0.0033	1.91	0.0016	0.0034	0.10	0.008
21	Inv. ex.	0.77	2.46	0.71	0.0031	0.0054	1.95	0.0013	0.0038	0.03	0.003
22	Inv. ex.	0.76	2.35	0.64	0.0051	0.0072	1.72	0.0016	0.0044	0.07	0.007
23	Inv. ex.	0.73	2.36	0.73	0.0033	0.0073	1.80	0.0019	0.0033	0.08	0.006
24	Inv. ex.	0.76	3.20	0.72	0.0076	0.0038	2.10	0.0010	0.0037	0.06	—
	Chemical compositions (mass %)
Ex.	Mo	W	Mg	Zr	Ca	Mn + V	Cr + V	Si—Cr	Mo + W
1	Inv. ex.	—	—	—	—	—	—	—	0.90	—
2	Inv. ex.	—	—	—	—	—	—	—	0.35	—
3	Inv. ex.	—	—	—	—	—	—	—	0.85	—
4	Inv. ex.	—	—	—	—	—	—	—	0.79	—
5	Inv. ex.	—	—	—	—	—	—	—	0.61	—
6	Inv. ex.	—	—	—	—	—	—	—	0.66	—
7	Inv. ex.	—	—	—	—	—	—	—	1.16	—
8	Inv. ex.	—	—	—	—	—	—	—	0.46	—
9	Inv. ex.	—	—	—	—	—	—	—	0.70	—
10	Inv. ex.	—	—	—	—	—	—	—	0.83	—
11	Inv. ex.	—	—	—	—	—	—	—	1.07	—
12	Inv. ex.	—	—	—	—	—	0.74	1.74	0.84	—
13	Inv. ex.	—	—	—	—	—	0.93	2.10	0.32	—
14	Inv. ex.	0.12	—	—	—	—	—	—	0.64	0.12
15	Inv. ex.	—	0.15	—	—	—	—	—	0.66	0.15
16	Inv. ex.	0.12	0.16	—	—	—	—	—	0.52	0.27
17	Inv. ex.	—	—	0.0005	—	—	—	—	0.87	—
18	Inv. ex.	—	—	—	0.0002	—	—	—	0.86	—
19	Inv. ex.	—	—	—	—	0.0011	—	—	0.37	—
20	Inv. ex.	0.12	0.17	0.0002	0.0003	0.0011	0.79	2.00	0.44	0.28
21	Inv. ex.	0.15	0.16	0.0003	0.0002	0.0003	0.74	1.98	0.51	0.31
22	Inv. ex.	0.11	0.17	0.0005	0.0003	0.0006	0.71	1.79	0.64	0.28
23	Inv. ex.	0.11	0.15	0.0004	0.0001	0.0012	0.81	1.88	0.56	0.27
24	Inv. ex.	0.10	—	—	—	—	0.77	2.16	1.10	0.10

	TABLE 1-2
	Chemical compositions (mass %)
Ex.	C	Si	Mn	P	S	Cr	Al	N	V	Nb
25	Inv. ex.	0.73	2.05	0.80	0.0062	0.0052	1.71	0.0011	0.0033	0.08	0.009
26	Inv. ex.	0.76	2.52	0.71	0.0069	0.0080	1.69	0.0024	0.0039	0.05	—
27	Inv. ex.	0.75	2.30	0.82	0.0074	0.0055	1.73	0.0018	0.0038	0.08	0.010
28	Inv. ex.	0.73	2.24	1.20	0.0052	0.0074	1.52	0.0027	0.0033	0.09	—
29	Inv. ex.	0.73	2.44	0.74	0.0076	0.0046	1.73	0.0024	0.0032	0.05	0.010
30	Inv. ex.	0.74	2.24	0.76	0.0052	0.0076	1.41	0.0016	0.0030	0.05	0.005
31	Inv. ex.	0.75	2.28	0.89	0.0065	0.0043	1.70	0.0011	0.0043	0.03	0.010
32	Inv. ex.	0.74	2.23	0.89	0.0043	0.0049	1.50	0.0022	0.0032	0.04	—
33	Inv. ex.	0.77	2.86	0.77	0.0057	0.0057	2.48	0.0015	0.0040	0.09	0.005
34	Inv. ex.	0.77	2.29	0.79	0.0055	0.0072	1.85	0.0016	0.0032	0.09	0.004
35	Inv. ex.	0.73	2.52	0.76	0.0078	0.0067	2.00	0.0022	0.0036	0.09	0.009
36	Inv. ex.	0.77	2.31	0.89	0.0046	0.0080	1.91	0.0021	0.0032	0.06	0.000
37	Inv. ex.	0.73	2.53	0.69	0.0046	0.0060	2.01	0.0022	0.0042	—	0.005
38	Inv. ex.	0.76	2.35	0.81	0.0034	0.0036	1.80	0.0016	0.0040	0.04	0.006
39	Inv. ex.	0.74	2.38	0.72	0.0031	0.0064	1.82	0.0010	0.0043	0.07	0.009
40	Inv. ex.	0.75	2.32	0.72	0.0034	0.0038	1.56	0.0017	0.0033	0.09	0.002
41	Inv. ex.	0.76	2.37	0.69	0.0056	0.0067	1.63	0.0028	0.0041	0.04	—
42	Inv. ex.	0.76	2.48	0.73	0.0050	0.0035	1.97	0.0016	0.0034	0.09	—
43	Inv. ex.	0.72	2.26	0.68	0.0053	0.0043	1.51	0.0027	0.0032	0.05	—
44	Inv. ex.	0.76	2.38	0.86	0.0077	0.0039	1.61	0.0020	0.0052	—	—
45	Inv. ex.	0.77	2.23	0.76	0.0060	0.0061	1.69	0.0029	0.0053	0.05	0.006
46	Inv. ex.	0.76	2.35	0.86	0.0060	0.0067	1.87	0.0025	0.0035	0.05	0.010
47	Inv. ex.	0.71	2.28	0.77	0.0050	0.0058	1.88	0.0030	0.0033	0.07	0.003
	Chemical compositions (mass %)
Ex.	Mo	W	Mg	Zr	Ca	Mn + V	Cr + V	Si—Cr	Mo + W
25	Inv. ex.	0.15	0.16	—	—	—	0.88	1.79	0.34	0.31
26	Inv. ex.	0.28	—	0.0002	0.0003	0.0015	0.76	1.74	0.84	0.28
27	Inv. ex.	0.13	0.15	0.0004	0.0001	0.0005	0.91	1.82	0.56	0.28
28	Inv. ex.	0.11	—	0.0003	0.0001	0.0011	1.29	1.62	0.72	0.11
29	Inv. ex.	0.12	0.15	0.0001	0.0003	0.0013	0.79	1.78	0.71	0.27
30	Inv. ex.	0.11	0.16	0.0003	0.0001	0.0011	0.81	1.46	0.83	0.27
31	Inv. ex.	0.14	0.16	0.0004	0.0002	0.0012	0.93	1.73	0.58	0.31
32	Inv. ex.	0.11	—	0.0001	0.0001	0.0002	0.93	1.54	0.73	0.11
33	Inv. ex.	0.11	0.16	0.0002	0.0002	0.0005	0.86	2.57	0.38	0.27
34	Inv. ex.	0.23	0.15	0.0004	0.0003	0.0006	0.88	1.93	0.45	0.38
35	Inv. ex.	0.18	0.28	0.0005	0.0001	0.0005	0.84	2.09	0.52	0.46
36	Inv. ex.	0.13	0.15	0.0001	0.0002	0.0006	0.95	1.97	0.40	0.28
37	Inv. ex.	0.14	0.17	0.0005	0.0002	0.0006	—	—	0.52	0.31
38	Inv. ex.	0.12	0.09	0.0004	0.0001	0.0009	0.85	1.84	0.55	0.21
39	Inv. ex.	0.11	0.28	0.0002	0.0002	0.0007	0.80	1.89	0.55	0.39
40	Inv. ex.	0.13	0.15	0.0004	0.0002	0.0012	0.81	1.65	0.76	0.28
41	Inv. ex.	0.14	0.17	0.0002	0.0002	0.0008	0.73	1.67	0.74	0.30
42	Inv. ex.	0.10	0.16	—	—	—	0.82	2.06	0.51	0.26
43	Inv. ex.	0.14	0.14	0.0002	0.0003	0.0002	0.73	1.56	0.75	0.29
44	Inv. ex.	0.14	0.15	—	—	—	—	—	0.77	0.28
45	Inv. ex.	0.11	0.15	—	—	—	0.81	1.74	0.54	0.26
46	Inv. ex.	0.10	0.15	0.0005	0.0003	0.0010	0.90	1.92	0.48	0.25
47	Inv. ex.	0.10	0.17	0.0005	0.0003	0.0006	0.84	1.95	0.41	0.27

	TABLE 1-3
	Chemical compositions (mass %)
Ex.	C	Si	Mn	P	S	Cr	Al	N	V	Nb
48	Comp. ex.	0.95	2.47	0.61	0.0033	0.0078	1.54	0.0022	0.0043	0.09	0.006
49	Comp. ex.	0.58	2.59	0.62	0.0079	0.0057	1.81	0.0021	0.0066	0.09	0.005
50	Comp. ex.	0.71	3.80	0.85	0.0049	0.0035	2.03	0.0017	0.0032	0.05	0.002
51	Comp. ex.	0.73	1.86	0.83	0.0034	0.0060	1.32	0.0013	0.0031	0.09	0.004
52	Comp. ex.	0.72	2.46	1.54	0.0076	0.0066	2.02	0.0024	0.0043	0.06	0.008
53	Comp. ex.	0.72	2.22	0.21	0.0057	0.0064	1.58	0.0024	0.0036	0.09	0.005
54	Comp. ex.	0.73	3.13	0.71	0.0048	0.0030	2.72	0.0014	0.0030	0.04	0.007
55	Comp. ex.	0.78	2.21	0.68	0.0051	0.0041	1.02	0.0013	0.0031	0.09	0.011
56	Comp. ex.	0.75	2.42	0.74	0.0034	0.0067	1.78	0.0017	0.0033	0.09	0.005
57	Comp. ex.	0.73	2.45	0.81	0.0046	0.0061	1.75	0.0027	0.0046	0.11	0.010
58	Comp. ex.	0.73	2.30	0.64	0.0044	0.0034	1.80	0.0011	0.0037	—	0.001
59	Comp. ex.	0.74	2.23	0.80	0.0060	0.0046	1.72	0.0026	0.0032	0.46	0.007
60	Comp. ex.	0.77	2.43	0.83	0.0062	0.0072	1.96	0.0025	0.0076	0.08	—
61	Comp. ex.	0.75	2.26	0.78	0.0069	0.0058	1.99	0.0013	0.0085	0.07	0.010
62	Comp. ex.	0.73	2.50	0.62	0.0049	0.0062	1.73	0.0020	0.0053	0.05	0.035
63	Comp. ex.	0.72	2.21	0.72	0.0074	0.0051	1.86	0.0026	0.0036	—	0.024
64	Comp. ex.	0.74	2.50	1.18	0.0053	0.0056	2.02	0.0022	0.0042	0.09	0.006
65	Comp. ex.	0.74	2.52	0.51	0.0064	0.0046	1.73	0.0019	0.0039	0.06	0.003
66	Comp. ex.	0.75	2.76	0.72	0.0048	0.0058	2.45	0.0014	0.0032	0.09	0.004
67	Comp. ex.	0.77	2.20	0.79	0.0059	0.0059	1.31	0.0016	0.0035	0.03	0.002
68	Comp. ex.	0.76	2.12	0.66	0.0068	0.0075	2.31	0.0026	0.0033	0.09	—
69	Comp. ex.	0.74	2.10	0.88	0.0056	0.0056	2.23	0.0023	0.0043	0.06	0.005
70	Comp. ex.	0.78	2.23	0.73	0.0067	0.0075	2.41	0.0023	0.0048	0.11	0.000
71	Comp. ex.	0.71	3.12	0.76	0.0068	0.0044	1.54	0.0021	0.0035	0.07	0.005
72	Comp. ex.	0.76	3.45	0.66	0.0071	0.0052	1.66	0.0017	0.0036	0.04	0.005
73	Comp. ex.	0.78	2.48	0.68	0.0076	0.0045	1.57	0.0022	0.0031	—	—
74	Comp. ex.	0.73	2.36	0.73	0.0033	0.0073	1.80	0.0019	0.0033	0.08	0.006
	Chemical compositions (mass %)
Ex.	Mo	W	Mg	Zr	Ca	Mn + V	Cr + V	Si—Cr	Mo + W
48	Comp. ex.	0.12	0.14	0.0004	0.0001	0.0013	0.70	1.63	0.93	0.26
49	Comp. ex.	0.15	0.15	0.0003	0.0001	0.0014	0.71	1.90	0.77	0.29
50	Comp. ex.	0.14	0.16	0.0002	0.0003	0.0008	0.90	2.09	1.77	0.31
51	Comp. ex.	0.14	0.14	0.0004	0.0003	0.0005	0.92	1.41	0.54	0.29
52	Comp. ex.	0.14	0.15	0.0003	0.0003	0.0014	1.60	2.08	0.44	0.29
53	Comp. ex.	0.15	0.17	0.0001	0.0002	0.0011	0.30	1.67	0.64	0.31
54	Comp. ex.	0.14	0.15	0.0003	0.0002	0.0011	0.75	2.76	0.41	0.29
55	Comp. ex.	0.14	0.17	0.0002	0.0001	0.0012	0.77	1.11	1.19	0.31
56	Comp. ex.	0.42	0.07	0.0004	0.0003	0.0005	0.83	1.87	0.63	0.49
57	Comp. ex.	0.12	0.50	0.0003	0.0001	0.0008	0.92	1.86	0.70	0.62
58	Comp. ex.	0.26	0.27	0.0005	0.0002	0.0007	—	—	0.50	0.53
59	Comp. ex.	0.11	0.16	0.0004	0.0001	0.0015	1.26	2.18	0.51	0.27
60	Comp. ex.	0.15	0.17	0.0002	0.0001	0.0004	0.91	2.04	0.47	0.32
61	Comp. ex.	0.10	0.16	0.0003	0.0002	0.0005	0.85	2.06	0.27	0.27
62	Comp. ex.	0.13	0.17	—	—	—	0.67	1.78	0.77	0.29
63	Comp. ex.	0.12	0.17	0.0002	0.0001	0.0012	—	—	0.34	0.29
64	Comp. ex.	0.10	0.16	—	—	—	1.27	2.11	0.48	0.26
65	Comp. ex.	0.12	0.16	0.0004	0.0002	0.0007	0.57	1.79	0.79	0.28
66	Comp. ex.	0.11	0.15	0.0002	0.0002	0.0015	0.81	2.54	0.31	0.26
67	Comp. ex.	0.14	0.16	0.0003	0.0001	0.0011	0.82	1.34	0.89	0.30
68	Comp. ex.	0.15	0.16	0.0003	0.0001	0.0013	0.75	2.40	−0.19	0.31
69	Comp. ex.	0.14	0.16	—	—	—	0.94	2.29	−0.13	0.30
70	Comp. ex.	0.12	0.17	0.0004	0.0003	0.0013	0.84	2.52	−0.18	0.28
71	Comp. ex.	0.12	0.15	—	—	—	0.84	1.61	1.58	0.26
72	Comp. ex.	0.11	0.16	0.0004	0.0002	0.0002	0.70	1.70	1.79	0.26
73	Comp. ex.	—	—	—	—	—	—	—	0.90	—
74	Comp. ex.	0.11	0.15	0.0004	0.0001	0.0012	0.81	1.88	0.56	0.27

	TABLE 1-4
	Chemical compositions (mass %)
Ex.	C	Si	Mn	P	S	Cr	Al	N	V	Nb
101	Inv. ex.	0.78	2.48	0.68	0.0076	0.0045	1.57	0.0022	0.0031	—	—
102	Inv. ex.	0.77	2.41	0.68	0.0034	0.0047	2.05	0.0011	0.0042	—	—
103	Inv. ex.	0.68	2.38	0.87	0.0045	0.0061	1.53	0.0013	0.0033	—	—
104	Inv. ex.	0.88	2.50	0.87	0.0063	0.0071	1.71	0.0018	0.0035	—	—
105	Inv. ex.	0.78	2.11	0.84	0.0057	0.0039	1.50	0.0010	0.0031	—	—
106	Inv. ex.	0.73	2.35	0.70	0.0046	0.0033	1.91	0.0016	0.0034	0.10	0.008
107	Inv. ex.	0.77	2.46	0.71	0.0031	0.0054	1.95	0.0013	0.0038	0.03	0.003
108	Inv. ex.	0.76	2.35	0.64	0.0051	0.0072	1.72	0.0016	0.0044	0.07	0.007
109	Inv. ex.	0.73	2.36	0.73	0.0033	0.0073	1.80	0.0019	0.0033	0.08	0.006
110	Comp. ex.	0.78	2.48	0.68	0.0076	0.0045	1.57	0.0022	0.0031	—	—
111	Comp. ex.	0.73	2.36	0.73	0.0033	0.0073	1.80	0.0019	0.0033	0.08	0.006
	Chemical compositions (mass %)
Ex.	Mo	W	Mg	Zr	Ca	Mn + V	Cr + V	Si—Cr	Mo + W
101	Inv. ex.	—	—	—	—	—	—	—	0.90	—
102	Inv. ex.	—	—	—	—	—	—	—	0.35	—
103	Inv. ex.	—	—	—	—	—	—	—	0.85	—
104	Inv. ex.	—	—	—	—	—	—	—	0.79	—
105	Inv. ex.	—	—	—	—	—	—	—	0.61	—
106	Inv. ex.	0.12	0.17	0.0002	0.0003	0.0011	0.79	2.00	0.44	0.28
107	Inv. ex.	0.15	0.16	0.0003	0.0002	0.0003	0.74	1.98	0.51	0.31
108	Inv. ex.	0.11	0.17	0.0005	0.0003	0.0006	0.71	1.79	0.64	0.28
109	Inv. ex.	0.11	0.15	0.0004	0.0001	0.0012	0.81	1.88	0.56	0.27
110	Comp. ex.	—	—	—	—	—	—	—	0.90	—
111	Comp. ex.	0.11	0.15	0.0004	0.0001	0.0012	0.81	1.88	0.56	0.27

TABLE 1-5
	Billet	Patenting	Quenching	Wire break etc.	Max.	Prior
	heating temp.	heating temp.	heating temp.	[good = no	spherical carbide	austenite grain	Residual austenite
Ex.	(° C.)	(° C.)	(° C.)	abnormality]	diameter μm	size (γ#)	(vol %)
1	Inv. ex.	1250	930	950	good	0.06	10	9
2	Inv. ex.	1250	930	950	good	0.13	12	9
3	Inv. ex.	1250	930	950	good	0.06	11	8
4	Inv. ex.	1250	930	950	good	0.09	12	13
5	Inv. ex.	1250	930	950	good	0.13	13	11
6	Inv. ex.	1250	930	950	good	0.10	12	10
7	Inv. ex.	1200	930	950	good	0.08	10	10
8	Inv. ex.	1250	930	950	good	0.02	10	6
9	Inv. ex.	1250	930	950	good	0.07	13	11
10	Inv. ex.	1200	930	1010	good	0.11	13	9
11	Inv. ex.	1250	930	1010	good	0.13	10	10
12	Inv. ex.	1250	930	1010	good	0.10	12	7
13	Inv. ex.	1250	930	950	good	0.10	13	13
14	Inv. ex.	1250	930	950	good	0.07	13	10
15	Inv. ex.	1250	930	950	good	0.10	11	12
16	Inv. ex.	1250	930	950	good	0.01	12	9
17	Inv. ex.	1250	930	950	good	0.12	11	13
18	Inv. ex.	1250	930	950	good	0.09	12	9
19	Inv. ex.	1250	930	950	good	0.11	10	7
20	Inv. ex.	1250	930	950	good	0.09	10	8
21	Inv. ex.	1250	930	950	good	0.04	12	12
22	Inv. ex.	1250	—	970	good	0.04	10	12
23	Inv. ex.	1250	930	950	good	0.06	12	11
24	Inv. ex.	1200	930	950	good	0.10	11	11
				Notch		Internal	Hardness of
	Tensile			bending deg.	Nakamura type rotary	hardness after	nitrided
Ex.	strength (MPa)	0.2% proof stress	Yield ratio (%)	(deg.)	bending (MPa)	nitriding (HV)	layer (HV)
1	Inv. ex.	2214	1887	85	36	927	597	788
2	Inv. ex.	2288	1939	85	39	918	638	787
3	Inv. ex.	2383	1945	82	37	924	599	819
4	Inv. ex.	2217	1811	82	39	915	621	799
5	Inv. ex.	2305	1844	80	36	918	627	817
6	Inv. ex.	2241	1899	85	39	918	623	816
7	Inv. ex.	2243	1978	88	39	911	618	793
8	Inv. ex.	2301	1811	79	38	914	600	789
9	Inv. ex.	2293	1880	82	35	928	607	806
10	Inv. ex.	2263	1801	80	37	912	612	795
11	Inv. ex.	2283	1939	85	37	929	608	808
12	Inv. ex.	2330	1844	79	38	913	616	780
13	Inv. ex.	2169	1904	88	39	925	627	819
14	Inv. ex.	2346	1854	79	35	919	612	808
15	Inv. ex.	2189	1917	88	36	924	599	804
16	Inv. ex.	2285	1822	80	38	916	625	805
17	Inv. ex.	2248	1950	87	40	929	595	787
18	Inv. ex.	2193	1904	87	38	925	616	804
19	Inv. ex.	2180	1949	89	36	928	598	786
20	Inv. ex.	2374	1891	80	39	924	612	817
21	Inv. ex.	2368	1846	78	36	911	638	806
22	Inv. ex.	2254	1853	82	40	914	629	795
23	Inv. ex.	2311	1852	80	39	923	625	783
24	Inv. ex.	2250	1956	87	35	911	616	793

TABLE 1-6
	Billet	Patenting	Quenching	Wire break etc.	Max.	Prior
	heating temp.	heating temp.	heating temp.	[good = no	spherical carbide	austenite grain	Residual austenite
Ex.	(° C.)	(° C.)	(° C.)	abnormality]	diameter μm	size (γ#)	(vol %)
25	Inv. ex.	1250	930	950	good	0.10	10	8
26	Inv. ex.	1250	930	950	good	0.12	11	13
27	Inv. ex.	1250	930	950	good	0.11	11	12
28	Inv. ex.	1200	930	950	good	0.05	11	13
29	Inv. ex.	1250	930	950	good	0.09	13	9
30	Inv. ex.	1250	930	950	good	0.04	11	7
31	Inv. ex.	1250	930	950	good	0.07	12	10
32	Inv. ex.	1250	930	950	good	0.01	11	12
33	Inv. ex.	1200	930	950	good	0.02	10	10
34	Inv. ex.	1250	930	950	good	0.03	11	11
35	Inv. ex.	1250	930	950	good	0.01	12	9
36	Inv. ex.	1250	930	950	good	0.05	12	13
37	Inv. ex.	1250	—	970	good	0.01	10	9
38	Inv. ex.	1250	930	950	good	0.07	11	6
39	Inv. ex.	1250	930	950	good	0.15	12	6
40	Inv. ex.	1250	930	950	good	0.06	13	8
41	Inv. ex.	1250	930	950	good	0.01	10	8
42	Inv. ex.	1250	930	950	good	0.06	13	6
43	Inv. ex.	1250	930	950	good	0.08	12	12
44	Inv. ex.	1250	930	950	good	0.06	12	8
45	Inv. ex.	1250	930	950	good	0.09	12	6
46	Inv. ex.	1250	930	950	good	0.13	12	12
47	Inv. ex.	1250	930	950	good	0.03	11	7
				Notch		Internal	Hardness of
	Tensile			bending deg.	Nakamura type rotary	hardness after	nitrided
Ex.	strength (MPa)	0.2% proof stress	Yield ratio (%)	(deg.)	bending (MPa)	nitriding (HV)	layer (HV)
25	Inv. ex.	2357	1953	83	39	916	619	792
26	Inv. ex.	2195	1865	85	35	930	628	794
27	Inv. ex.	2226	1872	84	38	922	592	795
28	Inv. ex.	2181	1854	85	38	910	631	784
29	Inv. ex.	2277	1890	83	37	916	635	786
30	Inv. ex.	2330	1938	83	37	922	593	798
31	Inv. ex.	2305	1867	81	38	921	630	791
32	Inv. ex.	2216	1880	85	35	919	592	813
33	Inv. ex.	2285	1929	84	39	911	597	797
34	Inv. ex.	2329	1944	83	39	916	601	794
35	Inv. ex.	2310	1824	79	35	911	609	793
36	Inv. ex.	2153	1951	91	37	927	615	810
37	Inv. ex.	2220	1904	86	35	919	606	816
38	Inv. ex.	2178	1971	91	38	921	605	803
39	Inv. ex.	2268	1820	80	39	920	590	784
40	Inv. ex.	2302	1860	81	39	919	604	814
41	Inv. ex.	2190	1896	87	37	924	611	806
42	Inv. ex.	2218	1893	85	37	927	639	787
43	Inv. ex.	2382	1949	82	40	929	624	819
44	Inv. ex.	2269	1869	82	35	918	618	806
45	Inv. ex.	2155	1880	87	37	919	627	782
46	Inv. ex.	2314	1964	85	38	912	639	813
47	Inv. ex.	2220	1892	85	37	919	610	798

TABLE 1-7
	Billet	Patenting	Quenching	Wire break etc.	Max.	Prior
	heating temp.	heating temp.	heating temp.	[good = no	spherical carbide	austenite grain	Residual austenite
Ex.	(° C.)	(° C.)	(° C.)	abnormality]	diameter (μm)	size (γ#)	(vol %)
48	Comp. ex.	1250	930	950	good	0.31	12	9
49	Comp. ex.	1250	930	950	good	0.02	13	7
50	Comp. ex.	1250	930	950	good	0.04	13	8
51	Comp. ex.	1250	930	950	good	0.12	11	7
52	Comp. ex.	1250	930	950	good	0.08	12	21
53	Comp. ex.	1250	930	950	good	0.10	11	2
54	Comp. ex.	1250	930	950	good	0.26	13	10
55	Comp. ex.	1250	930	950	good	0.03	10	7
56	Comp. ex.	1250	930	950	wire Break	—	—	—
57	Comp. ex.	1250	930	950	wire break	—	—	—
58	Comp. ex.	1250	930	950	wire break	—	—	—
59	Comp. ex.	1250	930	950	good	0.42	12	7
60	Comp. ex.	1250	930	950	good	0.25	12	8
61	Comp. ex.	1250	930	950	good	0.13	12	6
62	Comp. ex.	1250	930	950	wire break	—	—	—
63	Comp. ex.	1250	930	950	wire break	—	—	—
64	Comp. ex.	1250	930	950	good	0.03	11	17
65	Comp. ex.	1250	930	950	good	0.11	11	3
66	Comp. ex.	1250	930	950	good	0.33	11	10
67	Comp. ex.	1250	930	950	good	0.03	10	9
68	Comp. ex.	1250	930	950	good	0.28	12	11
69	Comp. ex.	1250	930	950	good	0.22	13	6
70	Comp. ex.	1250	930	950	good	0.26	12	8
71	Comp. ex.	1250	930	950	Decarburized	—	12	9
72	Comp. ex.	1250	930	950	Decarburized	—	10	12
73	Comp. ex.	1100	930	950	good	0.26	10	9
74	Comp. ex.	1100	930	950	good	0.28	10	6
				Notch		Internal	Hardness of
	Tensile			bending deg.	Nakamura type rotary	hardness after	nitrided
Ex.	strength (MPa)	0.2% proof stress	Yield ratio (%)	(deg.)	bending (MPa)	nitriding (HV)	layer (HV)
48	Comp. ex.	2320	1857	80	24	915	617	807
49	Comp. ex.	2002	1806	90	39	812	554	813
50	Comp. ex.	2338	1841	79	23	926	612	812
51	Comp. ex.	2236	1954	87	36	915	568	728
52	Comp. ex.	2318	1680	75	37	820	612	816
53	Comp. ex.	2395	1971	82	25	923	637	801
54	Comp. ex.	2363	1909	81	23	920	629	795
55	Comp. ex.	2382	1840	77	39	919	608	731
56	Comp. ex.	—	—	—	—	—	—	—
57	Comp. ex.	—	—	—	—	—	—	—
58	Comp. ex.	—	—	—	—	—	—	—
59	Comp. ex.	2227	1885	85	21	929	635	815
60	Comp. ex.	2250	1881	84	17	914	640	816
61	Comp. ex.	2384	1830	77	17	916	606	817
62	Comp. ex.	—	—	—	—	—	—	—
63	Comp. ex.	—	—	—	—	—	—	—
64	Comp. ex.	2233	1560	70	21	786	561	783
65	Comp. ex.	2232	1860	83	24	911	606	815
66	Comp. ex.	2311	1838	80	21	916	613	783
67	Comp. ex.	2276	1977	87	43	789	561	732
68	Comp. ex.	2297	1934	84	24	923	612	809
69	Comp. ex.	2243	1838	82	22	914	629	813
70	Comp. ex.	2329	1813	78	21	922	596	787
71	Comp. ex.	2359	1898	80	43	774	609	780
72	Comp. ex.	2371	1830	77	43	781	631	797
73	Comp. ex.	2114	1827	86	25	911	591	793
74	Comp. ex.	2251	1822	81	22	910	602	751

TABLE 1-8
	Billet	Patenting	Quenching	Wire break etc.	Max.	Prior
	heating temp.	heating temp.	heating temp.	[good = no	spherical carbide	austenite grain	Residual austenite
Ex.	(° C.)	(° C.)	(° C.)	abnormality]	diameter μm	size (γ#)	(vol %)
101	Inv. ex.	1250	—	—	—	0.03	—	—
102	Inv. ex.	1200	—	—	—	0.15	—	—
103	Inv. ex.	1250	—	—	—	0.08	—	—
104	Inv. ex.	1250	—	—	—	0.04	—	—
105	Inv. ex.	1250	—	—	—	0.09	—	—
106	Inv. ex.	1250	—	—	—	0.03	—	—
107	Inv. ex.	1250	—	—	—	0.02	—	—
108	Inv. ex.	1250	—	—	—	0.03	—	—
109	Inv. ex.	1200	—	—	—	0.14	—	—
110	Comp.	1100	—	—	—	0.21	—	—
	ex.
111	Comp.	1100	—	—	—	0.22	—	—
	ex.
				Notch		Internal	Hardness of
	Tensile			bending deg.	Nakamura type rotary	hardness after	nitrided
Ex.	strength (MPa)	0.2% proof stress	Yield ratio (%)	(deg.)	bending (MPa)	nitriding (HV)	layer (HV)
101	Inv. ex.	—	—	—	—	—	—	—
102	Inv. ex.	—	—	—	—	—	—	—
103	Inv. ex.	—	—	—	—	—	—	—
104	Inv. ex.	—	—	—	—	—	—	—
105	Inv. ex.	—	—	—	—	—	—	—
106	Inv. ex.	—	—	—	—	—	—	—
107	Inv. ex.	—	—	—	—	—	—	—
108	Inv. ex.	—	—	—	—	—	—	—
109	Inv. ex.	—	—	—	—	—	—	—
110	Comp.	—	—	—	—	—	—	—
	ex.
111	Comp.	—	—	—	—	—	—	—
	ex.

INDUSTRIAL APPLICABILITY

The present invention can be utilized for the production of steel wire for high strength spring use. The high strength spring material can be utilized in many industrial fields starting from the automotive industry.

REFERENCE SIGNS LIST

• 1 spherical carbides
• 2 punch
• 3 test piece
• 4 notch
• 5 pusher
• 6 load use fixture
• P load
• L distance between supports
• θ notch bending angle

Claims

1. Pre-drawn steel wire for high strength spring use characterized by containing, by mass %,

C: 0.67% to less than 0.9%,

Si: 2.0 to 3.5%,

Mn: 0.5 to 1.2%,

Cr: 1.3 to 2.5%,

N: 0.003 to 0.007%, and

Al: 0.0005% to 0.003%,

having Si and Cr satisfying the following formula: 0.3%≦Si−Cr≦1.2%,

having a balance of iron and unavoidable impurities,

having P and S as impurities comprising

P: 0.025% or less and

S: 0.025% or less, and,

furthermore,

having a circle equivalent diameter of undissolved spherical carbides of less than 0.2 μm.

2. Pre-drawn steel wire for high strength spring use as set forth in claim 1 characterized by, further, containing, by mass %, one or more of

V: 0.03 to 0.10%,

Nb: 0.015% or less

Mo: 0.05 to 0.30%,

W: 0.05 to 0.30%

Mg: 0.002% or less,

Ca: 0.002% or less, and

Zr: 0.003% or less,

when containing V

satisfying 1.4%≦Cr+V≦2.6% and 0.70%≦Mn+V≦1.3%, and,

when containing Mo and W,

satisfying 0.05%≦Mo+W≦0.5%.

3. Drawn heat treated steel wire for high strength spring use characterized by containing, by mass %,

C: 0.67% to less than 0.9%,

Si: 2.0 to 3.5%,

Mn: 0.5 to 1.2%,

Cr: 1.3 to 2.5%,

N: 0.003 to 0.007%, and

Al: 0.0005% to 0.003%,

having Si and Cr satisfying the following formula: 0.3%≦Si−Cr≦1.2%, and

having a balance of iron and unavoidable impurities,

having P and S as impurities comprising

P: 0.025% or less and

S: 0.025% or less,

furthermore,

having a metal structure comprised of at least residual austenite in a volume rate of over 6% to 15%,

having prior austenite grain size number of #10 or more, and

having a circle equivalent diameter of undissolved spherical carbides of less than 0.2 μm.

4. Drawn heat treated steel wire for high strength spring use as set forth in claim 3 characterized by, further, containing, by mass %, one or more of

V: 0.03 to 0.10%,

Nb: 0.015% or less

Mo: 0.05 to 0.30%,

W: 0.05 to 0.30%

Mg: 0.002% or less,

Ca: 0.002% or less, and

Zr: 0.003% or less,

when containing V

satisfying 1.4%≦Cr+V≦2.6% and 0.70%≦Mn+V≦1.3%, and,

when containing Mo and W,

satisfying 0.05%≦Mo+W≦0.5%.

5. Drawn heat treated steel wire for high strength spring use as set forth in claim 3 characterized in that said drawn heat treated steel wire for high strength spring use has a tensile strength of 2100 to 2400 MPa.

6. Drawn heat treated steel wire for high strength spring use as set forth in claim 5 characterized in that said drawn heat treated steel wire for high strength spring use has a yield strength of 1600 to 1980 MPa.

7. Drawn heat treated steel wire for high strength spring use as set forth in claim 3 characterized in that said drawn heat treated steel wire for high strength spring use has a surface Vicker's hardness of HV750 or more and an internal Vicker's hardness of HV570 or more aftersoft nitriding of keeping at 500° C. for 1 hour.

8. A method of production of pre-drawn steel wire for high strength spring use characterized by taking a bloom containing, by mass %,

C: 0.67% to less than 0.9%,

Si: 2.0 to 3.5%,

Mn: 0.5 to 1.2%,

Cr: 1.3 to 2.5%,

N: 0.003 to 0.007%, and

Al: 0.0005% to 0.003%,

having Si and Cr satisfying the following formula:

0.3%≦Si−Cr≦1.2%,

having a balance of iron and unavoidable impurities,

having P and S as impurities comprising

P: 0.025% or less and

S: 0.025% or less,

heating the bloom to 1250° C. or more, then hot rolling the bloom to produce a billet and heating the billet to 1200° C. or more, then hot rolling to produce pre-drawn steel wire.

9. A method of production of pre-drawn steel wire for high strength spring use as set forth in claim 8 characterized by the bloom further, containing, by mass %, one or more of

V: 0.03 to 0.10%,

Nb: 0.015% or less

Mo: 0.05 to 0.30%,

W: 0.05 to 0.30%

Mg: 0.002% or less,

Ca: 0.002% or less, and

Zr: 0.003% or less,

when containing V

satisfying 1.4%≦Cr+V≦2.6% and 0.70%≦Mn+V≦1.3%, and,

when containing Mo and W,

satisfying 0.05%≦Mo+W≦0.5%.

10. A method of production of pre-drawn steel wire for high strength spring use characterized by further heating pre-drawn steel wire as set forth in claim 8 to 900° C. or more, then patenting it at 600° C. or less.

11. A method of production of heat treated steel wire for high strength spring use characterized by

drawing said pre-drawn steel wire which was produced by the method of production of pre-drawn steel wire as set forth in claim 8,

heating it at a heating rate of 10° C./sec or more up to an A3 point,

holding it at a temperature of the A3 point or more for 1 minute to 5 minutes,

then cooling it at a cooling rate of 50° C./sec or more down to 100° C. or less.

12. A method of production of heat treated steel wire for high strength spring use characterized by

drawing said pre-drawn steel wire which was produced by the method of production of pre-drawn steel wire as set forth in claim 10,

heating it at a heating rate of 10° C./sec or more up to an A3 point,

holding it at a temperature of the A3 point or more for 1 minute to 5 minutes,

then cooling it at a cooling rate of 50° C./sec or more down to 100° C. or less.

13. A method of production of heat treated steel wire for high strength spring use as set forth in claim 11 characterized by further holding and tempering it at 400 to 500° C. for 15 minutes or less.

14. A method of production of heat treated steel wire for high strength spring use as set forth in claim 12 characterized by further holding and tempering it at 400 to 500° C. for 15 minutes or less.

15. Drawn heat treated steel wire for high strength spring use as set forth in claim 4 characterized in that said drawn heat treated steel wire for high strength spring use has a tensile strength of 2100 to 2400 MPa.

16. Drawn heat treated steel wire for high strength spring use as set forth in claim 4 characterized in that said drawn heat treated steel wire for high strength spring use has a surface Vicker's hardness of HV750 or more and an internal Vicker's hardness of HV570 or more aftersoft nitriding of keeping at 500° C. for 1 hour.

17. A method of production of pre-drawn steel wire for high strength spring use characterized by further heating pre-drawn steel wire as set forth in claims 9 to 900° C. or more, then patenting it at 600° C. or less.

18. A method of production of heat treated steel wire for high strength spring use characterized by

drawing said pre-drawn steel wire which was produced by the method of production of pre-drawn steel wire as set forth in claim 9,

heating it at a heating rate of 10° C./sec or more up to an A3 point,

holding it at a temperature of the A3 point or more for 1 minute to 5 minutes,

then cooling it at a cooling rate of 50° C./sec or more down to 100° C. or less.