Cold forging steel having improved resistance to grain coarsening and delayed fracture and process for producing same

A cold forging steel excellent in grain coarsening prevention and delayed fracture resistance and method of producing the same are provided that enable omission of a step of annealing or spheroidization annealing before cold forging and improvement of delayed fracture resistance of a high-strength component used with a heat-treated surface. The cold forging steel is a steel of a specified composition having dispersed in the matrix thereof particles of not greater than 0.2 .mu.m diameter of one or more of TiC, Ti(CN), NbC, Nb(CN) and (Nb, Ti)(CN) in a total number of not less than 20/100 .mu.m.sup.2. The method of producing a cold forging steel includes the steps of heating this steel to not lower than 1050.degree. C., hot-rolling the steel into steel wire or steel bar, and slowly cooling the steel at a cooling rate of not greater than 2 C./s during cooling to a temperature not higher than 600.degree. C. to obtain a steel having dispersed in the matrix thereof particles of not greater than 0.2 .mu.m diameter of one or more of TiC, Ti(CN), NbC, Nb(CN) and (Nb, Ti)(CN) in a total number of not less than 20/100 .mu.m.sup.2.

BACKGROUND OF THE INVENTION
 1. Field of the Invention
 The present invention relates to a cold forging steel excellent in grain
 coarsening prevention and delayed fracture resistance and a method of
 producing the same.
 2. Description of the Related Art
 Cold forging (including roll-forging) is utilized for bolts, gear
 components, shafts and numerous other products because it enables
 fabrication of products with excellent surface quality and dimensional
 precision, is lower in cost than hot forging, and is excellent in yield.
 In the cold forging of such products, use is made of medium-carbon machine
 structural carbon steels and alloy steels such as those specified by S G
 4051, JIS G 4052, JIS G 4104, JIS G 4105, JIS G 4106 and the like. The
 process usually includes a step of annealing or spheroidization annealing
 before the cold forging, in the manner of, for example: hot
 rolling--annealing--cold forging--quench-hardening--tempering. This is
 because the high as-rolled hardness of medium-carbon carbon steels and
 alloy steels like those listed above is a cause of various
 production-related problems, including high cost owing to heavy wear of
 the cold forging tool during the shaping of components such as bolts and
 occurrence of cracking during component shaping owing to the low ductility
 of the blank.
 As annealing involves considerable energy, labor and equipment costs,
 however, a need is felt for a material and process that enable omission of
 the annealing step. This has led to the development of numerous so-called
 low-carbon boron steels that enable omission of the annealing step by
 reducing the carbon and alloying element content of the steel to achieve
 lower as-hot-rolled hardness and improved ductility and that add a small
 amount of boron to make up for the degradation in quench-hardening
 performance caused by the reduced content of Cr, Mo and other alloying
 elements. Such steels are taught by, for example, JP-A-(unexamined
 published Japanese patent application)5-339676, JP-B-(examined published
 Japanese patent application)5-63624 and JP-A-61-253347. Although addition
 of a small amount of boron (B) improves the quench-hardening performance,
 this effect is lost when N is present in the steel in solid solution
 because the B combines with N to form BN. Ordinarily, therefore, Ti is
 added to fix the N in the steel as TiN and thereby suppress formation of
 BN.
 As the need for components with higher strength has increased, attempts
 have been made to apply such low-carbon boron steels to higher strength
 components. Since low-carbon boron steels are low in C and alloying
 elements, however, they sustain a decline in delayed fracture property
 when subjected to heat treatment for achieving a tensile strength of 1000
 MPa or higher. It is known that an attempt to obtain high strength by
 conducting low-temperature tempering results in degraded delayed fracture
 properties. However, when the amount of added C is increased or an SCR,
 SCM or other such alloy steel is used in order to secure high strength and
 bring the delayed fracture strength up to a practical level even with
 high-temperature tempering, the resulting increase in the steel hardness
 makes it impossible to eliminate the annealing step. Although low-carbon
 boron steels that enable omission of annealing are economical, they
 require the tempering temperature to be lowered for obtaining high
 strength. But this degrades the delayed fracture strength and causes
 problems from the practical aspect. Application to high-strength products
 is therefore difficult.
 In response to the call for application of boron steels to high-strength
 components, JP-A-8-60245, for example, teaches a steel reduced in impurity
 content so has to have delayed fracture property on a par with an alloy
 steel. When this boron steel was evaluated using a machined-surface test
 piece, it was in fact found to exhibit a delayed fracture property
 superior to an alloy steel. However, when the steel was used to fabricate
 a component on an actual production line, and the delayed fracture
 property was evaluated from the heat-treated surface condition, it was
 found that the boron steel component was inferior to an alloy steel in
 delayed fracture property. The technology taught by JP-A-8-60245 is
 therefore limited in its ability to respond to the need for higher
 strength components.
 In addition to the foregoing problems, a boron steel is also more likely
 than an annealed steel to sustain abnormal coarsening of specific
 austenite grains during heating for quench-hardening. A component that has
 experienced grain coarsening is liable to have low dimensional precision
 owing to quench-hardening distortion, reduced impact value and fatigue
 life, and, particularly in a high-strength component, degraded delayed
 fracture property. Application of a boron steel to a high-strength
 component therefore requires suppression of grain coarsening and crystal
 grain refinement. For suppressing the grain coarsening, it is effective to
 finely disperse a large quantity of particles that pin grain boundary
 movement.
 Methods have been proposed for preventing the aforesaid grain coarsening of
 boron steel. JP-A-61-217553, for example, aims to pin the grain boundaries
 by defining the Ti and N contents as 0.02&lt;Ti-3.42N so as to generate TiC.
 However, it is not possible to prevent grain coarsening merely by defining
 composition because the TiC cannot be finely dispersed. On the other hand,
 JP-B-63-64495, for instance, aims to prevent grain coarsening by keeping N
 content to a very low value of not greater than 0.0035% and subjecting the
 resulting composition having an excess of Ti relative to N to rolling
 under low-temperature heating. However, prevention of grain coarsening
 cannot be achieved unless the TiC, Ti(CN) precipitation condition is
 optimized before heating for quench-hardening.
 JP-A-52-114545, for example, puts TiC into solid solution at the material
 stage so that fine precipitation of TiC will first occur during heating
 for quench-hardening. When pinning particles precipitate during heating
 for quench-hardening, however, the amount of TiC precipitation is affected
 by the heating rate during heating for quench-hardening or heating for
 carburization. As this makes the expression of the pinning effect unstable
 and, even when the same material is used, a high probability arises of the
 coarsening prevention being degraded by a mere change in component size or
 the heat-treatment furnace. A problem therefore persists regarding quality
 stability in actual production.
 The aforesaid conventional methods cannot achieve a delayed fracture
 property of the actual component equal to or better than that of an alloy
 steel when the annealing or spheroidization annealing step before cold
 forging is omitted and heat treatment is conducted for imparting high
 strength.
 SUMMARY OF THE INVENTION
 An object of this invention is to overcome the aforesaid problems of the
 prior art and to provide a cold forging steel excellent in grain
 coarsening prevention and delayed fracture resistance and method of
 producing the same.
 During their research for achieving this object, the inventors discovered
 the following facts (A)-(D) regarding the effects of various factors on
 the delayed fracture property at the heat-treated surface of an actual
 component.
 (A) That the surface properties of an actual component strongly affect its
 delayed fracture property, specifically that an actual bolt with adhered
 heat-treatment scale (heat-treated surface) and a test piece removed of
 the surface layer by cutting, grinding or other such machining (machined
 surface) exhibit markedly different properties when subjected to delayed
 fracture testing under identical conditions, with the actual component
 with adhered heat-treatment scale exhibiting inferior delayed fracture
 property.
 (B) That delayed fracture property at the heat-treated surface can be
 improved by adding Cr within a certain optimum range so as to cause the
 scale formed during heat treatment of the component to become a dense
 scale enriched in Cr, thereby increasing corrosion resistance so as to
 reduce the amount of hydrogen produced in the process of corrosion of the
 scale and the steel surface inside the scale.
 (C) That when a boron steel is applied to a high-strength component such as
 a bolt having a tensile strength of 1000 MPa or higher, improvement of
 delayed fracture property requires the P and S contents to be limited to
 not more than prescribed values and requires prevention of grain
 coarsening.
 (D) That fine TiC, Ti(CN), NbC, Nb(CN) and (Nb, Ti)(CN) particles are
 effective as pinning particles for preventing grain coarsening, that the
 grain coarsening property is very closely related to the size and
 dispersion state (number of precipitated particles) of these precipitates,
 and that for stably securing the pinning effect of the precipitates it is
 necessary to finely precipitate at least a prescribed amount of particles
 of one or more of TiC, Ti(CN), NbC, Nb(CN) and (Nb, Ti)(CN) before heating
 for quench-hardening.
 The present invention is based on this new knowledge.
 In a first aspect, the present invention enables a marked improvement of
 delayed fracture property after production into an actual component by
 defining content of C as 0.10-0.40%, Si as not more than 0.15% and Mn as
 0.30-1.00% to secure component strength after quench-hardening and
 tempering, limiting content of P to not more than 0.015% (including 0%)
 and S to not more than 0.015% (including 0%) to improve delayed fracture
 property, limiting content of B to 0.0003-0.0050% to secure
 quench-hardenability, and defining content of Cr as 0.50-1.20% to improve
 delayed fracture property at the heat-treated surface. Further, N content
 can be limited to not more than 0.0100% (including 0%) and Ti content be
 defined as 0.020-0.100% to produce TiC and Ti(CN) utilized as pinning
 particles for preventing grain coarsening. By making the total number of
 particles of not greater than 0.2 .mu.m diameter of one or both of TiC and
 Ti(CN) in the matrix not less than 20/100 .mu.m.sup.2, the pinning effect
 can be maximized to provide a cold forging steel enabling prevention of
 grain coarsening during heating for quench-hardening and refinement of old
 austenite grains.
 In a second aspect, the present invention defines, in addition to the
 components of the first aspect, a Nb content of 0.003-0.100% and makes the
 total number of particles of not greater than 0.2 .mu.m diameter of one or
 more of TiC, Ti(CN), NbC, Nb(CN) and (Nb, Ti)(CN) in the matrix not less
 than 20/100 .mu.m.sup.2, thereby providing a cold forging steel enabling
 prevention of grain coarsening.
 In a third aspect, the present invention defines, in addition to the
 components of the first and second aspects, one or both of a V content of
 0.05-0.30% and a Zr content of 0.003-0.100%, thereby enabling further
 refinement of old austenite grains, and makes the total number of
 particles of not greater than 0.2 .mu.m diameter of one or more of TiC,
 Ti(CN), NbC, Nb(CN) and (Nb, Ti)(CN) in the matrix not less than 20/100
 .mu.m.sup.2, thereby providing a cold forging steel enabling prevention of
 grain coarsening.
 In a fourth aspect, the present invention provides a method of producing a
 cold forging steel comprising the steps of heating a steel having the
 composition components of the first, second or third aspect to not lower
 than 1050.degree. C., thereby once causing TiC, Ti(CN), NbC, Nb(CN) and
 (Nb, Ti)(CN) to enter solid solution in the matrix, hot-rolling the steel
 into steel wire or steel bar, softening the steel by slow cooling at a
 cooling rate of not greater than 2.degree. C./s during cooling to a
 temperature not higher than 600.degree. C., and dispersing fine particles
 of not greater than 0.2 .mu.m diameter of one or more of TiC, Ti(CN), NbC,
 Nb(CN) and (Nb, Ti)(CN) in the matrix in a total number of not less than
 20/100 .mu.m.sup.2.

DESCRIPTION OF THE PREFERRED EMBODIMENTS
 The reasons for the limitations on the composition components in the
 present invention will now be explained.
 Carbon (C) is an element effective for imparting strength to the steel.
 When the C content is less than 0.10%, the required tensile strength
 cannot be obtained, and when the C content is greater than 0.40%, the cold
 forgeability is degraded and the annealing or spheroidization annealing
 step before cold forging cannot be omitted. Moreover, since the component
 ductility and toughness are degraded and the delayed fracture property
 also tends to be degraded, the C content must be in the range of
 0.10-0.40%. It is preferably 0.20-0.30%.
 Silicon (Si) is an element effective for deoxidization as well as for
 imparting a required strength and quench-hardenability to the steel and
 improving resistance to temper-softening. However, when present in excess
 of 0.15%, it degrades toughness and ductility. It also degrades cold
 forgeability by increasing hardness. Si content must therefore be kept to
 not greater than 0.15% and is preferably not greater than 0.10%.
 Manganese (Mn) is an element effective for deoxidization as well as for
 imparting a required strength and quench-hardenability to the steel. At a
 content of less than 0.30%, its effect is insufficient, and at a content
 greater than 1.00%, it degrades cold forgeability by increasing hardness.
 Mn content must therefore be in the range of 0.30-1.00% and is preferably
 in the range of 0.40-0.70%.
 Phosphorus (P) is an element that, by increasing resistance to deformation
 and degrading toughness during cold forging, degrades cold forgeability.
 As it also degrades delayed fracture property by embrittling the grain
 boundaries of the component after quench-hardening and tempering, its
 content is preferably made as low as possible. P content must therefore be
 limited to not more than 0.015% and is preferably not more than 0.010%.
 Sulfur (S) is an element that promotes cracking during cold forging and
 therefore degrades cold forgeability. As, like P, it also degrades delayed
 fracture property by embrittling the grain boundaries of the component
 after quench-hardening and tempering, its content is preferably made as
 low as possible. S content must therefore be limited to not more than
 0.015% and is preferably not more than 0.010%.
 Chromium (Cr) is an element effective for imparting strength and
 quench-hardenability to the steel and for improving resistance to
 temper-softening. It is particularly an element that markedly improves
 delayed fracture property at the heat-treated surface. Cr has the effect
 of making the scale formed during heat treatment a dense scale enriched in
 Cr, thereby increasing corrosion resistance so as to reduce the amount of
 hydrogen produced in the process of corrosion of the scale and thus
 improve the delayed fracture property. The effect of Cr content on delayed
 fracture property is shown in FIG. 1 for the case of heat-treatment for
 obtaining a tensile strength of around 1350 MPa.
 Although FIG. 1 shows the test results in 0.1N HCl, substantially the same
 pattern is exhibited in 1% H.sub.2 SO.sub.4. As is clear from FIG. 1, the
 effect of Cr content on delayed fracture property at the heat-treated
 surface is great. A sufficient improvement in delayed fracture property is
 not obtained when the content is less than 0.50%, and when the content
 exceeds 1.2%, the cold forgeability is degraded owing to increased
 hardness, while the delayed fracture property is degraded rather than
 improved owing to promotion of grain boundary oxidation of the surface
 layer formed during heat treatment. This tendency increases with
 increasing component strength. The amount of added Cr must therefore be in
 the range of 0.50-1.20% and is preferably in the range of 0.60-0.90%.
 Boron (B) is an element effective for imparting quench-hardenability to the
 steel when added in a small amount. This effect is insufficient at a
 content of less than 0.0003% and saturates when the content exceeds
 0.0050%. The content must therefore be in the range of 0.0003-0.0050%. The
 preferable range is 0.0010-0.0030%.
 Nitrogen (N) combines with B to form BN. This is deleterious in the case of
 a B-added steel such as that of the present invention because it lowers
 the quench-hardenability improving effect of B. Moreover, when N combines
 with Ti, coarse TiN contributing substantially no pinning effect is formed
 and the amount of Ti available for forming Ti-containing carbonitrides is
 reduced. As this reduces the amount of fine precipitate, the N content is
 preferably made as low as possible. Thus the main aim in keeping the N
 content as low as possible is to control grain coarsening and, as pointed
 out later, the amount of Ti added can be reduced when the N content is
 low. As it is difficult to completely remove N in an actual production
 process, however, the N content is defined as not greater than 0.0100%.
 The preferable range is not greater than 0.0050%.
 Ti (titanium) is an element that, by combining with C and N to form TiC and
 Ti(CN), is effective for grain refinement and suppression of grain
 coarsening. When it is added together with B, formation of BN is
 suppressed because N enters the steel in solid solution in the form of TiN
 and Ti(CN). Ti is therefore an element effective for enhancing the
 quench-hardenability improving effect of B. However, these effects are
 insufficient at a content of less than 0.020% and saturate at a content
 exceeding 0.100%. A content exceeding 0.100% also degrades cold
 forgeability by increasing hardness. The Ti content must therefore be in
 the range of 0.020-0.100%. The preferable range is 0.025-0.50%.
 In order to fix all sol N in the steel in the form of TiN, it is necessary
 to increase the Ti content in accordance with the N content, and in order
 to secure an adequate amount of fine TiC and Ti(CN) effective for grain
 boundary pinning, it is necessary to increase the amount of Ti in
 accordance with the N content. Ti must be added in excess of at least 3.4N
 %.
 Niobium (Nb) is an element that by combining with C and N to form NbC,
 Nb(CN) and (Nb, Ti)(CN) is effective for grain refinement and suppression
 of grain coarsening. When Nb is added together with Ti, almost all of it
 forms stable (Nb, Ti)(CN), whereby a stable pinning effect can be
 obtained. This effect is insufficient at a content of less than 0.003% and
 saturates at a content exceeding 0.100%. A content exceeding 0.100% also
 degrades cold forgeability by increasing hardness. The Nb content must
 therefore be in the range of 0.003-0.100%. The preferable range is
 0.005-0.030%.
 Vanadium (V) is an element that by combining with C and N to form VC and VN
 is effective for grain refinement. This effect is insufficient at a
 content of less than 0.05% and saturates at a content exceeding 0.30%. A
 content exceeding 0.30% also degrades cold forgeability by increasing
 hardness. The V content must therefore be in the range of 0.05-0.30%. The
 preferable range is 0.10-0.20%.
 Zr (zirconium) is an element that by combining with C and N to form ZrC and
 ZrN is effective for grain refinement. This effect is insufficient at a
 content of less than 0.003% and saturates at a content exceeding 0.100%. A
 content exceeding 0.100% also degrades cold forgeability by increasing
 hardness. The Zr content must therefore be in the range of 0.003-0.100%.
 The preferable range is 0.005-0.030%.
 Although V and Zr are not required elements in the present invention, they
 can be added as required for the purpose of grain refinement.
 Although the present invention does not define an amount of Al to be added,
 Al is an element effective for deoxidization of the steel and can
 therefore be included in an amount normally used for deoxidization.
 Ordinarily, the Al content is about 0.010-0.050%. When one or more other
 elements (Si, Mn, Ti, Zr etc.) are added as deoxidizers in place of Al,
 however, addition of Al is not absolutely necessary.
 The dispersed state of TiC, Ti(CN), NbC, Nb(CN) and (Nb, Ti)(CN) in the
 matrix will now be explained.
 For suppressing the grain coarsening, it is effective to finely disperse a
 large quantity of particles for pinning the grain boundaries. A smaller
 particle diameter and larger particle quantity is preferable because it
 increases the number of pinning particles. The relationship between fine
 TiC, Ti(CN) and grain coarsening temperature is shown in FIG. 2. The
 relationship of FIG. 2 also holds for NbC, Nb(CN) and (Nb, Ti)(CN), which
 have similar effect.
 As seen in FIG. 2, the grain coarsening property is very closely related to
 the number of finely precipitated particles. When particles of not greater
 than 0.2 .mu.m diameter of one or more of TiC, Ti(CN), NbC, Nb(CN) and
 (Nb, Ti)(CN) are dispersed in the matrix in a total number of not less
 than 20/100 .mu.m.sup.2, no grain coarsening occurs in the practical
 temperature range of heating for quench-hardening or heating for
 carburization and excellent grain coarsening prevention is obtained. It is
 therefore necessary for particles of not greater than 0.2 .mu.m diameter
 of one or more of TiC, Ti(CN), NbC, Nb(CN) and (Nb, Ti)(CN) to be
 dispersed in the matrix in a total number of not less than 20/100
 .mu.m.sup.2.
 The invention production method will now be explained.
 A steel comprising the aforesaid invention composition components is melted
 in a converter, electric furnace or the like, adjusted in composition, and
 passed through a casting step and, if necessary, a slab rolling step to
 obtain a rolled material. Further improved characteristics can be obtained
 by subjecting the casting to soaking and dispersion treatment before the
 slab rolling step by holding it at a temperature of about
 1,200-1,350.degree. C. for several hours. This is because this treatment
 reduces segregation of P and other impurity elements, thereby further
 improving the delayed fracture property of the actual component, and also
 enables coarse precipitates precipitated in the casting step to be once
 put into solid solution, thereby making it easier for precipitates to
 enter the matrix in solid solution in the following step.
 Next, the rolled material is heated to a temperature of 1050.degree. C. or
 higher. Under heating conditions of a temperature lower than 1050.degree.
 C., TiC, Ti(CN), NbC, Nb(CN) and (Nb, Ti)(CN) cannot once be put into
 solid solution in the matrix, making it impossible to obtain a steel
 having one or more of TiC, Ti(CN), NbC, Nb(CN) and (Nb, Ti)(CN) finely
 precipitated therein after hot rolling. Moreover, when much coarse TiC,
 Ti(CN), NbC, Nb(CN) or (Nb, Ti)(CN) that could not enter solid solution
 remains, it degrades the ductility of the component and has an adverse
 effect on the delayed fracture property.
 When many coarse precipitates are present, moreover, they further promote
 coarsening by acting as precipitation nuclei during cooling after rolling.
 This makes fine dispersion of pinning particles in the matrix difficult.
 The heating temperature is therefore preferably made as high as possible.
 The preferable range is 1150.degree. C. and higher.
 Next, the rolled material heated to 1050.degree. C. or higher is hot-rolled
 into steel wire or steel bar and then slowly cooled at a cooling rate of
 not greater than 2.degree. C./s during cooling to a temperature not higher
 than 600.degree. C. Under cooling conditions exceeding 2.degree. C./s, the
 time period of passage through the precipitation temperature ranges of
 TiC, Ti(CN), NbC, Nb(CN) and (Nb, Ti)(CN) is too short to obtain a
 sufficient amount of precipitation and, as a result, it becomes impossible
 to obtain a steel containing a large quantity of finely precipitated TiC,
 Ti(CN), NbC, Nb(CN) and/or (Nb, Ti)(CN) effective as pinning particles.
 In addition, a rapid cooling rate increases the hardness of the rolled
 material. As this degrades the cold forgeability, the cooling rate is
 preferably made as slow as possible. The preferable range is not greater
 than 1.degree. C./sec. After hot-rolling, cooling to a still lower
 temperature range (500.degree. C. or below) is preferably conducted slowly
 at a cooling rate of 2.degree. C./s. When slow cooling is conducted to a
 low temperature range, the rolled material is further softened and
 improved in cold forgeability.
 EXAMPLE
 The present invention will now be further explained with reference to an
 example.
 Each of molten converter steels of the compositions shown in Table 1 was
 continuously cast, subjected to soaking and dispersion treatment as
 required, and slab-rolled into a 162 mm square rolled material. The rolled
 material was then heated to a temperature not lower than 1050.degree. C.
 and hot-rolled into steel bar or steel wire of a diameter of 5-50 mm. For
 comparison, the heating of a portion was conducted at temperature below
 1050.degree. C. Next, slow cooling was conducted using a heat-retention
 cover installed after the rolling line. For comparison, a portion was not
 subjected to slow cooling.
 To examine the dispersed state of TiC, Ti(CN), NbC, Nb(CN) and/or (Nb,
 Ti)(CN) effective as pinning particles, precipitates present in the steel
 bar or steel wire matrix were sampled by the extraction replica method and
 observed with a transmission electron microscope. Around 20 fields were
 observed at 15,000 magnifications, the total number of 0.2 .mu.m and
 smaller diameter particles of TiC, Ti(CN), NbC, Nb(CN) and (Nb, Ti)(CN)
 per field was counted and converted to number per 100 .mu.m.sup.2.
 The grain coarsening temperature of the steel bar or steel wire produced by
 the foregoing process was determined. The rolled material was drawn at an
 area reduction of 70%, heated for 30 min to 840-1200.degree. C. and
 water-quenched. A cut surface was polished/corroded and the old austenite
 grain diameter was observed to determine the coarse grain forming
 temperature (grain coarsening temperature).
 Quench-hardening of bolts and other actual components is usually conducted
 in the A.sub.C3 -900.degree. C. temperature range. A material with a
 coarse grain forming temperature below 900.degree. C. was therefore
 evaluated as inferior in grain coarsening property. The old austenite
 granularity was measured in conformity with JIS G 0551. About 10 fields
 were observed at 400 magnifications and coarsening was judged to have
 occurred if even one coarse grain of a granularity number of 5 or below
 was present.
 The delayed fracture property of the materials was then investigated. After
 70% cold drawing, the material was machined to obtain a delayed fracture
 test piece with an annular V-notch. The test piece was then imparted with
 1350 MPa class tensile strength by 900.degree. C..times.30 min
 heating/quench-hardening followed by tempering to fabricate a delayed
 fracture test piece with a heat-treated surface closely resembling the
 surface of an actual component. This delayed fracture test piece was
 soaked in 0.1N HCl and the time to fracture under different load stresses
 was measured. The test was continued for a maximum of 200 h and the
 maximum load stress at which fracture did not occur within 200 h was
 determined. The value obtained by dividing the maximum load at which
 fracture did not occur within 200 h by the fracture stress in air was
 defined as the "delayed fracture strength ratio" and used as an index of
 the delayed fracture property.
 The delayed fracture strength ratio of SCM435 currently commonly used for
 1000-1400 MPa class tensile strength components is around 0.5. A material
 having a delayed fracture strength ratio of less than 0.5 was therefore
 evaluated as inferior in delayed fracture property. The granularity of the
 test pieces subjected to the delayed fracture test was investigated. In
 the case of uniform grains, the average granularity of the matrix was
 measured. In the case of mixed grains or when coarse grains were present,
 the granularity number of the largest grain in the observed field was also
 determined. Measurement of old austenite granularity was measured by the
 same method as used to determine the grain coarsening temperature.
 The results of the tests are shown in Tables 2, 3 and 4.
 Symbols N and O in Table 2 indicate comparative examples whose Ti or N
 content is outside the range of the present invention and that are
 therefore inferior in grain coarsening property owing to a deficiency in
 the number of finely precipitated particles of TiC, Ti(CN), NbC, Nb(CN)
 and/or (Nb, Ti)(CN). Symbols V, X and Y indicate comparative examples in
 which TiC, Ti(CN), NbC, Nb(CN) and/or (Nb, Ti)(CN) failed to once enter
 the matrix in solid solution owing to low heating temperature for rolling
 and that are therefore inferior in grain coarsening property because a
 steel having fine precipitates precipitated during cooling after hot
 rolling could not be obtained.
 Symbols W and Z indicate comparative examples that are inferior in grain
 coarsening property owing to a deficiency of fine precipitates caused by
 too high a cooling rate after rolling.
 The delayed fracture properties of the rolled materials of Table 2 when
 adjusted to around 1350 MPa and 1200 MPa are shown in Tables 3 and 4,
 respectively. Symbols P, Q and T in Table 3 indicated comparative examples
 that are inferior in grain coarsening property because the amount of added
 Cr is outside the range of the present invention. Symbols R and S indicate
 comparative examples that are inferior in grain coarsening property
 because the P or S content is outside the range of the present invention.
 The materials that are inferior in grain coarsening property (Symbols N, O,
 V, W, X, Y and Z) are inferior in delayed fracture property owing to the
 formation of coarse particles in the delayed fracture test piece. As the
 tensile strength of the materials in Table 4 is in the neighborhood of
 1200 MPa, their delayed fracture property is better than those in Table 3.
 Steel No. 21 in Table 1 and the material indicated by Symbol U in Tables 2
 and 3 are examples of widely used alloy steels that do not permit
 annealing to be omitted. As can be seen from the tables, the materials
 that satisfy all of the conditions prescribed by the present invention
 exhibit grain coarsening prevention and delayed fracture resistance
 superior to those of the comparative examples.
 When the cold forging steel and the production method of the present
 invention are adopted, the annealing step before cold forging can be
 omitted and the degree of degradation of dimensional precision and the
 amount of reduction of impact value and fatigue strength owing to
 quench-hardening distortion caused by grain coarsening during heat
 treatment are less than in the prior art. In addition, materials can be
 provided for bolts, gear components, shafts and the like that are
 especially superior in delayed fracture property in the actual component
 used with a heat-treated surface.
 TABLE 1
 Steel
 No. C Si Mn P S Cr B Al
 Ti N Others
 Invention 1 0.23 0.05 0.50 0.007 0.004 0.70 0.0020
 0.027 0.036 0.0033
 2 0.24 0.10 0.80 0.001 0.010 0.50 0.0012
 0.020 0.100 0.0037
 3 0.19 0.07 0.48 0.010 0.005 0.89 0.0023
 0.035 0.036 0.0036
 4 0.11 0.15 0.30 0.008 0.001 1.05 0.0050
 0.017 0.032 0.0037
 5 0.38 0.09 0.99 0.005 0.015 0.61 0.0003
 0.043 0.020 0.0013
 6 0.14 0.01 0.35 0.015 0.005 1.20 0.0025
 0.011 0.040 0.0050
 7 0.24 0.08 0.45 0.007 0.007 0.77 0.0015 --
 0.034 0.0031
 8 0.20 0.06 0.44 0.005 0.004 0.66 0.0019
 0.025 0.027 0.0036 Nb: 0.003
 9 0.25 0.06 0.39 0.014 0.002 0.74 0.0025
 0.030 0.026 0.0038 Nb: 0.019
 10 0.19 0.05 0.35 0.009 0.008 0.82 0.0010
 0.035 0.039 0.0032 Nb: 0.010