JP7628445B2 - Cold forging steel - Google Patents

Description

The present invention relates to a hypoeutectoid steel for cold forging. In particular, the present invention relates to a steel that has excellent cold forgeability when cold forging is used for part forming, because spheroidized carbides are uniformly dispersed throughout the steel as a result of spheroidizing annealing.

When processing hypo-eutectoid steel, a low-carbon steel, into parts, cold forging is sometimes chosen from the standpoint of improving yield and production efficiency. In addition, when cold forging, spheroidizing annealing is usually performed to reduce deformation resistance.

The spheroidizing annealing of hypo-eutectoid steel usually involves the following steps: (Step 1) Raise the temperature to above the _A1 point. (Step 2) Hold the temperature above the _A1 point to create a structure consisting of two phases, ferrite and austenite. (Step 3) After that, slowly cool from just above the _A1 point to just below the _A1 point to suppress the austenite to pearlite transformation.
At this time, spheroidal cementite precipitates at the austenite/ferrite interface of the two-phase structure, and after slow cooling to just below point _A1 is completed, a structure is formed in which spheroidal cementite precipitates within the ferrite grains.

Now, in the spheroidized annealed structure of hypo-eutectoid steel, due to the process (step 2) of maintaining the structure in the two-phase region consisting of ferrite and austenite, uneven distribution of spheroidized cementite after slow cooling can inevitably occur. Such uneven distribution of cementite is undesirable from the viewpoint of workability, as it promotes uneven deformation during cold forging and is one of the causes of cracking.

In conventional steel materials with excellent cold forgeability, for example, case-hardened steel materials have been proposed that aim to improve cold forgeability by appropriately controlling the area ratios of ferrite, bainite, and pearlite before spheroidizing annealing to balance the shape, density, and dispersion state of the spheroidized carbides after spheroidizing annealing (see Patent Document 1).

In addition, a steel wire rod or bar for case hardening has been proposed in which fine carbides caused by bainite are dispersed after spheroidizing annealing by ensuring an appropriate amount of bainite area ratio before spheroidizing annealing, thereby improving cracking (see Patent Document 2).

In addition, a case-hardened steel for cold forging has been proposed that, in a structure consisting of two phases, ferrite and pearlite, before spheroidizing annealing, the ferrite and pearlite grains are made smaller, increasing the number of ferrite/pearlite grain boundaries and promoting cementite precipitation at the interfaces, thereby shortening the spheroidizing annealing time and uniformly dispersing cementite, thereby ensuring cold forgeability (see Patent Document 3).

JP 2014-185389 A JP 2005-220377 A Japanese Patent Application Publication No. 11-012684

In the spheroidized annealed structure of hypo-eutectoid steel, uneven distribution of spheroidized cementite after slow cooling can be unavoidable due to the process (step 2) described above, in which the structure is maintained in the two-phase region consisting of ferrite and austenite. Such uneven distribution of cementite promotes uneven deformation during cold forging, which is one of the causes of cracking, so from the viewpoint of workability, it is desirable to have a uniform distribution of cementite.

The above-mentioned proposals in the prior art attempt to optimize the shape and dispersion state of spheroidized cementite after spheroidizing annealing, but all of them are merely proposals to appropriately control the structure before spheroidizing annealing. In other words, when looking at the inside of ferrite grains in the two-phase structure consisting of ferrite phase and austenite phase in the normal spheroidizing annealing process (step 2), spheroidized cementite cannot be precipitated, so it cannot be said that they adequately address the issue of avoiding non-uniform distribution of spheroidized cementite.

Therefore, the problem that this invention aims to solve is to provide a steel that can suppress cracking during cold forging by uniformly dispersing spheroidized carbides throughout the structure.

The inventors of the present application have conducted extensive research focusing on the C diffusion behavior just below the austenitizing temperature. In conventional hypoeutectoid steels, solute C in ferrite grains stably precipitates as _M3C -type carbides (cementite) just below the austenitizing temperature. Therefore, C exists as M3C-type carbides in both pearlite grains (structures composed of ferrite and lamellar cementite) and bainite grains (structures composed of ferrite and cementite) and ferrite grains (cementite is stably _precipitated ). As a result, it is difficult to create a difference in C diffusion behavior, and sufficient driving force for C diffusion from pearlite grains and bainite grains to ferrite grains cannot be obtained.

On the other hand, by using a high Cr content, the solute C in the ferrite grains stably precipitates as _M7C3 type (M = a mixture of Fe and Cr) carbides just below the austenitization temperature. In other words, the difference in C diffusion behavior between pearlite grains (structures composed of ferrite and lamellar cementite) and bainite _grains (structures composed of ferrite and _cementite ) and ferrite grains ( _M7C3 carbides are stable) becomes large. As a result, it was found that the C diffusion from pearlite grains and bainite grains to ferrite grains is sufficiently promoted.

The inventors discovered that by adjusting the composition of the steel, the stable carbides within the ferrite grains just below the austenitizing temperature are made to be of the M ₇ C ₃ type, which makes it easier to uniformly disperse spheroidized carbides.

In this way, the inventors of the present application have discovered that by subjecting steel with a specified composition to a specified spheroidizing annealing process, spheroidized carbides are uniformly dispersed throughout the structure, resulting in a steel that can suppress cracking during cold forging.

Then, when the structure before the heat treatment is a two-phase structure of ferrite and pearlite, a two-phase structure of ferrite and bainite, or a three-phase structure of ferrite, pearlite, and bainite, (a) during holding just below the austenitizing temperature, the lamellar cementite constituting the pearlite and the cementite constituting the bainite are dissolved and spheroidized, (b) the dissolved C flows into the ferrite grains, (c) _M7C3 type carbides precipitate in the ferrite grains, and (d) these grow into spheroidized carbides, thereby making it possible to obtain a steel in which spheroidized carbides are uniformly dispersed throughout. Here, _{the M7C3 type} carbides are a mixture of M ₌ Fe and Cr.

In view of the above, the first means for solving the problems of the present invention is a cold forging steel consisting, in mass%, of 0.14-0.45% C, 0.05-1.00% Si, 0.10-0.90% Mn, 1.30-3.50% Cr, 0.020-0.200% Al, 0.0040-0.0300% N, with the balance being Fe and unavoidable impurities, with P and S as the unavoidable impurities being P: 0.030% or less and S: 0.030% or less, with an _A1 point of 750°C or more and having _{M7C3 type} carbides (wherein M in the _{M7C3 type} carbides is Fe or Cr).

The second means is a cold forging steel which contains, in addition to the chemical components described in the first means, one or more of Ni: 0.02-2.00%, Mo: 0.02-2.00%, with the balance being Fe and unavoidable impurities, in which the unavoidable impurities are P: 0.030% or less and S: 0.030% or less, has an _A1 point of 750°C or more, and has _M7C3 type _carbides (wherein M in _{M7C3 type} carbides is Fe or Cr).

The third means is a cold forging steel which contains, in addition to the chemical components described in the first or second means, one or more of Nb: 0.02-0.10%, Ti: 0.020-0.200%, B: 0.0010-0.0050%, V: 0.010-0.500%, with the balance being Fe and unavoidable impurities, in which the unavoidable impurities of P and S are P: 0.030% or less and S: 0.030% or less, has an _A1 point of ₇₅₀ °C or more, and has _M7C3 type carbides (wherein M in _{M7C3 type} carbides is Fe or Cr).

In the cold forging steel of the present invention, _{M7C3 carbides} are precipitated in the prior ferrite grains, so that the spheroidized carbides are uniformly dispersed throughout the structure. In addition, the area ratio of the ferrite grains in which the spheroidized carbides are not precipitated is within 6.0% of the entire structure, and the area ratio of the remaining pearlite grains and bainite grains after spheroidizing annealing is also 5.0% or less, so that the spheroidization of the carbides is sufficiently advanced. As a result, the steel does not crack even under high compression by cold forging. Therefore, the cold forging steel of the present invention has a hardness of 83 HRB or less, and can be cold worked without cracking even when cold forged at a cold upsetting rate of 75%, so that the cold forging steel of the present invention shows excellent cold forgeability.

FIG. 2 is a schematic diagram for explaining the temperature history of the holding temperature T (° C.), holding time (hr), and air cooling in the spheroidizing annealing conditions. The photographic image taken with an optical microscope is shown, along with the area ratio of ferrite grains without precipitation of spheroidized carbides, and the area ratio of remaining pearlite and bainite.

Before describing the embodiment of the steel according to the present invention, _the reasons for specifying the composition of the steel, the reasons for specifying the lower limit of the _A1 point, and the reasons for having _M7C3 type carbides will be described. Note that % in the following composition is mass %.

C: 0.14-0.45%
C is an element that improves the hardness of the material. If the C content is less than 0.14%, the core hardness after carburization decreases, resulting in insufficient strength. On the other hand, if the C content exceeds 0.45%, the material hardness increases too much, resulting in a decrease in workability, and thus the machinability and cold forgeability become poor.
Therefore, the C content is set to 0.14 to 0.45%, and preferably 0.14 to 0.30%.

Si: 0.05-1.00%
Silicon is an element useful for deoxidization. If the silicon content is less than 0.05%, there will be a shortage of deoxidizing material. On the other hand, if the silicon content exceeds 1.00%, the material hardness will increase too much, resulting in a decrease in workability, and carburization will be inhibited.
Therefore, the Si content is set to 0.05 to 1.00%, and preferably, the Si content is set to 0.30 to 0.80%.

Mn: 0.10-0.90%
Mn is an element that improves hardenability. If the Mn content is less than 0.10%, hardening is insufficient. On the other hand, if the Mn content exceeds 0.90%, workability is reduced.
Therefore, the Mn content is 0.10 to 0.90%, and preferably 0.15 to 0.50%.

Cr: 1.30-3.50%
Cr is an element that stabilizes _{M7C3 type} carbides. If the Cr content is less than 1.30%, _{the M7C3} type carbides do not precipitate, so the amount of C that flows from pearlite grains to ferrite grains during spheroidizing annealing is insufficient, and the distribution of spheroidized carbides becomes non-uniform, resulting in a decrease in cold forgeability. In addition, hardenability is insufficient. On the other hand, if the Cr content exceeds 3.50%, the material hardness increases too much, resulting in a decrease in workability. In addition, carburization inhibition occurs. Therefore, the Cr content is set to 1.30 to 3.50%. Desirably, the Cr content is 1.40 to 2.50%.

Al: 0.020-0.200%
Al is a useful element for deoxidization. If Al is less than 0.020%, it will be insufficient as a deoxidizer. In addition, since there will be a shortage of fine nitrides, there will be a shortage of precipitation nuclei for _M7C3 type carbides, and the toughness and _fatigue properties will decrease due to the coarsening of crystal grains. On the other hand, if Al exceeds 0.200%, coarse nitrides will be formed, which will decrease the fatigue properties and workability. Therefore, Al is set to 0.020 to 0.200%. Desirably, Al is set to 0.020 to 0.050%.

N:0.0040~0.0300%
N is an element that forms nitrides. If N is less than 0.0040%, fine nitrides are insufficient, resulting in a shortage of precipitation nuclei _{for M7C3} , and the toughness and fatigue properties are reduced due to the coarsening of crystal grains. On the other hand, if N exceeds 0.0300%, coarse carbonitrides are formed, resulting in reduced fatigue properties and workability. Therefore, N is set to 0.0040-0.0300%. Desirably, N is set to 0.0040-0.0200%.

P: 0.030% or less as an inevitable impurity If the P content exceeds 0.030%, the toughness decreases due to grain boundary segregation. Therefore, even when P is contained as an inevitable impurity, the P content is set to 0.030% or less.

S: 0.030% or less as an inevitable impurity If S exceeds 0.030%, MnS is formed, which reduces toughness and fatigue strength. Therefore, even when S is contained as an inevitable impurity, the content is set to 0.030% or less.

The following components describe optional ingredients that may be optionally added.
Ni: 0.02-2.00%
Ni is an element that improves hardenability and toughness. If Ni is less than 0.02%, the effect of improving hardenability is small, and the effect of improving toughness is also small. On the other hand, if Ni is added in an amount exceeding 2.00%, not only will the cost increase, but the material hardness will also increase, resulting in a decrease in workability. Therefore, when Ni is added, it is set to 0.02 to 2.00%. Desirably, the amount of Ni added is 0.02 to 1.80%.

Mo: 0.02-2.00%
Mo is an element that improves hardenability. If Mo is less than 0.02%, the effect of improving hardenability is small. On the other hand, if Mo exceeds 2.00%, not only does the cost increase, but the material hardness also increases, resulting in a decrease in workability. Therefore, when Mo is added, it is set to 0.02 to 2.00%. Desirably, the amount of Mo added is 0.02 to 0.50%.

Nb: 0.02-0.10%
Nb is an element that generates carbonitrides and suppresses grain coarsening. If the Nb content is less than 0.02%, the fine carbonitrides are insufficient, so the effect of suppressing grain coarsening is reduced, and toughness and fatigue strength are insufficient. On the other hand, if the Nb content exceeds 0.10%, the amount of carbonitrides becomes excessive and workability is reduced. Therefore, when Nb is added, it is set to 0.02 to 0.10%. Desirably, the amount of Nb added is 0.02 to 0.08%.

Ti: 0.020-0.200%
Ti is an element that generates carbonitrides and suppresses grain coarsening. If Ti is less than 0.020%, the amount of fine nitrides is insufficient. In addition, N is not fixed and forms BN, which reduces hardenability. In addition, the effect of suppressing grain coarsening is reduced.
On the other hand, if the Ti content exceeds 0.200%, the amount of carbonitrides becomes excessive, resulting in a decrease in workability.
Therefore, when Ti is added, the content is set to 0.020 to 0.200%, and preferably, the content of Ti added is set to 0.020 to 0.100%.

B: 0.0010-0.0050%
B is an element that improves hardenability. If B is less than 0.0010%, the effect of improving hardenability is small. If B exceeds 0.0050%, the workability decreases due to an increase in material hardness. Therefore, when B is added, it is set to 0.0010 to 0.0050%. Desirably, the amount of B added is 0.0010 to 0.0030%.

V:0.010~0.500%
V is an element that generates carbonitrides and suppresses grain coarsening. If V is less than 0.010%, fine carbonitrides are insufficient and the effect of suppressing grain coarsening is small, resulting in insufficient toughness and fatigue strength. On the other hand, if V exceeds 0.500%, the amount of carbonitrides becomes excessive and workability decreases. Therefore, when V is added, it is set to 0.010 to 0.500%. Desirably, the amount of V added is 0.010 to 0.400%.

_A1 point: 750°C or higher If _A1 point is less than 750°C, the holding temperature for spheroidizing annealing will be low, and the spheroidization of the carbides that compose pearlite and bainite will be insufficient. As a result, the cold workability will decrease due to insufficient softening.
Therefore, the _A1 point is set to be 750° C. or higher. More preferably, the _A1 point is set to be 750 to 800° C.
Incidentally, A ₁ in the present invention can be calculated using Ac ₁ =723° C.-14 Mn [%]+22 Si [%]-14.4 Ni [%]+23.3 Cr [%].

Presence of _M7C3 type carbides If the diffusion of carbon from pearlite and bainite grains to _ferrite grains during spheroidizing annealing is insufficient, carbides cannot precipitate in the ferrite grains, resulting in non-uniform carbide distribution and reduced cold forgeability. Therefore, in the present invention, _{M7C3 type} carbides are allowed to exist.
This results in a large difference in C diffusion behavior between pearlite grains (a structure composed of ferrite and lamellar cementite) and bainite grains (a structure composed of ferrite and _cementite ) on the one hand, and ferrite grains (in which _M7C3 carbide is stable) on the other. As a result, C diffusion from pearlite grains and bainite grains to ferrite grains is sufficiently promoted, making it easier for spheroidized carbides to be uniformly dispersed, resulting in excellent cold forgeability.

(Spheroidizing Annealing Conditions)
1 is a schematic diagram for explaining the holding temperature T (°C) and holding time t (hr) of the temperature history which are the spheroidizing annealing conditions. The holding temperature T and holding time t conditions for spheroidizing annealing in the present invention shall satisfy the following [Formula 1] to [Formula 3].
[Formula 1]: (A ₁ point - 30°C) ≦T ≦ (A ₁ point - 5°C)
[Formula 2]: t≧120/(TA ₁ +50)
[Formula 3]: A ₁ = 723°C - 14Mn [%] + 22Si [%] - 14.4Ni [%] + 23.3Cr [%]

(Manufacturing process of test material)
Steels having the chemical compositions shown in the invention steels No. 1 to 22 and the comparative steels No. 1 to 11 in Table 1, with the balance being Fe, were melted in a 100 kg vacuum induction furnace (VIM) to produce steel ingots, which were then forged to φ32 at 1250°C, normalized at 925°C for 1 hour, and then cut to 100 mm to obtain test specimens.

(Spheroidizing annealing process)
The test materials prepared above were subjected to spheroidizing annealing using a Kanthal furnace according to the following procedure: The test materials were placed in a furnace set to the holding temperature for spheroidizing annealing as shown in Tables 3 and 4, and the test materials were allowed to heat up for 30 minutes, held for a specified time, and then air-cooled.

(Check for the presence of _M7C3 type _carbides )
For each of the spheroidized annealed test materials, a TEM sample was prepared from the D/4 position by the extraction replica method, and the samples were then observed using a transmission electron microscope (TEM).

(TEM Observation Conditions)
For each test material sample, a total of 10 fields of view (10 μm × 10 μm) were photographed, and 5 carbides present in each field of view were randomly selected and subjected to identification analysis (component analysis by EDS, structural analysis by diffraction pattern) to confirm the presence or absence of _{M7C3 carbide} . The results are shown in Tables 3 and 4.

(Measurement of the area ratio of ferrite grains in which no spheroidized carbides are precipitated)
After spheroidizing annealing, the test pieces were cut out from the D/4 position, embedded in resin, polished, and the surface was etched with a nital etching solution. Five fields of view were then observed under an optical microscope at 400x magnification, and the area ratio was calculated by image analysis.
In addition, when the number of spheroidized carbides having an equivalent circle diameter of 0.5 μm or more present in a ferrite grain is less than 5, the ferrite grain is defined as "a ferrite grain in which no spheroidized carbides are precipitated."
As a result of measuring the area ratio of ferrite grains in which no spheroidized carbides were precipitated in the observation field after spheroidizing annealing, it can be said that those in which the area ratio exceeded 6.0% of the entire structure were insufficiently spheroidized or were not uniformly dispersed. The results are shown in Tables 3 and 4.

(Measurement of area ratio of remaining pearlite grains and bainite grains)
Similarly, the test material after spheroidizing annealing was observed in five fields of view using an optical microscope at a magnification of 400 times, and calculations were made by image analysis.
When the area ratio of the remaining pearlite grains and bainite grains after spheroidizing annealing is 5.0% or less, it can be said that the spheroidization of the carbides has progressed sufficiently, but when the area ratio of the remaining pearlite grains and bainite grains exceeds 5.0%, it can be said that the spheroidization of the carbides is insufficient, and the cold forgeability is reduced due to insufficient softening. The results are shown in Tables 3 and 4.

Figure 2 shows optical microscope photographs of invention steel 6, invention steel 21, and comparative steel 6 after spheroidizing annealing, as well as the area ratio of ferrite grains in which no spheroidized carbides have precipitated, and the area ratios of remaining pearlite grains and bainite grains.

(Hardness measurement)
The hardness of the spheroidized annealed steel bar material at the midpoint between the center and the outer periphery (referred to as "D/4", where D represents the diameter) was measured using a Rockwell hardness tester and taken as the hardness (HRB).
If the spheroidized annealing hardness exceeds 83 HRB, the cold forging deformability decreases, and cracks are likely to occur when cold forging is performed at a high processing rate. Therefore, it is preferable that the spheroidized annealing hardness is 83 HRB or less. The results are shown in Tables 3 and 4.

(Presence or absence of cracks at 75% cold upsetting)
Cylindrical test pieces measuring φ14×21 mm were prepared from the center of the test material (the forging direction of the test material was parallel to the longitudinal direction of the test piece). Using these cylindrical test pieces, a cold upsetting test was carried out at room temperature. The upsetting speed was 10 mm/min, and the final upsetting rate (the reduction rate of the height of the test piece) was 75%, and the test pieces were subjected to four tests (n=4) to observe whether there were any surface cracks after processing. The results are shown in Tables 3 and 4.

The invention steels 1 to 22 are within the range of chemical composition specified by the present invention, have _A1 points of 750°C or higher, and _{M7C3 carbides} are confirmed within the spheroidized annealed grains, but the area ratio of ferrite grains without precipitated spheroidized carbides is within 6.0% of the entire structure, and the area ratio of remaining pearlite grains and bainite grains is also within 5.0%, so that cold forged steels with sufficient spheroidization and uniform dispersion are obtained. That is, the invention steels 1 to 22 all have hardness of 83HRB or less, and did not show any cracks in a 75% cold upset test, so that the invention steels have excellent cold workability.

Comparative steels 1 and 2 had an excessive amount of Mn and an insufficient amount of Cr, the _A1 point was below 750° C., no _{M7C3 carbide} was observed, and cracks were observed in the 75% cold upset test.
Comparative steels 3 and 5 to 7 had an insufficient Cr content, the A ₁ point was below 750° C., no M ₇ C ₃ carbide was observed, and cracks were observed in the 75% cold upset test.
Comparative steel 4 had insufficient Cr and Al, the A ₁ point was below 750° C., no M ₇ C ₃ carbide was observed, and cracks were observed in the 75% cold upset test.
Comparative Example 8 had an excessive amount of Cr, was hard at 84 HRB, and cracks were observed in the 75% cold upset test.
Comparative Example 9 had excessive amounts of Mn and Cr, was hard at 85 HRB, and cracks were observed in the 75% cold upset test.
Comparative Example 10 had an excessive amount of Si, was hard at 85 HRB, and cracks were observed in the 75% cold upset test.
Comparative Example 11 had excessive amounts of C and Mo, was hard at 86 HRB, and cracks were observed in the 75% cold upset test.

Claims (3)

In mass percent,
C: 0.14-0.45%,
Si: 0.05-1.00%,
Mn: 0.10-0.90%,
Cr: 1.30-3.50%,
Al: 0.020-0.200%,
N: 0.0040-0.0300%,
The balance is Fe and unavoidable impurities.
The P and S contents as inevitable impurities are P: 0.030% or less and S: 0.030% or less,
A ₁ point is 750°C or higher,
_M7C3 type carbide _- containing cold forging steel.
(However, M in _M7C3 type _carbide is Fe or Cr.) In addition to the chemical components described in claim 1, the alloy contains at least one of Ni: 0.02 to 2.00% and Mo: 0.02 to 2.00%,
The balance is Fe and unavoidable impurities.
The P and S contents as inevitable impurities are P: 0.030% or less and S: 0.030% or less,
A ₁ point is 750°C or higher,
_M7C3 type carbide _- containing cold forging steel.
(However, M in _M7C3 type _carbide is Fe or Cr.) In addition to the chemical components according to claim 1 or 2,
Contains one or more of Nb: 0.02 to 0.10%, Ti: 0.020 to 0.200%, B: 0.0010 to 0.0050%, and V: 0.010 to 0.500%,
The balance is Fe and unavoidable impurities.
The P and S contents as inevitable impurities are P: 0.030% or less and S: 0.030% or less,
A ₁ point is 750°C or higher,
_M7C3 type carbide _- containing cold forging steel.
(However, M in _M7C3 type _carbide is Fe or Cr.)

Images