U.S. patent application number 13/138837 was filed with the patent office on 2012-01-26 for case-hardened steel superiorin cold workability, machinability, and fatigue characteristics after carburized quenching and method of production of same.
Invention is credited to Masayuki Hashimura, Shuji Kozawa, Manabu Kubota, Kei Miyanishi, Tatsuro Ochi.
Application Number | 20120018063 13/138837 |
Document ID | / |
Family ID | 42935858 |
Filed Date | 2012-01-26 |
United States Patent
Application |
20120018063 |
Kind Code |
A1 |
Hashimura; Masayuki ; et
al. |
January 26, 2012 |
CASE-HARDENED STEEL SUPERIORIN COLD WORKABILITY, MACHINABILITY, AND
FATIGUE CHARACTERISTICS AFTER CARBURIZED QUENCHING AND METHOD OF
PRODUCTION OF SAME
Abstract
Cold worked, machined, and carburized quenched case-hardened
steel prevented from formation of coarse grains, that is,
case-hardened steel superior in cold workability, machinability,
and fatigue characteristics after carburized quenching
characterized by limiting, by mass %, S: 0.001 to 0.15%, Ti: 0.05
to 0.2%, Al: 0.04% or less, and N: 0.0050% or less, containing
other specific ingredients in specific ranges, furthermore
containing one or more of Mg: 0.003% or less, Zr: 0.01% or less,
and Ca: 0.005% or less, limiting the amount of precipitation of AlN
to 0.01% or less, and having a density d (/mm.sup.2) of sulfides
with a equivalent circle diameter of over 20 .mu.m and an aspect
ratio of over 3 and a content of S [S] (mass %) satisfying
d.ltoreq.1700[S]+20.
Inventors: |
Hashimura; Masayuki; (Tokyo,
JP) ; Miyanishi; Kei; (Tokyo, JP) ; Kozawa;
Shuji; (Tokyo, JP) ; Kubota; Manabu; (Tokyo,
JP) ; Ochi; Tatsuro; (Tokyo, JP) |
Family ID: |
42935858 |
Appl. No.: |
13/138837 |
Filed: |
October 14, 2009 |
PCT Filed: |
October 14, 2009 |
PCT NO: |
PCT/JP2009/068083 |
371 Date: |
October 5, 2011 |
Current U.S.
Class: |
148/624 ;
148/328 |
Current CPC
Class: |
C22C 38/00 20130101;
C22C 38/02 20130101; C22C 38/26 20130101; C22C 38/38 20130101; C22C
1/02 20130101; C22C 38/06 20130101; C21D 1/06 20130101; C22C 38/28
20130101; C22C 38/001 20130101; C22C 38/002 20130101; C22C 38/48
20130101; C22C 38/32 20130101; C21D 8/0226 20130101; C22C 1/002
20130101; C22C 38/46 20130101; C23C 8/80 20130101; C22C 38/04
20130101; C22C 38/22 20130101; C22C 38/24 20130101; C23C 8/32
20130101 |
Class at
Publication: |
148/624 ;
148/328 |
International
Class: |
C21D 8/00 20060101
C21D008/00; C22C 38/38 20060101 C22C038/38; C22C 38/26 20060101
C22C038/26; C22C 38/22 20060101 C22C038/22; C22C 38/58 20060101
C22C038/58; C22C 38/32 20060101 C22C038/32; C22C 38/44 20060101
C22C038/44; C22C 38/46 20060101 C22C038/46; C22C 38/50 20060101
C22C038/50; C22C 38/54 20060101 C22C038/54; C22C 38/28 20060101
C22C038/28; C22C 38/24 20060101 C22C038/24 |
Foreign Application Data
Date |
Code |
Application Number |
Apr 6, 2009 |
JP |
2009-092176 |
Claims
1. Case-hardened steel superior in cold workability, machinability,
and fatigue characteristics after carburized quenching
characterized by containing, by mass %, C: 0.1 to 0.5%, Si: 0.01 to
1.5%, Mn: 0.3 to 1.8%, S: 0.001 to 0.15%, Cr: 0.4 to 2.0%, and Ti:
0.05 to 0.2%, limiting Al: 0.04% or less, N: 0.0050% or less, P:
0.025% or less, O: 0.0025% or less, further having one or more of
Mg: 0.003% or less, Zr: 0.01% or less, and Ca: 0.005% or less,
having a balance of iron and unavoidable impurities, limiting an
amount of precipitation of MN to 0.01% or less, and having a
density d (/mm.sup.2) of sulfides of a equivalent circle diameter
of over 20 .mu.m and an aspect ratio of over 3 and a content of S
[S] (mass %) satisfying d.ltoreq.1700[S]+20.
2. Case-hardened steel superior in cold workability, machinability,
and fatigue characteristics after carburized quenching as set forth
in claim 1, characterized by further containing, by mass %, Nb:
less than 0.04%.
3. Case-hardened steel superior in cold workability, machinability,
and fatigue characteristics after carburized quenching as set forth
in claim 1, characterized by further containing, by mass %, one or
more of Mo: 1.5% or less, Ni: 3.5% or less, V: 0.5% or less, and B:
0.005% or less.
4. Case-hardened steel superior in cold workability, machinability,
and fatigue characteristics after carburized quenching as set forth
in claim 1, characterized by limiting a structural fraction of
bainite to 30% or less.
5. Case-hardened steel superior in cold workability, machinability,
and fatigue characteristics after carburized quenching as set forth
in claim 1, characterized in that a grain size number of ferrite is
8 to 11 as defined by JIS G 0551.
6. Case-hardened steel superior in cold workability, machinability,
and fatigue characteristics after carburized quenching as set forth
in claim 1, characterized in that a maximum size of Ti precipitates
is 40 .mu.m or less.
7. A method of production of case-hardened steel superior in cold
workability, machinability, and fatigue characteristics after
carburized quenching characterized by heating a steel material
comprised of the ingredients of claim 1 to 1150.degree. C. or more,
hot working it at a finishing temperature of 840 to 1000.degree.
C., and cooling it in a 800 to 500.degree. C. temperature range by
1.degree. C./s or less.
Description
TECHNICAL FIELD
[0001] The present invention relates to case-hardened-steel
produced by hot rolling, hot forging, or other hot working, then
cold forged, rolled, or otherwise cold worked, cut, etc., then
treated by carburized quenching and a method of production of the
same.
BACKGROUND ART
[0002] Gears, bearings, and other rolling parts and constant
velocity joints, shafts, and other rotation transmission parts
require surface hardness, so are treated by carburized quenching.
These carburized parts are, for example, produced by the process of
using medium carbon alloy steel for machine structures prescribed
by JIS G 4052, JIS G 4104, JIS G 4105, JIS G 4106, etc. and hot
forging, warm forging, cold forging, rolling, or otherwise plastic
working it or cutting it to obtain a predetermined shape, then
treating it by carburized quenching.
[0003] When producing carburized parts, the heat treatment strain
arising due to the carburized quenching sometimes causes the shape
precision of the parts to degrade. In particular, with gears,
constant velocity joints, or other parts, the heat treatment strain
becomes a cause of noise or vibration. Furthermore, it sometimes
causes a deterioration of fatigue characteristics at the contact
surfaces.
[0004] Further, with a shaft etc., if the distortion due to heat
treatment strain becomes large, the efficiency of transmission of
power or the fatigue characteristics are impaired. The biggest
reason for this heat treatment strain is the coarse grains formed
unevenly due to the heating at the time of carburized
quenching.
[0005] In the past, annealing was performed after forging and
before carburized quenching so as to suppress the formation of
coarse grains. However, if annealing, the increase in production
costs becomes an issue.
[0006] Further, gears, bearings, and other rolling parts are
subjected to high surface pressures, so are treated by deep
carburization. With deep carburization, to shorten the
carburization time, usually the 930.degree. C. or so carburization
temperature is raised to a 990 to 1090.degree. C. temperature
region. For this reason, with deep carburization, coarse grains
easily form.
[0007] To suppress the formation of coarse grains at the time of
carburized quenching, the quality of the case-hardened steel, that
is, the material before plastic working, is important.
[0008] To suppress coarsening of the crystal grains at a high
temperature, fine precipitates are effective. Case-hardened steel
utilizing Nb and Ti precipitates, AlN, etc. have been proposed (for
example, Patent Literatures 1 to 5).
CITATION LIST
Patent Literature
[0009] PTL 1: Japanese Patent Publication (A) No. 11-335777 [0010]
PTL 2: Japanese Patent Publication (A) No. 2001-303174 [0011] PTL
3: Japanese Patent Publication (A) No. 2004-183064 [0012] PTL 4:
Japanese Patent Publication (A) No. 2004-204263 [0013] PTL 5:
Japanese Patent Publication (A) No. 2005-240175
SUMMARY OF INVENTION
Technical Problem
[0014] However, if utilizing fine precipitates to suppress the
formation of coarse grains, precipitation strengthening will cause
the case-hardened steel to harden. Further, the addition of alloy
elements for forming precipitates will also cause the case-hardened
steel to harden. For this reason, with steel prevented from forming
coarse grains at a high temperature, the deterioration of cold
forgeability, cutting, and other cold workability became a new
issue.
[0015] In particular, cutting is working requiring a high precision
close to the final shape. A slight rise in hardness has a great
effect on the precision. Therefore, when using case-hardened steel,
it is extremely important not only to prevent the formation of
coarse grains, but to also consider the machinability (ease of
cutting of material).
[0016] In the past, to improve the machinability, it has been known
to be effective to add Pb, S, and other elements improving the
machinability.
[0017] However, Pb is a substance having an environmental load. Due
to the importance of environmentally friendly technology, addition
of Pb to steel materials is being limited.
[0018] Further, S forms MnS etc. in the steel to improve the
machinability, but the coarse MnS inclusions elongated by the hot
working become origin of fracture. For this reason, addition of a
large amount of S can easily become a cause of a deterioration of
cold forgeability or rolling contact fatigue or other mechanical
properties.
[0019] The present invention, in view of this situation, prevents
the formation of coarse grains in case-hardened steel which is
forged, rolled, or otherwise cold worked, cut, and treated by
carburized quenching such as in carburized parts in which fatigue
characteristics are demanded, in particular bearing parts, rolling
parts, etc. in which rolling contact fatigue characteristics are
demanded, and provides case-hardened steel superior in cold
workability, machinability, and fatigue characteristics after
carburized quenching and a method of production of the same.
Solution to Problem
[0020] If treating steel to which Ti has been added by carburized
quenching, Ti precipitates will form origin of fatigue fracture and
the fatigue characteristics, in particular the rolling contact
fatigue characteristic, will easily be degraded. However, if
limiting the content of N and raising the hot rolling temperature
etc. so as to cause the Ti precipitates to finely disperse,
achievement of both prevention of coarse grains and good fatigue
characteristics is possible. Furthermore, for improvement of the
machinability, it is important to add S and add one or more of Mg,
Zr, and Ca to control the size and shape of the sulfides.
[0021] The gist of the present invention is as follows.
[0022] (1) Case-hardened steel superior in cold workability,
machinability, and fatigue characteristics after carburized
quenching characterized by containing, by mass %, [0023] C: 0.1 to
0.5%, [0024] Si: 0.01 to 1.5%, [0025] Mn: 0.3 to 1.8%, [0026] S:
0.001 to 0.15%, [0027] Cr: 0.4 to 2.0%, and [0028] Ti: 0.05 to
0.2%, [0029] limiting [0030] Al: 0.04% or less, [0031] N: 0.0050%
or less, [0032] P: 0.025% or less, [0033] O: 0.0025% or less,
[0034] further having one or more of [0035] Mg: 0.003% or less,
[0036] Zr: 0.01% or less, and [0037] Ca: 0.005% or less, [0038]
having a balance of iron and unavoidable impurities, [0039]
limiting an amount of precipitation of AlN to 0.01% or less, and
[0040] having a density d (/mm.sup.2) of sulfides of a equivalent
circle diameter of over 20 .mu.m and an aspect ratio of over 3 and
a content of S [S] (mass %) satisfying
[0040] d.ltoreq.1700[S]+20.
[0041] (2) Case-hardened steel superior in cold workability,
machinability, and fatigue characteristics after carburized
quenching as set forth in the above (1), characterized by further
containing, by mass %, [0042] Nb: less than 0.04%.
[0043] (3) Case-hardened steel superior in cold workability,
machinability, and fatigue characteristics after carburized
quenching as set forth in the above (1) or (2), characterized by
further containing, by mass %, one or more of [0044] Mo: 1.5% or
less, [0045] Ni: 3.5% or less, [0046] V: 0.5% or less, and [0047]
B: 0.005% or less.
[0048] (4) Case-hardened steel superior in cold workability,
machinability, and fatigue characteristics after carburized
quenching as set forth in any one of the above (1) to (3),
characterized by limiting a structural fraction of bainite to 30%
or less.
[0049] (5) Case-hardened steel superior in cold workability,
machinability, and fatigue characteristics after carburized
quenching as set forth in any one of the above (1) to (4),
characterized in that a grain size number of ferrite is 8 to 11 as
defined by JIS G 0551.
[0050] (6) Case-hardened steel superior in cold workability,
machinability, and fatigue characteristics after carburized
quenching as set forth in any one of the above (1) to (5),
characterized in that a maximum size of Ti precipitates is 40 .mu.m
or less.
[0051] (7) A method of production of case-hardened steel superior
in cold workability, machinability, and fatigue characteristics
after carburized quenching characterized by heating a steel
material comprised of the ingredients of any of the above (1) to
(3) to 1150.degree. C. or more, hot working it at a finishing
temperature of 840 to 1000.degree. C., and cooling it in a 800 to
500.degree. C. temperature range by 1.degree. C./s or less.
Advantageous Effects of Invention
[0052] The case-hardened steel of the present invention is superior
in forgeability, machinability, and other workability. Even when
producing parts by the cold forging process, coarsening of the
crystal grains due to heating at the time of carburized quenching
is suppressed. Deterioration of the dimensional precision due to
quenching strain is much smaller than the past.
[0053] Further, according to the case-hardened steel of the present
invention, the problem of the deterioration of machinability due to
the prevention of formation of coarse grains in the past is solved.
Further, higher precision of part shapes is achieved. Furthermore,
the tool life also becomes longer.
[0054] Further, parts made of the case-hardened steel of the
present invention are kept from forming coarse grains even in high
temperature carburization, sufficient strength characteristics such
as rolling contact fatigue characteristics can be obtained, etc.
The contribution to industry is extremely remarkable.
BRIEF DESCRIPTION OF DRAWINGS
[0055] FIG. 1 is a view for explaining a balance of machinability
and cold workability of the present invention.
[0056] FIG. 2 is a view showing a position for measuring a cooling
rate at the time of solidification.
[0057] FIG. 3 is a view showing a test piece used for an upset
test.
DESCRIPTION OF EMBODIMENTS
[0058] Coarsening of crystal grains due to carburized quenching is
prevented by using precipitates as pinning particles to suppress
grain growth. In particular, making Ti precipitates mainly
comprised of TiC and TiCS precipitate finely at the time of cooling
after hot working is extremely effective for preventing the
formation of coarse grains. Furthermore, to prevent the formation
of coarse grains, it is preferable to make NbC and other Nb
precipitates finely precipitate in the case-hardened steel.
[0059] However, if the amount of N contained in the steel is great,
the coarse TiN formed at the time of casting will not be
solubilized by the heating of the hot rolling or hot forging and
will sometimes remain in large amounts. If coarse TiN remains, at
the time of carburized quenching, the TiN will act as precipitation
nuclei resulting in TiC, TICS, and furthermore NbC precipitating
and fine dispersion of the precipitates being inhibited. Therefore,
to enable fine Ti precipitates and Nb precipitates to prevent
formation of coarse grains at the time of carburized quenching, it
is important to reduce the amount of N and solubilize the Ti
precipitates and Nb precipitates at the time of heating in hot
working.
[0060] Further, if coarse AlN remains at the time of heating in hot
working, in the same way as TiN, formation of fine precipitates
acting as pinning particles is inhibited.
[0061] However, the temperature at which AlN forms a solid solution
is lower than that of TiN, so compared with TiN, it is easier to
solubilize at the time of heating in hot rolling. Furthermore,
during the hot working and at the time of cooling after that, AlN
precipitates and grows slower than Ti precipitates and Nb
precipitates. Therefore, by preventing AlN from remaining at the
time of heating in hot working, it is possible to limit the amount
of precipitation of the AlN contained in the case-hardened
steel.
[0062] Therefore, according to the case-hardened steel of the
present invention limited in amount of precipitation of AlN, it is
possible to utilize fine Ti precipitates and Nb precipitates to
prevent the formation of coarse grains at the time of carburized
quenching.
[0063] Furthermore, to enable the pinning effect of Ti precipitates
and Nb precipitates to be stably exhibited, it is effective to
cause Ti precipitates and Nb precipitates to precipitate by
interphase boundary precipitation in the process of cooling after
hot working and the diffusion and transformation from austenite.
However, if bainite forms in the cooling process after hot rolling,
interphase boundary precipitation of precipitates will become
difficult.
[0064] Therefore, it is preferable to control the structure of the
steel after hot rolling and suppress the formation of bainite and
is more preferable to obtain a structure substantially not
containing any bainite.
[0065] In the method of production, first, it is necessary to heat
the steel material so that the Al, Ti, and Nb precipitates solute.
In particular, it is important to raise the heating temperature of
hot rolling, hot forging, or other hot working and cause the Ti
precipitates and Nb precipitates to solute.
[0066] Next, after hot working, that is, after hot rolling or after
hot forging, it is necessary to slow the cooling in the temperature
region of precipitation of Ti precipitates and Nb precipitates. As
a result, it is possible to make the Ti precipitates and Nb
precipitates finely disperse in the case-hardened steel.
[0067] Further, if the ferrite grains of the steel material before
carburized quenching are excessively fine, at the time of heating
for carburization, coarse grains will easily form. For that reason,
it is necessary to control the finishing temperature of the hot
rolling or hot forging to prevent formation of fine ferrite.
[0068] Further, when working the case-hardened steel of the present
invention into a gear etc., the teeth are formed by forging and
gear cutting before carburized quenching. At that time, MnS and
other sulfides cause the cold forgeability to drop, but are
extremely effective for gear cutting. That is, sulfides exhibit the
effect of suppressing changes in tool shape due to wear of the
cutting tools and extending so-called tool life.
[0069] In particular, in the case of precision shapes such as
gears, if the cutting tool life is short, stable formation of gear
shapes is not possible. For this reason, the cutting tool life has
an effect not simply on the production efficiency or cost, but also
the shape precision of the parts.
[0070] Therefore, to improve the machinability, it is desirable to
cause formation of sulfides in the steel.
[0071] On the other hand, in hot rolling or hot forging, in
particular the coarse MnS or other sulfides are often elongated.
Furthermore, if the sulfides increase in length, the probability of
their appearing as defects in the parts also becomes higher and the
performance of the parts is lowered. Therefore, not only the size
of the sulfides, but also control of the shape so as not to
elongate is important.
[0072] Note that, to suppress coarsening of the sulfides, it is
preferable to control the solidification speed at the time of
casting.
[0073] To reduce the MnS and other soft sulfides, it is also
effective to add Ti and cause the formation of TiCS and other Ti
sulfides. However, if the soft MnS is reduced, the added S will no
longer contribute to the improvement of the machinability.
[0074] Therefore, to improve the machinability, it is important to
not only add S, but also control the soft sulfides in the molten
steel to which Ti is added.
[0075] Therefore, it is preferable to control the shape of sulfides
by control of the AlN required for suppressing coarse grains,
addition of Ti, control of the amount of S, and, furthermore,
addition of Zr, Mg, and Ca.
[0076] The machinability and cold workability will be further
explained.
[0077] At the time of cold working, the sulfides mainly comprised
of MnS deform and become origin of fracture. In particular, the
coarse MnS lowers the limit compression rate and other aspects of
cold forgeability. Further, if the MnS in the steel is coarse,
anisotropy of the material characteristics will occur due to the
shape of the MnS.
[0078] To apply case-hardened steel to various complicated parts,
stable mechanical properties are demanded in all directions. For
this reason, in the case-hardened steel of the present invention,
it is preferable to make the sulfides mainly comprised of MnS finer
and make their shapes substantially spherical. Further, it is more
preferable that the change in shape be small even after forging and
other cold working.
[0079] Addition of Zr, Mg, and Ca is effective for causing
dispersion of fine sulfides. Furthermore, if Zr, Mg, Ca, etc.
solute in the MnS, the resistance to deformation becomes higher and
the sulfides no longer easily deform. Therefore, the addition of
Zr, Mg, and Ca is extremely effective for suppression of
elongating.
[0080] On the other hand, from the viewpoint of the machinability,
increase of the amount of S is important. Due to the addition of S,
the tool life at the time of cutting is improved. This effect is
determined by the total amount of S. The effect of the shape of the
sulfides is small. For this reason, by increasing the amount of
addition of S and controlling the shape of the sulfides, it is
possible to achieve both cold forgeability and machinability (tool
life).
[0081] In case-hardened steel, not only the prevention of formation
of coarse grains at the time of carburized quenching, but also
securing cold workability and machinability is important. If
increasing the amount of S, the machinability is improved, but a
deterioration of cold workability is invited. Therefore, it is also
important to secure a good cold workability when compared by the
same amount of S.
[0082] FIG. 1 compares the relationship of machinability and cold
workability for case-hardened steel with a good coarse grain
characteristic suppressed in formation of coarse grains at the time
of carburized quenching. In the present invention, it is possible
to maintain a good coarse grain characteristic (coarse grain
formation temperature >970.degree. C.) while achieving both cold
workability (limit compression rate) and machinability
(drillability VL1000). In FIG. 1, the further to the top right, the
better the balance of machinability and cold workability of the
material.
[0083] Below, the present invention will be explained in
detail.
[0084] First, the composition of ingredients will be explained.
Below, "mass %" will be simply described as "%".
[0085] C is an element raising the strength of steel. In the
present invention, to secure the tensile strength, 0.1% or more of
C is added. An amount of C of 0.15% or more is preferable. On the
other hand, if the content of C exceeds 0.5%, the steel remarkably
hardens and the cold workability is degraded, so the upper limit is
made 0.5%. Further, to secure toughness of the core part after
carburization, the amount of C is preferably made 0.4% or less. An
amount of C of 0.3% or less is more preferable.
[0086] Si is an element effective for deoxidation of steel.
[0087] In the present invention, 0.01% or more is added. Further,
Si is an element strengthening steel and improving the
quenchability. Addition of 0.02% or more is preferable.
Furthermore, Si is an element effective for increasing the grain
boundary strength. Furthermore, in bearing parts and rolling parts,
it is an element effective for extending lifetime by suppressing
structural changes and deterioration of quality in the process of
rolling contact fatigue. For this reason, when aiming at increasing
the strength, addition of 0.1% or more is more preferable. In
particular, to raise the rolling contact fatigue strength, addition
of 0.2% or more of Si is preferable.
[0088] On the other hand, if the amount of Si exceeds 1.5%, the
hardening causes the cold forging and other cold workability to
deteriorate, so the upper limit is made 1.5%. Further, to raise the
cold workability, it is preferable to make the amount of Si 0.5% or
less. In particular, when stressing cold forgeability, the amount
of Si is preferably 0.25% or less.
[0089] Mn is effective for deoxidation of steel. Furthermore, it is
an element improving the strength and quenchability of steel. In
the present invention, 0.3% or more is added. On the other hand, if
the amount of Mn exceeds 1.8%, the rise in hardness causes the cold
forgeability to be degraded, so 1.8% is made the upper limit. The
preferable range of the amount of Mn is 0.5 to 1.2%. Note that,
when stressing the cold forgeability, it is preferable to make the
upper limit of the amount of Mn 0.75%.
[0090] S is an element forming MnS in steel and improving the
machinability. In the present invention, to improve the
machinability, the content of S is made 0.001% or more. The
preferable lower limit of the amount of S is 0.1%. On the other
hand, if the amount of S is over 0.15%, grain boundary segregation
causes grain boundary embrittlement to be invited, so the upper
limit is made 0.15%. Further, if considering the fact that the
parts require high strength, the amount of S is preferably 0.05% or
less. Furthermore, when considering the strength or cold
workability and, furthermore, the stability of the same, the amount
of S is preferably made 0.03% or less.
[0091] Note that, in the past, in bearing parts and rolling parts,
it was considered necessary to reduce the S since MnS caused
deterioration of the rolling fatigue life. However, the inventors
etc. discovered that for improvement of the machinability, the
content of S has a large effect, while for improvement of the cold
workability, the shape of the sulfides has a large effect. In the
present invention, one or more of Mg, Zr, and Ca are added to
control the shape of the sulfides, so it is possible to make the
amount of S 0.01% or more. When stressing the machinability, the
amount of S is preferably made 0.02% or more.
[0092] Cr is an element effective for improving the strength and
quenchability of steel. In the present invention, 0.4% or more is
added. Furthermore, in bearing parts and rolling parts, it is
effective for increasing the residual amount of .gamma. of the
surface layer after carburization and increasing lifetime by
suppressing changes in structure and degradation of quality in the
process of rolling contact fatigue, so addition of 0.7% or more is
preferable. The more preferable amount of Cr is 1.0% or more. On
the other hand, if adding Cr over 2.0%, the rise in hardness causes
the cold workability to be degraded, so the upper limit is made
2.0%. To improve the cold forgeability, the amount of Cr is
preferably made 1.5% or less.
[0093] Ti is an element forming carbides, carbosulfides, nitrides,
and other precipitates in the steel. In the present invention, to
utilize the fine TiC and TICS to prevent the formation of coarse
grains at the time of carburized quenching, 0.05% or more of Ti is
added. The preferable lower limit of the amount of Ti is 0.1%. On
the other hand, if adding over 0.2% of Ti, precipitation hardening
causes the cold workability to remarkably degrade, so the upper
limit of the amount of Ti is made 0.2%. Further, to suppress
precipitation of TiN and improve the rolling contact fatigue
characteristic, it is preferable to make the amount of Ti 0.15% or
less.
[0094] Al is a deoxidizing agent. Addition of 0.005% or more is
preferable, but the invention is not limited to this. On the other
hand, if the amount of Al exceeds 0.04%, the AlN will remain
without being solubilized by the heating of the hot working. For
this reason, the coarse AlN will form precipitation nuclei for
precipitates of Ti and Nb and formation of fine precipitates will
be inhibited. Therefore, to prevent coarsening of the crystal
grains at the time of carburized quenching, the amount of Al has to
be made 0.04% or less.
[0095] N is an element forming nitrides. In the present invention,
to suppress the formation of coarse TiN and AlN, the upper limit of
the amount of N is made 0.0050%. This is because coarse TiN and AlN
form precipitation nuclei for Ti precipitates mainly comprised of
TiC and TiCS and Nb carbonitrides mainly comprised of NbC etc. and
inhibit the dispersion of fine precipitates.
[0096] P is an impurity. It is an element which raises the
resistance to deformation at the time of cold working and degrades
the toughness. If excessively included, the cold forgeability is
degraded, so the content of P has to be limited to 0.025% or less.
Further, to suppress embrittlement of the crystal grain boundaries
and improve the fatigue strength, the content of P is preferably
made 0.015% or less.
[0097] O is an impurity. It forms oxide inclusions in the steel and
impairs the workability, so the content is limited to 0.0025% or
less. Further, the case-hardened steel of the present invention
includes Ti, so oxide inclusions including Ti are formed and act as
precipitation nuclei causing TiC to precipitate. If the oxide
inclusions increase, the formation of fine TiC is sometimes
suppressed at the time of hot working.
[0098] Therefore, to make the Ti precipitates mainly comprised of
TiC and TiCS finely disperse and suppress the coarsening of crystal
grains at the time of carburized quenching, the upper limit of the
amount of O is preferably made 0.0020%.
[0099] Furthermore, in bearing parts and rolling parts, the oxide
inclusions sometimes serve as origin of rolling contact fatigue
fracture. For this reason, when used for bearing parts and rolling
parts, to improve the rolling life, the O content is preferably
limited to 0.0012% or less.
[0100] Furthermore, in the case-hardened steel of the present
invention, to control the form of the sulfides, it is necessary to
add one or more of Mg, Zr, and Ca. Mg, Zr, and Ca form roughly
spherical sulfides and further raise the deformation ability of MnS
to suppress elongating due to hot working. In particular, Mg and Zr
exhibit remarkable effects even when included in very small
amounts, so care is preferably exercised in secondary materials
etc. Furthermore, to stabilize the amounts of addition of Mg and
Zr, it is preferable to use refractories containing Mg and Zr to
control the content.
[0101] Mg is an element forming oxides and sulfides. Due to the
inclusion of Mg, composite sulfides (Mn,Mg)S with MgS or MnS etc.
are formed, so it is possible to suppress elongating of MnS. A very
small amount of Mg is effective for control of the form of the MnS.
To improve the workability, addition of 0.0002% or more of Mg is
preferable.
[0102] Further, oxides of Mg finely disperse and form the nuclei
for formation of MnS and other sulfides. To 35, utilize oxides of
Mg to suppress the formation of coarse sulfides, addition of
0.0003% or more of Mg is preferable. Furthermore, if adding Mg, the
sulfides become somewhat hard and become harder to elongate due to
hot working.
[0103] For control of the shape of the sulfides to contribute to
improvement of the machinability and prevent the cold workability
from being detracted from, addition of 0.0005% or more of Mg is
preferable. Note that, hot forging has the effect of causing fine
sulfides to uniform disperse and is effective for improvement of
the cold workability.
[0104] On the other hand, oxides of Mg easily float up in molten
steel, so the yield is low. From the viewpoint of the production
costs, the upper limit of the content of Mg is preferably 0.003%.
Further, if excessively adding Mg, large amounts of oxides are
formed in the molten steel and deposition on the refractories,
clogging of nozzles, and other trouble in steelmaking are sometimes
caused. Therefore, the amount of addition of Mg is more preferably
made 0.001% or less.
[0105] Zr is an element forming oxides, sulfides, and nitrides. If
adding a very small amount of Zr, it combines with the Ti in the
molten steel to form fine oxides, sulfides, and nitrides.
Therefore, in the present invention, the addition of Zr is
extremely effective for the control of inclusions and precipitates.
To control the form of the inclusions and improve the workability,
addition of 0.0002% or more of Zr is preferable, but the invention
is not limited to this.
[0106] Oxides, sulfides, and nitrides including Zr and Ti form
precipitation nuclei for MnS at the time of solidification. The Zr
and Ti dissolve into the MnS precipitated around these oxides,
sulfides, and nitrides including Zr and Ti resulting in a
deterioration of the deformation ability. Therefore, to suppress
the deformation of MnS and prevent elongating due to hot working,
addition of 0.0003% or more of Zr is preferable.
[0107] On the other hand, Zr is an expensive element, so from the
viewpoint of the production costs, the upper limit of the amount of
Zr is preferably made 0.01%. The more preferable amount of Zr is
0.005% or less, still more preferably 0.003% or less.
[0108] Ca is an element forming oxides and sulfides. To control the
form of the inclusions and improve the workability, 0.0002% or more
of Ca is preferably added. The CaS and (Mn,Ca)S and the composite
sulfides with Ti formed by the addition of Ca act as precipitation
nuclei for MnS at the time of solidification.
[0109] In particular, the Ca and Ti dissolve in the MnS
precipitated around the oxides and sulfides containing Ca and Ti
resulting in a deterioration of the deformation ability. Therefore,
to suppress deformation of MnS and prevent elongating due to hot
working, addition of 0.0003% or more of Ca is preferable.
[0110] On the other hand, in the same way as Mg, if excessively
adding Ca, deposition of the oxides on the refractories, clogging
of nozzles, and other trouble in steelmaking are sometimes caused.
Therefore, the amount of Ca is preferably made 0.005% or less.
[0111] Further, addition of two or more of Mg, Zr, and Ca is more
preferable. It is possible to make roughly spherical sulfides
finely disperse. When adding two or more of Mg, Zr, and Ca, it is
preferable to make the total content 0.0005% or more. Further, to
prevent deposition on the refractories etc. even when adding two or
more of Mg, Zr, and Ca, it is preferable to make the total content
0.006% or less, more preferable to make it 0.003% or less.
[0112] Furthermore, to suppress the formation of coarse grains at
the time of carburized quenching, in the same way as Ti, addition
of Nb forming carbonitrides is preferable. Nb, in the same way as
Ti, is an element bonding with C and N in the steel to form
carbonitrides. Due to the addition of Nb, the effect of suppression
of formation of coarse grains due to the Ti precipitates becomes
more remarkable. Even if the amount of Nb added is very small,
compared with the case of not adding Nb, the addition is extremely
effective for prevention of coarse grains.
[0113] This is because the Nb forms a solid solution in the Ti
precipitates and suppresses coarsening of the Ti precipitates. To
suppress the formation of coarse grains at the time of heating in
carburized quenching, addition of 0.01% or more of Nb is
preferable, but the invention is not limited to this. On the other
hand, if adding Nb in an amount of 0.04% or more, the steel hardens
and the cold workability, in particular the cold forgeability and
machinability, and, furthermore, the carburization characteristics
are sometimes degraded. Therefore, the amount of addition of Nb is
preferably made less than 0.04%. When stressing the cold
forgeability or other cold workability and machinability, the
preferable upper limit of the amount of Nb is less than 0.03%.
Further, when stressing the carburization ability in addition to
the workability, the preferable upper limit of the amount of Nb is
less than 0.02%.
[0114] Further, to achieve both prevention of coarse grains and
workability, it is preferable to adjust the total of the amount of
addition of Nb and the amount of addition of Ti. The preferable
range of Ti+Nb is 0.07% to less than 0.17%. In particular, in high
temperature carburization or cold forged parts, the preferable
range of Ti+Nb is over 0.09% to less than 0.17%.
[0115] Furthermore, to improve the strength and quenchability of
the steel, one or more of Mo, Ni, V, B, and Nb may be added.
[0116] Mo is an element improving the strength and quenchability of
steel. In the present invention, it is effective for increasing the
amount of residual .gamma. at the surface layer of carburized parts
and further to increase the lifetime by suppression of structural
changes and quality changes in the process of rolling contact
fatigue. However, if adding over 1.5% of Mo, the rise in hardness
causes the machinability and cold forgeability to be degraded in
some cases.
[0117] Therefore, making the content of Mo 1.5% or less is
preferable. Mo is an expensive element. From the viewpoint of the
production costs, making the amount 0.5% or less is more
preferable.
[0118] Ni, in the same way as Mo, is an element effective for
improving the strength and quenchability of the steel. However, if
adding Ni over 3.5%, the rise in the hardness causes the
cuttability and cold forgeability to deteriorate in some cases, so
making the content of Ni 3.5% or less is preferable. Ni is also an
expensive element. From the viewpoint of the production costs, the
preferable upper limit is 2.0%. The further preferable upper limit
of the amount of Ni is 1.0%.
[0119] V is an element improving the strength and quenchability if
forming a solid solution in the steel. If the amount of V is over
0.5%, the rise in the hardness causes the machinability and cold
forgeability to deteriorate in some cases, so making the upper
limit of content 0.5% is preferable. The preferable upper limit of
the amount of V is 0.2%.
[0120] B is an element effective for raising the quenchability of
steel with addition in a very fine amount. Further, B forms
boron-iron carbides in the cooling process after hot rolling,
increases the growth rate of ferrite, and promotes softening.
Furthermore, it is also effective for improving the grain boundary
strength of carburized parts and for improving the fatigue strength
and impact strength. However, if adding B in over 0.005%, the
effect becomes saturated and the impact strength is degraded, so
the upper limit of the content is preferably 0.005%. The preferable
upper limit of the amount of B is 0.003%.
[0121] Note that, the effect of the addition of Si and Cu and,
furthermore, the addition of Mo in suppressing structural changes
and quality changes in bearing parts and rolling parts in the
process of rolling contact fatigue is particularly large when the
residual austenite (residual .gamma.) at the surface layer after
carburization is 30 to 40%. To control the residual amount of
.gamma. of the surface layer to 30 to 40% in range,
carbonitridation treatment is effective. Carbonitridation treatment
is treatment for carburization, then nitridation in the process of
diffusion treatment.
[0122] To make the residual amount of .gamma. of the surface layer
30 to 40%, it is preferable to perform carbonitridation so that the
nitrogen concentration of the surface layer becomes 0.2 to 0.6% in
range. Note that, in this case, it is preferable to make the carbon
potential at the time of carburization 0.9 to 1.3% in range.
[0123] Further, in the case-hardened steel of the present
invention, the carbon and nitrogen penetrating the surface layer at
the time of carburized quenching and the solute Ti react and fine
Ti(C,N) precipitate in large amounts at the carburized layer. In
particular, at the bearing parts and rolling parts, the Ti(C,N) at
the surface layer causes the rolling fatigue life to be
improved.
[0124] Therefore, to improve the rolling fatigue life, it is
preferable to set the carbon potential at the time of carburization
to 0.9 to 1.3%. Further, with carburization, then nitridation in
the process of diffusion treatment, that is, carbonitridation
treatment, it is preferable to set the conditions so that the
nitrogen concentration of the surface becomes 0.2 to 0.6% in
range.
[0125] Next, among the precipitates included in the case-hardened
steel of the present invention, AlN and sulfides will be
explained.
[0126] AlN forms the precipitation nuclei for Ti precipitates and
Nb precipitates and inhibits the formation of fine precipitates.
Therefore, in the present invention, it is necessary to limit the
amount of precipitation of AlN included in the case-hardened steel.
If the amount of precipitation of AlN is excessive, coarse grains
are liable to be formed at the time of carburized quenching, so the
amount of precipitation of AlN in the case-hardened steel is
limited to 0.01% or less. The preferable upper limit of the amount
of precipitation of AlN is 0.005%.
[0127] To suppress the amount of precipitation of AlN of the
case-hardened steel, it is necessary to raise the hot working
heating temperature and promote solubilization. The case-hardened
steel of the present invention is limited in amount of N, so if
heating it to a temperature where AlN is solubilized, the Ti
precipitates and Nb precipitates can also be solubilized.
[0128] Note that, the amount of precipitation of AlN can be
measured by chemical analysis of the extraction residue.
[0129] The extraction residue is obtained by etching the steel by a
bromine methanol solution and filtering by a 0.2 .mu.m filter. Note
that, even if using a 0.2 .mu.m filter, in the process of
filtration, the precipitates cause the filter to clog, so
extraction of 0.2 .mu.m or smaller fine precipitates is also
possible.
[0130] MnS is useful for the improvement of the machinability, so
it is necessary to secure the density. On the other hand, elongated
coarse MnS impairs the cold workability, so the size and form have
to be controlled.
[0131] The inventors etc. studied the relationship between the
content of S, the size and shape of MnS inclusions, and the
machinability and cold workability.
[0132] As a result, it was learned that when MnS inclusions
observed under an optical microscope have a equivalent circle
diameter of over 20 .mu.m and an aspect ratio of over 3, they
become origin of fracture at the time of cold working.
[0133] The equivalent circle diameter of an MnS inclusion is the
diameter of a circle having an area equal to the area of the MnS
inclusion and can be found by image analysis. The aspect ratio is
the ratio of the length of the MnS inclusion divided by the
thickness of the MnS.
[0134] Next, the inventors etc. studied the effects of the
distribution of sulfides. The MnS inclusions of a hot rolled
material of a diameter of 30 mm were observed under a scanning
electron microscope and analyzed for the relationship of size,
aspect ratio and density, and cold workability and machinability.
The MnS inclusions are examined at a part of 1/2 radius from the
surface of the cross-section parallel to the rolling direction. Ten
fields of 1 mm.times.1 mm area were examined and the equivalent
circle diameters, aspect ratios, and numbers of the sulfide
inclusions present were found. Note that, the fact that the
inclusions are sulfides was confirmed by an energy dispersive X-ray
spectrometer attached to a scanning electron microscope.
[0135] The number of MnS inclusions with a equivalent circle
diameter over 20 .mu.m and an aspect ratio over 3 was counted and
divided by the area to find the density d. It was learned that the
density d of sulfides is influenced by the amount of S, so to
achieve both machinability and cold workability, the following
relation must be satisfied:
d.ltoreq.1700[S]+20 (/mm.sup.2)
[0136] Here, [S] indicates the content (mass %) of S. Furthermore,
if coarse Ti precipitates are present in the steel, they become
origin of contact fatigue fracture and the fatigue characteristics
deteriorate in some cases.
[0137] The contact fatigue strength is a required characteristic of
a carburized part and is the rolling contact fatigue characteristic
or surface fatigue strength. To raise the contact fatigue strength,
making the maximum size of the Ti precipitates less than 40 .mu.m
is preferable.
[0138] The maximum size of the Ti precipitates is found by
statistics of extremes measured in the cross-section of the
longitudinal direction of the case-hardened steel using a standard
inspection area of 100 mm.sup.2, inspection of 16 fields, and a
prediction area of 30000 mm.sup.2.
[0139] The method of measurement of the maximum size of
precipitates using statistics of extremes is, for example, as
described in Yukitaka Murakami, "Metal Fatigue--Effects of Small
Defects and Nonmetallic Inclusions", Yokendo, pp. 233 to 239
(1993), a two-dimensional test method of estimating the largest
precipitates obtained in a fixed area, that is, a prediction area
(30000 mm.sup.2).
[0140] The values are plotted on an extreme probability paper, the
primary function of the maximum precipitate size and statistics of
extremes standardized variable is found, and the maximum
precipitate distribution line is extrapolated to predict the size
of the largest precipitate in the prediction area.
[0141] Next, the structure of the case-hardened steel of the
present invention will be explained.
[0142] The structural fraction of bainite in the case-hardened
steel is preferably limited to 30% or less. This is because to
prevent the formation of coarse grains at the time of carburized
quenching, it is preferable to form fine precipitates at the grain
boundary. That is, if the structural fraction of bainite formed at
the time of cooling after hot working exceeds 30%, it becomes
harder for the Ti precipitates and the Nb precipitates to be made
to precipitate by interphase boundary precipitation.
[0143] Suppressing the structural fraction of bainite to 30% or
less is also effective for improving the cold workability.
[0144] In the case of high temperature carburization or otherwise
when the conditions for prevention of coarse grains are severe, the
upper limit of the structural fraction of bainite is preferably
made 20%, more preferably 10% or less. Furthermore, when cold
forging, then performing high temperature carburization etc., the
upper limit of the structural fraction of bainite is preferably
made 5% or less.
[0145] If the ferrite grains of the case-hardened steel of the
present invention are excessively fine, coarse grains easily form.
This is because at the time of carburized quenching, the austenite
grains become excessively fine. In particular, if the grain size
number of the ferrite exceeds 11 as defined by JIS G 0551, coarse
grains easily are formed. On the other hand, if the grain size
number of ferrite of the case-hardened steel becomes less than 8 as
defined by JIS G 0551, the ductility falls and the cold workability
is impaired in some cases. Therefore, the grain size number of
ferrite of the case-hardened steel is preferably 8 to 11 in range
as defined by JIS G 0551.
[0146] Next, the method of production of case-hardened steel of the
present invention will be explained.
[0147] Steel is produced by a converter, electric furnace, or other
usual method, adjusted in ingredients, and passed through a casting
process and, if necessary, a blooming process, to obtain a steel
material. The steel material is hot worked, that is, hot rolled or
hot forged, to produce steel rails or steel bars.
[0148] The sulfides of the steel material often precipitate in the
molten steel or at the time of solidification. The size of the
sulfides is greatly influenced by the cooling rate at the time of
solidification. Therefore, to prevent the coarsening of the
sulfides, it is important to control the cooling rate at the time
of solidification.
[0149] The cooling rate at the time of solidification is defined as
the cooling rate at the part of 1/2 of the distance from the
surface to the centerline in the thickness direction on the
centerline of the cast bloom width W in the cross-section of the
cast bloom shown in FIG. 2 (position from the surface of T/4 from
the surface with respect to the cast bloom thickness T). To
suppress coarsening of the sulfides, the cooling rate at the time
of solidification is preferably made 3.degree. C./min or more.
Preferably it is made 5.degree. C./min or more, more preferably
10.degree. C./min or more. Note that, the cooling rate at the time
of solidification can be confirmed by the secondary dendrite arm
spacing.
[0150] The cast bloom is reheated as it is and hot worked to
produce case-hardened steel or the material obtained by a blooming
process is reheated and hot worked to produce case-hardened steel.
In general, a cast bloom is bloomed to form a billet, cooled to
room temperature, then reheated to produce case-hardened steel.
Furthermore, in the production of gears or other parts, hot forging
is sometimes applied. At that time, in blooming, it is preferable
to hold the steel at a 1150.degree. C. or more high temperature for
10 minutes or more and cause the Ti and Nb precipitates to
solute.
[0151] To produce case-hardened steel, the steel material is
heated. If the heating temperature is less than 1150.degree. C., it
is not possible to make the Ti precipitates, Nb precipitates, and
AlN solute in the steel, and coarse Ti precipitates, Nb
precipitates, and AlN will remain.
[0152] To cause the fine Ti precipitates or Nb precipitates to
disperse in the case-hardened steel after hot working and suppress
the formation of coarse grains at the time of carburized quenching,
it is necessary to make the heating temperature 1150.degree. C. or
more. The preferable lower limit of the heating temperature is
1180.degree. C. or more.
[0153] The upper limit of the heating temperature is not
prescribed, but if considering the load of the heating furnace,
1300.degree. C. or less is preferable. To make the steel material
uniform in temperature and cause the precipitates to solute, a
holding time of 10 minutes or more is preferable. The holding time
is preferably 60 minutes or less from the viewpoint of
productivity.
[0154] If the finishing temperature of the hot working is less than
840.degree. C., the ferrite crystal grains become fine and coarse
grains easily form at the time of carburized quenching. On the
other hand, if the finishing temperature exceeds 1000.degree. C.,
hardening occurs and the cold workability deteriorates. Therefore,
the finishing temperature of hot working is made 840 to
1000.degree. C. Note that, the preferable range of the finishing
temperature is 900 to 970.degree. C., and the more preferable range
is 920 to 950.degree. C.
[0155] The cooling conditions after the hot working are important
for causing the Ti precipitates and Nb precipitates to finely
disperse. The temperature range at which precipitation of Ti
precipitates and Nb precipitates is promoted is 500 to 800.degree.
C. Therefore, the cooling is performed slowly by 1.degree. C./s or
less from a 800.degree. C. to 500.degree. C. temperature range to
promote the formation of Ti precipitates and Nb precipitates.
[0156] If the cooling rate exceeds 1.degree. C./s, the time of
passage through the region of the precipitation temperature of Ti
precipitates and Nb precipitates becomes shorter and the formation
of fine precipitates becomes insufficient. Further, if the cooling
rate becomes faster, the structural fraction of bainite becomes
larger. Further, if the cooling rate is large, the case-hardened
steel hardens and the cold workability deteriorates, so the cooling
rate is preferably 0.7.degree. C./s or less.
[0157] Note that, as the method for reducing the cooling rate, the
method of setting a heat retaining cover or heat retaining cover
with a heat source after the rolling line and thereby slowing the
cooling may be mentioned.
[0158] The case-hardened steel of the present invention can be
applied to parts produced by a cold forging process or parts
produced by hot forging. The hot forging process, for example, may
comprise hot forging of steel bar, normalization or other heat
treatment if necessary, cutting, carburized quenching, and grinding
or polishing if necessary.
[0159] By using the case-hardened steel of the present invention,
hot forging it at for example a 1150.degree. C. or more heating
temperature, then, as necessary, treating it by normalization, it
is possible to suppress the formation of coarse grains even if
applying high temperature carburization in a 950 to 1090.degree. C.
temperature region. For example, in the case of bearing parts or
rolling parts, even if treating them by high temperature
carburization, superior rolling contact fatigue characteristics can
be obtained.
[0160] The carburized quenching is not particularly limited, but
when aiming at a high rolling fatigue life in bearing parts and
rolling parts, it is preferable to set the carbon potential at 0.9
to 1.3%. Further, carburization, then nitridation in the process of
diffusion treatment, that is, carbonitridation treatment, is also
effective. Conditions whereby the nitrogen concentration of the
surface becomes 0.2 to 0.6% in range are suitable. By selecting
these conditions, fine Ti(C,N) precipitates in large amounts at the
carburized layer and the rolling life is improved.
Example 1
[0161] Steels having the compositions of ingredients shown in
Tables 1 to 3 were produced and cast at solidification cooling
rates of 10 to 11.degree. C./min. The blank fields in the
ingredients of Tables 1 to 3 mean the elements are deliberately not
added, while the underlines indicate the figures are outside the
ranges of the present invention.
[0162] The solidification cooling rate was adjusted in advance
based on data analyzing the relationship between the cooling
conditions and solidification cooling rate when casting various
sizes of cast blooms. The solidification cooling rate of some of
the cast blooms was confirmed by secondary dendrite arm spacing to
be 10 to 11.degree. C./min in range. Some of the cast blooms were
bloomed in accordance with need.
TABLE-US-00001 TABLE 1 Chemical ingredients (mass %) No. C Si Mn P
S Cr Ti Al N O Zr 1 0.21 0.19 1.30 0.018 0.011 1.06 0.13 0.026
0.0030 0.0011 0.0024 2 0.20 0.20 0.38 0.022 0.014 1.10 0.14 0.024
0.0047 0.0014 3 0.21 0.19 0.98 0.014 0.015 1.20 0.06 0.035 0.0033
0.0014 4 0.19 0.18 0.84 0.014 0.014 1.28 0.08 0.027 0.0045 0.0012
0.0007 5 0.19 0.21 0.88 0.005 0.016 1.22 0.08 0.038 0.0026 0.0015
0.0013 6 0.20 0.19 0.58 0.014 0.013 1.13 0.06 0.018 0.0029 0.0014 7
0.18 0.24 0.70 0.015 0.010 1.22 0.07 0.038 0.0029 0.0012 0.0025 8
0.20 0.19 0.41 0.021 0.030 1.23 0.10 0.026 0.0045 0.0014 9 0.21
0.21 1.23 0.011 0.026 1.10 0.12 0.037 0.0035 0.0015 10 0.19 0.21
1.04 0.017 0.031 1.23 0.11 0.038 0.0028 0.0014 0.0005 11 0.19 0.25
1.63 0.018 0.029 1.05 0.07 0.020 0.0031 0.0012 12 0.22 0.21 0.81
0.016 0.028 1.22 0.11 0.016 0.0032 0.0011 13 0.20 0.19 1.60 0.009
0.026 1.15 0.14 0.028 0.0026 0.0012 0.0016 14 0.19 0.19 0.99 0.018
0.029 1.15 0.15 0.034 0.0027 0.0010 0.0018 15 0.32 0.22 0.38 0.018
0.048 1.22 0.06 0.030 0.0032 0.0010 16 0.21 0.25 0.32 0.024 0.026
1.16 0.10 0.034 0.0026 0.0012 0.0018 17 0.22 0.18 1.77 0.009 0.015
1.21 0.12 0.022 0.0028 0.0011 18 0.21 0.20 0.54 0.025 0.013 1.21
0.12 0.014 0.0034 0.0014 19 0.19 0.23 0.86 0.005 0.012 1.22 0.09
0.027 0.0035 0.0012 0.0004 20 0.21 0.22 1.31 0.023 0.016 1.28 0.11
0.023 0.0034 0.0011 21 0.21 0.25 0.57 0.016 0.013 1.13 0.14 0.037
0.0047 0.0015 22 0.19 0.19 1.19 0.011 0.011 1.22 0.08 0.021 0.0041
0.0010 0.0008 23 0.22 0.19 0.57 0.013 0.013 1.13 0.05 0.019 0.0025
0.0014 0.0030 24 0.18 0.24 0.74 0.016 0.011 1.16 0.12 0.017 0.0032
0.0011 25 0.21 0.23 1.15 0.019 0.015 1.18 0.05 0.018 0.0032 0.0014
0.0027 26 0.22 0.21 0.48 0.013 0.013 1.27 0.07 0.025 0.0031 0.0014
0.0017 27 0.20 0.20 0.45 0.015 0.010 1.15 0.09 0.037 0.0036 0.0010
0.0010 28 0.20 0.22 1.11 0.022 0.017 1.12 0.13 0.024 0.0048 0.0015
0.0006 29 0.22 0.20 1.19 0.016 0.025 1.26 0.09 0.034 0.0029 0.0013
30 0.21 0.24 1.08 0.008 0.025 1.08 0.15 0.036 0.0030 0.0011 31 0.21
0.25 1.16 0.011 0.031 1.28 0.05 0.039 0.0028 0.0010 0.0022 32 0.19
0.23 1.73 0.009 0.040 1.23 0.06 0.016 0.0041 0.0015 0.0014 33 0.22
0.25 0.74 0.007 0.025 1.18 0.10 0.008 0.0026 0.0010 34 0.21 1.22
1.22 0.009 0.030 1.13 0.15 0.009 0.0038 0.0015 35 0.18 0.22 1.35
0.011 0.032 1.25 0.14 0.013 0.0039 0.0011 0.0020 Chemical
ingredients (mass %) No. Mg Ca Nb Mo Ni V B Remarks 1 Inv. ex. 2
0.0005 3 0.0025 4 0.0006 5 0.0020 6 0.0008 0.0014 7 0.0018 0.0013 8
9 10 11 0.0015 12 0.0012 13 0.0015 0.0014 14 0.0011 15 0.0015
0.0013 16 0.0003 0.0019 17 0.024 18 0.021 19 0.012 20 0.0012 0.019
21 0.0006 0.013 22 0.0004 0.013 23 0.0015 0.016 24 0.0014 0.0015
0.025 25 0.0017 0.0009 0.014 0.13 26 0.0007 0.0004 0.014 0.30 27
0.0005 0.0011 0.020 28 0.0016 0.0013 0.012 0.0015 29 0.014 30 0.010
31 0.022 32 0.014 33 0.0011 0.016 34 0.0008 0.023 35 0.0015 0.0009
0.024
TABLE-US-00002 TABLE 2 Chemical ingredients (mass %) No. C Si Mn P
S Cr Ti Al N O Zr 36 0.19 0.19 1.72 0.009 0.029 0.55 0.12 0.039
0.0049 0.0015 0.0015 37 0.22 0.18 1.68 0.024 0.028 1.06 0.06 0.023
0.0030 0.0012 38 0.21 0.20 0.32 0.010 0.028 1.08 0.10 0.032 0.0039
0.0013 0.0013 39 0.20 0.21 1.02 0.018 0.030 1.05 0.09 0.010 0.0046
0.0011 0.0019 40 0.19 0.20 0.33 0.025 0.035 0.62 0.12 0.022 0.0045
0.0015 0.0013 41 0.19 0.20 1.16 0.013 0.028 1.20 0.09 0.032 0.0049
0.0015 0.0021 42 0.19 0.23 1.37 0.012 0.017 1.08 0.13 0.032 0.0035
0.0013 0.0017 43 0.21 0.18 1.00 0.016 0.013 1.07 0.11 0.019 0.0044
0.0014 44 0.20 0.25 1.69 0.020 0.016 1.15 0.05 0.035 0.0031 0.0012
45 0.21 0.20 0.76 0.019 0.017 1.06 0.08 0.033 0.0031 0.0013 0.0012
46 0.20 0.22 1.52 0.015 0.015 1.30 0.10 0.018 0.0048 0.0013 0.0017
47 0.19 0.25 1.34 0.012 0.027 1.21 0.12 0.012 0.0041 0.0011 48 0.22
0.22 0.64 0.014 0.027 1.11 0.13 0.032 0.0050 0.0014 48 0.19 0.21
0.45 0.010 0.027 1.28 0.13 0.019 0.0026 0.0010 0.0010 49 0.21 0.21
0.56 0.021 0.044 1.62 0.15 0.039 0.0033 0.0010 0.0020 50 0.20 0.18
1.02 0.023 0.054 1.15 0.11 0.019 0.0033 0.0013 51 0.22 0.23 0.75
0.022 0.026 1.25 0.06 0.019 0.0047 0.0010 52 0.21 0.18 0.38 0.017
0.028 0.72 0.09 0.028 0.0031 0.0012 0.0018 53 0.21 0.20 0.82 0.018
0.029 1.12 0.09 0.035 0.0040 0.0012 0.0025 54 0.21 0.23 0.56 0.011
0.031 1.08 0.09 0.013 0.0049 0.0013 Chemical ingredients (mass %)
No. Mg Ca Nb Mo Ni V B Remarks 36 0.0006 0.019 Inv. ex. 37 0.0018
0.0012 0.020 38 0.0006 0.0013 0.019 39 0.0012 0.0013 0.020 0.21 40
0.0004 0.0013 0.016 0.95 41 0.0010 0.0013 0.022 0.0016 42 0.14 43
0.0004 0.16 44 0.0010 0.14 45 0.0017 0.14 46 0.0007 0.12 47 0.020
0.13 48 0.011 0.16 48 0.022 0.16 49 0.019 0.13 50 0.0003 0.011 0.15
51 0.0005 0.014 0.16 52 0.0008 0.0012 0.013 0.92 53 0.0004 0.019
0.12 54 0.0010 0.0017 0.014 0.15
TABLE-US-00003 TABLE 3 Chemical ingredients (mass %) Re- No. C Si
Mn P S Cr Ti Al N O Zr Mg Ca Nb Mo Ni V B marks 55 0.19 0.24 1.72
0.013 0.012 1.10 0.035 0.0126 0.0014 Comp. 56 0.18 0.23 0.98 0.013
0.013 1.14 0.08 0.018 0.0031 0.0011 ex. 57 0.20 0.19 1.15 0.024
0.013 1.26 0.12 0.034 0.0036 0.0012 58 0.19 0.22 1.06 0.025 0.012
1.09 0.14 0.034 0.0036 0.0013 59 0.19 0.21 1.57 0.017 0.030 1.10
0.017 0.0043 0.0011 60 0.19 0.25 0.79 0.005 0.030 1.14 0.029 0.0030
0.0012 61 0.20 0.24 1.72 0.008 0.012 1.28 0.14 0.034 0.0040 0.0012
0.0010 0.0005 62 0.19 0.19 0.84 0.007 0.027 1.27 0.15 0.015 0.0048
0.0015 0.0014 0.0014 63 0.20 0.19 0.31 0.009 0.014 1.17 0.12 0.022
0.0045 0.0013 0.0009 64 0.20 0.22 0.75 0.017 0.030 1.07 0.13 0.016
0.0031 0.0010 0.0025 65 0.20 0.24 0.71 0.023 0.030 1.13 0.14 0.020
0.0032 0.0013 0.0010 66 0.19 0.20 1.52 0.022 0.011 1.25 0.13 0.037
0.0043 0.0015 0.0025 0.0017 0.017 67 0.20 0.22 1.52 0.009 0.026
1.10 0.09 0.015 0.0049 0.0014 0.0024 0.0016 0.023 68 0.18 0.19 0.42
0.025 0.015 1.23 0.13 0.029 0.0043 0.0011 0.0004 0.014 69 0.20 0.21
1.78 0.013 0.027 1.26 0.13 0.011 0.0043 0.0012 0.0018 0.012 70 0.20
0.20 1.11 0.019 0.031 1.24 0.10 0.033 0.0031 0.0014 0.0016 0.023 71
0.20 0.24 1.02 0.022 0.017 1.09 0.12 0.013 0.0124 0.0012 0.0011 72
0.21 0.22 0.87 0.018 0.017 1.25 0.11 0.014 0.0145 0.0012 0.0006 73
0.19 0.21 1.02 0.019 0.013 1.26 0.10 0.018 0.0086 0.0011 0.0011 74
0.18 0.20 0.34 0.015 0.026 1.15 0.12 0.031 0.0098 0.0015 0.0024 75
0.19 0.22 0.33 0.008 0.030 1.16 0.06 0.030 0.0146 0.0012 0.0015 76
0.20 0.19 1.74 0.009 0.028 1.25 0.12 0.006 0.0113 0.0010 0.0016 77
0.20 0.24 1.57 0.009 0.011 1.25 0.020 0.0031 0.0013 0.0009 78 0.20
0.24 0.32 0.020 0.013 1.07 0.30 0.040 0.0049 0.0010 0.0011 79 0.20
0.22 1.60 0.015 0.027 1.05 0.14 0.034 0.0026 0.0011 0.0011 0.120 80
0.20 0.24 0.82 0.021 0.032 1.05 0.14 0.012 0.0045 0.0031 0.0020 81
0.18 0.24 1.45 0.006 0.032 1.16 0.10 0.008 0.0049 0.0015 0.0005 82
0.21 0.20 0.77 0.023 0.031 1.10 0.05 0.007 0.0040 0.0015 0.0004 83
0.21 0.20 0.77 0.023 0.031 1.10 0.05 0.007 0.0040 0.0015 0.0004 84
0.21 0.22 1.31 0.007 0.016 1.25 0.15 0.016 0.0031 0.0012 0.14 85
0.22 0.24 1.79 0.018 0.010 1.10 0.11 0.020 0.0038 0.0011 0.15 86
0.20 0.22 0.79 0.018 0.014 1.23 0.10 0.009 0.0031 0.0014 0.020 0.13
87 0.22 0.23 0.78 0.009 0.010 1.07 0.035 0.0126 0.0012 0.15 88 0.20
0.19 0.94 0.017 0.026 1.25 0.06 0.034 0.0033 0.0014 0.14 89 0.20
0.22 0.89 0.025 0.031 1.28 0.07 0.032 0.0030 0.0013 0.020 0.13
[0163] Next, the steels were hot worked to produce steel bars of
diameters of 24 to 30 mm. The steels were observed under a
microscope, the bainite fractions were measured, and the ferrite
grain size numbers were determined based on the provisions of JIS G
0551. The Vickers hardnesses were measured based on JIS Z 2244 and
used as indicators of cold workability and machinability. The
amounts of precipitation of AlN were found by chemical
analysis.
[0164] Further, the statistics of extremes method was used to
predict the maximum sizes of the Ti precipitates. Table 4 to 6 show
the hot working heating temperatures, finishing temperatures,
cooling rates, bainite fractions, ferrite grain size numbers, AlN
precipitation, Ti precipitate maximum sizes, and Vickers
hardnesses. Note that, the cooling rate is the cooling rate in the
500 to 800.degree. C. range. This was found from the time required
for cooling from 800.degree. C. to 500.degree. C.
[0165] The maximum sizes of the Ti precipitates were found as
follows. An optical microscope was used to observe the metal
structures and contrast was used to differentiate the precipitates.
Note that, the contrast of the precipitates was confirmed using a
scanning electron microscope and energy dispersive X-ray
spectrometer.
[0166] In the longitudinal direction cross-section of each test
piece, 16 fields of regions of standard inspection areas of 100
mm.sup.2 (10 mm.times.10 mm region) were prepared in advance. The
largest Ti precipitates in each 100 square mm standard inspection
area was detected and photographed by an optical microscope by
1000.times..
[0167] This was repeated 16 times for the 16 fields of the standard
inspection areas of 100 mm.sup.2. In this way, the test was
conducted for 16 fields and the size of the largest precipitate in
each standard inspection area was measured from the obtained
photographs. Note that, in the case of an ellipse, the geometric
mean of the long axis and short axis is found and used as the size
of the precipitate.
[0168] The 16 sets of data of the obtained maximum precipitate
sizes were plotted on an extreme probability paper by the method
described in Yukitaka Murakami, "Metal Fatigue--Effects of Small
Defects and Nonmetallic Inclusions", Yokendo, pp. 233 to 239
(1993), the largest precipitate distribution line, that is, the
primary function of the maximum precipitate size and statistics of
extremes standardized variable, was found, the largest precipitate
distribution line was extrapolated, and the diameters of the
largest precipitates in the prediction area (30000 mm.sup.2) were
found.
[0169] Further, to evaluate the cold workability by cold forging,
the test piece was annealed, then subjected to an upset test. The
grooved test piece shown in FIG. 3 was obtained and measured for
the limit compression rate until fracture. The compression rate was
changed and 10 test pieces were used to find the probability of
fracture. The compression rate when the probability became 50% was
made the limit compression rate.
[0170] The higher this limit compression rate, the better the
forgeability evaluated. This test method is a method of evaluation
close to cold forging, but has also been considered an indicator
showing the effects of sulfides on forgeability in hot forging.
[0171] The machinability was evaluated by a test finding the
lifetime until a drill broke. Note that, the drilling was performed
using a high speed steel straight shank drill having a diameter of
3 mm at a feed of 0.25 mm, a hole depth of 9 mm, and a drill
projection of 35 mm using a water soluble cutting fluid.
[0172] The speed of the drill was fixed at 10 to 70 muffin in range
and the cumulative hole depth until breakage was measured while
drilling. Here, the cumulative hole depth is the product of the
depth of one hole and the number of drilled holes.
[0173] The speed of the drill was changed and similar measurements
conducted. The maximum value of the speed of the drill where the
cumulative hole depth exceeds 1000 mm was found as VL1000. The
larger the VL1000, the better the tool life and the more superior
the machinability the material is evaluated as.
[0174] Further, the coarse grain characteristic was evaluated by
taking a test piece from a steel bar after spheroidal annealing,
cold upset forging it by a reduction rate of 50%, then heat
treating it simulating carburized quenching (referred to as
"carburization simulation"), and measuring the old austenite grain
size.
[0175] The carburization simulation comprised heat treatment
heating a test piece to 910 to 1010.degree. C., holding it there
for 5 hours, then water cooling it. The old austenite grain size
was measured in accordance with JIS G 0551.
[0176] The old austenite grain size was measured and the
temperature at which coarse grains formed (coarsening temperature)
was found. Note that, the old austenite grain size was measured by
observation at 400.times. for about 10 fields. If even one coarse
grain of a grain size number of 5 or less was present, it was
judged that coarse grains were formed.
[0177] The heating temperature of the carburized quenching
treatment is usually 930 to 950.degree. C., so a test piece with a
coarsening temperature of 950.degree. C. or less was judged to be
inferior in crystal grain coarsening characteristic.
[0178] Next, the reduction rate was made 50%, the steel was cold
forged, and a cylindrical rolling contact fatigue test piece of a
diameter of 12.2 mm was obtained and treated by carburized
quenching. The carburized quenching was performed by heating the
steel in an atmosphere of a carbon potential of 0.8% to 950.degree.
C., holding it there fore 5 hours, and quenching it in oil of a
temperature of 130.degree. C. Furthermore, the steel was held at
180.degree. C. for 2 hours and tempered. These carburized quenched
materials were investigated for the .gamma. granularity (carburized
layer austenite grain size number) of the carburized layers based
on JIS G 0551.
[0179] Furthermore, a point contact type rolling contact fatigue
test rig (Hertz maximum contact stress 5884 MPa) was used to
evaluate the rolling contact fatigue characteristic. As a measure
of the fatigue life, the L.sub.10 life, defined as "the number of
cycles of stress to fatigue fracture at a probability of failure of
10% obtained by plotting the test results on a Weibull probability
paper", was used. However, materials with frequent breakage at a
reduction rate of 50% were not subjected to subsequent fatigue
tests.
[0180] The results of these investigations are summarized in Tables
4 to 6. The rolling fatigue life shows the relative value of the
L.sub.10 life of each material indexed to the L.sub.10 life of No.
55 (comparative example) as "1".
TABLE-US-00004 TABLE 4 Hot working Ferrite Heating Finishing
Cooling Bainite grain AlN Ti precipitate Sulfide temp. temp. rate
fraction size precipitation max. density No. (.degree. C.)
(.degree. C.) (.degree. C./s) (%) number (%) size .mu.m (/mm.sup.2)
1 1270 930 0.50 0 9.8 0.003 21 16.0 2 1260 950 0.53 0 9.0 0.004 23
29.5 3 1190 940 0.53 0 9.4 0.004 26 26.6 4 1210 940 0.53 0 9.4
0.004 25 13.2 5 1260 940 0.55 0 9.8 0.003 23 11.6 6 1220 930 0.53 0
9.2 0.004 27 27.9 7 1190 940 0.48 0 10.5 0.003 29 25.6 8 1180 940
0.57 0 9.2 0.004 26 47.5 9 1220 930 0.55 0 10.2 0.003 31 52.0 10
1250 940 0.49 0 9.5 0.004 27 36.2 11 1270 930 0.48 0 9.8 0.003 24
53.3 12 1230 950 0.56 0 9.8 0.003 30 37.3 13 1200 930 0.47 0 9.4
0.003 24 51.4 14 1270 930 0.46 0 10.2 0.002 32 39.2 15 1190 940
0.52 5 9.0 0.004 25 30.2 16 1240 930 0.48 0 9.3 0.003 24 41.8 17
1220 940 0.47 0 10.1 0.002 25 27.9 18 1250 950 0.46 0 10.4 0.003 31
23.2 19 1190 940 0.57 0 10.0 0.002 23 19.7 20 1270 940 0.56 0 10.0
0.003 29 10.9 21 1230 930 0.52 0 9.2 0.002 26 21.7 22 1190 950 0.45
0 9.5 0.004 27 22.5 23 1220 930 0.57 0 9.7 0.004 27 25.6 24 1230
940 0.50 0 10.5 0.004 27 16.4 25 1250 930 0.56 0 10.2 0.004 29 28.3
26 1190 930 0.53 0 9.0 0.003 21 20.2 27 1250 940 0.52 0 10.3 0.004
27 25.1 28 1230 940 0.51 0 10.4 0.003 30 10.9 29 1200 940 0.52 0
9.6 0.002 27 59.5 30 1200 940 0.46 0 9.6 0.003 29 46.8 31 1230 940
0.56 0 10.4 0.003 22 57.1 32 1270 930 0.48 0 9.8 0.003 25 60.6 33
1200 950 0.56 0 9.3 0.004 29 53.3 34 1200 940 0.45 0 9.8 0.003 28
50.0 35 1280 940 0.49 0 9.0 0.002 23 38.1 Carburized layer Limit
Fatigue Vickers Coarsening austenite comp. Machineability life
hardness temp. grain size rate VL1000 (rel. No. (HV) (.degree. C.)
number (%) (m/min) value) Remarks 1 180 >1050 9.8 58 48 3.5 Inv.
ex. 2 183 >1050 8.8 56 46 3.7 3 187 >1050 9.9 56 45 3.0 4 184
>1050 8.9 55 49 3.4 5 185 >1050 8.7 56 46 3.5 6 194 >1050
8.6 57 46 2.8 7 188 >1050 8.0 55 49 2.6 8 172 >1050 9.7 55 55
2.5 9 183 >1050 8.9 54 51 3.8 10 188 >1050 8.5 53 54 3.4 11
176 >1050 10.0 56 50 3.0 12 178 >1050 8.4 54 53 3.2 13 187
>1050 8.4 56 50 2.6 14 183 >1050 8.5 54 51 3.2 15 192
>1050 9.8 52 51 3.4 16 177 >1050 9.9 56 52 2.8 17 173
>1050 9.1 59 46 2.7 18 178 >1050 9.6 56 49 3.7 19 174
>1050 9.6 56 46 3.1 20 180 >1050 9.4 58 48 2.5 21 194
>1050 10.0 60 48 3.2 22 179 >1050 8.4 58 50 2.5 23 181
>1050 9.0 59 48 3.2 24 192 >1050 8.9 56 46 2.9 25 174
>1050 9.3 59 50 3.3 26 193 >1050 8.4 55 49 2.8 27 175
>1050 8.7 59 48 2.6 28 184 >1050 8.4 55 46 2.6 29 180
>1050 9.1 54 50 3.3 30 177 >1050 9.0 56 53 3.7 31 175
>1050 9.5 49 58 3.6 32 189 >1050 8.6 54 53 3.5 33 189
>1050 8.6 55 51 3.3 34 191 >1050 9.7 54 50 3.7 35 176
>1050 9.5 54 53 3.2
TABLE-US-00005 TABLE 5 Hot working Ferrite Heating Finishing
Cooling Bainite grain AlN Ti precipitate Sulfide Vickers temp.
temp. rate fraction size precipitation max. size density hardness
No. (.degree. C.) (.degree. C.) (.degree. C./s) (%) number (%)
.mu.m (/mm.sup.2) (HV) 36 1210 940 0.52 0 8.8 0.003 26 53.9 173 37
1270 950 0.48 0 10.4 0.003 27 41.6 178 38 1190 950 0.46 0 9.7 0.003
23 45.0 173 39 1260 940 0.56 0 8.9 0.003 27 36.9 194 40 1240 950
0.47 0 9.5 0.003 24 59.7 187 41 1200 940 0.46 0 9.6 0.004 30 40.3
174 42 1200 930 0.49 4 9.5 0.003 20 15.2 201 43 1280 950 0.50 4
10.2 0.003 29 15.2 193 44 1260 940 0.45 4 9.9 0.004 27 29.3 185 45
1260 940 0.57 5 9.5 0.003 22 15.1 188 46 1200 950 0.50 7 9.5 0.002
28 28.9 188 47 1240 950 0.47 6 9.1 0.002 23 47.9 202 48 1250 950
0.56 5 9.7 0.004 26 32.8 184 48 1280 940 0.49 5 9.1 0.002 32 44.6
196 49 1190 930 0.02 16 10.1 0.003 28 70.0 189 50 1280 950 0.51 3
10.1 0.003 26 64.7 198 51 1220 940 0.55 5 10.0 0.003 22 32.0 197 52
1200 940 0.47 14 9.9 0.004 24 48.8 205 53 1280 940 0.55 3 9.9 0.004
21 57.5 186 54 1200 950 0.50 4 9.0 0.002 27 40.5 194 Carburized
layer Limit Fatigue Coarsening austenite compression Machineability
life temp. grain size rate VL1000 (rel. No. (.degree. C.) number
(%) (m/min) value) Remarks 36 >1050 8.9 55 52 3.4 Inv. ex. 37
>1050 8.9 54 53 3.0 38 >1050 8.6 53 53 3.2 39 >1050 9.2 55
52 3.3 40 >1050 8.5 56 52 3.7 41 >1050 9.8 54 52 2.6 42
>1050 8.8 58 43 3.1 43 >1050 9.5 54 42 3.9 44 >1050 9.1 55
41 3.1 45 >1050 9.4 57 41 3.3 46 >1050 8.2 56 43 3.3 47
>1050 9.7 53 47 3.0 48 >1050 8.4 51 47 3.3 48 >1050 8.4 52
48 3.1 49 >1050 8.1 51 52 3.1 50 >1050 8.7 52 50 3.5 51
>1050 9.2 52 47 3.5 52 >1050 8.2 53 45 3.8 53 >1050 8.1 51
48 3.2 54 >1050 9.0 53 49 3.5
TABLE-US-00006 TABLE 6 Hot working Ferrite Heating Finishing
Cooling Bainite grain AlN Ti precipitate Sulfide temp. temp. rate
fraction size precipitation max. density No. (.degree. C.)
(.degree. C.) (.degree. C./s) (%) number (%) size .mu.m (/mm.sup.2)
55 1210 900 0.47 0 10.3 0.003 -- 70.5 56 1200 930 0.51 0 10.4 0.002
22 46.9 57 1220 930 0.45 0 9.8 0.003 27 45.1 58 1210 930 0.53 0 9.1
0.004 28 58.9 59 1190 950 0.56 0 9.1 0.003 -- 126.6 60 1220 950
0.52 0 8.9 0.004 -- 149.5 61 1000 930 0.52 0 10.3 0.003 52 22.8 62
980 940 0.46 0 9.7 0.003 54 57.7 63 1000 940 0.56 0 9.2 0.003 52
12.9 64 980 940 0.57 0 10.4 0.003 55 48.4 65 980 940 0.49 0 10.4
0.003 49 48.1 66 1000 950 0.46 0 9.3 0.003 52 21.2 67 980 940 0.50
0 10.5 0.003 53 54.4 68 1000 940 0.52 0 9.6 0.004 52 24.2 69 980
950 0.51 0 9.2 0.004 56 41.6 70 980 950 0.52 0 10.0 0.003 55 35.4
71 1210 940 0.54 0 8.9 0.003 61 23.7 72 1240 940 0.49 0 9.8 0.003
56 20.7 73 1260 940 0.53 0 9.3 0.003 36 17.0 74 1270 940 0.48 0 9.0
0.002 40 42.7 75 1210 940 0.51 0 9.7 0.003 70 42.8 76 1240 930 0.45
0 10.2 0.004 59 51.9 77 1270 930 0.55 0 10.3 0.003 -- 76.5 78 1180
950 0.47 0 9.7 0.003 76 25.3 79 1200 930 0.47 0 9.8 0.004 31 55.5
80 1200 940 0.50 0 10.1 0.003 24 34.7 81 1200 930 1.50 35 9.9 0.002
25 54.6 82 1200 1030 0.56 0 7.0 0.002 23 40.1 83 1200 850 0.56 0
12.0 0.002 23 40.1 84 1190 930 0.54 0 8.9 0.003 24 48.2 85 1280 940
0.56 0 10.0 0.003 23 56.9 86 1230 930 0.46 0 10.0 0.004 28 54.5 87
1200 900 0.46 0 10.5 0.003 -- 75.4 88 1230 940 0.52 0 9.9 0.003 23
132.5 89 1250 940 0.56 0 9.9 0.003 24 116.2 Carburized layer Limit
Fatigue Vickers Coarsening austenite comp. Machineability life
hardness temp. grain size rate VL1000 (rel. No. (HV) (.degree. C.)
number (%) (m/min) value) Remarks 55 165 950 3.7 58 40 1.0 Comp.
ex. 56 191 >1050 8.1 50 30 2.6 57 195 >1050 8.2 51 30 2.6 58
176 >1050 8.5 50 33 2.8 59 160 910 3.5 45 47 60 162 910 3.7 43
49 61 190 910 4.9 59 46 3.2 62 181 920 3.4 56 53 3.3 63 183 910 3.0
59 48 3.1 64 193 910 4.5 56 55 2.9 65 183 920 4.1 56 52 2.6 66 177
910 4.3 58 47 3.7 67 181 920 4.5 55 52 3.6 68 172 910 3.4 58 50 3.2
69 180 910 4.9 54 55 3.6 70 189 920 3.1 54 53 2.7 71 188 930 3.7 50
25 2.7 72 176 930 3.5 52 26 2.6 73 180 >1050 9.8 51 25 2.7 74
194 >1050 9.4 45 35 75 193 930 3.7 44 34 76 189 920 3.7 46 35 77
165 910 3.0 58 50 1.1 78 203 910 3.2 30 30 79 205 910 3.4 32 30 80
179 910 4.0 57 53 0.3 81 220 930 3.4 30 30 82 184 910 3.5 53 54 1.2
83 184 910 3.5 53 54 1.3 84 194 >1050 8.6 47 25 85 191 >1050
8.1 46 28 86 205 >1050 9.0 45 25 87 175 910 3.7 50 35 1.2 88 200
>1050 9.1 41 43 89 201 >1050 8.5 41 43
[0181] It is clear that the crystal grain coarsening temperature of
the invention examples is 990.degree. C. or more, the .gamma.
grains of a 950.degree. C. carburized material are fine, regular
grains, and the rolling contact fatigue characteristic is also
superior. Regarding the cold forgeability and machinability as
well, it is clear that they are superior compared with the
comparative examples of similar amounts of S.
[0182] On the other hand, the comparative example of No. 55
corresponds to SCr420 prescribed by the JIS. It does not contain
Ti, Mg, Zr, or Ca, so has a low coarsening temperature and coarse
.gamma. grains.
[0183] Further, Nos. 56 to 58 exhibit effects of prevention of
coarse grains by Ti, but do not contain Ti, Mg, Zr, or Ca, so have
inferior machinability and furthermore insufficient cold
forgeability.
[0184] Nos. 59 and 60 are examples where the S is increased to try
to improve the machinability, but do not contain Ti, Mg, Zr, or Ca,
so have elongated sulfides and inferior cold forgeabilities.
[0185] Nos. 84 to 89 are examples where Mo and Nb are added and the
quenchability is improved, while No. 87 corresponds to SCM420
prescribed by the JIS. However, No. 87 does not contain Ti, Mg, Zr,
or Ca, so has a low coarsening temperature and coarse .gamma.
grains. Further, Nos. 84 to 86, 88, and 89 exhibit effects of
prevention of coarse grains by Ti, but do not contain Ti, Mg, Zr,
or Ca, so have inferior machinability and, furthermore,
insufficient cold forgeability.
[0186] Nos. 71 to 76 have large contents of N, coarse Ti
precipitates, and remarkable formation of coarse grains. Further,
Nos. 71 to 73 have reduced rolling contact fatigue characteristics
of carburized parts, while Nos. 74 to 76 are examples inferior in
cold forgeability and not subjected to rolling contact fatigue
tests.
[0187] No. 80 has a large O content, formation of coarse grains,
and no good rolling contact fatigue characteristic as well.
[0188] No. 77 has a small Ti content and a small pinning effect of
Ti, so has a reduced coarsening temperature.
[0189] No. 78 has a large Ti content, coarse Ti precipitates,
reduced coarsening temperature, and degraded cold workability due
to TiC precipitation hardening. Further, No. 78 has insufficient
solubilization of Ti precipitates and reduced rolling contact
fatigue characteristic of carburized parts.
[0190] No. 79 has a large Nb content, degraded cold workability due
to precipitation hardening, and inferior prevention of coarse
grains.
[0191] Nos. 61 to 70 have low heating temperatures, insufficient
solid solutions of Ti precipitates and Nb precipitates, and
inferior effects of prevention of coarse grains.
[0192] No. 81 has a fast cooling rate after hot rolling, increased
bainite structural fraction after hot working, and formation of
coarse grains.
[0193] No. 82 has a high finishing temperature in hot working,
coarse ferrite crystal grain size, and degraded prevention of
coarse grains.
[0194] No. 83 has a low finishing temperature in hot working, a
fine ferrite crystal grain size, and inferior prevention of coarse
grains.
* * * * *