Note: Descriptions are shown in the official language in which they were submitted.
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DESCRIPTION
Title of Invention
CASE-HARDENED STEEL SUPERIOR IN COLD WORKABILITY,
MACHINABILITY, AND FATIGUE CHARACTERISTICS AFTER
CARBURIZED QUENCHING AND METHOD OF PRODUCTION OF SAME
Technical Field
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
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.
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.
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
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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.
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.
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 C or so
carburization temperature is raised to a 990 to 1090 C
temperature region. For this reason, with deep
carburization, coarse grains easily form.
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.
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
Citation List
Patent Literature
PTL 1: Japanese Patent Publication (A) No. 11-335777
PTL 2: Japanese Patent Publication (A) No. 2001-
303174
PTL 3: Japanese Patent Publication (A) No. 2004-
183064
PTL 4: Japanese Patent Publication (A) No. 2004-
204263
PTL 5: Japanese Patent Publication (A) No. 2005-
240175
Summary of Invention
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Technical Problem
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.
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 machineability (ease of cutting of
material).
In the past, to improve the machineability, it has
been known to be effective to add Pb, S, and other
elements improving the machineability.
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.
Further, S forms MnS etc. in the steel to improve
the machineability, 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.
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,
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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
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 machineability, 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.
The gist of the present invention is as follows.
(1) Case-hardened steel consisting of, by mass,
C: 0.1 to 0.22%,
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,
0: 0.0025% or less,
further consisting of one or more of
Mg: 0.0002 to 0.003%,
Zr: 0.0002 to 0.01%, and
Ca: 0.0002 to 0.005%,
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having a balance of iron and unavoidable impurities,
limiting an amount of precipitation of AIN to 0.01% or less,
having a density d (/mm2) of sulfides of a circle equivalent
diameter of over 20 lam and an aspect ratio of over 3 and a content
of S [S] (mass%) satisfying
d.1700 [S] +20, and
having a maximum size of Ti-based precipitates of 40 pm or
less.
(2) Case-hardened steel consisting of, by mass%,
C: 0.1 to 0.22%,
Si: 0.01 to 1.5%,
Mn: 0.3 to 1.8%,
S: 0.001 to 0.15%,
Cr: 0.4 to 2.0%,
Ti: 0.05 to 0.2%, and
Nb: less than 0.04,
limiting
Al: 0.04% or less,
N: 0.0050% or less,
P: 0.025% or less,
0: 0.0025% or less,
further consisting of one or more of
Mg: 0.0002 to 0.003%,
Zr: 0.0002 to 0.01%, and
Ca: 0.0002 to 0.005%,
having a balance of iron and unavoidable impurities,
limiting an amount of precipitation of AIN to 0.01% or less,
having a density d (/1um2) of sulfides of a circle equivalent
diameter of over 20 pm and an aspect ratio of over 3 and a content
of S [S] (mass%) satisfying
c:11700 [S] +20, and
having a maximum size of Ti-based precipitates of 40 pm or
less.
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(3) Case-hardened steel consisting of, by mass%,
C: 0.1 to 0.22%,
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,
0: 0.0025% or less,
further consisting of one or more of
Mg: 0.0002 to 0.003%,
Zr: 0.0002 to 0.01%, and
Ca: 0.0002 to 0.005%,
further consisting of 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,
having a balance of iron and unavoidable impurities,
limiting an amount of precipitation of AIN to 0.01% or less,
having a density d (/mm2) of sulfides of a circle equivalent
diameter of over 20 pm and an aspect ratio of over 3 and a content
of S [S] (mass%) satisfying
d_.1700[S]+20, and
having a maximum size of Ti-based precipitates of 40 [Amor
less.
(4) Case-hardened steel consisting of, by mass%,
C: 0.1 to 0.22%,
Si: 0.01 to 1.5%,
Mn: 0.3 to 1.8%,
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S: 0.001 to 0.15%,
Cr: 0.4 to 2.0%,
Ti: 0.05 to 0.2%, and
Nb: less than 0.04,
limiting
Al: 0.04% or less,
N: 0.0050% or less,
P: 0.025% or less,
0: 0.0025% or less,
further consisting of one or more of
Mg: 0.0002 to 0.003%,
Zr: 0.0002 to 0.01%, and
Ca: 0.0002 to 0.005%,
further consisting of 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,
having a balance of iron and unavoidable impurities,
limiting an amount of precipitation of AIN to 0.01% or less,
having a density d (Atra2) of sulfides of a circle equivalent
diameter of over 20 pm and an aspect ratio of over 3 and a content
of S [S] (mass%) satisfying
d<1700 [S] +20, and
having a maximum size of Ti-based precipitates of 40 pm or
less.
(5) Case-hardened steel as set forth in any one of (1) to (4),
wherein a structural fraction of bainite is limited to 30% or less.
(6) Case-hardened steel as set forth in any one of (1) to (5),
wherein a grain number of ferrite is 8 to 11 as defined by JIS G
0551.
(7) A method of production of a case-hardened steel, comprising
heating a steel material comprised of the ingredients of any of
(1) to (4) to 1150 C or more, hot working it at a finishing
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terrperature of 840 to 1000 C, and cooling it in a 800 to 500 C
temperature range by lcVs or less.
10
Advantageous Effects of Invention
The case-hardened steel of the present invention is
superior in forgeability, machineability, 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.
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.
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
FIG. 1 is a view for explaining a balance of
machineability and cold workability of the present
invention.
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FIG. 2 is a view showing a position for measuring a
cooling rate at the time of solidification.
FIG. 3 is a view showing a test piece used for an
upset test.
Description of Embodiments
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.
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 soiubilize the Ti precipitates and Nb
precipitates at the time of heating in hot working.
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.
However, the temperature at which AIN 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
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time of cooling after that, AIN 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 AIN contained in the case-
hardened steel.
Therefore, according to the case-hardened steel of
the present invention limited in amount of precipitation
of AIN, it is possible to utilize fine Ti precipitates
and Nb precipitates to prevent the formation of coarse
grains at the time of carburized quenching.
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.
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.
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.
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.
Further, if the ferrite grains of the steel material
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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.
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.
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.
Therefore, to improve the machinability, it is
desirable to cause formation of sulfides in the steel.
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.
Note that, to suppress coarsening of the sulfides,
it is preferable to control the solidification speed at
the time of casting.
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 machineability.
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Therefore, to improve the machineability, it is
important to not only add S, but also control the soft
sulfides in the molten steel to which Ti is added.
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.
The machineability and cold workability will be
further explained.
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.
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.
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.
On the other hand, from the viewpoint of the
machineability, 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
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sulfides, it is possible to achieve both cold
forgeability and machineability (tool life).
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
machineability is important. If increasing the amount of
S, the machineability 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.
FIG. 1 compares the relationship of machineability
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 C) while achieving both cold workability
(limit compression rate) and machineability (drillability
VL1000). In FIG. 1, the further to the top right, the
better the balance of machineability and cold workability
of the material.
Below, the present invention will be explained in
detail.
First, the composition of ingredients will be
explained. Below, "mass%" will be simply described as
"%".
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.
Si is an element effective for deoxidation of steel.
In the present invention, 0.01% or more is added.
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,
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.
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.
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%.
S is an element forming MnS in steel and improving
the machineability. In the present invention, to improve
the machineability, 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
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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.
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 machineability, the amount of S is
preferably made 0.02% or more.
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 7 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.
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
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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.
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 AIN 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.
N is an element forming nitrides. In the present
invention, to suppress the formation of coarse TiN and
AIN, the upper limit of the amount of N is made 0.0050%.
This is because coarse TiN and AIN 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.
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.
0 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
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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.
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 0
is preferably made 0.0020%.
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 0 content is preferably limited to 0.0012% or
less.
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.
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.
Further, oxides of Mg finely disperse and form the
nuclei for formation of MnS and other sulfides. To
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
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become somewhat hard and become harder to elongate due to
hot working.
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.
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.
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.
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.
On the other hand, Zr is an expensive element, so
from the viewpoint of the production costs, the upper
CA 02757393 2011-09-30
= - 17 -
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.
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.
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.
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.
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.
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
CA 02757393 2011-09-30
- 18 -
=
is very small, compared with the case of not adding Nb,
the addition is extremely effective for prevention of
coarse grains.
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%.
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%.
Furthermore, to improve the strength and
quenchability of the steel, one or more of Mo, Ni, V, B,
and Nb may be added.
Mo is an element improving the strength and
quenchability of steel. In the present invention, it is
effective for increasing the amount of residual y 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
CA 02757393 2011-09-30
= - 19 -
=
hardness causes the machinability and cold forgeability
to be degraded in some cases.
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.
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%.
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%.
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%.
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
CA 02757393 2011-09-30
- 20 -
=
and rolling parts in the process of rolling contact
fatigue is particularly large when the residual austenite
(residual y) at the surface layer after carburization is
30 to 40%. To control the residual amount of y 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.
To make the residual amount of y 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.
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.
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.
Next, among the precipitates included in the case-
hardened steel of the present invention, AIN and sulfides
will be explained.
AIN forms the precipitation nuclei for Ti
precipitates and Nb precipitates and inhibits the
formation of fine precipitates. Therefore, in the present
CA 02757393 2011-09-30
= - 21 -
invention, it is necessary to limit the amount of
precipitation of AIN included in the case-hardened steel.
If the amount of precipitation of AIN is excessive,
coarse grains are liable to be formed at the time of
carburized quenching, so the amount of precipitation of
AIN in the case-hardened steel is limited to 0.01% or
less. The preferable upper limit of the amount of
precipitation of AIN is 0.005%.
To suppress the amount of precipitation of AIN 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.
Note that, the amount of precipitation of AIN can be
measured by chemical analysis of the extraction residue.
The extraction residue is obtained by etching the steel
by a bromine methanol solution and filtering by a 0.2 pm
filter. Note that, even if using a 0.2 pm filter, in the
process of filtration, the precipitates cause the filter
to clog, so extraction of 0.2 pm or smaller fine
precipitates is also possible.
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.
The inventors etc. studied the relationship between
the content of S, the size and shape of MnS inclusions,
and the machinability and cold workability.
As a result, it was learned that when MnS inclusions
observed under an optical microscope have a equivalent
circle diameter of over 20 pm and an aspect ratio of over
3, they become origin of fracture at the time of cold
working.
The equivalent circle diameter of an MnS inclusion
CA 02757393 2011-09-30
= - 22 -
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.
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 mmx1 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.
The number of MnS inclusions with a equivalent
circle diameter over 20 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:
d5.1700[S]+20 (/rae)
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.
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 pm
is preferable.
CA 02757393 2011-09-30
- 23 -
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 mm2, inspection of
16 fields, and a prediction area of 30000 mm2.
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 mm2).
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.
Next, the structure of the case-hardened steel of
the present invention will be explained.
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.
Suppressing the structural fraction of bainite to
30% or less is also effective for improving the cold
workability.
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
CA 02757393 2011-09-30
- 24 -
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.
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.
Next, the method of production of case-hardened
steel of the present invention will be explained.
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.
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.
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
CA 02757393 2011-09-30
- 25 -
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 C/rain or more.
Preferably it is made 5 C/min or more, more preferably
10 C/min or more. Note that, the cooling rate at the time
of solidification can be confirmed by the secondary
dendrite arm spacing.
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 C
or more high temperature for 10 minutes or more and cause
the Ti and Nb precipitates to solute.
To produce case-hardened steel, the steel material
is heated. If the heating temperature is less than 1150 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.
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 C or more. The preferable lower
limit of the heating temperature is 1180 C or more.
The upper limit of the heating temperature is not
prescribed, but if considering the load of the heating
furnace, 1300 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.
CA 02757393 2011-09-30
- 26 -
If the finishing temperature of the hot working is
less than 840 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 C, hardening occurs and the cold
workability deteriorates. Therefore, the finishing
temperature of hot working is made 840 to 1000 C. Note
that, the preferable range of the finishing temperature
is 900 to 970 C, and the more preferable range is 920 to
950 C.
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 C. Therefore, the
cooling is performed slowly by 1 C/s or less from a 800 C
to 500 C temperature range to promote the formation of Ti
precipitates and Nb precipitates.
If the cooling rate exceeds 1 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 C/s or
less.
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 roiling
line and thereby slowing the cooling may be mentioned.
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,
CA 02757393 2011-09-30
- 27 -
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.
By using the case-hardened steel of the present
invention, hot forging it at for example a 1150 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 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.
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
Steels having the compositions of ingredients shown
in Tables 1 to 3 were produced and cast at solidification
cooling rates of 10 to 11 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.
The solidification cooling rate was adjusted in
advance based on data analyzing the relationship between
the cooling conditions and solidification cooling rate
CA 02757393 2011-09-30
- 28 -
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 C/min in range. Some of the cast blooms were bloomed
in accordance with need.
Table 1
Chemical ingredients (mass%)
No.
Remarks
C Si Mn P S Cr Ti Al N 0 Zr -t-,1g Ca Nb Mo Ni V B
_
,
1 0.210.19 1.30 0.018 0.011-1.06 0.13 0.026-0.0030 0.0011 0.0024
Inv. ex.
2 0.20 D.200.38 0.022 0.0141-1.10 0.14 0.024-0.0047 0.0014 0.0005
_ _ _
r
3 0.21 0.19 0.98 0.014 0.015 1.20 0.06 0.0350.0033 0.0014 0.0025
-4 0.19 0.18 0.84 0.014 0.014-1.28 0.08 0.027 0.0045 0.0012 0.0007-
0.0006 .
0.19 0.21 0.88 0.005 0.016 1.22 0.08 0.038-0.0026-0.0015 0.0013 0.0020 _
6 O.200.19 0.58 0.014 0.013 1.13 0.06 0.018 0.0029-0.0014 0.0008
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 0.0018
0.0013 .
-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
0.19-0.21 1.04 0.017 0.031 1.23 0.11_0.0380.0028_0.00140.0005 _
0
11 0.19 0.25 1.63 0.018 0.029 1.05_0.07 0.020 0.00310.0012 0.0015
,
12 0.22 0.21 0.81 0.016 0.028 1.22 0.11 0.016Ø0032-0.0011 0.0012-
=o
,
13 0.20 0.19 1.60Ø009 0.026_1.15 0.14 0.028Ø0026 0.0012-0.0016 0.0015
0.0014, n.)
-.3
14 0.19 0.19 0.99_0.018 0.029_1.15 0.15 0.034 0.0027-0.0010 0.0018_0.0011
oi
-.3
0.32 0.22 0.38_0.018 0.048 1.22Ø06 0.030 0.0032-0.0010
0.0015 0.0013 w
_
l0
16 0.21 0.25 0.32 0.024 0.026 1.160.1O 0.034 0.0026 0.0012 0.0018 0.0003-
0.0019 w
.
.
17 '0.22 0.18 1.77-0.009 0.015 1.21 0.12 0.022 0.0028 0.0011 0.024
_
, n.)
18 0.21 0.20-0.54-$0.025 0.0131.21 0.12 0.014 0.0034 0.0014 i 0.021
o
19 0.19 0.230.86 0.005 0.012 _1.22 0.09 0.027 0.0035 0.0012 0.0004 0.012
i-,
_
N CL.
0.21 0.22 1.31-0.023 0.016 1.28 0.11 0.023 0.00340.0011 0.0012
0.019 Lo
21 0.21 0.25 0.57 0.016 0.013 1.13 0.14 0.037 0.0047 0.0015- W 0.0006
0.013
22 -0.19 0.19-1.19 0.011 0.011-1.22 0.08 0.021 0.0041 0.0010 0.0008
_0.0004_0.013 o
23 0.22 O.190.57 0.013 0.013 -1.13 0.05 0.019_0.0025 0.0014-0.0030_0.0015
0.016 .
24 0.18 0.24 0.74 0.016-0.011 1.16-0.12 0.0170.0032 0.0011' 0.0014
0.0015 0.025 _
0.210.23 1.15 0.019 0.015 1.18 0.05 0.018-0.0032H0.0014 0.0027-0.0017 0.0009
0.014 0.13
_ _ ,
26 0.22 0.21 0.48 0.013 0.013 1.27 0.07 0.025_0.0031 0.0014 0.0017 0.0007
0.0004 0.014 0.30
_
-27 0.20 0.20 0.45 0.015-0.010-1.150.09 0.037 0.0036 0.0010 0.0010
0.00050.0011-0.020
, _
.
28 0.20 0.22 1.11 0.022_0.017_1.12 0.13 0.024 0.0048 0.0015 0.0006 0.0016
0.0013 0.012 0.0015
. . .
-29 0.22 0.20 1.19,0.016_0.025 1.26'0.09 0.034 0.00290.0013- 0.014
0.21 0.24 1.08 0.008 0.025-1.080.15- 0.036-0.0030 0.0011 0.010 .
_
-31 _0.21 0.25 1.16 0.011-0.031-1.28-0.05-0.039-0.0028 0.0010-0.0022- 0.022
,
32 0.19 0.23 1.73 0.009 0.040 1.23 0.06 0.016-0.0041 0.0015_0.0014 0.014
_
33 0.22 0.25 0.74 0.007 0.025 1.18 0.10-0.008-0.0026- _0.0010 0.0011
0.016
34 0.21 1.22-1.22 0.009 0.030 1.130.15 0.009- _0.0038
0.0015 0.0008 0.023 _
0.18 0.22-1.35 0.011-0.032 1.25 0.14-0.013-0.0039 0.0011 0.0020 0.0015 0.0009
0.024
.- ,...._
,
Table 2
Chemical ________________________________ ingredients (mass%)
No. Remarks
C Si Mn P s Cr Ti Al N 0 Zr Mg Ca Nb Mo Ni V B
-1
36 0.19 0.19 1.72 0.009 0.029 -0.55 0.12 0.039 -0.0049 0.0015Ø0015 0.0006
0.019 r Inv. ex.
37 0.22 0.18 1.68 0.024 0.028-1.06 0.06 0.023 -0.0030 0.0012
0.0018 0.0012 0.020 _
_
_______________________________________________________________________________
________
38:0.21 0.20 0.32 0.010 0.028 1.08 0.10 0.032 -0.0039 0.00130.0013
0.0006,0.0013 0.019 ,
39 0.20 0.21 1.02 0.018 0.030 1.05 0.09 -0.010 0.0046 0.0011 0.0019 0.0012
0.0013 0.020 0.21
40 10.19 0.20 0.33 0.025 0.035 0.620.12 0.022'0.0045 0.0015 0.0013 0.0004
0.0013-0.016 ' 0.95-
_
41 0.19 0.20
1.16 -0.013 0.028 1.20'0.09 0.032 '0.0049 0.0015 0.0021 -0.0010 '0.00130.022-
0.0016
_
_
42 0.19 0.23 1.37 -0.012 0.017 1.08:0.13 0.032 0.0035 0.0013 0.0017 0.141-
T
43 '0.21 0.18 1.00 0.016 0.013 1.07_0.11 _0.019 0.0044 0.0014
0.0004 0.16 _
44 0.20 0.25 1.69 0.020 0.016 1.15 0.05 0.035 0.0031 0.0012 0.0010 0.14
1 _
45 0.21 0.20 -0.76 -0.019 0.017 1.06-0.08 0.033 0.0031 0.0013 0.0012
10.0017 0.14 _
46 0.20 ,0.22 1.52 0.015 0.015 1.30 0.10 0.018 0.0048 0.0013 0.0017 0.0007
0.12
47-0.19 0.25 1.34 0.012 0.027 1.21_ _0.12 0.012-0.0041 0.0011
0.020 0.13 0
-,.
48 0.22 O.220.64 0.014 0.027 1.11_0.13 ,0.032 _
_0.0050 0.0014 0.011 0.16 _
48 0.19'0.21-0.45 0.010 -0.027 1.28 0.13 0.019 0.0026 0.0010 0.0010 0.022
0.16 0
n.)
49 0.21 0.21-0.56 0.021 0.044 1.620.15 0.039 0.0033 0.0010 0.0020 0.019
0.13 --3
oi
500.20'0.18 .-1.02 0.023 0.054 1.15 0.11
0.019 0.0033 0.0013 0.0003 0.011 0.15 --3
,
_______________________________________________________________________________
________
51 0.22 0.23 -0.75 0.022 0.026 1.2-0.06 0.019 0.0047 0.0010-
0.0005 0.014 0.16 w
.
_______________________________________________________________________________
________ _ l0
52 0.21-0.18 0.38 0.017 0.028 0.72 0.09 0.028 0.0031 0.0012 0.0018 0.0008
0.0012-0.013-0.92- 1 w
53 0.21'O.20 -0.82 -0.018 0.029 1.120.090.035 0.0040 0.0012 0.0025 0.0004
'0.019 0.12 n.)
54 0.21 0.23 -0.56 0.011 ,0.031
1.08 0.09,0.013 0.0049-0.0013 0.0010 0.0017 0.014_ 0.15
i-,
O
I
ko
W
o
-- -
Table 3
Chemical ingredients (mass%)
No.
Remarks
C Si Mn P S Cr Ti Al N 0
Zr Mg Ca Nb Mo Ni V B
_
55 0.19 0.24 1.72 0.013 0.012 1.10 0.035 0.0126 0.0014Comp. ex.
56 0.18 0.23 0.98 0.013 0.013 1.14 0.08 -0.018 0.0031 0.0011 ___
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 O.241.72 0.008 0.012 1.26 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
0
65 0.20 0.24 0.71 0.023 0.030 1.13 0.14 0.020 0.0032 0.0013 0.0010
o
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 n.)
67 0.20 0.22 1.52 0.009 0.026 1.10,0.09 0.015 0.0049 0.0014 0.0024
0.00160.023 oi
=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 w
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 w
70 0.20 0.20 1.11 0.019 0.031 1.24,0.10 0.033 0.00310.0014 0.0016
0.023 n.)
I
71 0.20 0.24 1.02 0.022 0.017 i.090.12 0.013 0.0124'0.0012 0.0011
o
i-,
1
72 0.21,0.22 0.87 0.018 0.017 1.25 0.11 0.014 0.0145 0.0012
0.0006 1--,
73 0.19 0.21 1.02 0.019 0.013 1.26 0.10 0.018 0.0086 0.0011 0.0011
o
1
74 0.18 0.20 0.34 0.015 0.026 1.15 0.12 0.031 0.0098 0.0015 0.0024_ =.
_ w
75 0.19 0.22 0.33 0.008 0.030 1.16-0.06 0.030 0.0146,0.0012
0.0015 o
_
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
r
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 0.00400.0015"0.0004
83 0.21 0.20 0.77 0,023 0.031 1.10 :0.05 0.007 -0.0040-0.00150.0004 _
84 0.21 O.22'1.31 0.007 0.016 1.25 0.15 0.016 0.0031 0.00120.14
_ ___ ___
_
85 0.22 0.24 1.79 0.018 0.010 1.10 0.11 0.020 0.00380.0011- 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 1.070.035 0.0126 0.0012 -0.15
-88 -0.200.19 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
___ ___
CA 02757393 2011-09-30
- 32 -
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
machineability. The amounts of precipitation of AIN were
found by chemical analysis.
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, AIN precipitation, Ti
precipitate maximum sizes, and Vickers hardnesses. Note
that, the cooling rate is the cooling rate in the 500 to
800 C range. This was found from the time required for
cooling from 800 C to 500 C.
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.
In the longitudinal direction cross-section of each
test piece, 16 fields of regions of standard inspection
areas of 100 mm2 (10 mmx10 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 1000X.
This was repeated 16 times for the 16 fields of the
standard inspection areas of 100 mm2. 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
CA 02757393 2011-09-30
- 33 -
and short axis is found and used as the size of the
precipitate.
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 mm2) were found.
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.
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.
The machineability 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.
The speed of the drill was fixed at 10 to 70 m/min
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.
CA 02757393 2011-09-30
- 34 -
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 machineability
the material is evaluated as.
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.
The carburization simulation comprised heat
treatment heating a test piece to 910 to 1010 C, holding
it there for 5 hours, then water cooling it. The old
austenite grain size was measured in accordance with JIS
G 0551.
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 400X 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.
The heating temperature of the carburized quenching
treatment is usually 930 to 950 C, so a test piece with a
coarsening temperature of 950 C or less was judged to be
inferior in crystal grain coarsening characteristic.
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 C, holding it there fore 5
hours, and quenching it in oil of a temperature of 130 C.
Furthermore, the steel was held at 180 C for 2 hours and
CA 02757393 2011-09-30
- 35 -
tempered. These carburized quenched materials were
investigated for the y granularity (carburized layer
austenite grain size number) of the carburized layers
based on JIS G 0551.
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 Ln
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.
The results of these investigations are summarized
in Tables 4 to 6. The rolling fatigue life shows the
relative value of the Ln life of each material indexed to
the L10 life of No. 55 (comparative example) as "1".
Table 4
-No. Hot working Bainite Ferrite AIN Ti precipi- Sulfide Vickers
Coarsen- Carburized Limit Machine-Fatigue
Heating Finishing Cooling fraction grain precipi- tate max. density hardness
ing layer comp. ability life
temp. temp. rate (%) size tation size gm (iRmlz)
(Hy) temp. austenite rate VL1000 (rel. Remarks
( C) ( C) ( C/s) number (%) ( C) grain
size (%) (m/min) value)
number
.
.
1 1270 930 0.50 o -9.8 0.003 21 16.0 160
>1050 9.8 58 48 ,3.5 Inv. ex.
-2 -1260 950 Ø53 0 9.0 0.004 23
29.5 183 .>1050 8.8 56 46 3.7
3 7E190 940 0.53 0 .9.4 0.004 26 26.6 _ 187
>1050 9.9 56 45 3.0
4 -1210 940 0.53 o 9.4 0.004 25 13.2 164
>1050 8.9 55 -49 3.4
.
-
1260 =940 0.55 0 ,.9.8 0.003 23
11.6 185 >1050 6.7 56 46 3.5
.6 1220 930 0.53 0 _9.2 0.004 27 27.9 194
>1050 8.6 57 46 2.8
7 -1190 940 0.48 0 10.5 0.003 29
25.6 188 r>1050 8.0 55 .49 2.6
-
8 -1180 940 0.57 0 .9.2 0.004 26
47.5 172 >1050 9.7 55 55 2.5
9 1220 930 0.55 0 10.2 0.003 31
52.0 183 .>1050 8.9 54 51 3.8
_
1250 940 0.49 0 '-.9.5 0.004 27 36.2 188
>1050 8.5 53 54 3.4
11 -1270 .930 0.48 0 9.8 0.003 24 53.3 176
>1050 10.0 56 ,50 3.0 0
12 1230 950 Ø56 0 .9.8 0.003 30 37.3 178
'>1050 8.4 54 53 3.2
'
13 -1200 930 0.47 o 9.4 0.003 24 51.4 167
,>1050 8.4 56 50 2.6 =c)
t\.)
14 1270 930 0.46 o .10.2 0.002 32
39.2 163 >1050 8.5 54 .51 3.2 --3
_
1190 940 0.52 -5 .9.0 0.004 25 30.2 192
>1050 9.8 52 51 3.4 tn
--3
16 1240 930 0.48 0 9.3 0.003 24 41.8 177
>1050 ,9.9 56 52 2.8 w
17 1220 940 0.47 -0 .10.1 0.002 25
27.9 173 >1050 9.1 59 46 2.7 t.D
w
18 1250 950 0.46 0 10.4 0.003 31
23.2 178 >1050 9.6 56 49 3.7 I
19 1190 940 0.57 '..0 10 0 0.002 23
_ - 19.7 174 >1050 9.6
56 46 3.1 t\.)
_
c)
1270 940 0.56 0 10.0 0.003 -29
10.9 1 W
80 >1050 9.4 58 48 2.5 61 l-
21 1230 930 0.52 .0 9.2 0.002 .26 21.7 194
>1050 10.0 60 48 3.2
1
22 1190 950 0.45 0 9.5 0.004 27 22.5 179 -
>1050 8.4 58 50 2.5 I 0
23 1220 930 h I 0.57 0 ,9.7 0.004 27 25.6 181
>1050 9.0 59 48 2.2 .
24 1230 940 0.50 0 10.5 0.004 27
16.4 192 >1050 8.9 56 46 2.9 w
.
c)
1250 930 0.56 0 .10.2 0.004 29
28.3 174 >1050 9.3 59 50 3.3
,
26 -1190 930 0.53 0 9.0 0.003 ,21 20.2 193
.>1050 6.4 55 49 2.6
27 1250 940 0.52 -0 10.3 0.004 ,27
125.1 175 >1050 8.7 59 48 _
2.6
28 1230 940 _0.51 0 10.4 0.003 30
10.9 .184 >1050 8.4 '55 46 2.6
29 1200 940 0.52 0 9.6 0.002 27 59.5 180
>1050 9.1 54 50 72.3
20 1200 =940 0.46 0 9.6 0.003 29 .46.8 177
>1050 9.0 56 53 3.7
31 1230 940 0.56 0 10.4 0.003 22
57.1 -175 >1050 9.5 49 58 2.6
_
32 1270 930 -0.48 0 9.8 0.003 25 60.6 189
>1050 _.8.6 -54 53 3.5
33 ,1200 950 0.56 0 9.3 0.004 29 53.3 189
>1050 8.6 -55 -51 3.3
--
34 1200 940 0.45 0 9.8 0.003 2 .8
50.0 191 >1050 9.7 54 .50 3.7
_
_
1280 940 0.49 0 9.0 0.002 .123 _38.1 ,176
>1050 -9.5 54 _53 3.2
.- -
Table 5
Hot working
Carburized
Ferrite AIN Ti precipi- Coarsen-
Limit Machine- Fatigue
Bainite Sulfide Vickers layer
Heating Finishing Cooling grain precipi- tate ing
compres- ability life
No. fraction density hardness temp.
austenite Remarks
temp. temp. rate size tation max. size
sion rate VL1000 (rel.
(%) (/rme) (HV) grain
size
( C) ( C) ( C/s) number (%) pm ( C)
(%) (m/min) value)
.number
36 1210 940 0.52 0 8.8 0.003 26 53.9 173 .>1050
8.9 55 52 3.4 Inv. ex.
_
37 1270 _950 0.48 0 10.4 0.003 27
41.6 178 >1050 8.9 54 53 3.0
38 1190 950 0.46 0 9.7 0.003 23 45.0 173 >1050 8.6
53 53 3.2
39 1260 .940 0.56 0 8.9 0.003 27 36.9 194
>1050 9.2 55 -52 3.3
40 1240 950 0.47 0 9.5 0.003 24 59.7 187 >1050 8.5
56 52 3.7
_
_
41 1200 940 0.46 _0 9.6 0.004 30 40.3 -174
>1050 9.8 54 52 2.6
42 1200 930 0.49 4 9.5 0.003 20 15.2 201 >1050 8.8
58 -43 3.1
43 -_1280 _950 0.50 4 10.2 0.003 29 15.2 193
>1050 9.5 54 42 3.9
44 1260 _940 0.45 4 9.9 0.004 27 29.3 185
>1050 9.1 55 41 3.1
45 1260 940 0.57 5 9.5 0.003 22 15.1 -188 >1050 9.4
57 41 3.3
0
46 1200 .950 0.50 7 9.5 0.002 28 28.9 -188
>1050 8.2 56 43 3.3
4/ 1240 .950 0.47 6 _9.1 0.002 23 47.9 202
>1050 9.7 53 47 3.0
o
48 1250 950 0.56 5 9.7 0.004 26 32.8 -184 >1050 .8.4
51 47 3.3 t\.)
48 1280 940 0.49 5 9.1 0.002 32 44.6 -196 >1050 6.4
52 48 3.1
.
....3
(.n
49 1190 ,930 0.02 16 10.1 0.003 28
70.0 -189 >1050 8.1 51 -52 3.1 -4
50 1280 950 0.51 3 10.1 0.003 26 64.7 -198 >1050 8.7
52 50 3.5 W
t.D
51 1220 940 0.55 5 10.0 0.003 22 32.0 -.197 >1050 9.2
52 47 _3.5 w
-52 1200 940 0.47 14 9.9 0.004 24 48.8 205
>1050 8.2 53 45 3.8
53 1280 940 0.55 3 9.9 0.004 21 57.5 186 .>1050 8.1
51 48 3.2 0
_
(.,.) '-
54 1200 950 0.50 4 9.0 0.002 27 40.5 194 >1050 ,9.0
53 -49 .3.5 --.1
1
I
0
t.D
I
W
0
--
..
Table 6
Hot working
Carburized
Ferrite AIN Coarsen-
Limit Machine- Fatigue
Bainite Ti precipi- Sulfide Vickers ing
layer
Heating Finishing Cooling grain precipi-
comp. ability life
No. fraction tate max. density
hardness temp. austenite Remarks
temp. temp. rate
(%) size tation
size m (/mm2) (HV) grain
sire rate VL1000 (rel.
( C) ( C) m/s) number (%) 0'0)
number
00 cm/min) value)
,
-
55 1210 900 0.4/ 0 10.3 0.003 - 70.5 165
,950 3.7 58 40 1.0 Comp. ex.
56 1200 930 U.51 0 10.4 0.002 22 -46.9 191>1050 6.1
50 30 12.6
57 1220 930 0.45 0 9.8 -0.003 27 45.1 195 >1050 8.2
51 30 2.6
_
58 1210 930 0.53 0 9.1 0.004 28 58.9 176 >1050 8.5
50 33 2.8
..
59 1190 950 0.56 0 9.1 0.003 - 126.6 160 910 3.5
45 47
60 1220 950 -0.52 0 8.9 1.004 - 149.5 162 910 3.7
43 49
61 1000 930 0.52 0 10.3 0.003 52 22.8 190 910 4.9
59 46 3.2
62 980 940 0.46 0 9.7 0.003 54 57.7 181 920 3.4
56 53 3.3
63 1000 940 0.56 0 9.2 0.003 52 12.9 183 910 3.0
59 48 3.1
64 980 940 0.57 0 10.4 0.003 55 48.4 193 910 4.5
56 55 2.9
0
65 980 940 0.49 0 10.4 0.003 49 48.1 183 920 4.1
36 52 2.6
66-1000 _950 0.46 0 9.3 0.003 52 21.2 177
910 4.3 .58 47 3.7
940
o
67 960 0.50 0 10.5 0.003 53 54.4 181 920 4.5
55 52 3.6 . t\.)
68 -11000 940 Ø52 0 9.6 0.004 52 -24.2 ,172
910 -3.4 58 50 3.2 ....3
.
(.n
69 980 950 Ø51 ,0 9.2 0.004 56 41.6 180 910 4.9
54 55 3.6
_
....3
70 980 950 0.52 0 10.0 0.003 55 35.4 189 920 3.1
54 -53 2.7 . w
tb
71 1210 940 Ø54 0 8.9 0.003 61 23.7 188 930 3.7
50 25 2.7 w
72 1240 940 t 0.49 0 9.8 0.003 56 . 20.7 176
930 -3.5 52 26 2.6
. _
73 1260 940 0.53 0 9.3 0.003 36 1/.0 180 >1050 9.8
51 25 2.7 (..,) o
74 1270 940 0.48 0 9.0 0.002 ,40 42.7 194 >1050 9.4
45 35
'-
75 _1210 940 0.51 0 9.7 0.003 _70 42.8 193
930 3.7 44 34 1
_
I 0
76 1240 930 0.45 0 10.2 0.004 59 51.9 189 920 3.7
46 35 t.D
=- _
77 1270 930 0.55 0 10.3 0.003 - 16.5 165 910 3.0
58 50 1.1 1
w
78 1180 950 0.47 0 9.7 0.003 76 25.3 203 910 3.2
30 30 o
79 1200 930 0.47 0 9.8 0.004 31 55.5 205 910 3.4
32 30 _
BO 1200 940 0.50 0 10.1 0.003 24 34.7 179 910 4.0
57 53 0.3
_
_
81 1200 930 1.50 35 9.9 1 ..002 25 54.6 220
930 3.4 30 30
82 1200 1030 0.56 0 7.0 0.002 23 -40.1 184 910 3.5
53 54 1.2
83 1200 850 0.55 0 12.0 0.002 23 40.1 184 910 3.5
53 54 1.3
.
_
84 1190 -930 0.54 0 -8.9 0.003 24 48.2 194 _>1050 8.6
47 25
85 1280 940 0.56 0 10.0 0.003 23 56.9 191 ...,>1050 8.1
46 28 .
86 1230 930 0.46 0 10.0 0.004 28 _54.5 205 >1050 9.0 -
45 -25
87 1200 900 0.46 0 10.5 1.003 - _75.4 .175 910 3.7
50 35 1.2 .
_._
88 1230 940 0.52 0 9.9 -0.003 23 132.5 200 ->1050 -9.1
41 43
89 1250 940 0.56 0 9.9 0.003 24 _116.2 201 >1050 8.5
41 43
=
CA 02757393 2011-09-30
- 39
It is clear that the crystal grain coarsening
temperature of the invention examples is 990 C or more,
the y grains of a 950 C carburized material are fine,
regular grains, and the rolling contact fatigue
characteristic is also superior. Regarding the cold
forgeability and machineability as well, it is clear that
they are superior compared with the comparative examples
of similar amounts of S.
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 y grains.
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 machineability and furthermore
insufficient cold forgeability.
Nos. 59 and 60 are examples where the S is increased
to try to improve the machineability, but do not contain
Ti, Mg, Zr, or Ca, so have elongated sulfides and
inferior cold forgeabilities.
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 y 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 machineability and, furthermore,
insufficient cold forgeability.
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.
No. 80 has a large 0 content, formation of coarse
CA 02757393 2011-09-30
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grains, and no good rolling contact fatigue
characteristic as well.
No. 77 has a small Ti content and a small pinning
effect of Ti, so has a reduced coarsening temperature.
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.
No. 79 has a large Nb content, degraded cold
workability due to precipitation hardening, and inferior
prevention of coarse grains.
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.
No. 81 has a fast cooling rate after hot rolling,
increased bainite structural fraction after hot working,
and formation of coarse grains.
No. 82 has a high finishing temperature in hot
working, coarse ferrite crystal grain size, and degraded
prevention of coarse grains.
No. 83 has a low finishing temperature in hot
working, a fine ferrite crystal grain size, and inferior
prevention of coarse grains.
