U.S. patent application number 11/706687 was filed with the patent office on 2007-08-30 for steel part having long rolling contact fatigue life and method for producing the same.
This patent application is currently assigned to JFE Steel Corporation,a corporation of Japan. Invention is credited to Masao Goto, Hisashi Harada, Takashi Iwamoto, Hideto Kimura, Hisato Nishisaka, Kunikazu Tomita, Takaaki Toyooka.
Application Number | 20070199632 11/706687 |
Document ID | / |
Family ID | 38180041 |
Filed Date | 2007-08-30 |
United States Patent
Application |
20070199632 |
Kind Code |
A1 |
Iwamoto; Takashi ; et
al. |
August 30, 2007 |
Steel part having long rolling contact fatigue life and method for
producing the same
Abstract
A steel part having a long rolling contact fatigue life and
capable of further increasing the life of a bearing under severer
using condition than usual conditions. The steel part includes
steel having a composition containing 0.7% by mass to 1.1% by mass
of C, 0.5% by mass to 2.0% by mass of Si, 0.4% by mass to 2.5% by
mass of Mn, 1.6% by mass to 5.0% by mass of Cr, 0.1% by mass to
less than 0.5% by mass of Mo, 0.010% by mass to 0.050% by mass of
Al, less than 0.0015% by mass of Sb as an impurity, and the balance
composed of Fe and inevitable impurities, the steel being hardened
and tempered. In the steel structure of a portion from the surface
to a depth of 5 mm, residual cementite has a grain diameter of 0.05
to 1.5 .mu.m, prior austenite has a grain diameter of 30 .mu.m or
less, and the ratio by volume of the residual austenite is less
than 25%.
Inventors: |
Iwamoto; Takashi;
(Kurashiki, JP) ; Tomita; Kunikazu; (Kurashiki,
JP) ; Kimura; Hideto; (Kurashiki, JP) ;
Toyooka; Takaaki; (Kurashiki, JP) ; Nishisaka;
Hisato; (Osaka, JP) ; Goto; Masao; (Osaka,
JP) ; Harada; Hisashi; (Osaka, JP) |
Correspondence
Address: |
IP GROUP OF DLA PIPER US LLP
ONE LIBERTY PLACE, 1650 MARKET ST, SUITE 4900
PHILADELPHIA
PA
19103
US
|
Assignee: |
JFE Steel Corporation,a corporation
of Japan
Tokyo
JP
JTEKT Corporation, a corporation of Japan
Osaka
JP
|
Family ID: |
38180041 |
Appl. No.: |
11/706687 |
Filed: |
February 15, 2007 |
Current U.S.
Class: |
148/653 ;
148/334 |
Current CPC
Class: |
Y10S 148/906 20130101;
C22C 38/06 20130101; F16C 33/34 20130101; F16C 33/62 20130101; F16H
55/32 20130101; C22C 38/60 20130101; C21D 1/26 20130101; C22C 38/02
20130101; C21D 1/32 20130101; C22C 38/22 20130101; C21D 6/002
20130101; F16C 33/32 20130101; C22C 38/34 20130101; C21D 1/28
20130101; F16C 33/30 20130101 |
Class at
Publication: |
148/653 ;
148/334 |
International
Class: |
C22C 38/22 20060101
C22C038/22 |
Foreign Application Data
Date |
Code |
Application Number |
Feb 28, 2006 |
JP |
2006-052688 |
Claims
1. A steel part having a long rolling contact fatigue life,
comprising steel having a composition containing: C: about 0.7% by
mass to about 1.1% by mass; Si: about 0.5% by mass to about 2.0% by
mass; Mn: about 0.4% by mass to about 2.5% by mass; Cr: about 1.6%
by mass to about 5.0% by mass; Mo: about 0.1% by mass to less than
about 0.5% by mass; Al: about 0.010% by mass to about 0.050% by
mass; less than about 0.0015% by mass of Sb as an impurity, and the
balance composed of Fe and inevitable impurities, the steel being
hardened and tempered, wherein a portion from the surface to a
depth of about 5 mm has a steel structure in which residual
cementite has a grain diameter of about 0.05 to about 1.5 .mu.m,
prior austenite has a grain diameter of about 30 .mu.m or less, and
the ratio by volume of the residual austenite is less than about
25%.
2. The steel part according to claim 1, wherein the steel further
comprises at least one selected from the following: Ni: about 0.5%
by mass to about 2.0% by mass; V: about 0.05% by mass to about
1.00% by mass; and Nb: about 0.005% by mass to about 0.50% by
mass.
3. A method for producing a steel part having a long rolling
contact fatigue life, comprising: hot-working steel, spheroidizing
annealing the steel by maintaining the steel at about 800.degree.
C. to about 850.degree. C. for about 5 hours or more, cooling the
steel to about 700.degree. C. or less at a rate of about
0.01.degree. C./s or less, and hardening and tempering the steel,
wherein the steel has a composition comprising: C: about 0.7% by
mass to about 1.1% by mass; Si: about 0.5% by mass to about 2.0% by
mass; Mn: about 0.4% by mass to about 2.5% by mass; Cr: about 1.6%
by mass to about 5.0% by mass; Mo: about 0.1% by mass to less than
about 0.5% by mass; Al: about 0.010% by mass to 0.050% by mass;
less than about 0.0015% by mass of Sb as an impurity, and the
balance composed of Fe and inevitable impurities.
4. A method for producing a steel part having a long rolling
contact fatigue life, comprising: hot-working steel, cooling the
steel to about 200.degree. C. at a cooling rate of about
0.5.degree. C./s or less, spheroidizing annealing the steel by
maintaining at about 750.degree. C. to about 850.degree. C.,
cooling the steel to about 700.degree. C. or less at a rate of
about 0.015.degree. C./s or less, and hardening and tempering the
steel, wherein the steel has a composition comprising: C: about
0.7% by mass to about 1.1% by mass; Si: about 0.5% by mass to about
2.0% by mass; Mn: about 0.4% by mass to about 2.5% by mass; Cr:
about 1.6% by mass to about 5.0% by mass; Mo: about 0.1% by mass to
less than about 0.5% by mass; Al: about 0.010% by mass to about
0.050% by mass; less than about 0.0015% by mass of Sb as an
impurity, and the balance composed of Fe and inevitable
impurities.
5. The method according to claim 3, wherein the steel further
comprises at least one selected from the following: Ni: about 0.5%
by mass to about 2.0% by mass; V: about 0.05% by mass to about
1.00% by mass; and Nb: about 0.005% by mass to about 0.50% by
mass.
6. The method according to claim 4, wherein the steel further
comprises at least one selected from the following: Ni: about 0.5%
by mass to about 2.0% by mass; V: about 0.05% by mass to about
1.00% by mass; and Nb: about 0.005% by mass to about 0.50% by mass.
Description
RELATED APPLICATION
[0001] This application claims priority of Japanese Patent
Application No. 2006-052688, filed Feb. 28, 2006, herein
incorporated by reference.
TECHNICAL FIELD
[0002] The technology in this disclosure relates to steel parts
which have a long rolling contact fatigue life and which are used
as components of rolling bearings such as roller bearings and ball
bearings, and toroidal continuously variable transmissions. In
particular, it relates to steel parts which have a long life until
peculiar damage has occurred in a severe environment of bearing
using, i.e., until microstructural change (damage) has occurred
below a rolling contact plane due to cyclic load, and a method for
producing the steel parts.
BACKGROUND
[0003] As materials for steel parts constituting rolling bearings
used for automobiles and industrial machines, high-carbon chromium
bearing steel defined in JIS-SUJ2 is most frequently used. In
general, an important property of bearing steel is a long rolling
contact fatigue life, and a possible main factor which influences
the rolling contact fatigue life is a non-metallic inclusion in
steel. Therefore, as a commonly employed countermeasure, the oxygen
content in the high-carbon chromium steel is decreased to control
the amount, shape, and size of a non-metallic inclusion, thereby
improving a bearing life (refer to, for example, Japanese
Unexamined Patent Application Publication No. 1-306542 and Japanese
Unexamined Patent Application Publication No. 3-126839).
[0004] However, in order to produce bearing steel containing a
small amount of non-metallic inclusion, it is necessary to install
expensive refining equipment or significantly improve conventional
equipment. Therefore, there is the problem of a high economic
load.
[0005] Accordingly, research was conducted to resolve the problem.
As a result, it was found that even when the amount of a
non-metallic conclusion is simply decreased, in many cases, a large
effect cannot be obtained on improvement in the rolling life of a
bearing, particularly the bearing life under a severe condition
such as a high load or a high temperature. This led to the finding
that as a factor which determines the rolling life, there is a
factor other than the presence of a "non-metallic inclusion" which
has been conventionally discussed. Specifically, a microstructural
change layer composed of a white etched constituent occurs in a
lower layer (surface layer) of a contact plane due to shear stress
in contact between inner and outer rings and a rolling element of a
bearing as the environment of bearing using becomes severe. In
addition, the microstructural change layer is gradually grown as
the number of cycles increases, and finally spalling occurs by
rolling contact fatigue in the microstructural change portion to
determine the bearing life. It was also found that the severe
environment of bearing using, i.e., a higher plane pressure
(reduction in size) and an elevated using temperature, decrease the
number of cycles until a microstructural change has occurred,
resulting in a significant decrease in the bearing life. Such a
decrease in the bearing life in the severe environment of using
cannot be sufficiently suppressed only by controlling the amount of
a non-metallic inclusion as in related art. Therefore, it is
thought to be necessary to retard the microstructural change.
[0006] As a countermeasure, bearing steel containing 0.5 to 1.5% by
mass of C, over 2.5 to 8.0% by mass of Cr, 0.001 to 0.015% by mass
of Sb, 0.002% by mass or less of 0, and the balance composed of Fe
and inevitable impurities has been proposed, and bearing steel
containing these elements and further containing over 0.5 to 2.5%
by mass of Si, 0.05 to 2.0% by mass of Mn, 0.05 to 0.5% by mass of
Mo, and 0.005 to 0.07% by mass of Al has been developed (refer to
Japanese Unexamined Patent Application Publication No.
6-287691).
[0007] As a result, the microstructural change due to cyclic load
in rolling contact under high load was retarded, and so-called
"B.sub.50 high-load rolling contact fatigue life (total number of
cycles until a white portion of a microstructural change layer
spalls at a cumulative failure probability of 50% in a rolling
contact fatigue test)" was improved.
[0008] However, the environment of bearing using has been recently
made severer than that at the time of filing of Japanese Unexamined
Patent Application Publication No. 6-287691, and thus the
development of steel having a long rolling contact fatigue life has
been desired ardently.
[0009] Therefore, a steel was developed having a long rolling
contact fatigue life as steel capable of further increasing a
bearing life even under severe using conditions, the steel having a
composition containing 0.7 to 1.1% by mass of C, 0.5 to 2.0% by
mass of Si, 0.4 to 2.5% by mass of Mn, 1.6 to 4.0% by mass of Cr,
0.1 to less than 0.5% by mass of Mo, 0.010 to 0.050% by mass of Al,
and the balance composed of Fe and inevitable impurities, being
subjected to hardening and tempering, and having a microstructure
including residual cementite with a grain diameter of 0.05 to 1.5
.mu.m and prior austenite with a grain diameter of 30 .mu.m or less
(refer to Japanese Unexamined Patent Application Publication No.
2004-315890).
[0010] With that steel, the average grain diameter of residual
cementite is properly controlled to retard the microstructural
change and increase the n number of cycles until the microstructure
spalls. Furthermore, the grain diameter of prior austenite in the
microstructure after hardening and tempering is refined to suppress
the development of fatigue cracking and further improve the rolling
contact fatigue life.
[0011] However, when the steel is applied to a component of an
actual bearing, a sufficient rolling contact fatigue life may not
be exhibited, thereby causing the need for further improvement.
SUMMARY
[0012] We provide steel parts which have a long rolling contact
fatigue life and which are capable of further increasing the life
of a bearing under severer using conditions than usual conditions,
and provide a useful method for producing the steel parts.
[0013] We provide in particular: [0014] (1) A steel part having a
long rolling contact fatigue life, comprising steel having a
composition containing: [0015] C: 0.7% by mass to 1.1% by mass;
[0016] Si: 0.5% by mass to 2.0% by mass; [0017] Mn: 0.4% by mass to
2.5% by mass; [0018] Cr: 1.6% by mass to 5.0% by mass; [0019] Mo:
0.1% by mass to less than 0.5% by mass; [0020] Al: 0.010% by mass
to 0.050% by mass; less than 0.0015% by mass of Sb as an impurity,
and the balance composed of Fe and inevitable impurities, the steel
being hardened and tempered, wherein a portion from the surface to
a depth of 5 mm has a steel structure in which residual cementite
has a grain diameter of 0.05 to 1.5 .mu.m, prior austenite has a
grain diameter of 30 .mu.m or less, and the ratio by volume of the
residual austenite is less than 25%. [0021] (2) The steel part
having a long rolling contact fatigue life described above in (1),
the steel further containing at least one selected from the
following: [0022] Ni: 0.5% by mass to 2.0% by mass; [0023] V: 0.05%
by mass to 1.00% by mass; and [0024] Nb: 0.005% by mass to 0.50% by
mass. [0025] (3) A method for producing a steel part having a long
rolling contact fatigue life, the method including hot-working
steel, spheroidizing annealing the steel by maintaining at
800.degree. C. to 850.degree. C. for 5 hours or more and cooling to
700.degree. C. or less at a rate of 0.01.degree. C./s or less, and
hardening and tempering the steel, the steel having a composition
containing: [0026] C: 0.7% by mass to 1.1% by mass; [0027] Si: 0.5%
by mass to 2.0% by mass; [0028] Mn: 0.4% by mass to 2.5% by mass;
[0029] Cr: 1.6% by mass to 5.0% by mass; [0030] Mo: 0.1% by mass to
less than 0.5% by mass; [0031] Al: 0.010% by mass to 0.050% by
mass; less than 0.0015% by mass of Sb as an impurity, and the
balance composed of Fe and inevitable impurities. [0032] (4) A
method for producing a steel part having a long rolling contact
fatigue life, the method including hot-working steel, cooling the
steel to 200.degree. C. at a cooling rate of 0.5.degree. C./s or
less, spheroidizing annealing the steel by maintaining at
750.degree. C. to 850.degree. C. and cooling to 700.degree. C. or
less at a rate of 0.015.degree. C./s or less, and hardening and
tempering the steel, the steel having a composition containing:
[0033] C: 0.7% by mass to 1.1% by mass; [0034] Si: 0.5% by mass to
2.0% by mass; [0035] Mn: 0.4% by mass to 2.5% by mass; [0036] Cr:
1.6% by mass to 5.0% by mass; [0037] Mo: 0.1% by mass to less than
0.5% by mass; [0038] Al: 0.010% by mass to 0.050% by mass; less
than 0.0015% by mass of Sb as an impurity, and the balance composed
of Fe and inevitable impurities. [0039] (5) The method for
producing the steel part having a long rolling contact fatigue life
described above in (3) or (4), the steel further containing at
least one selected from the following: [0040] Ni: 0.5% by mass to
2.0% by mass; [0041] V: 0.05% by mass to 1.00% by mass; and [0042]
Nb: 0.005% by mass to 0.50% by mass.
[0043] According to the present invention, a microstructural change
in a rolling environment under a high load is retarded, and thus a
steel part having a high-load rolling contact fatigue life
represented by so-called B.sub.50 can be provided. Therefore, a
steel part required to have the resistance to rolling contact
fatigue as a constituent part of, for example, a roller bearing,
can be reduced in size, and a steel part usable in an environment
at a higher speed and higher load can be provided.
BRIEF DESCRIPTION OF THE DRAWINGS
[0044] FIG. 1 is a partial sectional view showing a toroidal
continuously variable transmission.
[0045] FIG. 2 is a schematic drawing showing a durability life test
rig used in a rolling contact fatigue test.
[0046] In the drawings, reference numerals denote the following:
[0047] 5 input disk [0048] 5b orbital plane [0049] 6b output disk
[0050] 6c orbital plane [0051] 7 roller [0052] 7b peripheral
surface [0053] 10 durability life test rig [0054] 11 disk [0055] 13
first roller [0056] 14 second roller [0057] 15 drive unit
DETAILED DESCRIPTION
[0058] We researched damage to a microstructure of steel used for a
bearing under a severe rolling environment. As a result, it was
found that the damage is mainly caused by stress concentration in a
hard portion of steel and diffusion of carbon (symbol: C) in the
periphery thereof. In other words, a microstructural change in the
steel can be retarded by suppressing C diffusion in the steel under
a using environment.
[0059] Therefore, a method for realizing the finding was further
researched. As a result, it was found that a method for suppressing
C diffusion in steel, the austenite grains (represented) by .gamma.
hereinafter) present in the metallic structure of the steel are
refined in a heating process for hardening, and the grain diameter
of retained cementite after hardening and tempering is controlled
to 0.05 to 1.5 .mu.m.
[0060] In high-carbon bearing steel represented by JIS-SUJ2, coarse
carbide with a grain diameter of 5 .mu.m or more, which is referred
to as "eutectic carbide", may remain in steel after hardening and
tempering due to the influence of coarse carbide crystallized when
melted steel is cast and solidified. Such coarse carbide is removed
of course, and the spheroidal carbide produced in spheroidizing
annealing functions as a stress concentrator with coarsening of the
spheroidal carbide to promote a microstructural change. As a
measure against this, the steel disclosed in Japanese Unexamined
Patent Application Publication No. 2004-315890 was developed.
[0061] In other words, in the technique disclosed in Japanese
Unexamined Patent Application Publication No. 2004-315890, in order
to retard a microstructural change, the average grain diameter of
residual cementite is controlled in a proper range, specifically
0.05 to 1.5 .mu.m, in which stress concentration in a boundary
between the residual cementite and a matrix can be suppressed while
promoting dissolution of C into the matrix. In addition, as a
method for controlling the average grain diameter of residual
cementite in the proper range, a method of spheroidizing annealing
by maintaining at 750.degree. C. to 850.degree. C. and then cooling
to 700.degree. C. or less at a rate of 0.015.degree. C./s or less
is used.
[0062] Furthermore, in an attempt to apply the steel to an actual
part, we worked the steel into the part shape by hot-working such
as hot-casting or the like and then spheroidized and annealed the
part under the conditions descried above to produce a desired
microstructure. In this attempt, the average grain diameter of
residual cementite in a surface layer was not necessarily
controlled in the range of 0.05 to 1.5 .mu.m, thereby causing the
problem of failing to obtain an expected rolling contact fatigue
characteristic. Furthermore, when finishing was performed by
cutting after spheroidizing annealing (before hardening and
tempering), the problem of very low machinability due to the hard
surface layer occurred.
[0063] As a result of intensive research on the causes of the
problems, it was found that when large amounts of Cr and Mo are
contained, the microstructure of the surface layer after
hot-working becomes a structure containing bainite or martensite,
not containing cementite, and thus the growth of cementite does not
sufficiently proceed even by subsequent spheroidizing annealing.
Since the growth of cementite is insufficient after spheroidizing
annealing, dissolution of C in an austenite phase as a mother phase
excessively proceeds during heating for hardening, and thus the
amount of residual austenite after hardening and tempering is
excessively increased, thereby causing an adverse effect on the
rolling contact fatigue life. It was further found that since the
microstructure of the surface layer becomes a structure containing
bainite or martensite after hot-working, softening does not
sufficiently proceed after spheroidizing annealing, thereby causing
difficulty in cutting.
[0064] As a result of further research, it was found that when
hot-working conditions or spheroidizing annealing conditions are
optimized, the microstructure of the surface layer of the steel
part can be optimized, and the steel part having a long rolling
contact fatigue life while maintaining machinability after
spheroidizing annealing can be obtained after hardening and
tempering.
[0065] Examples of the part include orbital parts such as an inner
ring and an outer ring, and a rolling element, which constitute a
rolling bearing, and a disk and a roller which constitute a
toroidal continuously variable transmission. FIG. 1 shows the
structure of a toroidal continuously variable transmission as an
example.
[0066] FIG. 1 is a schematic drawing showing a variator 1 of a full
toroidal continuously variable transmission which is a type of
toroidal continuously variable transmission. The variator 1
includes an input shaft 3 rotated by an output shaft 2 of an
engine, input disks 5 being supported near both ends of the input
shaft 3.
[0067] In each of the input disks 5, a concavely curved orbital
plane 5b is formed in one of the sides, and a plurality of spline
holes 5a is formed in the inner periphery. The spline holes 5a are
engaged to a splined shaft 3a provided on the input shaft 3 to
integrally rotatably attach each of the input disks 5 to the input
shaft 3. In addition, the opposite movements of the input disks 5
are restricted by anchor rings 51 fixed to the input shaft 3.
[0068] Further, output units 6 each including an output part 6a and
an output disk 6b integrally rotatably supported by the output part
6a are relatively rotatably provided at the center of the input
shaft 3 in the axial direction thereof. In addition, a concavely
curved orbital plane 6c is formed on one of the sides of each
output disk 6b which faces the orbital plane 5b of each input disk
5. Further, sprocket gears 6e are formed in the outer periphery of
each output part 6a so as to engage with a chain 6d so that power
is transmitted to the outside through the chain 6d.
[0069] Each of the output disks 6b is attached to allow slight
movement in the axial direction of the output part 6a, and a backup
plate 6h is disposed at the back of each output disk 6b with a gap
6g therebetween. The gap 6g is sealed with a casing 6f and a seal
(not shown in the drawing). When hydraulic pressure is supplied to
the gap 6g from a hydraulic power source, the output disk 6b is
urged toward the opposing input disk 5 to apply a predetermined
terminal load.
[0070] In addition, a toroidal gap is formed between the orbital
plane 5b of each input disk 5 and the orbital plane 6c of each
output disk 6b, which are opposed to each other. A lubricant
(traction oil) is supplied to the toroidal gap, and three disk
rollers 7 are disposed at peripheral positions at equal intervals
so as to rotate in contact with the orbital planes 5b and 6c
through an oil film. The portions of contact between the orbital
planes 5b and 6c are the peripheral surfaces 7bb of the rollers 7.
Each of the rollers 7 is rotatably supported by a carriage 8 so
that the rotational shaft 7a thereof can be tilted. In addition,
driving forces is applied to the carriages 8 by hydraulic pressure
in the direction crossing the drawing of FIG. 1.
[0071] In the variator 1, when the pair of input disks 5 is
rotated, torque is transmitted from the input disks 5 to the output
disks 6b by shearing force of the oil film through the three
rollers 7 on each of the right and left sides. The rollers 7
supported by the carriages 8 incline the rotational shafts 7a to
remove unbalance between the reaction force generated in the
carriages 8 due to torque transmission and the torque necessary for
driving the output disks 6b. Consequently, the positions of the
rollers 7 are changed as shown by two-dot chain lines in FIG. 1 to
continuously change the change gear ratio between the disks 5 and
6b.
[0072] In such a toroidal continuously variable transmission, the
portions of rolling contact between the input/output disks and the
rollers are subjected to a high temperature (100.degree. C. or
more) and high surface pressure (maximum contact surface pressure 4
to 4.5 GPa or more). In addition, for example, when three rollers
are disposed, large vertical stress is repeatedly applied to the
orbital planes of the input/output disks from the three points, and
high shearing stress is repeatedly applied due to the traction
during power transmission. Therefore, the orbital planes of the
input/output disks are under specific severe contact conditions as
compared with a rolling surface of a usual rolling bearing in which
only vertical stress is mainly applied to the surface. Therefore,
the orbital plane of each disk is required to have high fatigue
strength in order to prevent the occurrence of spalling starting at
the surface due to metal contact and spalling due to a structural
change even when used under such specific severe contact
conditions.
[0073] Next, the reasons for providing the ranges of the components
of the steel part will be described below.
C: about 0.7% by Mass to about 1.1% by Mass
[0074] In a metal structure of steel, C is dissolved in a matrix.
However, in order to strengthen martensite grains to cause an
effective function to secure hardness of steel after hardening and
tempering and improve the rolling contact fatigue life, a C content
of about 0.7% by mass or more is required. However, at an
excessively high C content, the formation of coarse carbide such as
eutectic carbide is accelerated, and a microstructural change due
to C diffusion in steel is promoted, thereby degrading the rolling
contact fatigue life. Therefore, the upper limit is about 1.1% by
mass.
Si: about 0.5% by Mass to about 2.0% by Mass
[0075] Si functions as a deoxidizer in refining of steel and is
solid-dissolved in a matrix to suppress a decrease in strength of
steel in tempering after hardening. Further, Si is an effective
element for retarding a microstructural change in an environment
under rolling load. In order to sufficiently exhibit these effects,
a Si content of about 0.5% by mass or more is required. On the
other hand, a content over about 2.0% by mass causes significant
deterioration in the castability and machinability of steel.
Therefore, the upper limit is about 2.0% by mass.
Mn: about 0.4% by Mass to about 2.5% by Mass
[0076] Mn functions as a deoxidizer in refining of steel and is an
effective element for decreasing the oxygen content of steel. In
addition, Mn effectively functions to improve the hardenability of
steel to improve toughness and strength of martensite constituting
a matrix, and improve the rolling contact fatigue life.
Furthermore, Mn has the effect of stabilizing a cementite phase and
retard a microstructural change. In order to sufficiently exhibit
these effects, a Mn content of about 0.4% by mass or more is
required. On the other hand, a content over about 2.5% by mass
causes significant deterioration in the castability and
machinability of steel. Therefore, the upper limit is about 2.5% by
mass.
Cr: about 1.6% by Mass to about 5.0% by Mass
[0077] Cr has the function to stabilize a cementite phase in steel
to suppress C diffusion and the function to suppress coarsening of
cementite grains to prevent stress concentration. Also, Cr is an
element effectively functioning to improve the rolling contact
fatigue life. In order to achieve the sufficient effect, a Cr
content of about 1.6% by mass or more is required. On the other
hand, at a content of over about 5% by mass, the amount of C
dissolved in martensite is decreased to decrease hardness after
hardening and tempering, thereby degrading the rolling contact
fatigue life. Therefore, the upper limit is about 5% by mass.
Mo: about 0.1% by Mass to Less than about 0.5% by Mass
[0078] Mo is solid-dissolved in a matrix and has the function to
suppress a decrease in strength of steel during tempering after
hardening. Further, Mo improves hardness of steel after hardening
and tempering and improves the rolling contact fatigue life. In
addition, Mo has the function to stabilize carbide to retard a
microstructural change. However, at a content of less than about
0.1% by mass, the effect cannot be sufficiently obtained, while at
a content of about 0.5% by mass or more, the effect is saturated,
thereby increasing the cost. Therefore, the content is about 0.1%
by mass to less than about 0.5% by mass.
Al: about 0.010% by Mass to about 0.050% by Mass
[0079] Al is necessary as a deoxidizer in refining of steel and is
positively added as an element functioning to refine prior
austenite grains by bonding to N in steel to effectively improve
the rolling contact fatigue life. In order to obtain the sufficient
effect, a content of about 0.010% by mass or more is required. On
the other hand, at a high content over about 0.050% by mass, the
rolling contact fatigue life is degraded by AlN precipitated in a
large amount in steel. Therefore, the content is about 0.010% by
mass to about 0.050% by mass.
Sb: Less than about 0.0015% by Mass
[0080] Sb is an element possibly mixed from an iron source such as
scraps, but segregates at austenite grain boundaries in hot-working
to degrade hot-workability, toughness, and the rolling contact
fatigue life of steel due to mixing of Sb. Therefore, it is
necessary to control the amount of Sb mixing to a low value by
appropriately selecting an iron source. The above problem generally
becomes significant when about 0.0010% by mass or more Sb is mixed.
However, when the grain diameter of prior austenite is achieved,
the allowable upper limit of the amount of Sb mixing can be
increased by increasing the grain boundary area. However, the
amount of Sb mixing in steel must be controlled to less than about
0.0015% by mass.
[0081] In addition to the above-described basic components, at
least one selected from about 0.5% by mass to about 2.0% by mass of
Ni, about 0.05% by mass to about 1.00% by mass of V, and about
0.005% by mass to about 0.50% by mass of Nb can be further
contained.
Ni: about 0.5% by Mass to about 2.0% by Mass
[0082] Since Ni is solid-dissolved in a matrix to suppress a
decrease in strength of steel after tempering, Ni is added
according to demand. In order to obtain the sufficient effect, a
content of about 0.5% by mass is required. On the other hand, a
content of over about 2.0% by mass causes the formation of a large
amount of residual austenite, thereby decreasing the strength after
hardening and tempering. Therefore, the upper limit of the content
is about 2.0% by mass.
V: about 0.05% by Mass to about 1.00% by Mass
[0083] V has the function to form stable carbide and improve
hardness of steel and the function to suppress a microstructural
change to improve the rolling contact fatigue life. Therefore, V is
added according to demand. In this case, at a content of less than
about 0.05% by mass, the sufficient effect is not obtained, while
at an excessively high content, the amount of dissolved C is
decreased to decrease hardness of steel after hardening and
tempering. Therefore, the upper limit of the content is about 1.00%
by mass.
Nb: about 0.005% by Mass to about 0.50% by Mass
[0084] Like V, Nb has the function to form stable carbide and
improve hardness of steel and the function to suppress a
microstructural change to improve the rolling contact fatigue life.
Therefore, Nb is added according to demand. In this case, at a
content of less than about 0.005% by mass, the sufficient effect
cannot be obtained, while at a content over about 0.05% by mass,
the effect is saturated. Therefore, the content is about 0.005% by
mass to about 0.50% by mass.
[0085] Next, the microstructure of the steel part will be described
below.
[0086] It was found that in a steel part required to have a rolling
contact fatigue life, the microstructure of a surface layer from
the surface to a depth of about 5 mm is particularly important.
Therefore, the steel part after hardening and tempering is required
to have a portion from the surface to a depth of about 5 mm which
satisfies the microstructure described below.
[0087] First, in the steel part after hardening and tempering, the
cementite grain diameter in the steel structure of a portion from
the surface to a depth of about 5 mm is controlled to about 0.05
.mu.m to about 1.5 .mu.m for the following reasons: [0088] When the
steel having the above-described C content is hardened and
tempered, cementite present before hardening remains in the
microstructure of the steel. Therefore, we repeatedly studied with
attention to the fact that a distribution form of the residual
cementite strongly influences the properties of a microstructural
change. As a result, it was found that when the average grain
diameter of residual cementite is smaller than about 0.05 .mu.m,
the ratio of the surface area of cementite to the volume is
increased to promote dissolution of C into the matrix. On the other
hand, when coarse residual cementite having an average grain
diameter over about 1.5 .mu.m is present, stress concentration in
the boundaries between the residual cementite and the matrix is
accelerated, and thus the number of cycles of stress application
until the occurrence of a microstructural change and spalling of
the microstructure is decreased. From this viewpoint, it was found
that the grain diameter of residual cementite after hardening and
tempering is preferably specified to about 0.05 .mu.m to about 1.5
.mu.m. [0089] In the steel structure of a portion from the surface
to a depth of about 5 mm, the grain diameter of prior austenite is
specified to about 30 .mu.m or less. The reason for this is that in
the microstructure of the steel after hardening and tempering, when
the grain diameter of prior austenite is about 30 .mu.m or less,
propagation of fatigue cracks within a crystal grain can be stopped
at the grain boundary, and further progress can be retarded. [0090]
Furthermore, in the steel structure of a portion from the surface
to a depth of about 5 mm, the ratio by volume of the residual
austenite is specified to less than about 25%. In other words, a
residual austenite phase at a residual austenite ratio of about 25%
or more is transformed to martensite accompanying volume expansion
in the using environment, thereby changing the dimensions of the
steel part. The dimensional change causes a stress concentration
portion to adversely affect the rolling contact fatigue.
[0091] The steel part having a long rolling contact fatigue life is
produced through the following steps: [0092] First, molten steel
having the above-described chemical composition is refined in a
steel making process and then continuously cast to form a cast
slab. The steel cast slab is formed into a steel material (for
example, a steel bar) by a hot-rolling process. The steel material
is then formed in a steel part such as a bearing race by
hot-working such as hot casting or the like. After spheroidizing
annealing, if required, the steel part is cut and hardened and
tempered to produce a steel part.
[0093] In the method for producing the steel part, it is necessary
to use the following conditions (I) or (II): [0094] (I) After the
hot working, the steel is spheroidized and annealed by maintaining
at about 800.degree. C. to about 850.degree. C. for about 5 hours
or more and then cooling at a rate of about 0.01.degree. C./s or
less, and then hardened and tempered. [0095] (II) After hot working
at abut 900.degree. C. or more, the steel is cooled to about
200.degree. C. at a cooling rate of about 0.5.degree. C./s or less,
spheroidized and annealed by maintaining at about 750.degree. C. to
about 850.degree. C. and then cooling to about 700.degree. C. at a
cooling rate of about 0.015.degree. C./s or less, and then hardened
and tempered.
[0096] In the method (I), the spheroidizing annealing conditions
are controlled to control the average grain diameter of cementite
in the surface layer of the steel part. When the usual hot-working
conditions and subsequent cooling conditions are not particularly
limited, in the steel composition, the structure of the surface
layer may become a bainite or martensite structure after
hot-working and cooling. Therefore, it is necessary to produce
cementite for achieving the above-described final structure by
subsequent spheroidizing annealing of the steel having a bainite or
martensite structure. As a condition for this purpose, a
spheroidizing annealing condition is required, in which the steel
is maintained at about 800.degree. C. to about 850.degree. C. for
about 5 hours or more and then cooled to about 700.degree. C. or
less at a cooling rate of about 0.01.degree. C./s or less.
[0097] Namely, when the retention temperature is lower than about
800.degree. C. or the retention time is less than about 5 hours,
the sufficient growth of cementite in the surface layer cannot be
expected, and the cementite grain diameter in the final surface
layer structure cannot be controlled in the above-described range.
Also, solid-dissolution of C in the austenite phase, which is a
matrix in hardening heating, excessively proceeds, and thus the
amount of the residual austenite in the final surface layer
structure is excessively increased. On the other hand, when the
retention temperature exceeds about 850.degree. C., cementite after
spheroidizing is coarsened, and, consequently, the residual
cementite after hardening and tempering is also coarsened. In
cooling after the retention, when the cooling rate to about
700.degree. C. exceeds about 0.01.degree. C./s, cementite
precipitation during cooling proceeds in the form of reproduction
of pearlite, not the growth of spheroidized cementite. Therefore,
softening after spheroidizing annealing does not sufficiently
proceed to degrade workability. Further, solid dissolution in
hardening heating is excessively accelerated to form a large amount
of residual austenite after hardening.
[0098] After the above-described spheroidizing annealing, hardening
and tempering is performed. In order to obtain a desired residual
cementite distribution and prior austenite grain diameter after
hardening and tempering, the heating temperature of hardening is
preferably about 800.degree. C. to about 950.degree. C. This is
because the most desirable microstructure is produced in this
temperature range. Although the fraction of the cementite structure
after hardening and tempering changes mainly depending on the C
content, the volume ratio is about 3 to about 25% in the
composition range of the present invention. In addition, a cutting
work may be performed before the hardening and tempering. The
above-mentioned spheroidizing annealing conditions have the effect
of improving machinability because the surface layer is
sufficiently softened.
[0099] Next, the method (II) will be described.
[0100] In the method (II), the hot-working conditions and
subsequent cooling conditions are regulated to control the surface
layer structure after the hot-working and cooling. When the
hot-working temperature is lower than about 900.degree. C., the
mold life is decreased due to the high deformation resistance of
steel, and cracks occur in casting due to the low deformation
ability. Therefore, the hot-working is performed at about
900.degree. C. or more. Further, in subsequent cooling, the
condition is controlled so that the above-described final structure
can be obtained in the surface layer.
[0101] When the cooling rate exceeds about 0.5.degree. C./s, the
sufficient growth of cementite in the surface layer cannot be
expected by subsequent spheroidizing annealing, and the cementite
grain diameter in the final surface layer structure cannot be
controlled in the above-described range. Also, solid-dissolution of
C in the austenite phase, which is a mother phase in hardening
heating, excessively proceeds, and thus the amount of the residual
austenite in the final surface layer structure is excessively
increased.
[0102] Then, spheroidizing annealing is performed by retention at
about 750.degree. C. to about 850.degree. C. and then cooling to
about 700.degree. C. at a cooling rate of about 0.015.degree. C./s
or less. When the retention temperature exceeds about 850.degree.
C., care must be taken because cementite after spheroidizing is
coarsened, and, consequently, the residual cementite after
hardening and tempering is also coarsened. At the same time,
layered cementite is newly formed in cooing after the retention,
thereby causing difficult in obtaining desired spheroidal
cementite.
[0103] On the other hand, when the retention temperature is lower
than about 750.degree. C., decomposition of cementite present as
pearlite before spheroidizing annealing does not sufficiently
proceed, and thus a desired residual cementite distribution cannot
be obtained.
[0104] In cooling after the retention, the cooling rate to about
700.degree. C. or less must be about 0.015.degree. C./s or less.
When the cooling rate exceeds about 0.015.degree. C./s, cementite
precipitation during cooling proceeds in the form of reproduction
of pearlite, not the growth of spheroidized cementite. Therefore,
softening after spheroidizing annealing does not sufficiently
proceed to degrade workability. Further, solid dissolution in
hardening heating is excessively accelerated to form a large amount
of residual austenite after hardening.
[0105] After the above-described spheroidizing annealing, hardening
and tempering is performed. In order to obtain a desired residual
cementite distribution and prior austenite grain diameter after
hardening and tempering, the heating temperature of hardening is
preferably about 800.degree. C. to about 950.degree. C. This is
because the most desirable microstructure is produced in this
temperature range. Although the fraction of the cementite structure
after hardening and tempering changes mainly depending on the C
content, the volume ratio is about 3 to about 25% in the
composition range of the present invention. In addition, a cutting
work may be performed before the hardening and tempering. The
above-mentioned spheroidizing annealing conditions have the effect
of improving machinability because the surface layer is
sufficiently softened.
EXAMPLES
[0106] Molten steel having each of the chemical compositions shown
in Table 1 was refined by a converter and then continuously cast to
form a case slab. The resulting cast slab was diffusion-annealed at
1200.degree. C. for 30 hours and then rolled to a steel bar of 64
mm or 90 mm in diameter.
[0107] The steel bar of 64 mm in diameter and the steel bar of 90
mm in diameter were hot-cast into a disk-like roller shape and a
disk shape, respectively, at least the temperature shown in Table 2
and then cooled at various cooling rates. These disk and roller
were normalized and then spheroidized and annealed. The
spheroidizing annealing was performed by cooling to 650.degree. C.
from various retention temperatures at various cooling rates shown
in Table 2 and then standing to cool. Then, in order to remove a
decarbonized layer, a cutting work was performed to form a test
piece with a final shape.
[0108] Further, after hardening, the tempering temperature was
changed from the heating temperature shown in Table 2 according to
the steel used, and the hardness HRc after tempering was controlled
to 60 to 62, followed by polishing and lapping finishing.
[0109] Next, the resulting test piece was cut in the height
direction of a column, and a section was corroded with a picric
acid alcohol solution and then corroded with a nitric acid alcohol
solution. Then, the microstructure was observed to measure the
average grain diameter of residual cementite and the average grain
diameter of prior austenite by image analysis.
[0110] A rolling contact fatigue test was conducted using a
durability life test rig 10 of a traction transmission part shown
in FIG. 2. The durability life test rig shown in FIG. 2 includes a
disk 11, a disk support 12 for supporting the disk 11, a first
roller 13 to be in rolling contact with one of the sides of the
disk 11, a second roller 14 to be in rolling contact with the other
side of the disk 11, a drive unit 15 for rotating the first roller
13, a differential rate mechanism 16 giving a peripheral speed
difference to the second roller 14 relative to the first roller 13,
and a pressure unit 17 for pressing the first roller 13 and the
second roller 14 on the disk 11, all of which are disposed on a
pedestal A.
[0111] In the durability life test rig having the above-described
constitution, in the state in which the disk 11 is supported by the
disk support 12, and the peripheral side of the disk 11 is held
between the first roller 13 and the second roller 14, the disk can
be caused to follow the rollers 13 and 14 by the pressure unit 17
to cause rolling contact each of the rollers 13 and 14 and the disk
11. In this case, a peripheral speed difference can be given to the
second roller 14 by the differential speed mechanism 16 relative to
the first roller 13. Therefore, slipping can be caused between the
disk 11 and the second roller 14.
[0112] The durability life test can be conducted by reproducing a
use state in an actual traction transmission part. In the
durability life test rig, the above-described test piece was used
for the disk 11, the first roller 13, and the second roller 14. The
time taken until failure had occurred due to rolling slide with the
roller 14 was measured as a durability life, and the durability of
the disk 11 was evaluated on the basis of the durability life.
[0113] The test conditions of the durability life test rig 10 were
determined as follows: [0114] (1) Roller rotational speed: 3000 rpm
[0115] (2) Slip ratio between roller and disk: 14% [0116] (3)
Maximum contact surface pressure: 4.2 GPa [0117] (4) Lubricant:
traction oil for toroidal continuously variable transmission [0118]
(5) Oil film parameter (.LAMBDA.): 1.8 The results of the
durability life test are shown in Table 2. In Table 1, steel No. 1
is conventional steel corresponding to JIS-SUJ2, and steel Nos. 2,
3, 4, 5, 6, 6, and 8 are comparative steel containing C, Si, Mn,
Cr, Mo, Al, and Sb, respectively, at contents out of the range of
the present invention.
[0119] In steel part Nos. 2 to 8 (comparative steel) containing
essential elements at contents out of the specified range, B.sub.50
is inferior to steel part No. 1 but is substantially the same value
as steel No. 1. In particular, in steel part Nos. 2 and 7 using
steels having a low C content and a low Al content, respectively,
the microstructure cannot be controlled in the specified range, and
the value of B.sub.50 is excessively decreased. In steel part Nos.
9, 11, 20, and 21 each using steel in which the chemical
composition is within the range, but the microstructure differs
from the specified structure, B.sub.50 is superior to steel part
No. 1 (conventional steel), but the improvement in B.sub.50 is only
small.
[0120] On the other hand, in steel part Nos. 10, 12 to 19, and 21
in which both the chemical composition and the microstructure are
within the specified ranges, B.sub.50 is 10 times or more superior
to steel part No. 1 (conventional steel). These test results
indicate that a damage preventing effect is particularly large in a
full-toroidal continuously variable transmission having a large
spin component in a contact portion between the roller and the disk
and severe contact conditions.
TABLE-US-00001 TABLE 1 Steel Chemical composition of steel (% by
mass) No. C Si Mn Cr Mo Al Ni V Nb Sb Remarks 1 1.00 0.25 0.50 1.50
0.01 0.033 -- -- -- 0.0014 Conventional steel 2 0.65 0.98 0.59 4.22
0.44 0.033 -- -- -- 0.0008 Comparative steel 3 0.90 0.13 0.68 2.51
0.22 0.030 -- -- -- 0.0009 '' 4 1.07 1.47 0.30 2.02 0.54 0.028 --
-- -- 0.0008 '' 5 0.86 1.22 0.53 1.20 0.51 0.033 -- -- -- 0.0009 ''
6 0.84 1.54 0.57 4.34 0.03 0.028 -- -- -- 0.0010 '' 7 0.88 1.04
0.69 2.65 0.31 0.005 -- -- -- 0.0011 '' 8 1.02 1.65 0.63 2.34 0.38
0.033 -- -- -- 0.0038 '' 9 1.00 1.00 0.45 3.50 0.45 0.035 -- -- --
0.0010 Steel of this invention 10 1.00 1.50 0.70 5.00 0.45 0.033 --
-- -- 0.0011 '' 11 0.95 1.10 1.50 2.00 0.30 0.027 -- -- -- 0.0009
'' 12 0.91 1.23 0.70 4.25 0.40 0.026 1.03 -- -- 0.0009 '' 13 0.90
0.91 0.70 2.55 0.50 0.033 -- 0.31 -- 0.0011 '' 14 0.85 1.28 0.78
4.49 0.56 0.029 -- -- 0.045 0.0009 '' 15 1.03 0.73 0.58 3.87 0.28
0.028 0.70 -- 0.031 0.0010 '' 16 1.07 0.74 0.62 4.34 0.25 0.032
0.62 0.15 -- 0.0008 ''
TABLE-US-00002 TABLE 2 Hot-working Spheroidizing Surface layer
structure after (casting) condition annealing condition hardening
and tempering Cooling Cooing Average Hot- rate after rate Hardening
Average prior .gamma. Amount of Steel working hot- Retention
Retention (retention Retention residual .theta. grain residual part
Steel temp. working temp. time temp. ~700.degree. C.) temp.
diameter diameter .gamma. (% by B.sub.50 No. No. (.degree. C.)
(.degree. C./s) (.degree. C.) (h) (.degree. C./s) (.degree. C.)
(.mu.m) (.mu.m) mass) life ratio Remarks 1 1 1000 0.4 810 8 0.007
900 0.47 11.1 13.1 1.0 *1) 2 2 1100 0.4 810 8 0.007 900 0.02 14.8
10.9 0.5 *2) 3 3 1050 0.4 810 8 0.007 900 0.30 11.5 12.3 0.9 '' 4 4
1000 0.4 810 8 0.007 900 0.40 13.4 11.0 1.1 '' 5 5 1050 0.4 810 8
0.007 900 0.34 11.0 12.6 1.0 '' 6 6 1050 0.4 810 8 0.007 900 0.30
13.5 12.5 1.2 '' 7 7 1050 0.4 810 8 0.007 900 0.34 36.2 10.3 0.4 ''
8 8 1000 0.4 820 7 0.007 900 0.30 14.4 14.0 1.3 '' 9 9 1000 1.0 780
7 0.008 900 0.03 19.3 30.0 2.6 '' 10 9 1000 1.0 830 7 0.006 900
0.38 13.2 12.0 12.6 *3) 11 9 1000 1.0 830 7 0.006 1000 0.06 45.0
24.0 2.3 *2) 12 9 1000 0.2 790 7 0.006 900 0.20 11.6 15.0 10.8 *1)
13 10 1000 0.4 810 8 0.006 950 0.26 16.4 12.9 14.9 '' 14 11 1000
0.4 820 7 0.007 900 0.43 13.0 13.6 3.3 '' 15 12 1050 0.3 830 8
0.005 950 0.25 18.3 11.6 16.4 '' 16 13 1050 0.2 780 8 0.005 900
0.31 12.5 11.6 9.2 '' 17 14 1050 2.0 845 7 0.003 950 0.25 17.6 11.4
16.9 '' 18 15 1000 0.4 820 7 0.007 900 0.40 14.9 10.5 12.8 '' 19 16
1000 2 790 3 0.005 900 0.03 21.0 35.0 3.3 *2) 20 16 1000 1.5 830 8
0.052 900 0.03 18.0 23.5 3.5 *2) 21 16 1000 1.5 830 8 0.005 900
0.35 13.5 11.5 14.4 *3) *1) Conventional example *2) Comparative
example *3) Example of this invention
* * * * *