Note: Descriptions are shown in the official language in which they were submitted.
1 2~4~6~9
TITLE OF THE INVENTION
Method of Making Steel for Springs
BACKGROUND OF THE INVENTION
The present invention relates to a method of
making steel for springs such as a diaphragm spring
provided in a clutch of a motor vehicle.
In recent years, the environm~antal temperature of
the spring used in a machine increases with increase of
the output power of the machine. :For example, clutch
torque of the clutch of the motor vehicle is increased
due to increase of the engine power of a motor vehicle
such as a four-wheel drive vehicle. As a result, the
environmental temperature of the clutch increases up to
250-~350°C from 150°C which is a maximum temperature of
a conventional motor vehicle.
The diaphragm spring is made of carbon tool steel
such as SK5 {Japanese Industrial Standard). However,
the spring of carbon steel relaxes quickly, namely it
becomes inoperative when the temperature thereof
increases to the above described high environmental
temperature.
It is known that if silicon content of steel is
increased, endurance of the spring, that is a property
of the spring resisting heat without settling,
z5 increases. However, conventional ;steel including a
2
large silicon content is liable to relax at a high
temperature.
SUMMARY OF THE INVENTION
An object of the present invention is to provide a
method of making steel having a property resisting high
temperature, whereby a spring made of the steel withstands
the relaxation thereof at high tempE~rature.
Another object of the present invention is to
provide steel which may be quickly quenched at a low
temperature, thereby preventing the steel from reducing
in endurance.
The inventors found that steel having excellent
endurance against the relaxation in spring at high
temperature could be made by properly controlling solid
solution and precipitation of carbide in the steel
including carbon (C), silicon (Si), Manganese (Mn).,
chromium (Cr), molybdenum (Mo) and others.
According to the present invention, the method for
making steel comprises hot-rolling steel material
consisting essentially by weight of from 0.4% to 0.80
carbon, from 0.50 to 2.50 silicon, from 0.3% to 2.Oo
manganese, from 0.1% to 1.5a chromium, from 0.1% to 0.50
molybdenum, from 0% to 0.50 vanadium, from Oo to 0.5%
niobium, from 0% to 0.0200 aluminum and remaining iron
and inevitable impurities to form a plate, annealing the
hot-rolled plate, cold-rolling the annealed hot-rolled
plate at rolling reduction of 10% to 80%, annealing the
20~~~~~
3
cold-rolled plate at a temperature below Acl critical
point, heating the annealed cold-rolled plate at a
temperature above Ac3 critical point for a time
sufficient to austenitize carbide, cooling the heated
cold-rolled plate at a speed higher than a lower
critical cooling speed, heating the cooled cold-rolled
plate for a time necessary for precipitating carbide
and then cooling it to a room temperature.
The lower critical cooling speed is a speed above
which the austenite is fully transformed to the
martensite.
In the last heating process, molybdenum carbide is
finely precipitated, thereby preventing the dislocation
migration which causes the relaxation of the spring at
high temperature. The heating is performed at a
temperature between 450°C and 600°C: for a time
sufficient to precipitate the carbide.
The silicon content and chromium content are
selected so as to satisfy the equation:
_~<4xSi(~)-lOxCr(~)<5.
The heating of the cooled cold',-rolled plate is
performed so as to provide an annealed hardness between
HV400 and HV550.
Furthermore, the annealing of the cold-rolled
plate is performed at a temperature between 550°C and
~0~4fi~9
4
730°C, thereby providing carbide having an average
grain diameter less than 2y~m.
The other object and features of this invention
will become understood from the following description
with reference to the accompanying drawings.
BRIEF DESCRIPTION OF DRAWINGS
The figure is a graph showing relationship between
heating temperature and hardness of steel.
DETAILED DESCRIPTION OF THE PREFERRED EMBODIMENTS
Quantity of each component, preparing condition
and reason for numerical limitation of the component
are described hereinafter.
(Carbon)
Carbon is effective in increasing the strength of
steel. In order to obtain a strength necessary for the
the spring; carbon content over 0.4 ~ by weight must be
included. However, if carbon is included in excess of
0.8 ~, quenching crack and reduction of toughness of
steel occur. Therefore the carbon of 0.4 ~ to 0.8 ~ by
weight is included in the steel.
(Silicon)
In the method of the present invention, the
material is tempered at high temperature. Silicon is
added to prevent the strength from reducing due to the
high temperature tempering. It is necessary to add
silicon over 0.5 ~ by weight. If ~cilicon content
2044639
exceeds 2.5 %, internal oxidation <~nd decarburization
which are unfavorable to the spring occur, and
graphitization is enhanced in the hot rolling and
annealing.
(Manganese)
Manganese is effective in deo:~idizing steel and in
increasing the hardenability of thE= steel, if the
manganese is included over 0.3 % by weight. If
manganese content exceeds 2.0 %, the toughness of steel
reduces exceedingly after quenching and tempering.
(Chromium)
Chromium acts to restrict the graphitization and
the internal oxidation which are enhanced by silicon,
and is effective in increasing the hardenability as is
effected by manganese, if chromium content is included
in excess of 0.1 % by weight. If chromium content
exceeds 1.5 %, the toughness of the steel reduces after
quenching and tempering.
Moreover, Si content and Cr content are determined
so as to satisfy the following equation, thereby
preventing decarburization and grad?hitization.
-7<4xSi(%)-lOxCr(%)<5
(Molybdenum)
The molybdenum included in the steel of the
present invention forms carbide in the steel after the
cold rolling and annealing thereof.. The carbon becomes
204409
6
solid solution in austenite when t:he steel is heated
over the Ac3 critical point. Consequently, the
austenite is transformed into mart~ensite after
quenching, and carbide separates finely upon tempering
at high temperature, thereby remarkably increasing
endurance withstanding against relaxation. In order to
obtain such an effect, it is necessary to include
molybdenum over 0.1 ~ and below 0.5 ~ by weight. If
molybdenum content exceeds 0.5 ~, .a large amount of
carbide remains without becoming solid solution in
austenite when the steel is heated above the Ac3
critical point.
(Vanadium, Niobium)
The vanadium and niobium included in the steel of
the present invention become carbide after the cold
rolling and annealing thereof. Remaining vanadium and
niobium without becoming solid solution in austenite
act to prevent austenite grain from growing. On the
other hand, solid solution of vanadium and niobium in
austenite are in solid solution in martensite when
quenching, and precipitate finely <~s carbide when
tempering, thereby enhancing endurance withstanding
against relaxation. In order to aittain these effects,
vanadium and niobium over 0.05 ~ a:re necessary. If the
content exceeds 0.5 $, quantity of un.dissolved carbide
in austenite increases when the stf~el is heated above
2044G~9
Ac3 point, thereby reducing fatigue strength of the
steel.
(Aluminum)
The spring is fatigued by repeated bending or
twisting. Existence of hard inclusions such as
aluminum aggravates th fatigue. In order to reduce the
influence of the hard inclusion, aluminum content is
limited below 0.020 weight percent.
Manufacturing conditions are described
hereinafter.
In the cold rolling, when rolling reduction is
smaller than 10 ~, the grain size of carbide becomes
coarse when annealed below the critical point Acl.
Consequently, a long time is required for transforming
the carbide to austenite when heated above the Ac3
critical point, which causes an increase in
decarburization and hence spring characteristic is
deteriorated. When the rolling reduction is larger
than 80 ~, work hardening due to t:he cold rolling is
2~ remarkably increased, causing deformation such as edge
crack. Therefore, an upper limit .is 80 ~.
If the annealing after the cold rolling is
performed at a temperature above 7.30°C (Ac1 critical
point), spheroidized grain of carbide becomes coarse.
Consequently, it takes a long time to transform the
carbide to austenite, resulting in increase of
2~44G~9
decarburization causing deterioration of spring
characteristic. Therefore, the annealing after the
cold rolling is carried out at a temperature below the
Acl point. If the annealing temperature is lower than
550°C, the hardness increases, so that the formability
of the material reduces. Therefore, the annealing
temperature is between 550°C and 730°C.
If the average grain diameter of carbide after the
annealing is less than 2~m, carbide is easily dissolved
austenite at quenching. Therefore, it is necessary
into maintain the average grain diameter of carbide to
a value smaller than 2~m for effectively performing the
quenching.
In order to increase the strength of the steel
made by the cold rolling and annealing to a value
necessary for the spring, the strip is heated at a
temperature higher than the critical point Ac3 for a
time sufficient for austenitizing 'the spheroidal
carbide, after which cooled at a speed higher than a
lower critical cooling speed, namely quenching.
Thereafter, the strip is heated at a temperature
between 450°C and 600°C for a time to precipitate fine
carbide and cooled to a room temperature (that is
tempering). At the quenching, the parent material is
austenitized by heating it over thE~ Ac3 point, and then
carbon and other elements are dissolved to martensite
20~4~~9
by cooling at a speed higher than the lower critical
cooling speed. By tempering the material at a
temperature higher than X50°C, carbide of Mo, V and Nb
is finely precipitated from the martensite, thereby
increasing the endurance withstanding against the
relaxation. If the tempering is carried out at a
higher temperature than 600°C, a carbide of Mo, V and
Nb becomes coarse which can not prevent the dislocation
migration. In addition, the strength of the steel
largely reduces. Therefore, the tempering is performed
at a temperature below 600°C.
Example 1
Table 1 shows contents of steels. In the table, A
to F are steels of the present invention, and G to L
are comparative steels.
Each of the steels A to F is :made into a
hot-rolled plate of 3.5 mmt by ordinary hot rolling and
then the plate is annealed and cold rolled at a rolling
reduction between 5 ~ and 90 ~. T:hereafter, the steel
is annealed at 700°C below the Acl paint for 10 hours,
and is soaked at 900°C above the Ac3 point for a period
necessary to provide remaining carlbide ratio below 1 ~
by weight. Thereafter, the steel :is quenched into oil.
Table 2 shows results of tests for edge crack and
depth of decarburization. When the rolling reduction
exceeds 80 ~, edge crack occurs. :If the rolling
l0 2044639
reduction is smaller than 10 ~, carbide becomes coarse.
Consequently, it takes a long tims~ to dissolve carbide
into austenite, so that the depth of decarburization
remarkably increases.
10
20
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~p~4639
Example 2
Each of the steels A to F is made into a
hot-rolled plate of 3.5 mmt by ordinary hot rolling and
annealed and cold rolled at rolling reduction of 35
to form a cold rolled plate of 2.3 mm. Thereafter, the
steel is annealed once at 700°C fo:r 10 hours, and is
heated at a temperature between 850°C and 900°C for 10
minutes. Thereafter, the steel is quenched into oil
and tempered at a temperature between 420°C and 63U°C
for 30 minutes.
A relaxation test was perform<=d in order to
estimate endurance against relaxat_~on. The test was
carried out at 350°C, initial 1.0 '~ strain, holding
time of 12 hours. Load reduction after the test was
regarded as relaxation rate.
Table 3 shows the result of the relaxation test.
Since comparative example G is smaller than the present
invention in carbon content, comparative example 1 is
smaller in silicon content, comparative example J is in
manganese content, and K is in chromium content, each
of these steels has low strength so that the relaxation
rate thereof is high. Although the comparative example
H has a large carbon content, the z-elaxation rate is
not largely reduced. Since the comparative example L
has not molybdenum, the carbide of which is effective
to increase the endurance, relaxation rate is very
14
high. Although each of comparative examples A', D' and
F' has the same ingredient content as the present
invention, the tempering temperature is out of the
range of the present invention. Consequently, the
relaxation rate is not largely reduced.
To the contrary each steel according to the
present invention has a very low relaxation rate
comparing with the comparative examples, which means
that the steel has a high endurance withstanding
against the relaxation.
20
15
TABLE
QuenchingTempering HardnessRelaxation
Steel Rate
Temp. Temp. (C) ( H V (%)
(C) )
900 480 496 16. 2
900 520 475 15. 1
present A -- ----~ --- -----
900 560 452 14. 4
in~~ention 8 5 0 5 2 0 4 6 2 1 5. 7
B 900 560 470 13. 5
C 900 520 479 1 5. 7
900 480 51 3 14. 2
900 520 492 1 3. 4
---- --_-. --~ -----
900 560 468 1 2. 6
'
850 560 452 1 3. 2
E 900 560 453 14. 1
F 900 580 472 1 1. 1
G 900 520 348 40. 2
H 900 560 473 20. 1
compara- -- ----- ----. .__.- .~_---
I 880 520 394 25. 2
tlVe --- ----.- --.--_ .--..~ -...._--
J 900 520 442 18. 2
example -- _.--- ----_ -_.- ---._
K 900 560 421 21. 7
L 880 520 427 32. 5
900 420 567 21. 5
A ' _.--- ----_ --- ----
900 630 41 3 1 9. 2
compara- -- .---- ----_ --- -.--.._
900 420 591 20. 1
tive D ' ---- ----_ --._ ----
900 630 426 18. 1
example -- ---- ---._ -._- --._-
900 420 625 19. 5
F ' ---- ----_ --- _.---
900 630 513 17. 3
16 2044G~9
Example 3
Alphabets A to G in Table 4 are steels of the
present invention and H to L are comparative steels.
S
15
25
,~ 2044~~9
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18
Each of the steels in the table was hot-rolled to
provide a hot-rolled plate having a thickness of 3.5
mmt, and then annealing the hot-rolled plate. The
plate was cold rolled at rolling reduction of 35 ~ to
prepare a cold-rolled plate of 2.3 mmt thickness. The
cold-rolled plate was annealed at a temperature between
650°C and 750°C for 10 hours to provide a test piece.
Hardenability test was performed in such a manner that
the test piece was rapidly heated to 850°C at the rate
of 140°C/sec, heated from 850°C to a test temperature
between 900°C and 1100°C at the rate of 30°C/sec, and
then rapidly cooled immediately after the heating
without taking a holding time. The hardenability was
estimated by the hardness of the test piece after the
quenching. Results of the test are shown in the
attached figure.
As is seen from the graph, the test piece A having
ingredient contents according to the present invention
has an average grain diameter of carbide less than 2~m
when annealed at 650°C and 700°C. Even if the test
piece A is heated to the lowest temperature 900°C, the
hardness becomes the higher value. However, if it is
annealed at 750°C so that the average grain diameter
exceeds 2~m, the hardness does not reach the highest
value unless the quenching temperature is elevated up
to 950°C.
19 ~04463~
The comparative example H has a SICR value of
-7.42 out of the range of the present invention.
Consequently a large amount of chromium remains in
carbide after the annealing. Accordingly, the steel
must be heated up to 1000°C in order to obtain the
higher hardness, although the average grain diameter is
smaller than 2um.
From the comparing, it will be seen that the
carbide is rapidly dissolved into austenite at a lower
temperature in accordance with the present invention.
Fatigue test and relaxation tc=st, piece are
estimated as follows. The cold-ro:Lled plate having 2.3
mm thickness is annealed at 680°C :Eor 10 hours, and
then heated at 900°C and quenched. Thereafter, a
plurality of the plates are tempered at various
temperatures for 30 minutes.
The fatigue test was performed in alternating
plane bending fatigue. The result of the test is shown
in Table 5.
25
20 204469
TABLE
Tempering Hardness Test Temp. Fatigue Strength
Steel
Temp. (C (HV) ( C ) (kgf/mm2)
)
25 52
.A 5 8 0 4 5 4 ---.-- --------
250 51
540 490 25 57
2 5 5 3
G 5 5 0 4 5 2 ----- -------
250 51
5 0 0 5 0 5 2 5 5 6
2 5 4 7
I 5 8 0 4 4 8 ---__ ._------
2 5 0 4 6
2 5 4 3
J 5 9 0 4 5 2 ----- -------
2 5 0 4 1
25
21 2044639
From the table, it will be seen that although the
steel A of the present invention has a hardness
appraximately equal to the comparative example I, the
steel A is superior to the comparative example I in
fatigue strength. This is caused by the fact that the
aluminum content of the steel A is less than 0.020
weight percent, which means hard inclusion causing
fatigue fracture is small. The steel G has the same
fatigue characteristic as steel A.
The comparative steel J has a small Cr content
compared with Si content, so that ~SICR value is 7.50
out of the range of the present invention, producing
graphite at annealing. In addition, since a long time
was required for austenitization, decarburization
increased. As a result, the fatigue characteristic is
inferior to the steels A and G.
The endurance withstanding against relaxation was
estimated by the relaxation test. Table 6 shows test
results.
25
22 204439
TABLE 6
Quenching Tempering Hardness Relaxation
Steel Rate
Temp. (C) Temp. (C) (1-IV) (/)
580 454 1 2. 4
A 9 0 0 ---- ---- ----
540 490 1 2. 2
580 461 13. 1
B 900 ---- --- ----
540 494 1 3. 1
C 900 520 4'71 1 6. 3
540 469 1 2. 0
D 9 0 0 --,_ ---- ----
510 496 11. 9
-- - ,._. --~.~ --- .-..-__
E 900 520 435 1'7. 3
F 900 560 455 1 2. 5
550 452 1 1. 5
G 9 0 0 --_ --- ----,-
500 G>05 1 1. 3
850 430 450 36. 7
K --.-_ --.T- -~.~ -~,.-
850 400 4'75 38. 5
L 900 550 x':59 3 1. 5
M 880 490 x:65 25. 2
25
23
Example 4
In table 7, A to G are steels of the present
invention, H to L are comparative steels.
10
20
2044639
24
r_ N r_
o I o I o I ~ I o I o ( o o I o I o I o ( o
O I O I O I O I O ( O I O O I O I O I O I O .
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0 0 0
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25 X044639
Each of the steels in the tab_Le was hot rolled to
provide a hot-rolled plate having a thickness of 3.5
mmt, and then annealing the hot-rolled plate. The
plate was cold rolled at rolling reduction of 35 ~ to
prepare a cold-rolled plate of 2.3 mmt thickness. The
cold-rolled plate was annealed at Ei80°C for 10 hours,
and then heated at a temperature between 850°C and
900°C for 10 minutes and quenched into oil. All plates
were tempered at various temperatures for 30 minutes.
The fatigue test was performed in alternating
plane bending fatigue. The result of the test is shown
in Table 8.
20
26
TABLE 8
Quenching TemperingHardnessTesting Fatigue
Steel Strength
Temp. (C Temp. ( H V Temp. (C (kgf/mm2)
) (G ) ) - )
5 4 0 4 6 3 2. 5 5 1 I
.A 900 480 508 2'.5 58 I
4 1 0 5 7 3 2 5 4 7 II
25 51 I
5 8 0 4 5 2 ----- ---- --
9 0 0 2 5 0 5 0 I
510 516 25 57 I
4 4 0 5 7 8 2 5 4 7 II
2 5 4 6 II
I 9 0 0 5 8 0 4 4 8 ----- ---- --
2 5 0 4 4 II
- note: I represents present invention steels, while II represents
comparative steels.
From the table, it will be seen that although the
steel E of the present invention has a hardness
approximately equal to the comparative example I, the
steel E is superior to the comparative example I in
fatigue strength because of small aluminum content.
Even if contents of ingredients are within the
range of the present invention, thc~ fatigue strength
reduces if the annealed hardness e:~ceeds HV550.
The endurance withstanding ag<~inst settling was
estimated by the relaxation test. Test temperature was
350°C, initial strain 1.0 ~, and holding time 12 hours.
Table 9 shows test results.
20
28
TABLE 9
Quenching TemperingHardnessRelaxation
Steel Rate Notes
Temp. (C Temp. ( H V ( % )
) (C ) )
480 496 16. 2
A 9 0 0 5 2 0 4 7 5 Tl 5.
1
560 452 14, 4
present
850 520 462 ~1 5.
7
B 9 0 0 5 6 0 4 7 0 1 3 .
5
invention
C 9 0 0 5 2 0 4 7 0 1 5. 7
440 509 16. 1
D 9 0 0 4 8 0 4 9 1 1 5 .
9
5 2 0 4 6 5 T1 5 .
8 5 0 5 2 0 4 4 9 2
1 6 .
2
900 580 453 ~~ 2~.
5
9 0 0 5 8 0 4 7 2 ~l 1 .
1
G 9 0 0 5 2 0 4 4 6 7. 3 ,
6
900 520 348 40. 2
'-'-
"-'-'- comparative
8 8 0 5 2 0 3 9 4 2 5 .
2
example
880 520 42? ~~2. 5
I
850 430 450 36. 7
A 900 650 391 22. 1
-- ---- ----
comparative
9 0 0 6 2 0 3 8 2 2 3 .
1
example
900 660 392 23. 2
II
G 900 620 378 2~5. 1
29
Comparative steels H and J have small C content
and Si content, and hence they have high relaxation
rates, respectively. Since comparative steel K has no
Mo, it has a high relaxation rate. Even if each of
steels A, D, E and G has the contents of the present
invention, the relaxation rate is not largely reduced
if tempering temperature increases and hardness is
lower than annealed hardness HV400, as shown in
comparative examples IT.
While the presently preferred embodiment of the
present invention has been shown and described, it is
to be understood that this disclosure is for the
purpose of illustration and that various changes and
modifications may be made without departing from the
scope of the invention as set forth. in the appended
claims.
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