Note: Descriptions are shown in the official language in which they were submitted.
20~01~
R~GROUND OF THE lNv~NlION
This invention relates to a method of manufacturing an
oriented silicon steel sheet having improved magnetic
characteristics and, more particularly, to an improved method
of preventing reduction of magnetic flux density
notwithstanding reduction of thickness of the silicon steel
sheet.
High magnetic flux density and a small core loss are
magnetic characteristics required in grain-oriented silicon
steel sheets. Recent progress in manufacture techniques has
made it possible to make, for example, a silicon steel sheet
having a magnetic flux density B8 (the value at a magnetizing
force of 800 A/m) of 1.92 T for a sheet having a thickness of
0.23 mm. It is also possible to manufacture, on an industrial
scale, an improved silicon steel sheet product having a core
loss characteristic W"/50 (value under a fully magnetized
condition: 1.7 T at 50 Hz) of 0.90 w/kg.
Silicon steel sheets having such improved magnetic
characteristics have crystalline structures in which the <001>
directions parallel to the axis of easy magnetization are
uniformly aligned in the direction of rolling of the steel
sheet. Such a texture is formed during finishing annealing
by a phenomenon called secondary recrystallization in which
crystal grains having a (110) ~001] direction called the Goss
directian are grown with priority into giant grains.
Fundamental requirements for effectively growing secondary
recrystallized grains include the existence of an inhibitor
for limiting the growth of crystal grains having undesirable
2048014
directions other than the (110) [001] direction in the
secondary recrystallization process and the formation of a
primary recrystallized crystalline structure suitable for
effectively developing secondary recrystallized grains in the
(110) [001] direction.
A fine precipitate of MnS, MnSe, AlN or the like is
ordinarily utilized as the inhibitor. The effect of the
inhibitor has been enhanced by adding a grain boundary
segregation type component such as Sb or Sn to the inhibitor.
Conventionally, methods in which NnS or MnSe is used as a main
inhibitor are advantageous in reducing the core loss of
certain sheets because they assist in reducing the sizes of
the secondary recrystallized grains. However, methods based
on laser irradiation or plasma jetting have recently been
provided to artificially form pseudo grain boundaries so that
the magnetic domains are fractionated and the core loss is
reduced. For this reason, the advantage of reducing the sizes
of the secondary recrystallized grains has been lost.
Further, the concept of increasing the magnetic flux density
of the steel sheet has become advantageous.
A method of manufacturing an oriented silicon steel sheet
having a large magnetic flux density is disclosed in Japanese
patent Publication 46-23820. According to this method, the
desired steel sheet can be manufactured by (a) introducing
Al into the steel as an inhibitor component, (b) quenching
to obtain cooling before final cold rolling to precipitate
AlN, and (c) increasing the rolling reduction of the final
cold rolling from a lower reduction to a higher reduction,
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like from 65 to 95 %.
The method of the Japanese Publication, however, entails
a problem in that the magnetic flux is abruptly reduced along
with the reduction of thickness of the product sheet. It is
very difficult or impossible to manufacture by the method of
the Japanese Publication the type of silicon steel sheet
presently in demand, e.g., a thin product having a thickness
of 0.25 mm or less and having a B8 value of 1.94 T or higher.
In Japanese patent Publication 46-23820, immersing a
steel sheet in hot water at 100C after annealing to quench
the sheet is disclosed, but there is no consideration or
mention of any phase of any carbides after quenching.
Ordinarily, in the case of slow cooling from 600C or lower,
carbides are precipitated from grain boundaries at a higher
temperature and are precipitated in crystal grains at a lower
temperature. Carbides precipitated are finer and have a
higher density if precipitation is started at a reduced
temperature. Accordingly, with respect to the first
embodiment of Japanese patent Publication 46-23820 in which
the time for cooling from 1,000 to 750C is about 10 seconds
and the time for cooling from 750 to 100C is about 25
seconds, it is not unreasonable to conclude that very fine
carbides having particle sizes of several tens of angstroms
are precipitated or that the extent of carbide precipitation
is limited and that the carbon is simply supersaturated in the
steel.
Japanese Patent Publication 56-3892 discloses a technique
for controlling carbides in other steels during cooling after
204801~
annealing. In this method, with respect to two-stage cold
rolling, the steel is cooled at a cooling speed of 150C/min
or higher from 600 to 300C during cooling after annealing
followed by final cold rolling so that the amount of solid
solution carbon after cooling is increased. This method is
intended to improve the magnetic characteristics of the steel
by increasing the amount of solid solution carbon in the steel
and by optimizing the aging effect between cold rolling paths.
Such an effect of solid solution carbon is well known in the
case of ordinary cold-rolled steel sheets. If the amount of
solid solution C or solid solution N before cold rolling is
increased, the (110) intensity in the recrystallized structure
formed by recrystallization annealing after cold rolling is
increased. In the case of oriented silicon steel sheets, the
(110) grains become nuclei for secondary recrystallization,
so that the number of secondary recrystallized grains is
increased, the secondary-recrystallized grains are finer, and
improved magnetic characteristics can be achieved. This
method, however, does not enable the magnetic flux density of
a thin oriented silicon steel sheet to be increased.
As a technique for controlling the form of C in steel to
increase the (110) intensity of the steel, a method of
precipitating many fine carbide grains during cooling after
intermediate annealing is disclosed in Japanese Patent Laid-
Open Publication 58-157917. In this method, quenching of the
steel to 300C is effected after intermediate annealing and
slow cooling is applied for 8 to 30 seconds through a
temperature range of 300 to 150C, thereby precipitating fine
2048014
carbides. The (110) intensity of the steel after
recrystallization is thereby increased so that the magnetic
characteristics of the steel are improved. However, the
magnetic characteristics achieved by these methods are at most
1.94 T with respect to Blo and 1.92 T with respect to B8 when
the sheet thickness is 0.3 mm, which value is not high enough
to be satisfactory.
Japanese Patent Laid-Open Publication 61-149432discloses
a technique based on setting the cooling speed of steel to
10C/s or higher at the time of cooling after intermediate
annealing, creating a work strain of 1 to 30 % during cooling
from 1,000 to 400C, and perorming finishing rolling at a
temperature in the range of 100C to 400C. According to this
method, a work strain of 1 to 30 % is created at a temperature
in the range of 1,000 to 400C in which the C diffusion speed
is very high to provide high-density dislocations, so that C
is finely precipitated at the dislocations and the (110)
intensity is increased. To finely precipitate C in
dislocations at a high density, the working is performed by
rolling, and a high cooling speed of 10C/s or higher is set
for the precipitation step. The core loss can be reduced to
a certain extent by this method but the magnetic flux density
achieved by this method is only 1.91 T with respect to B1o
(1.89 T with respect to B8), which is low.
20~80~4
OBJECTS OF THE INVENTION
It is an object of the present invention to provide a
method of manufacturing an oriented silicon steel sheet which
enables maintenance of high magnetic flux density
notwithstanding reduction of steel sheet thickness. Another
object is to achieve a high magnetic flux density with desired
stability while reducing the core loss of steel sheet.
SUMM~RY OF THE lNV~;N'l'ION
It has been discovered that, in an Al-containing oriented
silicon steel sheet in which Sb is also present, the
precipitation of carbides is greatly changed during cooling
for annealing before final cold rolling, and that such
precipitation is effective to increase the ultimate (lll)
intensity of the recrystallized structure after final cold
rolling of sheet rather than the (110) intensity, and that
carbides precipitated in crystal grains at a high temperature
in the range of about 200 to 500C under strain during cooling
for annealing before final cold rolling, which are
conventionally regarded as undesirable, surprisingly have the
effect of increasing the {111} <112> intensity while reducing
the {111} <uvw> intensity, more particularly the {111} <llO>
intensity, so that a very high magnetic flux density can be
obtained with stability irrespective of the thickness of the
final product.
That is, according to the present invention, there is
provided a method of manufacturing an oriented silicon steel
sheet having greatly improved magnetic characteristics in
which a hot-rolled steel sheet of a silicon steel containing
2048014
about 0.01 to 0.15 6 by wel~ht of acid-solub.Le Al and abo~lt
0.005 to 0.04 % by weight of Sb as lnhlbltor components ls
cold-rolled once or a plurality of times until its thickness
ls reduced to the desired predetermined final thickness. The
method further comprises
softening-annealing the steel sheet before final
cold-rolling;
successlvely guenchlng the steel sheet at a cooling
speed of about 15 to 500C/s to a temperature of about 500C
or lower;
creatlng upon the sheet a small strain ranglng from
about 0.005 to 3.0 % ln a temperature range from about the
temperature reached by quenching to about 200C;
controlling carblde preclpltation at an effective
coollng speed about 2~/s or less by cooling the steel sheet
durlng the stralnlng step or by coollng the steel sheet for
about 60 to 180 seconds after the strainlng step durinq whlch
the steel sheet ls malntalned ln essentlally the same
temperature range as the temperature range at the end of the
quenchlng step; and
thereafter performing final cold-rolling with a
rolllng reductlon of about 80 to 95 %.
Thls can be done in coniunctlon with additional
steps of effecting anneallng for primary recrystalllzatlon as
well as decarburization;
applying an anneallng separation agent; and
effectlng secondary recrystalllzatlon annealing and
purlflcatlon-annealing.
73461~27
E ~!
2048014
Other features and variations of the present
invention will become apparent from the following detailed
description of the invention.
BRIEF DESCRIPTION OF THE DRAWINGS
Figs. 1 to 4 are transmission-electron-microscopic
photographs of examples of structures of steel sheets after
8a
t ~ 73461-27
ii '.V
2048014
annealing followed by final cold rolling, showing forms of
carbides at a depth of one-tenth of the sheet thickness
measured from the surfaces of the steel sheets.
DETAILED DESCRIPTION OF THE INVENTION
First, the results of experiments on which the present
invention is based will be described below.
Al-containing oriented silicon steel sheets to which Sb,
Sn, Ge, Ni and Cu (well-known as additive components) were
separately added were provided. These sheets were rolled
different times to manufacture products; one group of these
steel sheets was cold-rolled only one time to obtain products
having a thickness of 0.30 mm, and another group was cold-
rolled twice to obtain products having a thickness of 0.23 mm.
The rolling reduction of the final cold rolling was set
at 88 %, and annealing immediately before final cold rolling
was performed at 1,150C for 90 seconds with respect to the
steel sheets cold-rolled one time, and at 1,100C for 90
seconds with respect to the steel sheets cold-rolled twice.
Cooling was performed by immersing each steel sheet in hot
water at 80C.
The results of this experiment are as shown in Table 1.
Each of the 0.30 mm thick steel sheets had a high magnetic
flux density while each of the 0.23 mm thick steel sheets had
a reduced magnetic flux density. The reduced sheet thickness
had seriously reduced the flux density in every case.
Table 1
Sample No. 1 2 3 4 5 6
Additive Constituent name No additive Ni Cu Sb Sn Ge
- Amount of additive (%) - 0.08 0.10 0.03 0.05 0.02
Magnetic flux Product 0.30mm 1.924 1.925 1.923 1.936 1.903 1.914
density B8 (T) thickness 0.23mm 1.885 1.885 1.887 1.894 1.882 1.884
O O
o
204801~
By ex~mi ning the results of Table 1 in detail, it is
evident that sample 4 in which Sb was present had a slightly
better magnetic flux density than the other five samples.
To ~x~mine the cause of this effect we examined the
textures of samples of decarburized primary recrystallized
sheets with respect to the samples having a product thickness
of 0.23 mm, and ex~m; ~ed the forms of precipitated carbides
in the steel of each sample after intermediate annealing. The
results of these ex~mi n~tions are shown in Table 2.
Table 2
Sample No. 1 2 3 4 5 6
Additive No additive Ni Cu Sb Sn Ge
constituent
(110) Intensity 0.15 0.16 0.18 0.12 0.22 0.25
(222) Intensity 7.3 7.5 7.0 8.8 6.4 6.8
Form of carbideMostly in solid Mostly in solid PrecipitatedPrecipitated finely O
precipitation in solution, partially solution, partially slightly coarselyand at a high ~,
intermediate precipitated finely precipitated finely in grainsdensity in grains
annealed sheet ~,
Precipitated sizeAbout 80~ About 80~ About 200~ About 50~
2048014
As can be understood from Table 2, no increase in the
(110) intensity is attributed to the presence of Sb as
observed in sample 4 containing Sb, unlike the effect that
might have been expected in view of conventional technical
concepts, but the (111) intensity (equivalent to (222)) was
remarkably increased in the sample containing Sb. Further,
different forms of carbides exist after annealing followed by
final cold rolling and, as a result of the addition of Sb, the
fine high-density precipitated state or the C solid solution
state was changed so that carbides were precipitated in the
form of slightly coarse grains (Table 2, column 4) having
particle sizes much greater than the others in the Table.
In contrast, in the case of addition of Sn or Ge,
carbides were finely precipitated at a high density, and the
(110) intensity of the primary recrystallized structure was
remarkably improved.
The cause of this special effect achieved by the presence
of Sb is not clear. However, it is speculated that the
tendency of Sb to strongly segregate at grain boundaries or
surfaces is related to the phenomenon leading to the
occurrence of specially precipitated forms of carbides.
With a view to positive utilization of such variations
of the forms of carbides before final cold rolling, and to
create new effects by varying cooling conditions, further
experiments were conducted. Tests were conducted on the same
Al-containing oriented silicon steel sheets as those used in
the above-described experiments to which only Sb was added,
and also on the same Al-contAining silicon steel sheets which
13
2048011
had no added component. The tested steels were processed by
ordinary two-stage rolling to product products each having a
thickness of 0.23 mm. In this experiment, the rolling
reduction of final cold rolling was set at 85 ~, annealing
before the final cold rolling (intermediate annealing) was
effected at 1,100C for 90 seconds, and cooling was effected
under the following different cooling conditions:
(a) Condition (a) wherein the steel sheet was quenched at a
rate of 50C/s until 500C was reached, and thereafter cooled
at a very low cooling speed of 0.5 to 2C/s by being inserted
in a heat maintaining furnace,
(b) Condition (b) wherein the steel sheet was quenched at a
rate of 50C/s until 350C was reached, and thereafter cooled
at a very low cooling speed of 0.5 to 2C/s by insertion into
a heat maintaining furnace,
(c) Condition (c) wherein the steel sheet was quenched at a
rate of 50C/s until 350C was reached, successively skin-
pass-rolled to reduce by O.S ~, and cooled at a very low
cooling speed of 0.5 to 2C/s by insertion into a heat
maintAi~in~ furnace,
(d) Condition (d) wherein the steel sheet was quenched at a
rate of 50C/s until 150C was reached, and thereafter cooled
at a very low cooling speed of 0.5 to 2C/s by insertion into
a heat maintaining furnace,
(e) Condition (e) wherein the steel sheet was immersed in hot
water at 80C so that the average cooling speed was 62C/s,
was maintained at 80C after being cooled to this temperature,
and was thereafter cooled naturally.
2048014
The products thereby manufactured were examined with
respect to magnetic flux density, (110) intensity and (222)
intensity of the decarburized primary recrystallized sheets
and the precipitated forms of carbides in the intermediate
annealed sheets. The results are shown in Table 3.
2048014
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16
2048011
Figs. 1 to 4 are transmission-electron-microscopic
photographs of the structures of steel sheets after annealing
followed by final cold rolling, showing forms of carbides at
a depth of 1/10 of the sheet thickness from the surfaces of
the steel sheets. Fig. 1 shows a sample to which Sb was added
and which was cooled under Condition (e), Fig. 2 shows a
sample (Table 3, column c, bottom) to which Sb was added and
which was cooled under Condition (c), Fig. 3 shows a sample
which had no additive component and which was cooled under
Condition (e) (Table 3, column 3, top), and Fig. 4 shows a
sample which had no additive component and which was cooled
under Condition (c) (Table 3, column c, top).
As is shown in Table 3, the magnetic flux density (B8)(T)
of the sample to which Sb was added (bottom half of Table 3)
and which was manufactured under the intermediate annealing
cooling condition (c) (Table 3, column c) was particularly
high. In this sample, carbide precipitates having a size
ranging from 300 to 500 A and sparsely precipitated were
observed after the intermediate annealing, and are shown in
Fig. 2, as heretofore noted. In contrast, in the sample which
had no additive component and which was manufactured under the
same cooling condition (c) (Table 3, column c, top), fine
carbide precipitates having a size of about 100 A were
undesirably precipitated at a high density, as shown in ~ig.
4.
With respect to the steel sheets which had no additive
- component, in the case of creating a work strain by skin-
pass-rolling in accordance with the condition (c), carbide
2048014
precipitation sites were increased during cooling so that
carbides were finely precipitated at a high density, as is
apparent from comparison with processing under the condition
(b). In contrast, with respect to the steel sheets to which
Sb was added, precipitation sites were not increased and
slightly coarse precipitates were observed. According to our
study after these experiments, such sparse precipitation of
carbides having a size ranging from 300 to 500 ~ increases the
(111) intensity of the structure primarily recrystallized by
decarburization annealing after final cold rolling and reduces
the {lll} <uvw>, in particular the {111} <110> intensity while
increasing the {111} <112> intensity. The {111} <110> grains
limit the growth of the (110) [001] secondary grains which
contribute to the increase in the magnetic flux density, while
the {111} <112> grains promote the growth of (110) [001]
secondary grains. It is thought that addition of Sb in the
particular process provides this effect and enables formation
of a product having a substantially high magnetic flux density
as in the case of Condition (c) as shown in the top portion
of Table 3.
It is thought that this effect of Sb in steel relates to
segregation of Sb, that Sb is segregated at base points in
crystal grains such as to form carbide precipitation sites,
and that this segregation results from the limitation of
precipitation carbides during cooling.
This action of Sb is particularly effective in a
temperature range of about 200 to 500C; the amount of strain
to be applied may be very small, e.g., about 0.005 to 3 %.
18
2~0i~
It has also been found that the aging effect at the time of
final cold rolling can also be improved according to this
invention because the amount of solid solution carbon is
increased by the carbide precipitation limiting effect of Sb.
It is known that a small strain of 0.5 % created by skin-
pass-rolling is concentrated at a surface-layer portion of the
steel sheet. In this work as well, the form of precipitated
carbides was changed according to the change in the amount of
strain in the thickness direction of the sheet, and the
density of precipitated carbides was reduced toward the center
of the sheet in the thickness direction.
The fact that the form of precipitated carbides was
changed in the sheet thickness direction is regarded as a
reason for the success of this work. To positively utilize
this effect, a similar experiment was also conducted by
creating a strain of 0.5 % by bending with a leveler, and
suitable effects were thereby obtained.
A carbide precipitation processing method is disclosed
in Japanese Patent Laid-Open 61-149432. In this method, high-
density dislocations uniform in the direction of sheet
thickness are provided by rolling at a high temperature of
1,000 to 400C, and the speed of cooling in a step of
precipitating carbon is high, such as 10C/s. This method is
intended to precipitate finely divided carbides and to
increase the (110) [001] intensity of the texture of the
product.
Japanese Patent Laid-Open 58-15797 also discloses a
technique for precipitating carbides of a size of 100 to
2048014
500 ~. In this case, however, the precipitation temperature
range is a range of low temperatures, i.e., 300 to 150C, and
the effect of Sb is not effectively utilized, and there is no
disclosure or suggestion of our special ideas relating to the
precipitation processing which constitutes a feature of the
present invention, including that of creating a strain during
precipitation. This technique is therefore sharply different
from the present invention with respect to the carbide
precipitation density and requires high-density precipitation
for increasing the (110) [001] intensity as in the case of the
method disclosed in Japanese Patent Laid-Open 61-149432.
In contrast, in accordance with the present invention,
it is important to precipitate carbides sparsely to reduce the
{111} <uvw> intensity, in particular the {111} <110> intensity
of the primary recrystallized structure while increasing its
{111} <112> intensity.
It is important to define the ranges of chemical
components of the composition of the oriented silicon steel
sheet in accordance with the present invention. Preferable
ranges of the components will be described below.
C is necessary for improving the hot-rolled structure of
the steel. However, if the C content is excessive, it is
difficult to decarburize the steel. It is therefore
preferable to limit the carbon content to a range of about
0.035 to 0.090 % by weight.
If the Si content is below a lower limit the desired core
loss characteristic cannot be obtained. If the Si content is
excessive it is difficult to perform cold rolling. It is
2048014
preferable to provide an Si content in the range of about 2.5
to 4.5 % by weight.
Mn can be utilized as an inhibitor component. In case of
an excessively large amount of Mn, Mn compound in the steel
cannot be dissolved during slab-reheating process, and it is
accordingly preferable to provide an Mn content in the range
of about 0.05 to 0.15 % by weight.
S or Se is effective when combined with Mn to form MnS
or MnSe which acts as an inhibitor. The range of S or Se
content for finely precipitating MnS or MnSe is preferably
about 0.01 to 0.04 % by weight in either case of whether used
alone or together.
It is specifically necessary for the steel sheet of the
present invention to contain acid-soluble Al or N as inhibitor
components for the purpose of achieving a high magnetic flux
density, and addition of certain amounts of acid-soluble Al
or N is required. However, if these contents are excessive
fine precipitation is difficult. It is preferable to maintain
the content of acid-soluble Al to a range of about 0.01 to
0.15 % by weight and the content of N to a range of about
0.0030 to 0.020 % by weight.
Further, according to the present invention, the presence
of Sb in the steel is indispensable, and it is possible to
limit precipitation of C at grain boundaries or in crystal
grains in the steel by providing a content of Sb. To enable
such an effect, about 0.005 % or greater by weight of Sb is
necessary. However, if the Sb content exceeds about 0.040 %
by weight, the problem of grain boundary embrittlement is
2048014
encountered, and it is difficult to perform cold rolling. The
Sb content is therefore maintained within a range of about
0.005 to 0.040 % by weight.
To improve magnetic properties, other inhibitor
strengthening components such as Cu, Cr, Bi, Sn, B, Ge and the
like may be added as desired. The content of each of such
components may be within well-known ranges. To prevent
occurrence of surface defects due to hot-rolling
embrittlement, it is preferable to add Mo in a range of about
0.005 to 0.020 % by weight.
Next, a process of manufacture in accordance with the
present invention will be described below.
Well-known manufacturing methods are applied for
manufacturing the steel sheet, and ingots or slabs are
reproduced as desired, adjusted to the desired size, and
thereafter heated and hot-rolled. The hot-rolled steel sheet
is processed by cold rolling one time or in a plurality of
stages until its thickness is reduced to a desired final
thickness.
For annealing before final cold rolling a high
temperature in a range of about 850 to 1,200C is required to
dissolve AlN, and, after this annealing, quenching to 500C
or lower is required to precipitate AlN and it is also
necessary to prevent precipitation of C at grain boundaries.
If the cooling speed is lower than 15C/s, C is pxecipitated
at grain boundaries, or, if the cooling speed exceeds 500C/s,
the shape of the steel sheet after the cooling is
deteriorated. The cooling rate is therefore maintained within
2048014
a range of about 15 to 500C/s.
Thereafter, a small strain ranglng from about 0.005
to 3.0 % is created ln a temperature range from the
temperature reached by ~uenching (about 500C at the maximum)
to about 200C. The steel sheet is cooled slowly at a coolln~
speed of about 2C/s or lower during this straining or after
the straining for a perlod of tlme of about 60 to 1~0 seconds
ln which the steel sheet is maintained in essentially the same
temperature range.
This step ls intended to preclpitate sparsely
arranged carbldes havlng a slze ranging from about 300 to
500 ~ ln grains, whlch effect relates to one of the most
lmportant features of the present lnvention. This processinq
ls performed withln a hlgh temperature range from the
temperature reached by cooling, i.e., about 500C at the
maximum to about 200C, and a strain is created in this
temperature range, a feature unknown before the present
invention. The precipitation of carbides is controlled to
provide the deslred slze and denslty by balanclng three
influenclng factors including (a) the fact that the C
diffusion speed is comparatively high so that carbides are
coarsely formed, (b) the fact that the carbide precipitation
polnts are lncreased by straining so that carbides precipltate
flnely at a hlgh density, and (c) the fact that precipitation
of carbides at graln boundarles and ln crystal gralns ls
llmlted by the segregation effect of the presence of Sb.
E
73461-~7
, . . _ .. ., .. ~
2048014
Carbide preclpltates have an excesslvely large slze
if the precipitation temperature exceeds about 500C. They
are
~, 23a
73461-27
204801~
excessively fine if the precipitation temperature is lower
than about 200C. Preferably the temperature at which
precipitation is performed is within the range of about 450C
to 300C.
If the maintenance time is shorter than about 60 seconds,
the carbides are not formed sufficiently coarsely. If it is
longer than about 180 seconds, carbides are formed excessively
coarsely, and the number of precipitation points is increased
and the amount of solid solution is considerably reduced, with
undesirable results.
When slow cooling is performed instead of the constant-
temperature maintenance step it is necessary to
set the cooling speed to about 2C/s or lower.
It is necessary to effect straining immediately after
quenching or in the temperature range of about 500 to 200C
before the carbon precipitation processing. It is thereby
possible to prevent carbides from precipitating excessively
coarsely. If the amount of strain provided is less than about
0.005 % by weight, the carbides are formed excessively
coarsely. If the strain is more than about 3.0 %, carbides
are finely precipitated at an excessively high density. The
amount of strain is therefore set within a range of about
0.005 to 3.0 ~. A range of 0.01 to 1.0 % is particularly
preferable.
Needless to say, straining may be performed by any
conventional straining method, e.g., a skin pass method based
on rolling, a bending method using a bending roll, a straining
method using a leveler roll, shot blasting, or the like.
24
20~8014
The steel sheet is then subjected to final cold rolling.
At this time, to obtain a high magnetic flux density, it is
necessary to set the rolling reduction to a range of about 80
to 95 %, as is well known.
Performing well-known aging or hot rolling treatment
during this final cold rolling is further effective in the
process of the present invention, because the amount of solid
solution C in the steel of the present invention is large.
The aging temperature is preferably adjusted to the range of
about 200 to 400C. If the aging temperature is higher than
about 400C the shapes of precipitated carbides are changed
so that the object of the present invention cannot be
achieved. If the aging temperature is lower than about 200C,
solid solution C or solid solution N is not sufficiently fixed
on dislocations, and further improvements in characteristics
cannot be expected.
It is necessary to set the rolling reduction to a range
of about 80 to 95 %, as is well known. If the rolling
reduction is less than about 80 %, a sufficiently high
magnetic flux density cannot be obtained. If the rolling
reduction exceeds about 95 %, it is difficult to develop
secondary recrystallization grains.
The steel sheet after final cold rolling is degreased and
is then annealed for decarburization and primary
recrystallization. An annealing separation agent having MgO
as a main component is thereafter applied to the steel sheet,
and the steel sheet is coiled to be subjected to finishing
annealing and is coated with an insulating material if
204801~
necessary. Needless to say, the steel sheet may also be
processed to fractionate magnetic domains by laser, plasma or
any other means.
(Examples)
Example 1
Eleven steel ingots B, D, E, F, G, H, I, J, X, L, and M
shown in Table 4 were provided in conformity with the present
invention. These steels and other two steels A, C provided
as comparative examples, thirteen steels in all were hot
rolled in a conventional manner to form hot-rolled coils each
having a thickness of 2.2 mm.
20'18014
n ~ I ~ ~
Z v ~ D ' V ~ D ,_, ,,, ,,, ,V ,V V ~ V v
z ~ ~ U~ ~ O ~ ~ ~ U~ ~ o ~
-
6 ~ ~ ~ ~ ~ N ~ ~ ~
-
0
O
V V V V V V V VV V J- .
o
N N ~N
eo o o o o oo oo o ~o o
U~ . .. .. . .. . .
o o o o o oo oo o oo o
~O O O O O OO OO O OO O
U
O O O O O OO OO O OO O
In
V V V V V VV VV . VV V
o
O O _I~ `1 ~ ~'1 t~
rn V
o o o oo o o o o o o o
:~ O O O OO O O O O O O O O
o O O O O OO O O O O O O O
~r
E- v. o ~ ~ ~ ~~ ~ o
~V V V V VV . VVVVV~
O
r a~ oa~co o N O ~ ~ a~
rn v v v
o o oo o o o o o o
_Io o o oo o o o o o o o
rn o o o oo o o o o o o o o
.
o o o o oo o o o o o o o
,~ O O O OO O O O O O O O O
O O O O OO O O O O O O O
O O O OO O O
O O O O OO O O O O O O O
O O O O OO O O O O O O O
.
O O O O OO O O O O O O O
rn
o~ u~ C O ~ ~ ~ ~
u o o o o oo o o o o o o o
- o o o o oo o o o o o o o
o ~
r~ u a ~~ r~
n
204801~
Each steel sheet was thereafter subjected to normal
annealing at l,000C for 90 seconds and was cold-rolled until
its thickness was reduced to an intermediate thickness of 1.50
mm. The reduced steel sheet was further annealed at 1,100C
for 90 seconds, quenched at a rate of 60C/s to 350C, and
passed through a slow cooling box having a bending roll and
was thereby strained to an extent of 1.5 ~ while being cooled
at a rate of 2C/s to 200C. The steel sheet was thereafter
cooled in atmospheric air.
- The steel sheet was then rolled until its thickness was
reduced to a final thickness of 0.22 mm, electrolytically
degreased, and subjected to decarburization/primary
recrystallization annealing at 850C for 2 minutes in a wet
hydrogen atmosphere. An MgO agent containing 5 % TiO2 was
then applied to the steel sheet, and the steel sheet was
subjected to finishing annealing at 1,200C for 10 hours.
Thereafter, the surfaces of the sheet were coated to give the
steel sheet tensile stress and were partially processed to
fractionate magnetic domains at 10 mm pitches by the plasma
jet method. Table 5 shows the magnetic characteristics before
and after the magnetic domain fractionating processing of the
steel sheets.
28
204801~
Table 5
Magnetic domain MagnetiC flux Core loss
Ingot fractionating density Wl7/so (W/kg)Note
symbol processing B8 (T)
A Unprocessed 1.875 1.15 Comparative
Processed 1.874 1.09 example
B -Unprocessed 1.935 0.92 Conformable
Processed 1.936 0.78 example
C Unprocessed 1.883 1.07 Comparative
Processed 1.883 1.02 example
D Unprocessed 1.938 0.95 Conformable
Processed 1.938 0.84 example
E Unprocessed 1.941 0.87 ditto
Processed 1.942 0.73
F Unprocessed 1.946 0.85 ditto
Processed 1.945 0.70
G Unprocessed 1.942 0.86 ditto
Processed 1.943 0.72
H Unprocessed 1.937 0.97 ditto
Processed 1.938 0.83
I Unprocessed 1.940 0.87 ditto
Processed 1.941 0.72
J Unprocessed 1.941 0.83 ditto
Processed 1.941 0.70
K Unprocessed 1.938 0.86 ditto
Processed 1.937 0.73
L Unprocessed 1.942 0.85 ditto
Processèd 1.943 0.71
M Unprocessed 1.939 0.88 ditto
Processed 1.938 0.75
Note: * Magnetic domain fractionating at lOmm pitches by plasma jet
method
29
20~8~4
As appears in Table 5, the conformable examples (all
except A and C) have characteristics improved in magnetic flux
density and core loss due to this invention, in comparison
with those of the comparative Examples A and C. The magnetic
flux density of the conformable examples was 1.946 T (Ingot
F) at the m~;mum with respect to B8, as compared to 1.875 and
1.883 for comparative Examples A and C. The magnetic domain
fractionating processing remarkably improved the core loss but
did not substantially adversely influence the magnetic flux
density.
Example 2
The steel ingot F shown in Table 4 was hot-rolled in a
conventional manner to provide hot-rolled steel sheets having
thicknesses of 2.4, 2.2, 2.0, and 1.5 mm.
The hot-rolled steel sheets having thicknesses of 2.4 and
2.2 mm were respectively annealed at 1,175C for 90 seconds
and at 1,150C for 90 seconds, then quenched to 400C at an
average cooling speed of 50C/s, strained to an extent of 2
% by a hot skin pass roller, slowly cooled to 250C at an
average cooling speed of 1.5C/s, and quenched in water.
Thereafter, these steel sheets were respectively cold-rolled
to final thicknesses of 0.30 and 0.28 mm. When the
thicknesses of these steel sheets were respectively reduced
to 1.3 and 1.0 mm, each sheet was separated into two. One of
them was successively cold-rolled and the other was aged at
300C for 2 minutes and cold-rolled to the final thickness.
The hot-rolled steel sheets having thicknesses of 2.0 and
1.5 mm were normalized at 1,000C for 90 seconds, naturally
20~801~
cooled, respectively cold-rolled to thicknesses of 1.4 and 1.1
mm, annealed at 1,100C for 90 seconds, and quenched to 350C
at an average speed of 60C/s. They were then strained to an
extent of 1.0 % by a hot leveler, maintained at 320C for 120
seconds, and taken out of the furnace and naturally cooled.
Thereafter they were respectively cold-rolled to final
thicknesses of 0.20 and 0.15 mm. When the thicknesses of
these steel sheets were respectively reduced to 0.7 and 0.55
mm, each sheet was separated into two. One of them was
successively cold-rolled and the other was aged at 300C for
2 minutes and cold-rolled to the final thickness. After final
cold rolling the steel sheets were degreased and subjected to
decarburization/primary recrystallization annealing at 850C
for 2 minutes in a wet hydrogen atmosphere. An MgO containing
2 % SrSO4 was then applied to the steel sheets, and the steel
sheets were subjected to finishing annealing at 1,200C for
10 hours. Thereafter the surfaces of the sheets were coated
to give a tensile stress to the sheets and processed to
fractionate magnetic domains by 5 mm pitch electron beam
irradiation. Table 6 shows the magnetic characteristics of
the steel sheets thus processed.
Table 6
Item Non-aged Aged *
Final
thickness (mm) B8 (T)Wl7/so (W~Kg) B8 (T)W17/50 (W/Kg)
0.30 1.9420.97 1.9450.90
0.28 1.9480.93 1.9440.88
0.20 1.9400.87 1.9420.82
w
N
0.15 1.934 0.86 1.930 0.77
Note: * Aged at 300C for 2 minutes during cold rolling
o
204801~
As appears in Table 6 the magnetic flux density was
improved even though the final thickness was substantially
reduced down to 0.15 mm, and the magnetic domain fractionating
processing during the cold rolling remarkably improved the
core loss but did not substantially influence the magnetic
flux density.
Example 3
The ingot G shown in Table 4 was hot-rolled in a
conventional manner to provide a hot-rolled coil having a
thickness of 2.0 mm. This steel sheet was normalized at
- 1,000C for 90 seconds and was cold-rolled to an intermediate
thickness of 1.50 mm. ThiS steel sheet was separated into
three pieces and all were subjected to intermediate annealing
at l,100C for 90 seconds. This cooling was performed under
three different sets of conditions.
The first set of conditions (I) was that the steel sheet
was cooled in hot water at 80C.
The second set of conditions (II) was that the steel
sheet was cooled to 350C at an average cooling speed of
60C/s, was slowly cooled to 300C for 2 minutes while being
strained to an extent of 0.5 % by a bending roll, and was
cooled in atmospheric air.
The third set of conditions (III) was that the steel
sheet was cooled to 400C at an average cooling speed of
60C/s, was cooled to 250C at a cooling speed of 2C/s, and
was cooled in atmospheric air.
Each of these three steel sheets was separated into two.
One of them was cold-rolled in a conventional manner to a
2~801~
final thickness of 0.20 mm, while the other was hot-rolled at
250C to a final thickness of 0.20 mm. After final cold
rolling, all the steel sheets were degreased and subjected to
decarburization/primary recrystallization annealing at 860C
for 2 minutes in a wet hydrogen atmosphere. An MgO separator
containing 10 % TiOz was then applied to the steel sheets, and
the steel sheets were subjected to finishing annealing at
1,200C for 10 hours. Thereafter the surfaces of the sheets
were tension-coated and the magnetic characteristics were
measured. Table 7 shows the results of this measurement.
34
Table 7
Item Normally rolled sheet Warm-rolled sheet *
Cooling Note
conditionB8 (T) Wl7/50 (W/Kg) B8 (T) Wl7/so (W/Kg)
(I) 1.882 1.08 1.888 0.97 Comparative
example
(II) 1.939 0.85 1.941 0.82 Conformable
example
w (III) 1.896 1.05 1.894 1.95 Comparative
example
Note: * Finishing-cold-rolled at 250C
o
2~01~
As shown in Table 7, the conformable example processed
under the cooling conditions (II) was improvçd in both
magnetic flux density and core loss in comparison with the
comparative examples processed under the cooling conditions
(I) and (III), and it was found that the creation of a small
strain in a temperature range of 500 to 200C during the
cooling for the annealing before the final cold rolling was
effective in improving the magnetic characteristics of the
sheet.
According to the present invention, a silicon steel sheet
cont~ining Al and Sb is used and cooling control and creation
of a small strain are effected during cooling for annealing
before final cold rolling, so that an oriented silicon steel
sheet having a high magnetic flux density can be manufactured
with stability even if the sheet thickness is reduced. The
oriented silicon steel sheet manufactured in accordance with
the present invention has excellent properties for use in
transformer cores and other products having high magnetic flux
density and good stability with reduced core loss.
