Note: Descriptions are shown in the official language in which they were submitted.
3 3 ~
METHOD ~F MA~lNG REGUI~R
GRAIN ORIENTED SILICO~ STEEL
WITHOUT A H()T BAND A~iJEAL
Jerr~ W. Schoen
ECHNICAL FIELD
The present invention relates to a process o
producing regular grain oriented silicon steel in
thicknesses ranging from about 18 mils (0.45 mm) to
about 7 mils ~0.18 mm) without a hot band anneal, and
to such a process wherein the intermediate anneal
following the first cold rolling stage has a very
lS short soak time and a two-part temperature-controlled
cooling cycle to control carbide precipitation.
BACRGROUND ART
The teachings of the present invention are
applied to silicon steel having a cube-on-edge
orientation, designated (110) [001~ by Miller's
Indices. Such silicon steels are generally referred
to as grain oriented silicon steels. Grain oriented
silicon steels are divided into two basic
categories: regular grain oriented silicon steel and
high permeability grain oriented silicon steel.
Regular grain oriented silicon steel utilizes
manganese and sulfur (and/or selenium) as the
0 principle grain growth inhibitor and generally has a
permeability at 796 A/m o less than 1870. High
permeability silicon steel relies on aluminum
nitrides, boron nitrides or other species known in
the art made in addit;on to or in place of manganese
sulphides and/or selenides as grain growth
- 2 ~ 3 ~ ~
1 inhibitors and has a permeability greater than 1870.
The teachings of the present invention are applicable
to regular grain oriented silicon steel.
Conventional processing of regular grain oriented
silicon steel comprises the steps of preparing a melt
of silicon steel in conventional facilities, refining
and casting the silicon steel in the form of ingots
or strand cast slabs. The cast silicon steel
preferably contains in weight percent less than about
0.1% carbon, ~bout 0.025% to about 0.25% manganese,
ahout 0.01% to 0.035% sulfur and/or selenium, about
2.5% to about 4.0~ silicon with an aim silicon
content of about 3.15%, less than about 50 ppm
nitrogen and less than about 100 ppm total aluminum,
the balance being essentially iron. Additions of
boron and/or copper can be made, if desired.
If cast into ingots, the steel is hot rolled into
slabs or directly rolled from ingots to strip. If
continuous cast, the slabs may be pre-rolled in
accordance with U.SO Patent 4,713,951. If developed
commercially, strip casting would also benefit from
the process of the present invention. The slabs are
hot rolled at 2550 F (1400 C) to hot band thickness
and are subjected to a hot band anneal of about
1850 F (1010 C) wi~h a soak of about 30 seconds.
The hot band i5 air cooled to ambient temperature.
Thereafter, the material is cold rolled to
intermediate gauge and subjected to an intermediate
anneal at a temperature of about 1740 F (950 C)
with a 30 second soak and i5 cooled as by air cooling
to ambient temperature. Following the intermediate
anneal, silicon steel is cold rolled to final gauge.
The silicon steel at final gauge is subjected to a
conventional decarburizing anneal which serves to
2 ~ 3 ~ ~
1 recrystallize the steel, to reduce the carbon content
to a non-aging level and to form a fayalite surface
oxide. The decarburizing anneal is generally
conducted at a temperature of from about 1525 F to
about 1550 ~ (about 830 C to about 845 C3 in a wet
hydrogeD bearing atmosphere for a time sufficient to
bring the carbon content down to a~out 0.003% or
lower. Thereafter, the silicon steel is coated with
an annealing separator such as magnesia and is bo~
annealed at a temperature of about 2200 F (1200 C~
for twenty-four hours. This final anneal brings
about secondary recrystalliæation. A forsterite or
"mill" glass coating is formed by reaction of the
fayalite layer with the separator coating.
Representative processes for producing regular
grain oriented (cube-on-edge) silicon steel are
taught in U.S. Patent Nos. 4,202,711; 3,764,406; and
3,843,422.
The present invention is based upon the discovery
that in the conventional routing given above, the hot
band anneal can be eliminated if the intermediate
anneal and cool;ng practice of the present invention
is ~oll~wed. The intermediate anneal and cooling
procedure of the present invention contemplates a
very short soak preferàbly at lower temperatures,
toyether with a temperature controlled, two-stage
cooling cycle~ as will be fully described
hereinafter
The teachings of the present invention yield a
number o~ advantages over the prior art. At all
final gauges within the above stated range, magnetic
guality is achieved which is at least equal to and
3 ~ ~
l often better than that achieved by the conventional
routing. The magnetic quality is also more
consistent. The teachings of the present invention
shorten the annealing cycle by from 20% or more,
thereby increasing line capacity. The process of the
present invention enables for the first time the
manufacture of thin gauge, typically about 9 mils
~0.23 mm) to about 7 mils (0.18 mm), regular grain
oriented silicon steel having good magnetic
characteristics without a hot band anneal following
hot rolling to hot band. This enables thin gauge
regular grain oriented silicon steel to be
manufactured where hot band annealing can not be
practic~d. The lower tempPrature of the intermediate
anneal of the present invention increases the
mechanical strength of the silicon steel during the
anneal, which previously was marginal at high
annealing temperatures.
European Patent 0047129 teaches the use of rapid
coolin~ from 1300 F to 400 F (705 C to 205 C) for
the production of high permeability electrical
steel. This rapid cooling enables the achievement of
smaller secondary grain size in the final product.
U.S. Patent 4,517,932 teaches rapid cooling and
controlled carbon loss in the intermediate ann~al for
the production of high permeability electrical steel,
including an aging treatment at 200 F to 400 F
(95 C to 205 C) for from lO to 60 seconds to
condition the carbide
These high permeability silicon steel references
employ a very low temperature and lengthy
intermediate anneal cycle having a 120 second soak at
1600 F (870~ C) followed by rapid ~ooling from
2 ~
1 1300 F (705 C) and an aging treatment to condition
the carbide precipitates. It has been found,
however, that in the intermediate anneal of the
present invention, rapid cooling from above about
1150 F (620 C) or higher produces poorer maqnetic
quality owing to the formation of martensite which
increases hardness, degrades mechanical properties
for subsequent cold rolling, and contributes to
poorer magnetic quality in the final product.
In the above-noted U.S. Patent 4,517,032, a low
temperature aging treatment following rapid cooling
is employed. This practice, if used for regular
grain oriented materials, has been found to produce
enlarged secondary grain size and poorer magnetic
quality in the final product since it impaires the
fine iron carbide precipitates. Lower tamperature
annealing at about 1640 F (895 C~ or lower, to
avoid the formation of austenite, could be used to
provide adequate solution of iron carbide without
forming a second phase which must be conditioned out
of the microstructure. However, this procedure
requires much longer annealing times to effect
carbide solution. Such a procedure would permit
direct rapid cooling from soak temperature without
the two-~tage cooling cycle of the present invention.
U.S. Patent 4,478,653 teaches that a higher
intermediate anneal temperature can be used to
produce 9 mil (0.23 mm~ regular grain oriented
silicon steel without hot band annealing. It has
been found, however, that 9 mil (0.23 mm) regular
grain oriented silicon steel made in accordance with
this patent has more variable magnetic quality than
when a routing utilizing a hot band anneal i.s used.
2 ~ 3 f' ~
1 It has further been found that the no hot band
anneal-high temperature intermediate anneal practice
taught in this reference provides generally poor
magnetic quality at thinner gauges of 9 mils (0.23
mm) or less, when compared to the above noted
practice employing a hot band anneal. Finally, the
very high temperature of the intermediate anneal of
U.S. Patent 4,478,653 results in low mechanical
strength of the silicon steel, making processing more
~ifficult.
DISCL~SUR~ OF THE I~VE~TIO~
According to the invention, there is provided a
method ~or processing regular grain oriented silicon
steel having a thickness in the range of from about
18 mils (0.45 mm) to about 7 mils (0.18 mm)
comprising the steps of providing silicon steel
consisting essentially of, in weight percent, of less
than about 0.1% carbon, about 0.025% to 0.25%
manganese, about 0.01% to 0.035% sulfur and~or
selenium, about 2.5% to 9.0% silicon, less than about
100 ppm total aluminum, less than about 50 ppm
nitrogen, the balance being essentially iron.
Additions o boron and/or copper can be made, i~
de~ired.
The silicon steel is cold rolled from hot band to
intermediate thickness without a hot band anneal.
The cold rolled intermediate thickness silicon steel
is subjected to an intermediat~ anneal at about
1650 F to about 2100 F (about 900 C to about
1150 C) and preferably from about 1650 F to about
1700 F (from about 900 C to about 930 C) for a
soak time of from about 1 to about 30 seconds, and
2 ~
1 preferably for about 3 to 8 seconds. Following this
soak, the silicon steel is cooled in two stages. The
first is a slow cooling stage from soak temperature
to a t~mperature of from 1000 F to 1200 F (540 C
to 650 C), and preferably to a temperature of
1100 F ~ 50 F (595 C ~ 30 C) at a rate less than
about 1500 F ~835 C) per minute, and preferably at
a rate of from about 500 F (280 C~ to 1050 F
(585 C) per minute. The second stage is a fast
cooling stage at a rate of greater than 1500 F
(835 C) per minute, and preferably at a rate of
2500 F to 3500 F (1390 C to 1945 C) per minute
followed by a water quench at about 600 F to about
700 F (about 315 C to about 370 C). Following the
intermediate anneal, the silicon steel is cold rolled
to final thickness,- decarburized, coated with an
annealing separator, and subjected to a final anneal
to effect secondary recrystallization.
o BRIEF_DESCRIPTI~ OF T~IE DRAWI~IG
The Figure is a graph illustrating th~
intermediate anneal time/temperature cycle of the
present invention and that of a typical prior art
intermediate anneal.
DESCRIPTIO~ O~ THE PRE:E'ERPcED EMBODIME~T~;
In the practice of the present invention, the
routing for the regular grain oriented silicon steel
is conventional and is the same as that given above
with two exceptions. The first e~ception is that
there is no hot band anneal. The second exception is
the development of the intermediate anneal and
cooling cycle of the present invention, following the
first stage of cold rolling.
2~J ~: 3 rJ ~
To this end, the starting material referred to as
"hot band" can be produced by a number of methods
known in the art such as ingot casting/continuous
casting and hot rolling, or by strip casting~ The
silicon steel hot band scale is removed, but no hot
band anneal prior to the first stage of cold rolling
is practiced.
Following the first sta~e of cold rolling, the
silicon steel is subjected to an intermediate anneal
in accordance with the teachings of the present
inv~ntion. Reference is made to the Figure, which is
a sch~matic of the time/temperature cycle for the
intermediate anneal of th~ present invention. The
Figure also shows, with a broken line, the
time/temperature cycle for a typical, prior art
intermediate anneal.
A primary thrust of the present invention is the
discovery that the intermediate anneal and its
cooling cycle can be adjusted to provide a fine
carbide dispersion. The refinement of the carbide
enables production of regular grain oriented silicon
steel over a wide range of melt carbon, even at final
gauges of 7 mils (O.lB mm) and less, having good and
consistent magnetic properties in the final product
without the necessity of a hot band annealing step.
During the heat-up portion of the intermediate
annealt recrystallization occurs at about 1250 F
(675 C~, roughly 20 seconds after entering the
furnace, after which normal grain growth occurs. The
start of recrystallization is indicated at "0" in the
Figure, Above about 1280 F (690 C) carbides will
3 ~ ~
1 begin dissolving, as indicated at "A" in the Figure.
This event continues and accelerates as the
temperature increases. Above about 1650 F (900 C~,
a small amount of ferrite transforms to austenite.
The austenite provides for more rapid solution o
carbon and restricts normal grain growth, thereby
establishing the intermediate annealed grain size.
Prior art intermediate anneal practice provided a
soak at about 1740 F (950 C~ for a period of from
25 to 30 seconds. The intermediate anneal procedure
of the present invention provides a soak time of from
about 1 to 30 seconds, and preferably from about 3 to
8 seconds. The soak temperature has been determined
not to be critical. The soak can be conducted at a
temperature of from about 1650 F (900 C~ to about
2100 F (1150 C). Preferably, the soak is conducted
at a temperature of from about 1650 F (900 C) to
about 1700 F (930 C), and more prefarably at about
1680 F (915 C). The shorter soak time and the
lower soak temperature are preferred because less
austenite is formed. The austenite present in the
form of dispersed islands at the prior ferrite grain
boundaries is finer. Thus, the austenite is easier
to decompose into ferrite with carbon in solid
solution for subsequent precipitation of fine iron
carbide. To e~tend either the soak temperature or
time results in the enlargement of the austenite
islands which rapidly become carbon-rich compared to
the prior ferrite matri~. Both growth and carbon
enrichment of the austenite hinder its decomposition
during cooling. The desired structure exiting the
furnace consists of a recrystallized matri~ of
ferrite having less than about 5~ austenite uniformly
dispersed throughout the material as fine islands.
At the end of the anneal, the carbon will
-- 10 --
2 ~
1 be in solid solution and ready ~or reprecipitation on
cooling. The primary reason behind the redesign of
the intermediate anneal time and temperature at soak
is the control of the growth of the austenite
islands. The lower temperature reduces the
equilibrium volume fraction of austenite which
forms. The shorter time reduces carbon diffusion,
thereby inhibiting growth and undue enrichment of the
austenite. The lower strip tempPrature, the reduced
volume fraction and the finer morphology of the
austenite makes it easier to decompose during the
cooling cycle.
Immediately after the soak, the cooling cycle is
initiated. The cooling cycle o~ the present
invention contemplates two stages. The first stage
e~tending from soak to the point "E" on the Figure is
a slow cool from soak temperatllre to a temperature of
from about 1000 F (540 C) to about 1~00 F (650 C)
and preferably to about 1100 F + 50 F
(595 C ~ 30 C). This first slow cooling stage
provides for the decomposition of austenite to
carbon-saturated ferrite. Under equilibrium
conditions, au~tenite decomposes to carbon-saturated
ferrite between from about 1650 F (900 C) and
1420 F (770 C). However, the kinetics of the
cooling process are such that austenite decomposition
does not begin in earnest until the mid 1500 F
(815 C) range and contînues somewhat below 1100 F
(595 C).
Failure to decompose the austenite in the first
cooling stage will result in the formation of
mart~nsite and/or pearlite. Martensite, i~ present,
will cause an enlargement of the secondary grain
1 size, and the deterioration of the quality of the
(110)[001~ orientation. Its presence adversely
affects energy storage in the second stage of cold
rolling, and results in poorer and more variable
magnetic quality of the final silicon steel product.
Lastly, martensite degrades the mechanical
properties, particularly the cold rolling
characteristics. Pearlite is more benign, but still
tieæ up carbon in an undesired form.
As indicated above, austenite decomposition
begins at about point "C" in the Figure and continues
to about point "E". At point "D" fine iron carbide
begins to precipitate from the carbon-saturated
ferrite. Under eguilibrium conditions, carbides
begin to precipitate rom carbon-saturated ferrite at
temperatures balow 1280 F (690 C). HoweYsr, the
actual process requires some undercooling to start
precipitation, which begins in earnest at about
1200 F (650 C). It wil:i be noted that the
austenite decomposition to carbon~rich ferrite and
carbide precipitation from the ferrite overlap
somewhat. The carbide is in two forms. It is
present as an intergranular film and as a fine
intragranular precipitate. The former precipitates
at temperatures above about 1060 F (570 C). The
latter precipitates below about 1060 F (570 C).
The slow cooling first stage, e~tending from point
"C" to point "E" of the Figure has a cooling rate of
less than 1500 F (835 C) per minute, and preferably
from about 500 F to abou~ 1050 F t280 C to 585 C)
per minute.
The second stage of the cooling cycle, a fast
cooling stage, begins at point "E" in the Figure and
- 12 -
1 extends to point "G" between 600 F and 1000 F
(315 C and 540 C) at which point the strip can be
water quenched to complete the rapid cooling stage.
The strip temperature after water quenching is 150 F
(65 C) or less, which is shown in the Figure as room
temperature (75 F or 25 C). During the second
cooling stage, the cooling rate is preferably from
about 2500 F to about 3500 F (1390 C to 1945 C)
per minute and more preferably greater than 3000 F
per minute (1665 C) per mlnute. This assures the
precipitation of fine iron carbide.
It will be evident from the above that the entire
intermediate anneal and cooling cycle oE the present
invention is required in the process of oh$aining the
desired microstructure, and precise controls are
critical. The prior art cycle time shown in the
Figure required at least 3 minutes, terminating in a
water bath, not shown, at a strip speed of about 220
feet pex minute (57 meters per minute)O The
intermediate anneal cycle time of the present
invention requires about 2 minutes, 10 seconds which
enabled a strip speed of about 260 feet per minute
(80 meters per minute) to be used. It will therefore
be noted that the annealing cycle of the present
invention enables greater productivity of the line.
No aging treatment after the anneal is either needed
or desired, since it has been found $o cause the`
formation of an enlarged secondary grain size which
degrades the magnetic quality of the final silicon
steel product.
The intermediate anneal is followed by the second
stage of cold rolling where the silicon steel is
reduced to the desired final gauge. The silicon
-- 13 --
2 ~ ,e 3
1 steel is thereafter decarburized, coated with an
annealing separator and subjected to a final anneal
to effect secondary recrystallization.
In the plant, two regular grain oriented silicon
steel heats having an aim silicon content of 3.15%,
were processed. The chemistries for these two heats
in weight percent are given in TABLE I below.
TA~LE I
C Mn S S; Al N ~_
A 0.02800.0592 0.02153.163 0.0016 0.0033 0.094
B 0.02880.0587 0.02163.175 0.0013 0.0029 0.083
The processing was without a hot band anneal and each
of the two heats were separated and processed to to
final gauges of 11 mils (0.28 mm), 9 mils (0-.23 mm)
and 7 mils (0.18 mm3 each using three different
intermediate gauges. The three intermediate gauges
for each of the 7, 9 and 11 mil (0.18 mm, 0.23 mm and
0.28 mm) materials are given in TABLE II below.
TABLE II
Final Intermediate Gau~e
Gauqe (inch)(mm~ _
7-mil (0.18 mm~ 0.019 O.g8
0.021 0.53
0.023 0.58
9-mil (0.23 mm) 0.021 0.53
0.023 0.58
0O025 0.63
ll-mil (O.28 mm) 0.022 0.56
0.024 0.61
0.026 0`.64
- 14 -
The standard prior art aim gauges for 7 mil (0.18
mm), 9 mil (0.23 mm) and 11 mil (0.28 mm) materials
were, respectively, 0.021 inch (0.53 mm), 0.023 inch
(0.58 mm), and 0.024 inch (0.61 n~). The silicon
steels were given an intermediate anneal and cooling
cycle according to the present invention. To this
end they were soaked for about 8 seconds at about
1680 F (915 C). Thereafter they were cooled to
about 1060 F (570 C~ at a rate of from about 850 F
to about 1200 F (from about 470C to about 670 C)
per minute. They were then cooled to about 600 F
(350 C) at a rate of about 1500~ F to about ~000 F
(about 830 C to about 1100 C) per minute, followed
by water quenching to less than 150 F (65 C). The
silicon steels were cold rolled to final gauge,
~ecarburized at 1525 F (830 C) in wet hydrogen
bear;ng atmosphere, magnesia coated, and given a
final box anneal at 2200 F (1200 C) for 24 hours in
wet hydrogen.
The coil front and back average results of both
heats A and B are summarized in TABLE III below.
O h~
_1 q`
:c ~
- o ~ o
3 o o O
_ ~_
e P~ ~ ,, 00 ~
E _~ ~o r-- r`
oo ~ ~ ~ ~
g _ O o o
_ o~
~ _~ ~ U~ ~o
a _ ~ ~
_ O o O
ei ~`1
1~ 3 o o
o o o
1~ CO ~
C ~ ~ ~
~1 u~ ~ ~ o
~1 ~ l ~,
o~ ~ ~ ~
C 000
_
*~ _~
_~
Z 01~
:9 ~ o o O
~a o ~ ~ ~1
O O O
1:~ ~ o o o
~ .
b ~1
3 u~
~ IQ . ~ ~
E _ OE~ ct~ oo
0: 31 t~
_ ~_ O o O
_ ~
I ~U. ~ ~ CO
r P ~ ~ u~ ~
C: o o
o o o
~-r .
H ~ O O O
1 Based upon prior art results, the aim 15 kGa core
loss values for the 7-mil (0.18 mm), 9-mil (0.23 mm)
and ll-mil (0.2B mm) material, respectively, were
.390 W/lb (0.867 W/Kg), .420 W/lb (0.933 W/Kg) and
.480 W/lb (1.067 W/Kg~. It will be noted that for
each of the 7, 9 and 11 mil (0.18 mm, 0~23 mm, and
0.28 mm) materials a slight core loss improvement was
achieved at the prior art intermediate gauges. Even
greatex improvement was achieved at heavier
intermediate gauyes. This clearly shows that the
optimum intermediate gauge has shifted upwardly with
the adoption of the intermediate anneal cycle of the
present invention. It will be noted that the H-10
permeability also improves at the heavier
intermediate gauges.
The present invention has thus far been described
in its application to partially austenitic grades o
regular grain oriented silicon steel. Fully ferritic
grades undergo no transformation from bcc type
crystal structure to fcc. This can be determined
from the ferrite stability index calculated as:
FSI = 2.54~40.53*(C+N)+0.43*(Mn+Ni~+0.22*Cu
-2.65*Al-3.95*P-1.26*(Cr~Mo)-Si
Compositions having a value equal to or less than
0.0 are fully ferritic. Increasing positive ferrite
stability inde~ values represent increasing volume
fractions of austenite will be present. For fully
ferritic compositions, rapid cooling can be initiated
directly at the end of the soak since there is no
austenite present, and thus a stage one slow cooling
is not required.
- 17 - ~ i3~
l Modifications may be made in the invention
wi~hout departing from the spirit of it.
