U.S. patent application number 13/389677 was filed with the patent office on 2012-06-14 for method for producing an iron-chromium alloy.
This patent application is currently assigned to THYSSENKRUPP VDM GMBH. Invention is credited to Heike Hattendorf, Osman Ibas.
Application Number | 20120145285 13/389677 |
Document ID | / |
Family ID | 43016657 |
Filed Date | 2012-06-14 |
United States Patent
Application |
20120145285 |
Kind Code |
A1 |
Hattendorf; Heike ; et
al. |
June 14, 2012 |
METHOD FOR PRODUCING AN IRON-CHROMIUM ALLOY
Abstract
The invention relates to a method for producing a component,
made of an iron-chromium alloy that precipitates Laves phases
and/or particles containing Fe and/or particles containing Cr
and/or particles containing Si and/or carbides, by subjecting a
semi-finished product made of the alloy to a thermomechanical
treatment, wherein in a first step, the alloy is solution heat
treated at temperatures.gtoreq.the solution heat treatment
temperature and is subsequently quenched in stationary protective
gas or air, moving (blown) protective gas or air, or water. In a
second step, a mechanical forming of the semi-finished product in a
range from 0.05 to 99% is performed, and in a subsequent step,
Laves phases Fe.sub.2(M, Si) or Fe.sub.7(M, Si).sub.6 and/or
particles containing Fe and/or particles containing Cr and/or
particles containing Si and/or carbides are precipitated in a
specific and finely distributed manner in that the component
produced from the formed semi-finished product is brought to an
application temperature between 550.degree. C. and 1000.degree. C.
by means of heating at 0.1.degree. C./min to 1000.degree.
C./min.
Inventors: |
Hattendorf; Heike; (Werdohl,
DE) ; Ibas; Osman; (Nachrodt, DE) |
Assignee: |
THYSSENKRUPP VDM GMBH
Werdohl
DE
|
Family ID: |
43016657 |
Appl. No.: |
13/389677 |
Filed: |
August 18, 2010 |
PCT Filed: |
August 18, 2010 |
PCT NO: |
PCT/DE2010/000975 |
371 Date: |
February 9, 2012 |
Current U.S.
Class: |
148/504 ;
148/325; 148/608; 148/714 |
Current CPC
Class: |
C21D 8/0226 20130101;
C21D 6/002 20130101; C22C 38/22 20130101; C22C 38/26 20130101; C21D
6/02 20130101; C21D 8/0236 20130101 |
Class at
Publication: |
148/504 ;
148/714; 148/608; 148/325 |
International
Class: |
C21D 11/00 20060101
C21D011/00; C22C 38/02 20060101 C22C038/02; C22C 38/38 20060101
C22C038/38; C22C 38/22 20060101 C22C038/22; C21D 8/00 20060101
C21D008/00; C22C 38/26 20060101 C22C038/26 |
Foreign Application Data
Date |
Code |
Application Number |
Sep 1, 2009 |
DE |
10 2009 039 552.0 |
Claims
1. Method for production of a component, from an iron-chromium
alloy precipitating Laves phases and/or Fe-containing particles
and/or Cr-containing particles and/or Si-containing particles
and/or carbides, in that a semifinished product produced from the
alloy is subjected to a thermomechanical treatment, wherein in a
first step the alloy is solution annealed at
temperatures.gtoreq.the solution-annealing temperature, followed by
quenching in stationary protective gas or air, moving (blown)
protective gas or air or in water, in a second step mechanical
working of the semifinished product in the range from 0.05 to 99%
is performed and in a subsequent step Fe.sub.2(M, Si) or
Fe.sub.7(M, Si).sub.6 Laves phases and/or Fe-containing particles
and/or Cr-containing particles and/or Si-containing particles
and/or carbides are precipitated purposefully and in finely
dispersed form, by the fact that the component made from the worked
semifinished product is brought to an application temperature
between 550.degree. C. and 1000.degree. C. by heating at
0.1.degree. C./min to 1000.degree. C./min.
2. Method for production of a component from an iron-chromium alloy
precipitating Laves phases and/or Fe-containing particles and/or
Cr-containing particles and/or Si-containing particles and/or
carbides, in that a semifinished product produced from the alloy is
subjected to a thermomechanical treatment, wherein in a first step
the alloy is solution annealed at temperatures.gtoreq.the
solution-annealing temperature, followed by quenching in stationary
protective gas or air, moving (blown) protective gas or air or in
water, in a second step mechanical working of the semifinished
product in the range from 0.05 to 99% is performed and in a
subsequent step Fe.sub.2(M, Si) or Fe.sub.7(M, Si).sub.6 Laves
phases and/or Fe-containing particles and/or Cr-containing
particles and/or Si-containing particles and/or carbides are
precipitated purposefully and in finely dispersed form by the fact
that the worked semifinished product is subjected for a time
between t.sub.min and t.sub.max to a heat treatment in the
temperature range between 550.degree. C. and 1060.degree. C. under
protective gas or air, followed by quenching in stationary
protective gas or air, moving (blown) protective gas or air or in
water or for heat treatments up to 800.degree. C. is quenched in
the oven, wherein and t.sub.min and t.sub.max are calculated
according to the following formulas:
t.sub.min=T.sub.a10.sup.(6740/Ta-9.216) and
t.sub.max=T.sub.a10.sup.(17960/Ta-15.72) where T.sub.a=T+273.15,
and wherein the desired component is made before or after this heat
treatment.
3. Method according to claim 1, wherein, in the first step, the
alloy is solution annealed at a temperature .gtoreq.1050.degree. C.
for longer than 6 minutes.
4. Method according to claim 1, wherein, in the first step, the
alloy is solution annealed at a temperature .gtoreq.1060.degree. C.
for longer than 1 minute.
5. Method according to claim 1, wherein semifinished product of the
following chemical composition (in % by weight) is
thermomechanically treated: Cr 12-30% Mn 0.001-2.5% Nb 0.1-2% W
0.1-5% Si 0.05-1% C 0.002-0.1% N 0.002-0.1% S max. 0.01% Fe
remainder as well as the usual melting-related impurities, wherein
a mechanical deformability at room temperature of >13% is
obtained, measured as plastic elongation in the tension test.
6. Method according to claim 1, wherein only little or even no
Fe.sub.2(M, Si) or Fe.sub.7(M, Si).sub.6 Laves phases and/or
Fe-containing particles and/or Cr-containing particles and/or
Si-containing particles and/or carbides are still present in the
semifinished product after solution annealing at temperatures the
solution-annealing temperature, preferably .gtoreq.1050.degree. C.
for longer than 6 minutes or .gtoreq.1060.degree. C. for longer
than 1 minute, followed by quenching in stationary protective gas
or air, moving (blown) protective gas or air or in water, in the
initial state before deformation.
7. Method according to claim 1, wherein the working of the
semifinished product takes place by hot working.
8. Method according to claim 1, wherein the hot working of the
semifinished product begins with a starting temperature
>1070.degree. C., wherein the last 0.05 to 90% of mechanical
deformation is applied between 1000.degree. C. and 500.degree.
C.
9. Method according to claim 1, wherein the hot working of the
semifinished product begins with a starting temperature
>1070.degree. C., wherein the last 0.05 to 95% of mechanical
deformation is applied between 1000.degree. C. and 500.degree.
C.
10. Method according to claim 1, wherein the hot working of the
semifinished product begins with a starting temperature
>1070.degree. C., wherein the last 0.05 to 99% of mechanical
deformation is applied between 1000.degree. C. and 500.degree.
C.
11. Method according to claim 1, wherein the hot working of the
semifinished product is followed by cold working.
12. Method according to claim 1, wherein the working of the
semifinished product is carried out by cold working.
13. Method according to claim 12, wherein the degree of cold
working of the semifinished product is 0.05 to 99%.
14. Method according to claim 12, wherein the cold working of the
semifinished product is 0.05 to 95%.
15. Method according to claim 12, wherein the cold working of the
semifinished product is 0.05 to 90%
16. Method according to claim 1, wherein the mechanical working of
the semifinished product is 20 to 99% and then the worked
semifinished product is subjected for a time between t.sub.min and
t.sub.max to a heat treatment in the temperature range between
950.degree. C. and 1060.degree. C. under protective gas or air,
followed by quenching in stationary protective gas or air, moving
(blown) protective gas or air or in water and after this the
desired component is made with
t.sub.min=T.sub.a10.sup.(6740/Ta-9.216) and
t.sub.max=T.sub.a10.sup.(17960/Ta-15.72) where T.sub.a=T+273.15 and
indication of t.sub.min and t.sub.max in minutes and of
heat-treatment temperature T in .degree. C.
17. Method according to claim 1, wherein the alloy additionally
contains (in % by weight) 0.02 to 0.3% La.
18. Method according to claim 1, wherein the alloy additionally
contains (in % by weight) 0.01 to 0.5% Ti.
19. Method according to claim 1, wherein the alloy additionally
contains 0.02 to 0.3% of one or more of the elements Ce, Pr, Ne,
Sc, Y, Zr or Hf.
20. Method according to claim 1, wherein the alloy additionally
contains (in % by weight) 0.001 to 0.5% Al.
21. Method according to claim 1, wherein the alloy additionally
contains (in % by weight) 2.0 to 6.0% Al.
22. Method according to claim 21, wherein the alloy additionally
contains (in % by weight) 2.5 to 5.0% Al.
23. Method according to claim 1, wherein the alloy additionally
contains one or more of the elements 0.0001 to 0.07% Mg, 0.0001 to
0.07% Ca, 0.002-0.03% P.
24. Method according to claim 1, wherein the alloy further contains
0.01 to 3.0% of one or more of the elements Ni, Co or Cu.
25. Method according to claim 1, wherein the alloy further contains
up to 0.005% B.
26. Method according to claim 1, wherein the iron-chromium alloy,
which is thermomechanically treated and which precipitates Laves
phases in finely dispersed form, has the following composition
containing (in % by weight) Cr 12-30% Mn 0.001-2.5% Nb 0.1-2% W
0.1-5% Si 0.05-1% C 0.002-0.03% N 0.002-0.03% S max. 0.005% Fe
remainder as well as the usual melting-related impurities.
27. Method according to claim 1, wherein the alloy additionally
contains (in % by weight) 0.02 to 0.2% of the element La.
28. Method according to claim 1, wherein the alloy additionally
contains (in % by weight) 0.02 to 0.2% Ti.
29. Method according to claim 1, wherein the alloy additionally
contains (in % by weight) 0.02 to 0.2% of one or more of the
elements Ce, Pr, Ne, Sc, Y, Zr or Hf.
30. Method according to claim 1, wherein the alloy additionally
contains (in % by weight) one or more of the elements 0.0001-0.05%
Mg, 0.0001-0.03% Ca, 0.002-0.03% P.
31. Method according to claim 1, wherein the alloy further contains
(in % by weight) up to 0.003% B.
32. Method according to claim 1, wherein (in % by weight) the Nb
content is 0.3 to 1.0% and the Si content is 0.15-0.5%.
33. Method according to claim 1, wherein the W content is replaced
entirely or partly by at least one of the elements Mo and/or
Ta.
34. Method according to claim 1, wherein the alloy contains (in %
by weight) max. 0.2% V and/or max. 0.005% S.
35. Method according to claim 1, wherein the alloy contains (in %
by weight) max. 0.01% O.
36. Method according to claim 1, wherein the alloy contains (in %
by weight) max. 0.01% of each of the elements Zn, Sn, Pb, Se, Te,
Bi and Sb respectively.
37. Method according to claim 1, wherein the semifinished product
is formed by sheet, strip, bar, forging, pipe or wire.
38. Method according to claim 1, wherein the heat treatment is
carried out only after finishing of the component.
39. Method according to claim 1, wherein, by the thermomechanical
treatment of the semifinished product, a particularly high creep
strength is produced in the semifinished product and/or in the
component with simultaneous elongation >13% in the tension test
at room temperature.
40. Metallic component or semifinished product, consisting of the
following chemical composition (in % by weight) Cr 12-30% Mn
0.001-2.5% Nb 0.1-2% W 0.1-5% Si 0.05-1% C 0.002-0.1% N 0.002-0.1%
S max. 0.01% Fe remainder as well as the usual melting-related
impurities, which at the end of a thermomechanical treatment has a
deformed microstructure, to the effect that Laves phase(s) is or
are embedded in finely dispersed form in the microstructural
dislocations of the microstructure, wherein, in a creep test with,
in particular, 35 MPa at 750.degree. C. and an elongation of at
least 18%, a time to break that exceeds the time to break of a
coarse-grained, completely recrystallized microstructure by a
factor of at least 1.5 is established in the microstructure.
41. Metallic component or semifinished product, consisting of the
following chemical composition (in % by weight) Cr 12-30% Mn
0.001-2.5% Nb 0.1-2% W 0.1-5% Si 0.05-1% C 0.002-0.1% N 0.002-0.1%
S max. 0.01% Fe remainder as well as the usual melting-related
impurities, which at the end of a thermomechanical treatment has a
deformed microstructure, to the effect that Laves phase(s) is or
are embedded in finely dispersed form in the microstructural
dislocations of the microstructure, wherein, in a creep test with,
in particular, 35 MPa at 750.degree. C. and an elongation of at
least 18%, a time to break hours that exceeds the time to break of
a coarse-grained, completely recrystallized microstructure by a
factor of at least 3 is established in the microstructure.
42. Use of a component produced according to claim 1 as
interconnector in a fuel cell.
43. Use of a component produced according to claim 1 as material in
a component, such as a reformer or a heat exchanger or in an
ancillary aggregate of a fuel cell.
44. Use of a component produced according to claim 1 in the
exhaust-gas line of a combustion engine.
45. Use of a component produced according to claim 1 for steam
boilers, superheaters, turbines and other parts of a power plant or
in the chemical process industry.
Description
[0001] The invention relates to a ferritic iron-chromium alloy
produced by melting metallurgy.
[0002] DE 100 25 108 A1 discloses a high-temperature material
comprising an iron alloy forming chromium oxide with up to 2% by
weight of at least one oxygen-affine element from the group Y, Ce,
Zr, Hf and Al, up to 2% by weight of an element M from the group
Mn, Ni and Co, which forms a spinel phase of MCr.sub.2O.sub.4 type
at high temperatures, up to 2% by weight of a further element from
the group Ti, Hf, Sr, Ca and Zr, which increases the electrical
conductivity of Cr-based oxides. The chromium content should lie
within a concentration range between 12 and 28%. Areas of use for
this high-temperature material are bipolar plates in a
high-temperature fuel cell.
[0003] EP 1 298 228 A1 relates to a steel for a high-temperature
fuel cell that has the following composition: not more than 0.2% C,
not more than 1% Si, not more than 1% Mn, not more than 2% Ni,
15-30% Cr, not more than 1% Al, not more than 0.5% Y, not more than
0.2% REM and not more than 1% Zr, the remainder being iron and
production-related impurities.
[0004] Low hot strength and inadequate creep strength at
temperatures of 700.degree. C. and above are common to these two
alloys. Precisely in the range above 700.degree. C. up to
approximately 900.degree. C., however, these alloys have excellent
oxidation and corrosion resistance.
[0005] A creep-resistant ferritic steel, comprising precipitates of
an intermetallic phase of Fe.sub.2(M, Si) or Fe.sub.7(M, Si).sub.6
type with at least one metallic alloying element M, which may be
formed by the elements niobium, molybdenum, tungsten or tantalum,
has become known from DE 10 2006 007 598 A1. The steel is
preferably intended to be used for a bipolar plate in a fuel-cell
stack.
[0006] EP 1 536 031 A1 discloses a metallic material for fuel
cells, containing C.ltoreq.0.2%, 0.02 to 1% Si, .ltoreq.2% Mn, 10
to 40% Cr, 0.03 to 5% Mo, 0.1 to 3% Nb, at least one of the
elements from the group Sc, Y, La, Ce, Pr, Nd, Pm, Sn, Zr and
Hf.ltoreq.1, the remainder being iron and unavoidable impurities,
wherein the composition is supposed to satisfy the following
equation: 0.1.ltoreq.Mo/Nb.ltoreq.30.
[0007] EP 1 882 756 A1 describes a ferritic chromium steel,
especially usable in fuel cells. The chromium steel has the
following composition: C max. 0.1%, Si 0.1-1%, Mn max. 0.6%, Cr
15-25%, Ni max. 2%, Mo 0.5-2%, Nb 0.2-1.5%, Ti max. 0.5%, Zr max.
0.5%, REM max. 0.3%, Al max. 0.1%, N max. 0.07%, the remainder
being Fe and melting-related impurities, wherein the content of
Zr+Ti is at least 0.2%.
[0008] In comparison with DE 100 25 108 A1 and EP 1 298 228 A2, all
of these alloys have better hot strength and elevated creep
strength at temperatures of 700.degree. C. and above, specifically
due to formation of precipitates, which hinder the dislocation
movements and thus plastic deformation of the material. In the case
of DE 10 2006 007 598 A1, for example, these precipitates consist
of a Laves phase, an intermetallic compound with the composition
Fe.sub.2(M, Si) or Fe.sub.7(M, Si).sub.6, wherein M may be niobium,
molybdenum, tungsten or tantalum. Therein a proportion by volume of
1 to 8%, preferably 2.5 to 5%, should be reached. However, there
may also be other precipitates such as Fe-containing particles
and/or Cr-containing particles and/or Si-containing particles, such
as described, for example, in EP 1 536 031 A1, or there may be
carbides containing Nb, W, Mo. It is common to all of these
particles that they make deformation of the material difficult.
[0009] From the state of the art described above, it is known that
small additions of Y, Zr, Ti, Hf, Ce, La and similar reactive
elements can influence the oxidation resistance of Fe--Cr alloys
very positively.
[0010] The alloys cited in DE 10 2006 007 598 A1, EP 1 536 031 A1
and EP 1 882 756 A1 are optimized for the application as
interconnector plates for the high-temperature fuel cells: By use
of a ferritic alloy containing 10 to 40% chromium, they have an
expansion coefficient adapted as well as possible to the ceramic
components anode and electrolyte.
[0011] Further requirements on the interconnector steel of a
high-temperature fuel cell are, besides the creep strength already
mentioned above, very good corrosion resistance, good conductivity
of the oxide layer and little chromium volatilization.
[0012] The requirements on the reformers and the heat exchangers
for the high-temperature fuel cell are the best possible creep
strength, very good corrosion resistance and little chromium
volatilization. The oxide for these components does not have to be
conductive.
[0013] The requirements for components for the exhaust-gas line of
a combustion engine, for example, or for steam boilers,
superheaters, turbines and other parts of a power plant, are best
possible creep strength and very good corrosion resistance. In
these cases chromium volatilization does not cause any poisoning
phenomena as in the fuel cell, and the protecting oxide does not
have to be conductive for such components.
[0014] In DE 10 2006 007 598 A1, for example, the excellent
corrosion resistance is achieved by formation of a chromium oxide
top layer. By the fact that a spinel containing Mn, Ni, Co or Cu is
additionally formed on the chromium oxide top layer, fewer volatile
chromium oxides or chromium oxyhydroxides that poison the cathode
are formed. By the fact that Si is stably bound in the Fe.sub.2(M,
Si) or Fe.sub.7(M, Si).sub.6 Laves phase, a nonconductive
subsurface layer of silicon oxide is also not formed under the
chromium oxide top layer. The corrosion resistance is further
improved by the fact that the Al content is kept low and so the
increase of the corrosion due to the internal oxidation of the
aluminum is avoided. A small Ti addition additionally favors
strengthening of the surface and thus prevents swelling of the
oxide layer and the inclusion of metallic zones in the oxide layer,
which increases the oxidation. In addition, the addition of
oxygen-affine elements such as La, Ce, Y, Zr or the like further
increases the corrosion resistance.
[0015] From the market, increased requirements are being imposed on
products, necessitating elevated hot strength and creep strength
together with an elongation of at least 18% at application
temperature for avoidance of brittle failure together with at least
equally good oxidation or corrosion resistance and a higher service
temperature of the alloy, specifically while retaining acceptable
deformability, measured as plastic deformation in the tension test
with an elongation of >13% at room temperature.
[0016] Furthermore, the following investigation methods are
used.
[0017] In a creep test, a specimen is subjected to a constant
static tensile force at a constant temperature. For the purpose of
comparability, this tensile force is expressed as an initial
tensile stress relative to the initial cross section of the
specimen. In the creep test, the time t.sub.B until break--the time
to break--of the specimen is measured in the simplest case. The
test can then be performed without measurement of the elongation of
the specimen in the course of the test. The elongation at break is
then measured after the end of the test.
[0018] The specimen is mounted at room temperature in the
creep-testing machine and heated to the desired temperature without
loading by a tensile force. After reaching the test temperature,
the specimen is maintained for one hour without loading for
temperature equilibration. Thereafter the specimen is loaded with
the tensile force and the test time begins.
[0019] The time to break can be taken as a measure of the creep
strength. The longer the time to break is at a specified
temperature and initial tensile stress, the greater the creep
strength of the material is. The time to break and the creep
strength decrease with increasing temperature and increasing
initial tensile stress (see, for example, "Burgel", page 100).
[0020] The deformability is determined in a tension test according
to DIN 50145 at room temperature. In the process, the offset yield
strength R.sub.p0.2, the tensile strength R.sub.M and the
elongation at break are determined. The elongation A is determined
on the broken specimen from the elongation of the original gauge
length L.sub.0:
A=(L.sub.u-L.sub.0)/L.sub.0100%=.DELTA.L/L.sub.0100%
Where L.sub.u=gauge length after break.
[0021] Depending on gauge length, the elongation at break is
denoted by subscripts:
[0022] A.sub.5, gauge length L.sub.0-5d.sub.0 or L.sub.0=5.65
S.sub.0
[0023] A.sub.10, gauge length L.sub.0=10d.sub.0 or
L.sub.0=11.3S.sub.0
or, for example, A.sub.L=100, for the freely chosen gauge length
L=100 mm. (d.sub.0 initial diameter, S.sub.0 initial cross section
of the flat specimen)
[0024] The magnitude of the elongation A in the tension test at
room temperature can be taken as a measure of the
deformability.
[0025] The Laves phase(s) or the Fe-containing particles and/or
Cr-containing particles and/or Si-containing particles and/or
carbides can be made visible on a metallographic ground section by
etching with V2A pickling fluid or electrolytic etching with oxalic
acid. During etching with V2A pickling fluid, the grains or grain
boundaries also are additionally etched visibly. Only particles
with a size of approximately 0.5 .mu.m and larger are visible by
viewing in an optical microscope. Smaller particles may not be
recognized, but are definitely present. Therefore metallography is
used only in support of the explanation, while the efficacy of a
measure is assessed more practically by the time to rupture or
creep strength.
[0026] In the Manual of High-Temperature Materials Technology, Ralf
Burgel, 3rd Revised Edition, Viehweg Verlag, December 2006,
hereinafter referred to as "Burgel", the "Possible measures for
increasing the creep strength of metallic materials" are presented
on pages 196 to 199 and in Table 3.7.
[0027] The measures [0028] "High melting point, face-centered cubic
material", [0029] "High modulus of elasticity", [0030] "Material
with low stacking fault energy", cannot be used for improvement of
the cited parameters, since they necessitate a change of material
type, which is not possible here and also is not the task.
[0031] The measures [0032] "Solid solution hardening" [0033]
"Particle hardening" [0034] "High particle volume fraction" [0035]
"Particles with small diffusion coefficients of the alloying
element in question" have already been employed in DE 10 2006 007
598 A1 and/or in EP 1 536 031 A1 and/or in EP 1 882 756 A1.
[0036] The measures [0037] "Particles with low solubility in the
matrix" [0038] "Coherent particles with low interfacial enthalpy
relative to the matrix", are not applicable for the precipitates
under consideration.
[0039] Likewise, the measures [0040] "Carbides or borides as
grain-boundary precipitates; avoid oxides and sulfides", [0041]
"Add positively active grain-boundary elements in precisely
controlled dosage, for example, B, C, Zr, Ce", [0042] "Higher
purity of the alloy" [0043] "Add getter elements (for example, for
S)", [0044] "High corrosion resistance" have already been described
in DE 10 2006 007 598 A1 and/or in EP 1 536 031 A1 and/or in EP 1
882 756 A1.
[0045] The measures [0046] "Small dendrite arm spacings in cast
microstructures", [0047] "Small grain structure in the main loading
direction", [0048] "Single crystal" [0049] "Low density of
components loaded by their own weight and rotating" cannot be
applied to this alloy type, or to the production route, or the
use.
[0050] For the task of improving the creep strength of the
precipitation-hardened iron-chromium alloy, the measures
1) "Coarse grain microstructure", 2) "Jaggedness of the grain
boundaries due to precipitates", 3) "Optimized heat treatment
(adjust optimum particle diameter, eliminate segregations in cast
microstructure, purposefully adjust possible grain boundary
roughness)", 4) "Avoid cold working", are to be considered.
[0051] The task of the invention is to provide a method for
production of a component made from a precipitation-hardened
iron-chromium alloy, by means of which the high hot strength or
creep strength of a precipitation-hardened ferritic alloy can be
further increased compared with the state of the art while
retaining acceptable deformability at room temperature.
[0052] It is also intended to provide a thermomechanically treated
component/semifinished product consisting of an iron-chromium
alloy, which can be used for achievement of high hot strength or
creep strength while retaining acceptable deformability at room
temperature.
[0053] Finally, it is intended that the component/semifinished
product produced in this way can be used for specific technical
applications in the temperature range above 550.degree. C.
[0054] This task is accomplished on the one hand by a method for
production of a component from an iron-chromium alloy precipitating
Laves phases and/or Fe-containing particles and/or Cr-containing
particles and/or Si-containing particles and/or carbides, in that a
semifinished product produced from the alloy is subjected to a
thermomechanical treatment, wherein in a first step the alloy is
solution annealed at temperatures.gtoreq.the solution-annealing
temperature, followed by quenching in stationary protective gas or
air, moving (blown) protective gas or air or in water, in a second
step mechanical working of the semifinished product in the range
from 0.05 to 99% is performed and in a subsequent step Fe.sub.2(M,
Si) or Fe.sub.7(M, Si).sub.6 Laves phases and/or Fe-containing
particles and/or Cr-containing particles and/or Si-containing
particles and/or carbides are precipitated purposefully and in
finely dispersed form by the fact that the component made from the
worked semifinished product is brought to an application
temperature between 550.degree. C. and 1000.degree. C. by heating
at 0.1.degree. C./min to 1000.degree. C./min.
[0055] This task is accomplished on the other hand by a method for
production of a component from an iron-chromium alloy precipitating
Laves phases and/or Fe-containing particles and/or Cr-containing
particles and/or Si-containing particles and/or carbides, in that a
semifinished product produced from the alloy is subjected to a
thermomechanical treatment, wherein in a first step the alloy is
solution annealed at temperatures the solution-annealing
temperature, followed by quenching in stationary protective gas or
air, moving (blown) protective gas or air or in water, in a second
step mechanical working of the semifinished product in the range
from 0.05 to 99% is performed and in a subsequent step Fe.sub.2(M,
Si) or Fe.sub.7(M, Si).sub.6 Laves phases and/or Fe-containing
particles and/or Cr-containing particles and/or Si-containing
particles and/or carbides are precipitated purposefully and in
finely dispersed form by the fact that the worked semifinished
product is subjected for a time between t.sub.aw, and t.sub.max to
a heat treatment in the temperature range between 550.degree. C.
and 1060.degree. C. under protective gas or air, followed by
quenching in stationary protective gas or air, moving (blown)
protective gas or air or in water or for heat treatments up to
800.degree. C. is quenched in the oven, wherein t.sub.min and
t.sub.max are calculated according to the following formulas:
t.sub.min=T.sub.a10.sup.(6740/Ta-9.216) and
t.sub.max=T.sub.a10.sup.(17960/Ta-15.72) where
T.sub.a=T+273.15,
and wherein the desired component is made before or after this heat
treatment.
[0056] The times t.sub.min and t.sub.max are expressed in minutes
and the heat treatment temperature T in .degree. C.
[0057] Advantageous further developments of the inventive method
are to be inferred from the associated dependent claims pertaining
to the method.
[0058] For the first step, the following temperature ranges and
times are practical for solution annealing:
>1050.degree. C. for longer than 6 minutes >1060.degree. C.
for longer than 1 minute
[0059] The resulting changes of the material characteristics are
explained in more detail in the course of the further
description.
[0060] Furthermore, the task is also accomplished by a metallic
component or semifinished product consisting of the following
chemical composition (in % by weight)
Cr 12-30%
Mn 0.001-2.5%
Nb 0.1-2%
W 0.1-5%
Si 0.05-1%
C 0.002-0.1%
N 0.002-0.1%
S max. 0.01%
[0061] Fe remainder as well as the usual melting-related
impurities, which at the end of a thermomechanical treatment has a
deformed microstructure, to the effect that Laves phase(s) is or
are embedded in finely dispersed form in the microstructural
dislocations of the microstructure, wherein, in a creep test with,
for example, 35 MPa at 750.degree. C. and an elongation of at least
18%, a time to break that exceeds the time to break of a
coarse-grained, completely recrystallized microstructure by a
factor of at least 1.5 is established in the microstructure.
[0062] A comparable result is achieved for creep tests with
different stresses and temperatures, wherein the temperatures for
the creep test preferably lie in the range between 500 and
1000.degree. C.
[0063] Measures 1 to 4 described above will now be considered.
[0064] Surprisingly it has been found in this connection that, in
contrast to measure 4 "cold deformation", preworking followed by an
adapted annealing treatment can bring about prolongations of the
times to break of the specimen in the creep test that go more than
1.5 times, preferably more than 3 times beyond the times to break
for a coarse-grained microstructure (measure 1).
[0065] Furthermore, it is proposed that, for the third step--the
precipitation of the Laves phase(s)--the worked semifinished
product or if applicable the component made therefrom, by a
combination of heating at 0.1.degree. C./min to 1000.degree. C./min
to a heat-treatment temperature between 550.degree. C. and
1060.degree. C. with subsequent heat treatment for a time between
t.sub.min and t.sub.max at this temperature under protective gas or
air, followed by quenching in stationary protective gas or air,
moving (blown) protective gas or air or in water or for heat
treatments up to 800.degree. C. is quenched in the oven, after
which the desired component is made if applicable, wherein
t.sub.min and t.sub.max are calculated according to the following
formulas:
t.sub.min=T.sub.a10.sup.(6740/Ta-9.216) and
t.sub.max=T.sub.a10.sup.(17960/Ta-15.72) where
T.sub.a=T+273.15.
[0066] In addition, the possibility exists that, for the third
step--the precipitation of the Laves phase(s)--the worked
semifinished product or the component made therefrom is subjected
for a time between t.sub.min and t.sub.max to a heat treatment in
the temperature range between 550.degree. C. and 1060.degree. C.
under protective gas or air, followed by quenching in stationary
protective gas or air, moving (blown) protective gas or air or in
water or for heat treatments up to 800.degree. C. is quenched in
the oven, after which the desired component is made if applicable,
wherein t.sub.min and t.sub.max are calculated according to the
following formulas: t.sub.min=T.sub.a10.sup.(6740/Ta-9.216) and
t.sub.max=T.sub.a10.sup.(17960/Ta-15.72) where T.sub.a=T+273.15 and
then the component that has been made is brought by heating at
0.1.degree. C./min to 1000.degree. C./min to an application
temperature between 550.degree. C. and 1000.degree. C.
[0067] According to a further concept of the invention, a
semifinished product from an alloy of the following composition (in
% by weight) is treated thermomechanically:
Cr 12 to 30%
Mn 0.001 to 2.5%
Nb 0.1 to 2%
W 0.1 to 5%
S 0.05 to 1%
C 0.002 to 0.03%
N 0.002 to 0.03%
S max. 0.01%
[0068] Fe remainder as well as the usual melting-related
impurities.
[0069] With the inventive method it is possible to produce
semifinished products in the form of sheets, strips, bars,
forgings, pipes or wire and to make components in the most diverse
forms needed for the respective application.
[0070] It is of special advantage that only little or even no
Fe.sub.2(M, Si) or Fe.sub.7(M, Si).sub.6 Laves phases and/or
Fe-containing particles and/or Cr-containing particles and/or
Si-containing particles and/or carbides are still present in the
semifinished product after solution annealing at
temperatures.gtoreq.the solution-annealing temperature, preferably
.gtoreq.1050.degree. C. for longer than 6 minutes or
>1060.degree. C. for longer than 1 minute, followed by quenching
in stationary protective gas or air, moving (blown) protective gas
or air or in water, in the initial state before deformation.
[0071] The working of the semifinished product can take place by
hot working. Alternatively, however, the forming can also be
brought about by cold working.
[0072] In the first case, the semifinished product is hot-worked
with a starting temperature .gtoreq.1070.degree. C., wherein the
last 0.05 to 95% of mechanical deformation is applied between
1000.degree. and 500.degree. C., advantageously the last 0.5 to 90%
between 1000.degree. C. and 500.degree. C.
[0073] In the second case, the degree of cold working of the
semifinished product is 0.05 to 99%, advantageously 0.05 to 95% or
0.05 to 90%.
[0074] As a further concept according to the invention, it is
proposed that the mechanical working of the semifinished product be
20 to 99% and then the worked semifinished product be subjected for
a time between t.sub.min and t.sub.max to a heat treatment in the
temperature range between 950.degree. C. and 1060.degree. C. under
protective gas or air, followed by quenching in stationary
protective gas or air, moving (blown) protective gas or air or in
water, after which the desired component be made, wherein t.sub.min
and t.sub.max are calculated according to the following
formulas:
t.sub.minT.sub.a10.sup.(6740/Ta-9.216) and
t.sub.max=T.sub.a10.sup.(17960/Ta-15.72) where T.sub.a=T+273.15
with t.sub.min and t.sub.max in minutes and the heat-treatment
temperature T in degrees Celsius. If the already indicated alloy is
used as an interconnector for a solid oxide fuel cell, then a
content of 0.001-0.5% aluminum is advantageous.
[0075] For other areas of use, such as, for example, in the
reformer or heat exchanger for the fuel cell, for which no
conductive oxide layer is necessary, a content of 2 to 6% aluminum
is advantageous, since then a closed aluminum oxide layer can form,
which once again has a much slower growth rate compared with a
chromium oxide layer and additionally has much less chromium oxide
volatilization than a chromium-manganese spinel.
[0076] For the areas of use that neither need a conducive oxide nor
have special requirements on chromium volatilization, both variants
may be considered. In this connection, it is to be kept in mind in
particular that the processability and weldability of the alloy
deteriorate with increasing aluminum content and so higher costs
are incurred. Therefore, when an oxide layer consisting of a
chromium oxide and a chromium-manganese spinel, adequate oxidation
resistance can be assured by use of 0.001-0.5% aluminum. If greater
oxidation resistance is necessary, as is assured, for example, by
the formation of an aluminum oxide layer, a content of 2.0-6.0%
aluminum is advantageous. These two alloy variants can be used, for
example, as components for the exhaust-gas line of a combustion
engine or for steam boilers, superheaters, turbines and other parts
of a power plant.
[0077] A preferred aluminum range is in particular the range from
2.5% to 5.0%, which is still characterized by good
processability.
[0078] In the already indicated alloy, the following elements may
be additionally used individually or in combination:
La 0.02 to 0.3%
Ti 0.01 to 0.5%
Mg 0.0001 to 0.07%
Ca 0.0001 to 0.07%
P 0.002 to 0.03%
Ni/Co/Cu 0.01 to 3%
B up to 0.005%.
[0079] The contents of the elements that can be additionally
introduced in the alloy may be adjusted as follows: Mg 0.0001 to
0.05%, Ca 0.0001 to 0.03%, P 0.002 to 0.03%.
[0080] Furthermore, the alloy (in % by weight) may contain one or
more of the elements Ce, La, Pr, Ne, Sc, Y, Zr or Hf in contents of
0.02-0.3%.
[0081] If necessary, the alloy (in % by weight) may contain one or
more of the elements Ce, Pr, Ne, Sc, Y, Zr or Hf in contents of
0.02-0.2%.
[0082] To achieve the desired effects, the Nb content is 0.3 to
1.0% and the Si content 0.15 to 0.5%.
[0083] If necessary, the element tungsten may be replaced entirely
or partly by at least one of the elements Mo or Ta.
[0084] If necessary, the alloy may also even contain max. 0.2% V
and/or max. 0.005% S. In this case the oxygen content should not be
greater than 0.01%.
[0085] If necessary, the alloy may also even contain max. 0.003%
boron.
[0086] Furthermore, the alloy should have a maximum of 0.01% of the
following elements respectively: Zn, Sn, Pb, Se, Te, Bi, Sb.
[0087] Components/semifinished products that on the one hand
consist of the cited alloy composition and on the other hand have
been produced by the inventive method may preferably be used as
interconnector in a fuel cell or as material in a component, such
as a reformer or a heat exchanger in an ancillary aggregate of the
fuel cell.
[0088] Alternatively, the possibility also exists of using the
component/semifinished product produced according to the inventive
method or the alloy itself as a structural element in the
exhaust-gas line of a combustion engine or for steam boilers,
superheaters, turbines and other parts of a power plant or in the
chemical process industry.
[0089] By means of the inventive method, Laves phases, by virtue of
the thermomechanical treatment, can be precipitated purposefully
and in fine dispersion at the dislocations of the microstructure in
alloys produced by melting metallurgy.
[0090] The details and the advantages of the invention will be
explained in more detail in the following examples.
[0091] In the following, the inventive method steps will be
subjected to closer examination.
[0092] The first step for the thermomechanical treatment of an
iron-chromium alloy precipitating Laves phases and/or Fe-containing
particles and/or Cr-containing particles and/or Si-containing
particles and/or carbides must be annealing above the solution
annealing temperature, so that the Laves phases and/or
Fe-containing particles and/or Cr-containing particles and/or
Si-containing particles and/or carbides are dissolved and are
available for precipitation in the subsequent thermomechanical
treatment. The solution annealing temperature is alloy-specific,
but preferably lies above 1050.degree. C. for a period of longer
than 6 minutes or above 1060.degree. C. for longer than 1 minute,
followed by quenching in stationary protective gas or air, moving
(blown) protective gas or air or in water. The exact temperature
control above this solution annealing temperature is not
determining for the characteristics. The annealing may be carried
out in air or under protective gas. It should lie under the melting
temperature, preferably <1350.degree. C. Also, for cost reasons,
the annealing times should preferably be <24 hours, but may also
be longer depending on performance. The solution annealing follows
quenching in stationary protective gas or air, moving (blown)
protective gas or air or in water, during which only little Laves
phase is newly formed.
[0093] In addition, care is to be taken that, especially for
thicker-walled components, all parts of the component reach the
required minimum annealing time at the specified temperature. This
is to be considered in the determination of the starting point of
the annealing time.
[0094] In a second step, an elevated dislocation density must be
introduced into the material. Elevated dislocation densities have
worked microstructure or recovered microstructure, wherein the
dislocations there are arranged at small-angle grain
boundaries.
[0095] The second step must therefore be working, so that the
dislocations are introduced into the material, which then, in the
subsequent annealing treatment, ensure a homogeneous dispersion of
the Laves phases and/or Fe-containing particles and/or
Cr-containing particles and/or Si-containing particles and/or
carbides.
[0096] This deformation may be cold working, but also hot working,
wherein it must be ensured during hot working that the
microstructure is not already recrystallized during rolling. This
is achieved by restricting the deformation range for the last
working and the temperature at which this is carried out. For
deformations above 1000.degree. C., the material already tends to
recrystallization or recovery during working, so that the working
must preferably be carried out below 1000.degree. C. At
temperatures below 500.degree. C., exist in the range of the
embrittlement that occurs in ferrites at 475.degree. C. There this
has a smaller elongation and an elevated working resistance, which
makes working less advantageous and reduces the economic
benefit.
[0097] Precipitates smaller than a certain size are less effective
(see, for example, "Burgel", page 141). Therefore the dislocation
density generated by the deformation should not be too high, since
then very many precipitates are indeed formed but are too fine, and
the excess dislocations can move freely and in this way the
preworking becomes harmful. This means that preferably the greatest
deformation is 90% for the part of the hot working
.ltoreq.1000.degree. C. and 90% for the cold working.
[0098] During working in the range of 20 to 99%, annealing between
950.degree. C. and 1050.degree. C. may cause recovery of the
microstructure. Thereby the dislocation density is reduced, so that
the positive effect on the dispersion of the Laves phases and/or
Fe-containing particles and/or Cr-containing particles and/or
Si-containing particles and/or carbides is established again.
[0099] The one possibility of introducing the Laves phases and/or
Fe-containing particles and/or Cr-containing particles and/or
Si-containing particles and/or carbides into the worked material is
to make the needed components from the semifinished product and
then to bring the component that has been made to the application
temperature between 550.degree. C. and 1000.degree. C. by heating
at 0.1.degree. C./min to 1000.degree. C./min. During heating, the
Laves phases and/or Fe-containing particles and/or Cr-containing
particles and/or Si-containing particles and/or carbides are
precipitated as a fine dispersion in the microstructure. The fine
dispersion is generated by nucleation in the lower temperature
range, followed by some growth of the nuclei at the higher
temperatures. Therefore the heating rate should not be slower than
1000.degree. C./min, because otherwise the time for this process is
too short. Heating rates slower than 0.1.degree. C./min are
uneconomical.
[0100] A second possibility is a separate heat treatment of the
material. For this purpose, the worked semifinished
product/component is subjected for a time between t.sub.min and
t.sub.max to a heat treatment in the temperature range between
550.degree. C. and 1060.degree. C. under protective gas or air,
followed by quenching in stationary protective gas or air, moving
(blown) protective gas or air or in water or for heat treatments up
to 800.degree. C. is quenched in the oven, wherein
t.sub.min=T.sub.a10.sup.(6740/Ta-9.216) and
t.sub.max=T.sub.a10.sup.(17960/Ta-15.72) where
T.sub.a=T+273.15,
with indication of t.sub.min and t.sub.max in minutes and heat
treatment temperature T in .degree. C. In this connection, the
desired component can be made before or after this heat
treatment.
[0101] In the annealing steps, care is to be taken that, especially
for thicker-walled semifinished products/components, all parts of
the component reach the required minimum annealing time at the
specified temperature. This is to be considered in determination of
the starting point of the annealing time. Likewise, care is to be
taken that no region of the semifinished product/component exceeds
the required maximum annealing time.
[0102] Times shorter than t.sub.min are not sufficient for
formation of the Laves phases and/or Fe-containing particles and/or
Cr-containing particles and/or Si-containing particles and/or
carbides. For times longer than t.sub.max, the danger exists of too
great coarseness of the precipitates, whereby the particles can no
longer contribute markedly to the creep strength. For times longer
than t.sub.max, the possibility exists in the upper temperature
range of 550.degree. C. and 1060.degree. C. that a recovered
microstructure will be formed, which certainly can still be
effective. However, with increasing recovery, the dislocation
density is further reduced, so that the dispersion of the
precipitates becomes increasingly inhomogeneous and the positive
effect on the creep strength ultimately vanishes. In addition, at
the lower temperatures in the range of 550.degree. C. and
1060.degree. C., times longer than t.sub.max are uneconomical.
[0103] The annealing step may be carried out under protective gas
(argon, hydrogen and similar atmospheres with reduced oxygen
partial pressure). For economic reasons, the quenching step is
carried out in stationary protective gas or air, moving (blown)
protective gas or air or in water, while oven quenching should be
avoided in particular for temperatures above 800.degree. C. but is
also possible at temperatures <800.degree. C.
[0104] The oxidation resistance and the thermal expansion
coefficient of the material are determined via the chromium
content. The oxidation resistance of the material is based on the
formation of a closed chromium oxide layer. Below 12%,
iron-containing oxides, which impair the oxidation resistance, are
formed increasingly, especially at higher operating temperatures.
The chromium content is therefore adjusted to >12%. Above 30%
chromium, the processability of the material and its usability are
impaired by increased formation of embrittling phases, especially
the sigma phase. The chromium content is therefore limited to
.ltoreq.30%. The expansion coefficient decreases with increasing
chromium content.
[0105] Especially for use in a fuel cell, therefore, the expansion
coefficient can be adjusted in a range that matches the ceramics in
the fuel cell. These are chromium contents around 22 to 23%. For
other applications, however, for example for reformers or in power
plants, this restriction does not exist.
[0106] The addition of manganese brings about formation of a
chromium-manganese spinel on the chromium oxide layer, which is
formed on the material at low aluminum contents below 2%. This
chromium-manganese spinel reduces the chromium volatilization and
improves the contact resistance. A manganese content of at least
0.001% is necessary for this. More than 2.5% manganese impairs the
oxidation resistance by formation of a very thick
chromium-manganese spinel layer.
[0107] Niobium, molybdenum, tungsten or tantalum can participate in
the formation of precipitates in iron-containing alloys, such as,
for example, carbides and/or the M in the Fe.sub.2(M, Si) or
Fe.sub.7(M, Si).sub.6 Laves phases. Molybdenum, tungsten or
tantalum are also good solid-solution hardeners and thus contribute
to the improvement of the creep strength. In this connection the
lower limit is determined in each case by the fact that a certain
content must be present in order to be effective, while the upper
limit is determined by the processability. Thus the preferred
ranges are, for
Nb 0.1-2%
W: 0.1-5%
[0108] W may also be replaced entirely or partly by Mo and/or Ta:
0.1-5%.
[0109] Silicon can participate in the formation of precipitates in
iron-containing alloys, for example in the Fe.sub.2(M, Si) or
Fe.sub.7(M, Si).sub.6 Laves phases. It favors the increased
precipitation and stability of these Laves phases and in this way
contributes to the creep strength. During formation of the Laves
phases, it is completely bound in these. Thus the formation of a
silicon dioxide layer no longer takes place under the chromium
oxide layer. At the same time, the incorporation of M in the oxide
layer is reduced, whereby the negative influence of M on the
oxidation resistance is prevented. At least 0.05% Si must be
present for the desired effect to occur. If the content of Si is
too high, the negative effect of the Si may reappear. The Si
content is therefore limited to 1%.
[0110] Aluminum in contents below 1% impairs the oxidation
resistance, since it leads to internal oxidation. However, an
aluminum content higher than 1% leads to formation, under the
chromium oxide layer, of an aluminum oxide layer, which is not
electrically conductive and thus reduces the contact resistance.
Therefore the aluminum content is limited to .ltoreq.0.5% when a
chromium oxide former is desired or its oxidation resistance is
sufficient. An example of this is, for example, for the application
as interconnector plate. However, a certain aluminum content of at
least 0.001% is necessary for deoxidation of the melt. If no
conductive oxide is necessary and at the same time the requirement
of much higher oxidation resistance than is given by a chromium
oxide layer is still required, the alloy may form a closed aluminum
oxide layer by a content of aluminum of at least 2% (DE 101 20
561). Aluminum contents above 6.0% lead to processing problems and
thus to increased costs.
[0111] Carbon leads to carbide precipitates and thus contributes to
the creep strength. The carbon content should be <0.1%, in order
not to impair the processability. However, it should be >0.002%,
so that an effect can occur.
[0112] The nitrogen content should be 0.1% maximum, in order to
avoid formation of nitrides, which impair the processability. It
should be higher than 0.002%, in order to assure the processability
of the material.
[0113] The contents of sulfur should be made as low as possible,
since this interface-active element impairs the oxidation
resistance. A maximum of 0.01% S is therefore stipulated.
[0114] Oxygen-affine elements such as Ce, La, Pr, Ne, Sc, Y, Zr, Hf
improve the oxidation resistance by reducing the oxide growth and
improving the adherence of the oxide layer. A minimum content of
0.02% of one or more of the elements Ce, La, Pr, Ne, Sc, Y, Zr, Hf
is practical in order to obtain the oxidation-resistance-increasing
effect of the Y. For cost reasons, the upper limit is set at 0.3%
by weight.
[0115] As with every oxygen-affine element, titanium is bound in
the oxide layer during oxidation. In addition, it also causes
internal oxidation. However, the resulting oxides are so small and
finely dispersed that they cause hardening of the surface and thus
prevent swelling of the oxide layer and inclusion of metallic zones
during the oxidation (see DE 10 2006 007 598 A1). These swellings
are unfavorable, since the resulting cracks cause an increase of
the oxidation rate. Thus Ti contributes to the improvement of the
oxidation resistance. For effectiveness of the Ti content, at least
0.01% Ti must be present, but not more than 0.5%, since this does
not improve the effect further but increases the costs.
[0116] The content of phosphorus should be lower than 0.030%, since
this interface-active element impairs the oxidation resistance. Too
low P content increases the costs. The P content is therefore
.gtoreq.0.002%.
[0117] The contents of magnesium and calcium are adjusted in the
spread ranges of 0.0001 to 0.05% by weight and 0.0001 to 0.03% by
weight respectively.
[0118] It has been found that cobalt contents of 3% and higher
impair the oxidation resistance. For cost reasons, the lower limit
is set at 0.01% by weight. For nickel and copper, the same applies
as for cobalt.
[0119] Boron is limited to max. 0.005%, since this element reduces
the oxidation resistance.
[0120] The subject matter of the invention will now be explained in
more details on the basis of exemplary embodiments.
[0121] The analyses of the batches used for the following examples
are presented in Table 1. These batches were melted in the arc
furnace in an amount of approximately 30 metric tons, thereafter
cast in a ladle and subjected to a decarburization and deoxidation
treatment as well as to a vacuum treatment in a VOD system and cast
to ingots. These were then hot-rolled and, depending on final
thickness, cold-rolled with intermediate annealing steps. After the
hot-rolling, the oxide layer was removed by pickling.
[0122] A material with an analysis as indicated in Table 1
precipitates mainly Fe.sub.2(M, Si) or Fe.sub.7(M, Si).sub.6 Laves
phases and, in much smaller contents, carbides.
EXAMPLE 1
[0123] In this example, material from the batch 161061 listed in
Table 1 was hot-rolled to 12 mm thick sheet after solution
annealing above 1070.degree. C. for a period of longer than 7
minutes followed by quenching in stationary air, wherein the
mechanical working was begun with a start temperature
>1070.degree. C. and the last 78% of mechanical deformation was
applied by rolling between 500.degree. C. and 1000.degree. C.
[0124] FIG. 1 shows the typical appearance of a microstructure
deformed in this way. In the microsections etched by means of
electrolytic etching with oxalic acid, it can be clearly seen that
only little Laves phase has been precipitated in microscopically
visible form.
[0125] When the material formed in this way is then annealed at
1075.degree. C. for 20 minutes with quenching in stationary air, a
microstructure is obtained with only few precipitates of Laves
phase and a grain size of approximately 137 .mu.m (FIG. 2), which
is a typical coarse-grained microstructure.
[0126] When a creep test as described above is performed on this
material with an initial stress of 35 MPa at a temperature of
750.degree. C., the specimen breaks after 12.8 hours at an
elongation A of 69.8%. (Table 2). At room temperature, the material
annealed in this way has an elongation of 35%, which is a very good
value for a ferrite.
[0127] If, in contrast, from the hot-rolled material, which is
synonymous with preworking, a specimen for a creep test is made as
simulation for a component and this is then heated at approximately
60.degree. C./minute to an application temperature of 750.degree.
C. and then a creep test is performed with an initial stress of 35
MPa at a temperature of 750.degree. C., the specimen surprisingly
breaks only after 255 hours at an elongation A of 29%, which means
a prolongation of the time to break by a factor of 20. Making the
component is very easily possible, since the hot-worked condition,
as was described above, has an elongation of 19% in the tension
test at room temperature, which is a good value and makes the
material readily processable.
[0128] This example clearly shows that the microstructure with the
preworking and the coarse-grained microstructure is superior with
respect to time to break or creep strength, which contradicts the
state of the art as described in "Burgel" pages 196 to 199 Table
3.7.
EXAMPLE 2
[0129] In this example, annealing steps were performed on the
hot-rolled material from Example 1 for 20 minutes each between
600.degree. C. and 1000.degree. C. or for some temperatures also
for 240 or 1440 minutes (see Table 3 for t.sub.min and t.sub.max
according to Equation 1 and 2) in air, followed by quenching in
stationary air. After the heat treatment, specimens were made from
the sheet and then the creep test was performed with a stress of 35
MPa at 750.degree. C. as described above. The results are compiled
in Table 3.
[0130] After 20 minutes at 1000.degree. C., a time to break of only
10.4 hours is reached at an elongation A of 79.5%. After 20 minutes
at 600.degree. C. to 950.degree. C., a time to break of longer than
100 hours, at least 7 times longer, is achieved at an elongation A
of greater than 22.7%. The longest time to break for annealing
steps of 20 minutes is achieved at 850.degree. C., with 564 hours.
The longest time to break for annealing steps of 240 minutes is
achieved at 800.degree. C., with 396 hours. After 1440 minutes at
700.degree. C., a time to break of 645 hours is even reached. FIG.
3 shows the microstructure after the various annealing steps for 20
minutes. The microstructures in FIG. 3 are not globularly
recrystallized. Up to 850.degree. C. (the maximum of the time to
break), the microstructure has the typical appearance of a deformed
microstructure. Starting from approximately 900.degree. C.,
recovery can be clearly recognized, but this means that the
dislocation density is still increased compared with a globularly
recrystallized microstructure. In a recovered microstructure, the
dislocations have become partly reordered at small-angle grain
boundaries. This has an effect similar to that of preworking. In
the microsections etched by means of electrolytic etching with
oxalic acid, it can be clearly recognized that the Laves phase is
precipitated in microscopically visible form starting from
approximately 750.degree. C., wherein it is precipitated
increasingly more densely and more homogeneously up to 850.degree.
C. (the maximum of the time to break). From approximately
900.degree. C. on, small-angle grain boundaries or grain boundaries
can also be recognized markedly besides the precipitates in the
grain, thus assuring jaggedness of the small-angle grain boundaries
or grain boundaries, which corresponds to measure 2 for increasing
the creep strength (see above). At 1000.degree. C., very large
grains have recognizably formed due to advancing recovery, such
that the dislocation density is greatly reduced and so no further
increase of the time to break occurs. The maximum of the time to
break occurs in the deformed microstructure with dense
homogeneously precipitated Laves phases.
[0131] At room temperature, the sheet annealed for 20 minutes at
all temperatures between 600.degree. C. and 950.degree. C. has an
elongation of at least 13%, which is still to be regarded as
satisfactory for a ferritic alloy and makes the material
processable. The elongation is smallest in the range of 700.degree.
C. to 800.degree. C. and is improved at the lower or higher
annealing temperatures respectively, because at the lower
temperatures Laves phase is certainly already precipitated but is
not yet microscopically visible and therefore has a smaller
proportion by volume, but in return is very finely dispersed. At
the higher temperatures, a larger proportion by volume is
precipitated, but in return is somewhat coarser and recognizable at
the small-angle grain boundaries and grain boundaries.
[0132] The annealing at 1000.degree. C. for an annealing time of 20
minutes exceeds t.sub.max=19.6 minutes. Thus it is not in the range
of the invention and is used as reference. Also, the time to break
is only 10.4 hours. The annealing time of 20 minutes at the
temperatures between 600.degree. C. and 950.degree. C. lies in the
inventive range between t.sub.min and t.sub.max. Accordingly, the
time to break was clearly prolonged according to the invention by
more than a factor of 7 compared with the coarse-grained,
globularly recrystallized condition from Example 1, which is
obtained after annealing at 1075.degree. C./20 minutes followed by
quenching in stationary air.
EXAMPLE 3
[0133] In this example, material from the batch 161061 listed in
Table 1 was hot-rolled to 12 mm thick sheet after solution
annealing above 1070.degree. C. for a period of longer than 7
minutes followed by quenching in stationary air, wherein the
working was begun with a start temperature >1070.degree. C. and
the last 60% of mechanical working was applied by rolling between
1000.degree. C. and 500.degree. C.
[0134] If the sheet worked in this way is then annealed
industrially in the continuous furnace at 920.degree. C. for 28
minutes in air and quenched in stationary air, a tension specimen
made from this material has a time to break of 391 hours at an
elongation A of 38% (Table 4) in the creep test with an initial
stress of 35 MPa at a temperature of 750.degree. C. The
microstructure is not globularly recrystallized but instead is
recovered. It has precipitates in the grain and at the small-angle
grain boundaries or grain boundaries (FIG. 4). The time to break is
30 times the time achieved in Example 1 after annealing at
1075.degree. C. for 20 minutes with a globularly recrystallized
coarse-grained microstructure with a grain size of 137 .mu.m. The
annealing at 920.degree. C. for an annealing time of 28 minutes
lies in the inventive range between t.sub.min=0.32 minutes and
t.sub.max=162.6 minutes.
[0135] At room temperature, the sheet treated in this way has a
very good elongation of 18%, an offset yield strength of 475 MPa
and a tensile strength of 655 MPa (see Table 4), which makes the
material readily workable.
EXAMPLE 4
[0136] In this example, material from batch 161061 and batch 161995
was cold-rolled to 1.5 mm thick sheet after solution annealing at
above 1070.degree. C. for a period longer than 7 minutes followed
by quenching in blown air and hot rolling as well as removal of the
oxide layer, wherein cold working of 53% was applied. Subsequently
annealing at 1050.degree. C. was carried out for 3.4 minutes under
protective gas in the continuous furnace with subsequent quenching
in the cold stream of protective gas. Thereafter both batch 161061
(FIG. 5) and batch 161995 exhibit a recovered microstructure with
elongated grains (FIG. 7) and precipitation of Laves phase,
although much less than recognizable in FIG. 4. Thereafter part of
the material was annealed once again at 1050.degree. C. for 20
minutes under air with subsequent quenching in stationary air.
After this, both batches were globularly recrystallized, batch
161061 with a grain size of 134 .mu.m (FIG. 6) and batch 161995
with a grain size of 139 .mu.m. Only slight precipitated Laves
phase can still be found.
[0137] Tables 5a and 5b show the results of the creep tests and of
the tension tests at room temperature. After the annealing at
1050.degree. C. for 3.4 minutes, batch 161061 has a time to break
of 25.9 hours at an elongation A of 50% in a creep test at
750.degree. C. with an initial load of 35 MPa, and after additional
annealing at 1050.degree. C. for 20 minutes, which produces very
coarse grain, a time to break of only one third, 7.9 hours, at an
elongation A of 83%.
[0138] Similarly, batch 161995 has a time to break of 33.5 hours
89% in a creep test at 750.degree. C. with an initial load of 35
MPa, and after additional annealing at 1075.degree. C. for 20
minutes, which produces very coarse grain, a time to break of only
one third, 7.9 hours, at an elongation A of 92%. The elongation of
28% in the tension test at room temperature for batch 161061 for
1050.degree. C. and 3.4 minutes of annealing time and of 26% for
batch 161995 is very good for a ferrite, which makes the material
very readily workable. For the coarse-grained structure, it is even
higher, with 31% for batch 161061 and 29% for batch 161995.
[0139] This shows the influence of the annealing time at
temperatures around 1050.degree. C., In short-time annealing steps
of a few minutes, dislocations (deformation) and adequate Laves
phase are present in the material, which in this example has the
consequence of a 3 to 4 times longer time to break in the creep
test. For longer annealing steps, the Laves phase is sufficiently
dissolved, as batch 161061 shows, and the microstructure
recrystallizes globularly with correspondingly short times to break
in the creep test.
[0140] The annealing at 1050.degree. C. for 20 minutes lies with an
annealing time of 20 minutes above t.sub.max=6.0 minutes. Thus it
does not fall within the range of the invention and is used as
reference, just as the annealing at 1075.degree. C. for 20 minutes.
The annealing at 1050.degree. C. for 3.4 minutes lies with an
annealing time of 3.4 minutes in the inventive range between
t.sub.min=0.32 minutes and t.sub.max=6.0 minutes and according to
the invention exhibits a clearly improved time to break in the
creep test.
EXAMPLE 5
[0141] In this example, material from the batch 161061 was
hot-rolled to 12 mm thick sheet after solution annealing above
1070.degree. C. for a period of longer than 7 minutes followed by
quenching in stationary air, wherein the working was begun with a
start temperature >1070.degree. C. and the last 70% of
mechanical deformation was applied by rolling between 1000.degree.
C. and 500.degree. C.
[0142] When the material worked in this way is then subjected to
solution annealing at 1075.degree. C. for 22 minutes with quenching
in stationary air, a very coarse-grained microstructure is obtained
with only few precipitates of Laves phase and a grain size of
approximately 134 to 162 .mu.m (FIG. 9). When a creep test is
performed on this material with an initial stress of 40 MPa at a
temperature of 700.degree. C., the specimen breaks after 228 hours
at an elongation A of 51%. (Table 6) When the creep test is
performed at 60 MPa, the specimen breaks after 8.1 hours, at an
elongation A of 43%. At room temperature, the material annealed in
this way has an elongation of 35%, which is a very good value for a
ferrite.
[0143] If the material solution annealed at 1075.degree. C. for 22
minutes is additionally subjected to annealing for 4 hours at
700.degree. C. with subsequent quenching in stationary air, Laves
phase dispersed in the microstructure is precipitated. (FIG. 10).
When the creep test is then performed at 700.degree. C. with an
initial stress of 40 MPa, the specimen already breaks after 104
hours, at an elongation A of 72.6%, therefore a much shorter time
than after the solution annealing at 1075.degree. C. for 22
minutes. When the creep test is performed at 60 MPa, the specimen
breaks after 6.3 hours, at an elongation A of 63%, therefore also
after substantially shorter time than after the solution annealing
at 1075.degree. C. for 22 minutes.
[0144] This is the proof that the precipitation of the Laves
phase(s) must take place in a microstructure with elevated
dislocation density, therefore in a worked or recovered
microstructure, in order to achieve prolongation of the time to
break. Precipitation in a solution annealed microstructure has
exactly the opposite effect, namely shortening of the time to
break. The cause of this is the more homogeneous dispersion of very
fine precipitates in the case of precipitation in a microstructure
with elevated dislocation density, or in other words a deformed or
recovered microstructure, in comparison with precipitation in a
dislocation-poor coarse-grained microstructure.
EXAMPLE 6
[0145] In this example, as in Example 2, annealing steps were
performed on the hot-rolled material from Example 1 for 20 minutes
each between 750.degree. C. and 1000.degree. C. or for some
temperatures also for 120 minutes, 240 minutes, 480 minutes, 960
minutes, 1440 minutes or 5760 minutes (see Table 7 for t.sub.min
and t.sub.maX according to Equation 1 and 2) in air, followed by
quenching in stationary air. After the heat treatment, specimens
were made from the sheets and then the creep test was performed
with a stress of 40 MPa at 750.degree. C. as described above. The
higher stress in comparison with Example 2 was chosen for
shortening of the test time. The objective was to find
heat-treatment times suitable for the annealing steps. The results
are compiled in Table 7.
[0146] After 20 minutes at 1000.degree. C., a time to break of only
8.8 hours is reached at an elongation A of 78.7%. In Example 2,
after 20 minutes at 1000.degree. C. and a creep test at 750.degree.
C. and 35 MPa, a time to break comparable with that after solution
annealing at 1075.degree. C. for 20 minutes with quenching in
stationary air was reached, and so this value can be taken as
reference for the time to break of the solution annealed condition.
Inventive variants should also exceed this break time once again by
a factor of at least 1.5.
[0147] After 20 minutes at 750.degree. C. to 900.degree. C., a time
to break of longer than 100 hours, at least 10 times longer, is
achieved at an elongation A of greater than 27%. The longest time
to break for annealing steps of 20 minutes is achieved at
850.degree. C. with 296 hours. The longest time to break for
annealing steps of 120 minutes is achieved at 800.degree. C. with
227 hours. The longest time to break for annealing steps of 240
minutes is achieved at 750.degree. C. with 182 hours, but in this
connection no value exists for 700.degree. C. The longest time to
break for annealing steps of 480 minutes is achieved at 800.degree.
C. with 169 hours. For 960 minutes, only one time to break was
determined, for 750.degree. C., with a value of 139 hours at an
elongation of 24.2%. After 1440 minutes and 5760 minutes at
750.degree. C. and 800.degree. C., only times to break clearly
shorter than the maximum times to break achieved at these
temperatures are still achieved. At 800.degree. C., for example,
the time to break after a treatment time of 480 minutes drops from
169 hours to a value of 46 hours after a treatment time of 1440
minutes, to a value of 17.5 hours after a treatment time of 5760
minutes, although this is still in the inventive range. A further
prolongation of the treatment time should shorten the time to break
further, so that t.sub.max of 7059 minutes is logically somewhat
longer than the time of 5750 minutes. All elongations for the heat
treatment temperatures from 750 to 900.degree. C. and times from 20
minutes to 5760 minutes lie between 24.2% and 43% and therefore are
larger than 18%, as required, in order to avoid brittle failure.
For the microstructure after 20 minutes of annealing, what was said
in Example 2 is applicable, since the annealing steps are the same.
Even at the higher stress of 40 MPa in the creep test, the maximum
of the time to break occurs in the deformed microstructure, with
dense homogeneously precipitated Laves phase.
[0148] At room temperature, the sheet annealed for 20 minutes at
all temperatures between 600.degree. C. and 900.degree. C. in
Example 2 has an elongation of at least 13%, which is still to be
regarded as satisfactory for a ferritic alloy and makes the
material processable.
EXAMPLE 7
[0149] In this example the 1.5 mm thick material of batch 161995,
which was annealed after cold working of 53% at 1050.degree. C. for
3.4 minutes under protective gas in the continuous furnace with
subsequent quenching in the stream of cold protective gas, was used
once again. In the same manner, 2.5 mm thick material from batch
161995 was produced by annealing it, after cold working of 40%, at
1050.degree. C. for 2.8 minutes under protective gas in the
continuous furnace with subsequent quenching in the stream of cold
protective gas. Even the 2.5 mm thick material then exhibits a
recovered microstructure with elongated grains (FIG. 11), just as
the material from Example 4 in FIG. 7, and precipitation of Laves
phase, albeit clearly less than recognizable in FIG. 4. Part of the
material was then annealed once more at 1050.degree. C. for 10
minutes under air with subsequent quenching in stationary air.
After this the material had globularly recrystallized structure,
with a grain size of 108 .mu.m. Only little precipitated Laves
phase is still to be found. The material with 1050.degree. C./2.8
minutes and the material with 1075.degree. C./10 minutes as the
last heat treatment was then rolled with degrees of working between
2.8 and 40%. Thereafter creep tests were performed at 750.degree.
C. and 35 MPa and tension tests were performed at room temperature.
The results are summarized in Table 8.
[0150] After the annealing at 1050.degree. C. for 3.4 minutes,
batch 161995, in a creep test at 750.degree. C. with an initial
load of 35 MPa, had a time to break of 33.5 hours at an elongation
A of 89% and, after the additional annealing at 1050.degree. C. for
10 minutes, which produces very coarse grain, it had a time to
break amounting to only one third, 10.8 hours, at an elongation A
of 50.4%.
[0151] If, from the material worked after 1050.degree. C./2.8
minutes, a specimen for a creep test is made as simulation for a
component and this is then heated at approximately 60.degree.
C./minute to an application temperature of 750.degree. C. and then
a creep test is performed with an initial stress of 35 MPa at a
temperature of 750.degree. C., the elongation at break for degrees
of working between 5 and 40% drops to values around the 10 hours
with elongations at break greater than 45%.
[0152] If, in contrast, from the material formed after 1050.degree.
C./10 minutes, a specimen for a creep test is made as simulation
for a component and this is then heated at approximately 60.degree.
C./minute to an application temperature of 750.degree. C. and then
a creep test is performed with an initial stress of 35 MPa at a
temperature of 750.degree. C., the elongation at break for degrees
of working between 2.9 and 40% increases to values between 49 and
137 hours, which means an increase by more than a factor of 4 in
the time to break compared with the material worked after
1050.degree. C./2.8 minutes, wherein a maximum occurs at 10% and
the elongations at break lie between 18.9 and 60%.
[0153] From 10% degree of working on, however, the elongation at
break in the tension test at room temperature becomes smaller than
8%, so that the material is increasingly more poorly processable.
In other words, preferred degrees of forming for cold shaping lie
between 0.05 and 10%. Thus this example shows that the increase of
the elongation at break after working does not occur if the
annealing before working was carried out at too low temperatures or
for too short times. (Here at 1050.degree. C. for 2.8 minutes) The
increase of the elongation at break occurred after annealing was
carried out at >1050.degree. C. for >6 minutes (here at
1075.degree. C. for 7 minutes).
[0154] The titles/descriptions of the tables/figures are reproduced
as follows: [0155] Table 1 Composition of the investigated alloy
(all values in % by weight) [0156] Table 2 Results of the creep
tests at 750.degree. C. with 35 MPa and of the tension tests at
room temperature for the hot rolling and the heat treatments in
Example 1 for a 12 mm thick sheet. (R: reference according to the
state of the art, I: according to the invention) [0157] Table 3
Results of the creep tests at 750.degree. C. with 35 MPa and of the
tension tests at room temperature for the hot rolling from Example
1 and the heat treatment from Example 2 for a 12 mm thick sheet.
(R: reference according to the state of the art, I: according to
the invention) [0158] Table 4 Results of the creep tests at
750.degree. C. with 35 MPa and of the tension tests at room
temperature for Example 3 for a 12 mm thick sheet. (R: reference
according to the state of the art, I: according to the invention)
[0159] Table 5 Results of the creep tests at 750.degree. C. with 35
MPa and of the tension tests at room temperature for Example 4 for
a 1.5 mm thick strip. (R: reference according to the state of the
art, I: according to the invention) [0160] Table 6 Results of the
creep tests at 700.degree. C. and of the tension tests at room
temperature for Example 5 on 12 mm thick sheet (R: reference
according to the state of the art, I: according to the invention)
[0161] Table 7 Results of the creep tests at 750.degree. C. with 40
MPa and of the tension tests at room temperature for the hot
rolling and the heat treatments in Example 6 for a 12 mm thick
sheet. (R: reference according to the state of the art, I:
according to the invention) [0162] Table 8 Results of the creep
tests at 750.degree. C. with 35 MPa and of the tension tests at
room temperature for Example 7 for 1.5 mm to 2.5 mm thick strip
from batch 161995. (R: reference according to the state of the art,
I: according to the invention)
[0163] FIG. 1 Microstructure of the hot-worked material in Example
1
[0164] FIG. 2 Microstructure of the hot-worked material in Example
1 after annealing at 1075.degree. C. for 20 minutes and quenching
in stationary air, grain size 137 .mu.m.
[0165] FIG. 3 Microstructure of the material in Example 2 after
annealing between 600.degree. C. and 1000.degree. C. for 20 minutes
in each case and quenching in stationary air.
[0166] FIG. 4 Microstructure of the material in Example 3 after
annealing at 920.degree. C. in the continuous furnace in air with
subsequent quenching in stationary air for 20 minutes in each case
and quenching in stationary air. (etching with V2A pickling
fluid)
[0167] FIG. 5 Microstructure of batch 161061 in Example 4 after
annealing at 1050.degree. C./3.4 minutes under, protective gas in
the continuous furnace with quenching in the stream of cold
protective gas.
[0168] FIG. 6 Microstructure of batch 161061 in Example 4 after
annealing at 1050.degree. C./3.4 minutes under protective gas in
the continuous furnace with quenching in the stream of cold
protective gas and annealing at 1050.degree. C./20 minutes under
air with subsequent quenching in stationary air, grain size 134
.mu.m (etching with V2A pickling fluid)
[0169] FIG. 7 Microstructure of batch 161995 in Example 5 after
annealing at 1050.degree. C./3.4 minutes under protective gas in
the continuous furnace with quenching in the stream of cold
protective gas.
[0170] FIG. 8 Microstructure of batch 161995 in Example 5 after
annealing at 1050.degree. C./3.4 minutes under protective gas in
the continuous furnace with quenching in the stream of cold
protective gas and annealing at 1075.degree. C./20 minutes under
air with subsequent quenching in stationary air, grain size 139
[0171] FIG. 9 Microstructure of the hot-worked material in Example
5 after annealing at 1075.degree. C. for 22 minutes and quenching
in stationary air, grain size 134 .mu.m to 162 .mu.m
[0172] FIG. 10 Microstructure of the hot-worked material in Example
5 after annealing at 1075.degree. C. for 22 minutes followed by
quenching in stationary air and subsequent annealing at 700.degree.
C. for 4 hours followed by quenching in stationary air. Grain size
136 .mu.m.
[0173] FIG. 11 Microstructure of batch 161995 in Example 7 after
annealing at 1050.degree. C. for 2.8 minutes under protective gas
in the continuous furnace with quenching in the stream of cold
protective gas.
[0174] FIG. 12 Microstructure of batch 161995 in Example 7 after
annealing at 1050.degree. C. for 2.8 minutes under protective gas
in the continuous furnace with quenching in the stream of cold
protective gas, followed by annealing at 1075.degree. C. for 10
minutes under air with subsequent quenching in stationary air.
Grain size 108 .mu.m.
TABLE-US-00001 TABLE 1 Batch Batch Element 161061 161995 C 0.007
0.009 S <0.002 <0.002 N 0.015 0.018 Cr 22.9 22.6 Ni 0.30 0.22
Mn 0.43 0.43 Si 0.21 0.24 Mo 0.02 0.02 Ti 0.07 0.06 Nb 0.51 0.49 Cu
0.02 0.02 Fe Remainder Remainder P 0.014 0.017 Al 0.02% 0.019 Mg
0.0006 <0.01 Pb <0.001 <0.001 Sn <0.01 <0.01 Ca
0.0002 0.01 V 0.05 0.02 Zr <0.01 <0.01 W 1.94 1.97 Co 0.04
0.02 La 0.08 0.05 Ce <0.01 O 0.004
TABLE-US-00002 TABLE 2 Creep test with 35 Tension test at room
temperature MPa at 750.degree. C. Offset yield Tensile Time
t.sub.B.sub.-- to Elongation strength strength Elongation Heat
treatment break in hours A in % Rp0.2 RM in MPa A 5 in % R
1075.degree. C./20 minutes 12.8 69.8 359 494 35 Hot-worked 511 604
19 I Component + heating at 255 29.0 60.degree. C./minute to
application temperature 750.degree. C.
TABLE-US-00003 TABLE 3 Limit times according Creep test with 35
Tension test at room temperature Heat treatment to equations 1 and
MPa at 750.degree. C. Offset yield Tensile Temperature Time in 2 in
minutes Time t.sub.B.sub.-- to Elongation strength strength
Elongation in .degree. C. minutes t.sub.min t.sub.max break in
hours A in % Rp0.2 RM in MPa A 5 in % Hot-worked 511 604 19 I 600
20 27.9 39.2*10.sup.6 278 37.6 I 650 20 11.26 31.8*10.sup.5 286
35.2 527 622 17 I 700 20 5.00 33.6*10.sup.4 260; 264 22.7; 39.1 537
676 13 I 750 20 2.41 44300 263 430.6 514 707 16 I 800 20 1.25 7059
344 25.0 505 699 13 I 850 20 0.69 1328 564 32.2 484 672 17 I 900 20
0.40 289 337 25.3 467 638 18 I 950 20 0.24 71.2 121; 72 31.1; 28.8
451 614 17 R 1000 20 0.15 19.6 10.4 79.5 n.m. n.m. n.m. I 650 240
11.26 31.8*10.sup.5 293 36.2 n.m. n.m. n.m. I 700 240 5.00
33.6*10.sup.4 233 32.8 n.m. n.m. n.m. I 750 240 2.41 44300 224 23.2
n.m. n.m. n.m. I 800 240 1.25 7059 396 43.2 n.m. n.m. n.m. I 850
240 0.69 1328 181 35.2 n.m. n.m. n.m. I 900 240 0.40 289 45.6 55.7
n.m. n.m. n.m. R 950 240 0.24 71.2 10.8 78.7 n.m. n.m. n.m. I 700
1440 5.00 33.6*10.sup.4 645 30.9 n.m. n.m. n.m. For comparison from
Example 1 R 1075 20 12.8 69.8 359 494 35 n.m. = not measured
TABLE-US-00004 TABLE 4 Limit times according Creep test with 35
Tension test at room temperature Heat treatment to equations 1 and
MPa at 750.degree. C. Offset yield Tensile Temperature Time in 2 in
minutes Time t.sub.B.sub.-- to Elongation strength strength
Elongation in .degree. C. minutes t.sub.min t.sub.max break in
hours A in % Rp0.2 RM in MPa A 5 in % I 920 28 0.32 162.6 391 38
475 655 18
TABLE-US-00005 TABLE 5a Limit times according Creep test with 35
Tension test at room temperature to equations 1 and MPa at
750.degree. C. Offset yield Tensile Batch 161061 2 in minutes Time
t.sub.B.sub.-- to Elongation strength strength Elongation Heat
treatment t.sub.min t.sub.max break in hours A in % Rp0.2 RM in MPa
A 5 in % I 1050.degree. C./3.4 minutes 0.1 6.0 25.9 50 385 541 27
380 537 28 R 1050.degree. C./20 minutes 0.1 6.0 7.9 83 344 494
31
TABLE-US-00006 TABLE 5b Limit times according Creep test with 35
Tension test at room temperature to equations 1 and MPa at
750.degree. C. Offset yield Tensile Batch 161995 2 in minutes Time
t.sub.B.sub.-- to Elongation strength strength Elongation Heat
treatment t.sub.min t.sub.max break in hours A in % Rp0.2 RM in MPa
A 5 in % I 1050.degree. C./3.4 minutes 0.1 6.0 33.5 89.0 399 538 26
R 1050.degree. C./20 minutes 7.7 92.1 331 475 29
TABLE-US-00007 TABLE 6 Limit times according Tension test at room
temperature to equations 1 and Creep test at 700.degree. C. Offset
yield Tensile Batch 161061 2 in minutes in hours Stress in Time
t.sub.B.sub.-- to Elongation strength strength Elongation Heat
treatment t.sub.min t.sub.max MPA break in hours A in % Rp0.2 RM in
MPa A 5 in % R 1075.degree. C./22 minutes 40 228 51 367 502 35 R
1075.degree. C./22 minutes 60 8.1 43.1 367 502 35 R 1075.degree.
C./22 minutes + 0.083 5601 40 104 72.6 700.degree. C./4 hours R
1075.degree. C./22 minutes + 0.083 5601 60 6.3 63.5 700.degree.
C./4 hours
TABLE-US-00008 TABLE 7a Limit times according Creep test with 40
Tension test at room temperature Heat treatment to equations 1 and
MPa at 750.degree. C. Offset yield Tensile Temperature Time in 2 in
minutes Time t.sub.B.sub.-- to Elongation strength strength
Elongation in .degree. C. minutes t.sub.min t.sub.max break in
hours A in % Rp0.2 RM in MPa A 5 in % I 750 20 2.41 44300 128 34.7
514 707 16 I 800 20 1.25 7059 189 27.2 505 699 13 I 850 20 0.69
1328 296 32.1 484 672 17 I 900 20 0.40 289 174 35.4 467 638 18 R
1000 20 0.15 19.6 8.8 78.7 n.m. n.m. n.m. I 750 120 2.41 44300 150
31.7 n.m. n.m. n.m. I 800 120 1.25 7059 227 26.2 n.m. n.m. n.m. I
850 120 0.69 1328 133 29.7 n.m. n.m. n.m. I 900 120 0.40 289 32.7
40 n.m. n.m. n.m. n.m. = not measured
TABLE-US-00009 TABLE 7b Limit times according Creep test with 35
Tension test at room temperature Heat treatment to equations 1 and
MPa at 750.degree. C. Offset yield Tensile Temperature Time in 2 in
minutes Time t.sub.B.sub.-- to Elongation strength strength
Elongation in .degree. C. minutes t.sub.min t.sub.max break in
hours A in % Rp0.2 RM in MPa A 5 in % I 750 240 2.41 44300 182 32.2
n.m. n.m. n.m. I 800 240 1.25 7059 163; 135 27.5; 26.5 n.m. n.m.
n.m. I 850 240 0.69 1328 57 33.9 n.m. n.m. n.m. I 750 480 2.41
44300 152 43.4 n.m. n.m. n.m. I 800 480 1.25 7059 169 26.5 n.m.
n.m. n.m. I 850 480 0.69 1328 35 33 n.m. n.m. n.m. I 750 960 2.41
44300 139 24.2 n.m. n.m. n.m. I 750 1440 2.41 44300 82 25.5 n.m.
n.m. n.m. I 800 1440 1.25 7059 46 46.1 n.m. n.m. n.m. I 750 5760
2.41 44300 54 52.9 n.m. n.m. n.m. I 800 5760 1.25 7059 17.5 50.3
n.m. n.m. n.m. n.m. = not measured
TABLE-US-00010 TABLE 8a Limit times according Creep test with 40
Tension test at room temperature Heat treatment to equations 1 and
Degree of MPa at 750.degree. C. Offset yield Tensile Temperature
Time in 2 in minutes working in Time t.sub.B.sub.-- to Elongation
strength strength Elongation in .degree. C. minutes t.sub.min
t.sub.max % break in hours A in % Rp0.2 RM in MPa A 5 in % I 1050
3.4 0.1 6.0 0.0 33.5 89 399 538 26 R 1050 2.8 0.1 6.0 5.0 10.6 45.8
581 616 16 R 1050 2.8 0.1 6.0 10.0 10.9 68.6 663 685 8 R 1050 2.8
0.1 6.0 20.0 9.4 71.8 725 745 6 R 1050 2.8 0.1 6.0 40.0 12.0 85.9
811 832 4 n.m. = not measured
TABLE-US-00011 TABLE 8b Creep test with 40 Tension test at room
temperature Degree of Heat treatment MPa at 750.degree. C. Offset
yield Tensile working in Temperature Time in Deformation Time
t.sub.B.sub.-- to Elongation strength strength Elongation % in
.degree. C. minutes in % break in hours A in % Rp0.2 RM in MPa A 5
in % I 0.0 1075 10 0.0 10.8 50.4 338 479 28 I 2.8 1075 10 2.8 86
23.3 506 545 17 I 2.8 1075 10 2.8 53 19.8 506 545 17 I 5.0 1075 10
5.0 75 18.9 548 571 15 I 5.0 1075 10 5.0 110 49.5 548 571 15 R 10.0
1075 10 10.0 137 31.7 650 555 8 R 20.0 1075 10 20.0 108 23.9 739
750 5 R 20.0 1075 10 20.0 73 60.6 739 750 5 R 40.0 1075 10 40.0 61
24.5 831 837 3 R 40.0 1075 10 40.0 49 34 831 837 3 n.m. = not
measured
* * * * *