Note: Descriptions are shown in the official language in which they were submitted.
CA 02801637 2012-12-05
WO 2011/154515 PCT/EP2011/059657
A METHOD FOR PRODUCING A TEMPERED MARTENSITIC HEAT RESISTANT STEEL FOR
HIGH TEMPERATURE APPLICATIONS
Global requirements for energy are forecast to double by 2030 with a projected
build
of 40 nuclear power stations alone in the EU. In addition to this significant
increase in
fossil fuelled stations will be required to meet the predicted demand. In
order to
minimise emissions of greenhouse gases improved generating efficiency, through
increased steam temperatures and pressures, and carbon capture technology will
be
required.
Currently martensitic steels used for turbine blades and casings, bolting and
boiler
tubes are limited to service temperatures of about 620 C and the best
commercially
available alloy is Steel 92 - 9%Cr, 0.5%Mo, 2%W. Over the last 15 years
significant
effort has been invested by international consortia, such as COST 522 and 536,
in
Europe to raise the operating temperatures of these martensitic alloys.
Attempts at
alloy modifications have not thus far yielded any solutions backed up by long
term
creep data.
It is an object of this invention to provide a method for producing a tempered
martensitic heat resistant steel for high temperature applications at an
application
temperature of up to 650 C.
A further object is to provide a tempered martensitic heat resistant steel for
high
temperature applications at an application temperature of up to 650 C.
The object is reached by a method as described in claim 1 and a steel as
described in
claim 9. Preferable embodiments are described in the dependent claims.
According to the invention, a method is provided for producing a tempered
martensitic
heat resistant steel for high temperature applications at an application
temperature of
up to 650 C, wherein the steel comprises, on the basis of percent by weight:
- 8.5 to 12% Cr,
- upto0.13%C,
- up to 0.5% Si,
- up to 2.0% W,
- upto3.0%Co,
- upto2%Cu,
- up to 0.8% Mn,
- up to 1.0% Mo,
- up to 0.7% Ni,
- up to 0.04% Al,
CA 02801637 2012-12-05
WO 2011/154515 PCT/EP2011/059657
-2-
- between 0.001 and 0.015 B,
- between 0.005 and 0.07 N,
- up to 0.25%V,
- at least one of from 0.01% up to 0.09% Nb and/or from 0.01% up to
0.14% Ta,
- balance iron and inevitable impurities;
wherein the C:N ratio is below 1.3 to favour formation of nano-scale carbo-
nitrides of the M(C,N) and M2(C,N)-type and to reduce the fraction of
M23(C,B)6
precipitates, the process comprising the steps of
- solution treating the steel in the austenite range at a temperature below
the transformation temperature to delta-ferrite to dissolve substantially all
precipitates including boron-nitrides and carbo-nitrides thereby bringing
the precipitating elements in solid solution;
- quenching the steel after solution treating to create a substantially fully
martensitic matrix and to suppress precipitation on cooling
- tempering the steel in one or more tempering treatments after quenching
to precipitate nano-scale particles M(C,N) or M2(C,N) particles, or mixtures
thereof, at a tempering temperature between 10 to 50 C higher than the
intended application temperature.
Tempered martensitic heat resistant steel derive their creep strength from
four
principal sources:
= Solid solution strengthening
= Dislocation substructure
= M23(C,B)6 precipitation at lath boundaries
M(C,N) and M2(C,N) precipitation at both inter- and intra-granular locations.
During exposure at elevated temperature the dislocation density is reduced and
precipitate coarsening occurs both of which reduce the resistance to creep
deformation. Coarsening of M23(C,B)6 is more rapid than M(C,N) and/or M2(C,N),
but
the rate can be reduced by boron additions. However the long term creep
strength of
these alloys is strongly dependent on the volume fraction and stability of the
M(C,N)
precipitate dispersion. The stability of the precipitates is enhanced by
tempering
above the application temperature and the volume fraction is controlled by the
solution treatment temperature and cooling rate prior to tempering. The use of
lower
tempering temperatures according to this invention in combination with a tuned
chemical composition will produce a high density of fine and stable
precipitates.
CA 02801637 2012-12-05
WO 2011/154515 PCT/EP2011/059657
-3-
High solution treatment temperatures increase the alloy in solution prior to
tempering, but in practice solution temperatures are restricted in order to
ensure a
fully martensitic microstructure and to prevent excessive grain growth which
might
lower creep ductility. Also, it is important to prevent formation of 6-ferrite
during the
solution treatment to ensure a substantially martensitic (a') matrix on
cooling
because the occurrence of a dual phase (a' + b) structure is detrimental for
creep
resistance. The presence of more or less austenite stabilising elements shift
the
transformation temperature to 6-ferrite up or down.
The solution treatment temperature should be high enough to dissolve
completely the
BN, M(2)(C,N) and M23(C,N)6 particles. The temperatures of complete
dissolution of BN
are usually higher than the transformation temperature to 6-ferrite. The same
applies
for some of the M(2)(C,N)-precipitates, depending on the composition and
amounts.
Therefore, not all particles will completely dissolve at the austenitization
temperatures used. M23(C,N)6 particles dissolve at lower austenitization
temperatures
and are normally completely dissolved.
On tempering, alloying elements dissolved during austenitization are
precipitated.
Lowering the tempering temperature from the conventionally used 780 C for
tubes or
710 C for blading (second step) to a temperature of e.g. 660 C increases the
volume
fraction of M(2)(C,N) and M23(C,N)6. Lowering the carbon level significantly
decreases
the volume fraction of M23(C,N)6 in the ferrite region. In the case of
M(2)(C,N) particles
lowering the C:N ratio resulted in a further increase in their volume
fraction.
The inventors found that control of the C: N ratio and the free boron in solid
solution
is essential in the control of the size and stability of the various
precipitate types. The
most important types with regard to high temperature creep properties are
M23(C,B)6
wherein M is mainly Fe, Cr, W or Mo (and mixtures thereof), and carbo-nitrides
of the
type M(C,N) and/or M2(C,N) wherein M is mainly Nb, V, Ta or Cr (and mixtures
thereof). The C:N ratio (as expressed in wt.%) must be below 1.3. Preferably
the C:N
is below 1.2.
Good, long term creep resistance requires a high density of small precipitates
which
are resistant to coarsening during long term exposure under stress at elevated
temperatures. The purpose is to stabilise the dislocation substructure and to
inhibit
dislocation mobility. The former is controlled mainly by M23(C,B)6
precipitates at lath
boundaries and the latter by the nano-scale carbo-nitride precipitation at
boundaries
and within laths. It is surprising that the tempering just above the
application
temperature of the steel in service results in a better creep resistance, even
after a
subsequent post-weld heat treatment where improvements of 170% over
conventionally heat treated samples are obtained.
CA 02801637 2012-12-05
WO 2011/154515 PCT/EP2011/059657
-4-
In an embodiment of the invention, a method is provided for producing a
tempered
martensitic heat resistant steel for high temperature applications at an
application
temperature of at least 620 C, or preferably between 620 to 650 C.
According to the invention the tempering treatment is performed at a
temperature in
the range of 10 to 50 C above the intended application temperature. For a
service or
application temperature of 650 C, the tempering treatment would be performed
at a
temperature of between 660 and 700 C. With intended application temperature
the
operating temperature is meant at which the heat resistant steel is used.
In an embodiment of the invention the solution treating is performed between
1150 C and 1250 C. These temperatures allow for a complete dissolution of the
M23(C,B)6, and the dissolution of the majority of the M(2)(C,N) particles.
The quenching after solution treatment should be as fast as possible to ensure
that
the dissolved elements remain in solid solution and to ensure formation of a
fully
martensitic microstructure. The quenching is preferably performed in oil. In
an
embodiment the oil is at ambient temperatures. However, the quenching could
also
be performed by other means such as forced air, (hot) mist or even (hot)
water, as
long as the martensitic microstructure is obtained, the dissolved elements
remain in
solid solution, and the stresses as a result of the quenching do not exceed
critical
levels so that no warping or cracking occurs.
In a preferable embodiment the tempering treatment is performed in at least
two
tempering treatments, and more preferably wherein the at least two separate
heat
treatments are performed at substantially the same tempering temperature for
substantially the same period of time. Preferably the period of time at the
tempering
temperature is between 1 and 5 hours, preferably between 2 and 4 hours. It is
clear
that the reheating to the tempering treatment should be as fast as possible to
prevent
undesirable reactions taking place during the reheating. By choosing different
tempering temperatures for the at least two tempering treatments within the
range of
between 10 and 50 C above the application temperature or service temperature
the
composition and size distribution of the precipitates which govern the creep
resistance
can be very effectively and reliably controlled.
In an embodiment of the invention a method is provided wherein the first of
the at
least two separate heat treatments is at the temperature range from 500 C up
to
10 C higher than the application temperature and second is or the following
are at
the temperature range from 10 to 50 C higher than the application temperature.
This
will result in even finer dispersion of precipitates. The latter higher
temperature
temper is particularly relevant where a stress relieving treatment after
welding is
CA 02801637 2012-12-05
WO 2011/154515 PCT/EP2011/059657
-5-
required. As an example for an application temperature of 650 C the tempering
can
be done in the first step at 500-660 C and the second step at 660-700 C.
In a preferable embodiment the steel comprises:
- between 1.0 and 2.0% W, and/or
- between 1 and 2% Co if Cr >_ 10%, and/or
- up to 1.5% Cu, and/or
- up to 0.6% Mn, and/or
- up to 0.8% Mo, and/or
- up to 0.5% Ni, and/or
- between 0.15 and 0.25%V, and/or
- between 0.03 and 0.09% Nb and/or
- between 0.05 and 0.12% Ta and/or
- C:N < 1.2.
It should be noted that the term "and/or" when used in this description or the
claims
must be interpreted in the sense that one, more or all of the preferable
ranges or
process conditions may be applicable.
In a preferable embodiment the steel comprises:
- 8.5 to 9.5% Cr, and/or
- between 0.07 and 0.13% C, and/or
- between 1.5 and 2.0% W, and/or
- between 0.30 and 0.60% Mn, and/or
- between 0.3 and 0.6 Mo, and/or,
- up to 0.4% Ni, and/or
- between 0.001 and 0.006% B and/or
- between 0.03 and 0.07% N and/or
- between 0.18 and 0.25%V, and/or
- between 0.04 and 0.07% Nb.
In a second aspect of the invention a tempered martensitic heat resistant
steel for high
temperature applications is provided having a chemical composition as
described
hereinabove and produced in accordance with the method as described
hereinabove
wherein the microstructure of the steel after tempering comprises
intragranular
precipitates having a size of at most 70 nm of the M(C,N) and/or M2(C,N) type
wherein M is one or more of Nb, V, Ta or Cr and wherein the microstructure of
the
steel after tempering comprises M23(C,B)6 precipitates wherein M is mainly
composed
of Cr and Fe on the lath, block, packets and/or prior austenite grain
boundaries.
CA 02801637 2012-12-05
WO 2011/154515 PCT/EP2011/059657
-6-
In an embodiment the M(C,N) and/or M2(C,N) precipitates have a size of at
least 1
nm and at most 70 nm, preferably at most 50 nm and more preferably at most 30
nm. In an embodiment the M(C,N) and/or M2(C,N) precipitates have a size of
between 10 and 70 nm, more preferably between 10 and 50 nm, even more
preferably between 10 and 30 nm.
In a preferable embodiment both Ta and V are present as an alloying element.
In a preferred embodiment the steel according to the invention is used in the
production of components for high temperature applications such as turbine
blades or
casings, bolting and boiler tubes, heat exchangers or other elements in power
generation systems, for use at an application temperature of up to 650 C.
In a preferred embodiment of the invention the C:N ratio of the steel is below
1.2.
Now the preferable embodiments as to the chemical composition will be
described:
Cr level should be selected according to the application temperature for steam
oxidation and corrosion resistance. Recommended Cr level is 9.0-11.0%.
Co is optionally added only to avoid the formation of delta-ferrite on high
solution
treatment temperatures and is not necessary if there is no risk of formation
of this
phase at the temperatures that allow the dissolution of precipitates. For
steels
containing 9.0% Cr the addition of cobalt is not necessary. For 10.0-11.0% Cr
the
addition of Co is recommended and the recommended amount of Co is 1.5%. A
suitable maximum content is 2%.
Cu is optionally added to influence the morphology of the Laves phase and to
avoid
formation of delta-ferrite in a similar way as Co. For 10.0-11.0% Cr the
recommended amount of Cu is 1.5%, for 9.0% Cr this addition can be lower.
Mn and Ni are added to help to avoid formation of delta-ferrite similar to Co.
Recommended levels are below 0.5% Mn and up to 0.6% Ni. In order to benefit
from
the effect of Ni and Mn a minimum amount of 0.1% for one or both elements is
preferable. A suitable minimum Si level is 0.1%, preferably at least 0.15%.
W and Mo are added for solid solution strengthening. Tungsten additionally
stabilizes
M23(C,B)6. recommended combination is 1.5% W and 0.5% Mo. A suitable minimum
W-content is 0.5%.
C:N ratio should be low in order to favour the formation of M(C,N) or M2(C,N)
particles rich in nitrogen and reduce M23(C,B)6fraction. Examples of
favourable C and
N contents are: 0.073% C and 0.065% N, 0.02% C and 0.06% N.
CA 02801637 2012-12-05
WO 2011/154515 PCT/EP2011/059657
-7-
V and Nb or V and Ta (or combination of V, Nb, Ta) is important for the nano-
scale
particles. The examples of favourable V and Nb or Ta contents are 0.18 to
0.25% V
and 0.04 to 0.07% Nb or 0.07 to 0.12% Ta.
Addition of B is important for stabilization of M23 (C,B)6 precipitates. It is
essential is
to optimize the B:N ratio in order to maximise the boron in solid solution
during
solution treatment. It is also important to dissolve MX or M2X particles in
the solution
treatment. In these particles M= V, Nb, Ta, Cr or mixtures thereof and X=N
and/or C.
Therefore solution treatment temperature should be as high as possible but not
resulting in formation of delta-ferrite. A recommended range of solution
treatment
temperatures is 1150 to 1250 C.
The use of low temperature tempering, up to 50 C higher than application
temperature, results in a very fine distribution of nano-scale M(C,N) or
M2(C,N)
particles. The tempering should be done in one or more, preferably one or two,
steps
at the same or different temperatures, up to 50 C higher than the application
temperature. Recommended temperatures are especially those in the range of 10
to
50 C higher than application temperature as they are regarded as especially
favourable for fine distribution of nano-scale MX(C,N) or M2(C,N) particles.
It is also
believed that low temperature tempering favours formation of M2(C,N) over
M(C,N).
The invention will now be further explained by means of the following, non-
limiting
examples.
A 50kg air induction melt of Steel 92 material was produced (Steel A). A
second 60kg
cast with higher nitrogen and lower carbon contents was produced by vacuum
induction melting (Steel B). Two additional casts (Steels C and D) of 50 kg
each were
also produced by vacuum induction melting. Cast C and D have the lowest C:N
ratio
to favour formation of M(C,N) and/or M2(C,N) particles. The chemical analyses
are
shown in Table 1. The ingots were forged to 50mm square bar then rolled to
19mm
diameter round bar for the production of test specimens.
CA 02801637 2012-12-05
WO 2011/154515 PCT/EP2011/059657
8
z N N m m
O .~ m M
U N O O
0 Lf1 LO
U
CR r,
Ln LO
U
co
0
N O O O
N N N N
O O O O
o LO 0
4- z 0 0 , O
O O
Lr) k.0
a-+ z 0000
V) (o O O O O
U
4 ~Lnii
C O O O
a) 0000
C) O O O
a)
x N LO N N
Q O O O O
0 0 0 0 O
V)
a)
N N to to
(a 0000
C
Q
.D~o00 N
( Lfl LfI Q
E O O O O E
(6
X
U vo000 a) 1-1 1-1
a, a, >
Lf) a, '-
0000 co
fo Ln OOOO c-
0o v v E
0
U
N O N N
-1 -1 O O C
o 0 0
O o v V
a)
00 00 M M
OOOO Q)
L
-- 41
m N Lfl Ln 0
M N N N
O O O O
U ^ N N 0
-4000 U
0000 ruo
4~
0
C
Q co u 0
CA 02801637 2012-12-05
WO 2011/154515 PCT/EP2011/059657
-9-
Samples from Steel A were given a standard heat treatment or a treatment
involving
a higher solution treatment temperature at 1150 C and lower tempering
temperature
of 660 C as outlined in Table 2. In order to study the effect of PWHT samples
with
the standard and inventive treatments were given a simulated PWHT of 1 hour at
740 C followed by air cooling. For steel B samples were prepared with the low
temperature tempering treatment and two solution treatments at 1150 C and
1200 C. Steel C and D are given the treatment as it is shown in Table 3.
Table 2: Heat Treatment Schedule for Steel A and B.
Treatment Purpose Austenitization Tempering
Standard Typical (commercial) Al 1060 C/1 h/AC Tl 780 C/2h/AC
High austenitization temperatures A2 1150 C/1 h/AC 660 C/3h/AC &
Inventive and double tempering T2
A3 1200 C/1 h/AC 660 C/3h/AC
Table 3: Heat Treatment Schedule for Steel C and D
Treatment Purpose Austenitization Tempering
600 C/3h/AC &
High austenitization temperatures, T3 660 C/3h/AC
fast quenching and low tempering A4 1200 C/1 h/oil
Inventive 660 C/3h/AC &
temperatures (double or single quenching T2
tempering) 660 C/3h/AC
T4 660 C/6h/AC
Plain and notched stress rupture tests in the temperature range 600-675 C were
carried out according to BS EN 10291:2000. Stress rupture results for Steel 92
(A)
are shown in Table 4 and 5 (b stands for broken and ub for unbroken samples).
Tests
were carried out at 600-675 C at stresses designed to give aim lives between
1,000
and 30,000 hours. In most cases for the standard heat treatment the alloy
failed
close to the aim lives. However at longer aim durations and higher
temperatures the
actual lives were shorter than intended.
The inventive heat treatment gave dramatic improvements in creep life compared
with material given the standard treatment. Aim lives were generally exceeded
by
significant margins, although the improvement decreased with increasing
temperature.
CA 02801637 2012-12-05
WO 2011/154515 PCT/EP2011/059657
- 10-
Table 4: Plain Stress Rupture Properties for Steel 92 (A)
Temp. Stress Aim Inventive HT Standard HT
( C) (MPa) (h) (h) (h)
600 187 1000 - 856 b
172 3000 50,584 b 2722 b
156 10,000 70,184 ub 8222 b
140 30,000 - 11,992 b
625 140 3000 16,967 b -
122 10,000 26,280 b -
107 30,000 41,808 b -
650 122 1000 - 807 b
110 3000 6179 b 1734 b
92 10,000 12,089 b 6109 b
78 30,000 - 14,729 b
675 81 3000 4192 b 2456 b
66 10,000 10,625 b -
Table 5: Notched Stress Rupture Properties for Steel 92 (A)
Temp. Stress Aim Plain Notched
( C) (MPa) (h) (h) (h)
600 172 3000 50,584 b 32,881 b
156 10,000 70 184 u b 29,401 u b
650 110 3000 6179 b 6305 b
92 10,000 12,089 b 16,133 b
The stress rupture data for the high nitrogen Steel B at 600 and 650 C are
shown in
Table 6 for material given the low temperature temper and at two solution
treatment
temperatures of 1150 and 1200 C. The performance is compared with the
conventional composition given the low temperature temper and the standard
treatment.
Table 6: Plain Stress Rupture Properties for Steels B, C and D - compared to
Steel A
Steel B Steel B Steel C Steel C Steel C Steel D Steel D
Steel A Steel A with with with with with with with
Aim with with inventive inventive inventive inventive inventive inventive
inventive
Temp. Stress life standard inventive
( C) (MPa) (h) HT (h) HT (h) HT HT HT HT HT HT HT
(h) (h) (h) (h) (h) (h) (h)
Al-T1 A2-T2 A2-T2 A3-T2 A4-T2 A4-T3 A4-T4 A4-T2 A4-T3
187 1000 856 b - 29,401 ub 38,795 b - - - - -
600 172 3000 2722 b 50,584 b 30,452 b 39,649 ub - - - - -
156 10,000 8222 b 70,184 ub 29,401 ub 39,649 ub - - - - -
122 1000 807 b 3378 up 7544 b 7973 b 5508 up 4868 b 5156 up 4001 b 3861 b
650 110 3000 1734 b 6179 b 10,255 b 17,118 b - - - - -
92 10,000 6109 b 12,089 b 16,920 b 27,908 b - - - - -
575 110 363 160 b 906 b - - 1655 b 1187 b 1409 b 2516 b 2858 b
The high nitrogen-low carbon Steel B is outperforming the conventional
composition
of Steel 92 (Steel A). The best properties were obtained using the higher
solution
treatment temperature of 1200 C. At 650 C and 92 MPa Steel A with the
inventive
heat treatment failed after 12,089 hours for a 10,000 hour aim. The high
nitrogen
CA 02801637 2012-12-05
WO 2011/154515 PCT/EP2011/059657
-11-
Steel B given a 1200 C solution treatment failed after 27,908 hours, meaning
that
Steel B outperforms Steel A by a factor 2.
For Steel 92 (Steel A) with conventional heat treatment the tests at 675 C
failed far
below the aim life. Better performance was achieved for the same steel but
with
higher solution treatment and lower tempering temperatures. However, the best
creep lives at this temperature and stress level were achieved for Steels C
and D
(Steel B was not tested) again with higher solution treatment and lower
tempering
temperatures than conventional heat treatment. The aim life was exceeded by up
to
a factor of seven. The aim life at 650 C/122 MPa was also exceeded by a factor
of
seven-eigth in the case of Steel B.
The stress rupture tests at 650 C for novel steels with inventive heat
treatments
demonstrate excellent results, especially for Steel B and C.
The results of a test programme for Steel A to assess notched properties
reveal that
the notched properties have been similar to, or in excess of, the plain
properties. The
effect of PWHT is that PWHT (Table 7) has little effect on the conventionally
heat
treated material as the tempering and PWHT temperatures are similar. For both
steels the inventive heat treated material has exceeded the base values for
normally
heat treated materials. The results show that the 740 C PWHT reduced the creep
life
to about between 63% and 76% of the non-PWHT value but still showed an
improvement of 170% compared with the standard heat treatment in the case of
Steel A.
Table 7: Effect of PWHT on Stress Rupture Properties for Steel A and B
Material Heat treatment 600 C / 172MPa 650 C / 11OMPa
+ PWHT - PWHT + PWHT - PWHT
Steel A 1060/780 C 3059 b 2722 b 1714 b 1734 b
Steel A 1150/660/660 C 19,041 ub 50,584 b 4688 b 6179 b
Steel B 1150/660/660 C 19,082 ub 30,452 b 7868 b 10,255 b
Steel B 1200/660/660 C 19,082 ub 39,649 ub 10,778 ub 17,118 b
The microstructure of the samples in as tempered condition has been
characterised
and the results thereof are presented in the figures 1 to 4.
Figure 1 and 2 show examples of the microstructure comprising very fine
M2(C,N) and
M(C,N) precipitates formed in Steel C A4-T2 sample (solutionised at 1200 C,
quenched in oil and tempered at 660 C/3h AC + 660 C/3h AC).
Figure la: TEM Bright Field (BF) micrograph of Steel C A4-T2 sample foil
showing
martensite lath interior precipitates.
Figure 1b: TEM Dark Field (DF) micrograph of Steel C A4-T2 with M2X
precipitates
taken in diffraction spot "df1" of Figure 1c.
CA 02801637 2012-12-05
WO 2011/154515 PCT/EP2011/059657
- 12-
Figure 1c: SAD pattern from wide sample area of Figure la corresponds to BCC
matrix, a=0.288 nm and M2X crystal type (s-Fe2N type, S.G. P-31m), a=0.492 nm,
c=0.447 nm.
Figure 1d: Indexation of SAD as M2X crystal type (8-Fe2N type, S.G. P-31m),
a=0.492 nm, c=0.447 nm; zone axis [236]. Lattice spacing indicated in
Angstroms.
Figure le: Indexation of SAD of Figure 1c as BCC lattice, a=0.286 nm; zone
axis
[101]. Lattice spacing indicated in Angstroms.
Figure 2a: TEM BF micrograph of Steel C A4-T2 with MX precipitates
Figure 2b: SAD pattern from wide sample area of Figure 2a corresponds to BCC
matrix, a=0.288 nm and MX crystal type (VN type, S.G. Fm-3m) a=0.425 nm,
Figure 2c: Indexation of SAD Figure 2b as BCC crystal type, a=0.288 nm; zone
axis
[210]. Lattice spacing indicated in Angstroms.
Figure 2d: Indexation of SAD Figure 2b as MX crystal type, a=0.425 nm; zone
axis
[112]. Lattice spacing indicated in Angstroms.
The example of EDX spectra from M(C,N) precipitate in Steel C A4-T2 is shown
in
Figure 3, where the M component is mainly V, Cr, Ta and some Mo.
The example of EDS spectra from M2(C,N) precipitate in Steel C A4-T3
(solutionised
at 1200 C, quenched in oil and tempered at 600 C/3h AC + 660 C/3h AC) is
presented in Figure 4. In this case the M2(C,N) precipitate has M component
composed mainly of Cr, V, Ta and some Mo.
