Note: Descriptions are shown in the official language in which they were submitted.
~9~ S
BACKGROUND OF THE INVENTI ON
2 'rhis invention relates to alumina-forming
3 nickel-based austenitic alloys having superior oxidation
4 and carburization resistance, superior creep strength,
goo~ high temperature ductility, and good microstructural
6 stability.
7 Various industrial processes, especially
8 chemical processes, create an insatiable demand for
g alloys which can withstand higher and higher tempèra-
tures and environments deleterious to the alloy. One such
11 deleterious environment is a carburizing environment, the
12 effects of which are known to significantly affect plant
13 performance and efficiency in many indus-trial processes.
14 These effects are evidenced in heat treatment equipment,
ethylene pyrolysis tubing, carbon dioxide and helium-
16 cooled nuclear reactors, coal processing plants, and
17 hydrocarbon reformers.
18 A variety of alloy steels exhibiting both
19 heat and carburization resistance have been developed
for use in pyrolysis furnaces for the thermal decom-
21 position of organic compounds, such as the steam cracking
22 of hydrocarbons. Generally, the pyrolysis furnace con-
23 tains a series of heat-resistant alloy steel tubes in
24 which the reaction occurs. It will be noted that the term
"tube" as used herein also includes fittings, pipes and
26 other parts used to contain carburizing materials.
27 It is well known that alloy steels containing
28 various amounts of nickel, chromium, and silicon, as
29 well as the addition of elements such as tungsten and/or
niobium, are frequently used in high temperature applica-
31 tions. ~owever, after extended use at high temperatures,
32 most, if not all, of these known alloys fail to accomplish
33 all of the above objectives.
3~ The major cause of failure, especially in
pyrolysis tubes, is creep rupture, brought about by the
36 combined effect of thermal stress and carburization.
37 Carburization of such tubes, which is the diffusion of
38 carbon into the alloy steel causing the formation of
:~9~ 35
1 additional carbides principally at the grain boundaries,
2 and the attendant depletion of the matrix in chromium,
3 brings about both a loss of creep strength and embrittle-
4 nent of the grain boundaries. Once the steel has become
embrittled, it is more susceptible to Eailure by creep
6 rupture at high temperatures, brittle fracture at low
7 temperatures, or both, because of ther~al stress.
8 Although progress has been rnade in the develop-
9 ment of alloy steels capable of withstanding the rigors of
high ternperature hydrocarbon environments, there still
11 exists a need in the art for the further development of
12 alloy steels having high temperature properties superior
13 to those known in the art
14 SU~MARY OE` THE INVENTION
I~ accordance with the present invention
16 there is provided alumina-forming nickel~based austenitic
17 alloys having improved oxidation and carburization
18 resistance, improved creep rupture properties, good high
19 temperature ductility, and good microstructural stability.
These alloys can be characterized by the following
21 composition (~ by weight):
22 Cr 20-25
23 Fe 10~15
24 Al 3-6
Hf 1-2
26 W 1.5-3.25
27 Nb 1-2
28 Y n.01-1
29 C 0.2-0.3
Ni balance
31 BRIEF DESCRIPTIO~_OF THE FIGURES
32 Figure 1 is a graph showing the superior creep
33 strellgth and ductility of the alloys of the present
34 invention compared with alloys commercially available for
hic~h temperature service.
36 Figure 2 is a graph which again shows the
37 superior creep and ductility properties of the alloys of
3~ the present invention over those conventionally employed.
~.~.96~3C15
1 The alloys represented on this ~raph were first aged at
2 1000C for 3,000 to 5,000 hours before testing.
3 DETAIL~D DESCRIPTION OF TIIE INVENTION
4 Tl~e alloys accordin~ to the present invention
are particularly adapted to constitute metallic parts to
6 be used on the inside or outside of reforming and steam
7 cracking furnaces wherein there exists an oxidizing and
~ carburizing atmosphere and operating temyeratures in the
9 range of about 900C to about 1100C.
The alloys of the present invention simulta-
11 neously possess the following properties:
12 (a) superior oxidation an~ carburization
13 resistance;
14 (b) good creep stren~th;
(c) good high temperature ductility; and
16 (d) good microstructural stability.
17 The composition of the alloys of the present
18 invention, b~ weight percent based on the total weight of
19 the alloy, can be characterized as follows:
20 Cr 20-25 preferably24-25
21 Fe 10-15 preferably12-1
22 Al 3-6 preferably~-5
23 ~f 1-2 preferably1.5-2
2~ W 1.5-3.25 preferably2.75-3.25
25 Nb 1-2 preferably1.25-1.75
26 Y 0.01-1 preferably0.5-1
27 C 0.2-0.3 preferably0.3
28 The rest being nickel with the usual minimum impurities.
29 In the composition o~ the alloys of the present
invention:
31 Chromium (Cr) and aluminum (Al) are jointly
32 responsible for hi~h temperature oxidation and car-
33 burization resistance of the alloys. Aluminum in the
34 ran~e oE about 3 to 6 wt.~, preferably from about ~ to
5 wt.~, leads to the development of protective A12O3
36 scales on the alloy surface, provided the chromium level
37 is in excess of 20 wt.%. At lower levels of chromium,
38 aluminwn will have a tendency to oxidize internally and a
:~3~ `5
protective A12O3 scale will not develop on the alloy
2 surface. Chromium levels in excess of 25 wt. ~ will lead
3 to the precipitation of alpha chromium and sigma phases
4 which lead to microstructural instability. A12O3 scales
5 have been found to be more stable than Cr2O3 scales
6 (which forms in Al free alloys) at temperatures in
7 excess of about 1050C.
8 Chromium also provides strength by virtue of its
g presence in solid solution and the formation of chromium
10 carbide particles.
11 Carbon (C) provides strength at elevated tem-
12 peratures in the presence of carbide forming elements
13 through the formation of finely dispersed alloy carbides
1~1 in the matrix and discontinuous blocky carbides in the
grain boundaries. The latter inhibit grain boundary
16 sliding and thereby constrain rnatrix deformation at a
17 carbon level of 0.2-0.3. Higher levels of carbon promote
18 the formation of a continuous layer of carbide at grain
19 boundaries which serve as an easy path for crack propa-
gation and thus impart poor ductility.
21 Hafnium (Hf) additions result in the formation
22 of highly stable hafnium carbides (HfC). These form in
23 preference to chromium carbides during solidification and
24 precipitate as discrete particles at grain boundaries.
This process removes carbon from solution and suppresses
26 the precipitation of chromium carbides and thereby pro-
27 motes the formation of discrete particles of haEnium
28 carbides in preference to the continuous carbide boundary
29 film formed in the absence of hafnium.
Yttrium (Y) in levels less than about 1 wt.9~
31 have been shown to significantly improve the adherence of
32 '91203 scales; therefore, a Y content of about 0.01 to
33 1 wt.%, preferably about 0.5 to 1 wt.% is emyloyed for use
34 in the instant alloys.
Tungsten (W) contributes to solid solution
36 strengthening at high temperatures and 1.5 to 3.25 wt.%,
37 preEerably 2.75 to 3.25 wt.~ is employed in the alloys of
38 the present invention.
~g~ 5
-- 5
Niobium (`l~b), when present in the alloys of
2 the present invention, will form fine niobium carbide
3 precipitates on dislocations in the alloy structure
4 and contributes to the strength of the alloy. Niobium
levels of about 1 to 2 wt.% are suitable for use herein,
6 preferred is a niobium content of about 1.25 to 1.75
7 wt.%.
a Nickel constitutes the balance of the alloys
9 with residual impurities at as low a concentration as
possible.
11 The following examples serve to describe,
12 more fully, the present invention. It is understood
13 that these examples in no way serve to limit the true
14 scope of this invention, but rather, are presented
for illustrative purposes.
16 _omparative Example A
17 A coupon measur ing 2 cm x 1 cm x 5 mm was
18 taken from a cast tube comprised of 0.55 wt.96 C, 2.31 wt.%
19 5i, 1.21 wt.% Mn, 29.75 wt.% Cr, 28.70 wt.~ Ni, balance
Fe. The coupon was pack carburized, that is, placed in a
21 carbon bed having access to air, at 1100C for 72 hours.
22 The coupon was then nickel-plated, cross-sectioned,
23 polished, and examined under a scanning electron micro-
24 scope. It was observed that carburization occurred
throughout the coupon.
26 Example 1
27 A coupon havin~ the same dimensions as that of
28 the above Comparative Example was taken from a cast tube
29 comprised of 23.û wt.% Cr, 12.2 wt.~ Fe, 4.8 wt.% Al,
1.23 wt.g Hf, 2.85 wt.~ W, 1.68 wt.% Nb, 0.48 wt.% Y,
31 0.2 wt.~6 C, and the balance Ni. The coupon was pack
32 carburized at ]100C for 72 hours. The coupon was then
33 prepared and analyzed as in the above Comparative Example
3~ and it was found that a protective A12o3 scale had
formed which enabled the coupon to resist carburization.
36 Exan~
37 A COUpOIl of the alloy of Example 1 above was
38 oxidized in air at 1100C for 100 hours. The coupon was
3~i8~5
- 6 -
1 then analyzed and it was found that a protective layer of
2 A12O3 had formed on its surface, and that some internal
3 A12O3 stringes were also present. At the chromium and
4 aluminum levels claimed herein, the alloys of the present
invention are selectively oxidized to form A12O3 scales
6 which resist further oxidation - thus evidencing the
7 oxidation resistance of the instantly claimed alloys.
8 Example 3
9 The microstructural stability of the alloy of
Example 1 above was determined by studying samples of
11 the alloy after one sample was e~posed to air at 1100C
l2 for 1000 hours and another sample was exposed to air at
13 1175C for 100 hours. Both samples were found to be
14 structurally stable, that is, the grain structure and
precipitates remained unchanged throughout the experimentO
16 Comparative Examples B and C and Example 4
17 Three alloy samples, shown in Table I below,
18 were prepared and tested for creep strength and ductility
19 at 1000C and 3000 psi for up to about 700 hours~ The
results were recorded and are illustrated in Figure 1
21 herein which shows strain-inch~inch versus time in hours.
22 It is evidenced by this Figure 1 that the alloy of the
23 instant invention is superior to the other two alloys
24 which are representative of those alloys conventionally
ernployed at elevated temperatures.
26 TABLE I
27 Comp. Ex. BComp. Ex. C Ex. 4
28 C - 0.11 0.43 0.2
29 Si - 0.88 1.1~ _
30 Cr - 24.5 24.0 23.8
31 Ni - 38.7 41.8 balance
32 Fe - 33 30.7 12.2
33 Mn - 1.05 1.45
34 ~l 1.47
35 Nb - 1.49 1.68
36 W _ 2.85
37 Y 0.48
38 Hf 1.23
___
39 Al _ 4.8
l Samples of the above alloys were first aged
2 before being tested at 1000C and 3000 psi. The alloy
3 corresponding to Comparative B and C were aged at 1000C
4 for 5000 hours whereas the alloy corresyonding to E~ample
5 4 was aged at 1000C or 3000 hours. The results, as
6 illustrated in Figure 2, again demonstrate the superior
7 creep strengt}l and ductility of the instantly claimed
8 alloys.
