Volume 13 - Corrosion Part 12 potx

absorption follows Sievert's law; that is, absorption is proportional to the square root of the partial pressure of nitrogen in
the sintering atmosphere. This nitrogen absorption provides significant strengthening (Fig. 3). Upon completion of
sintering, when the part enters the cooling zone of the furnace, the solubility of nitrogen decreases sharply with
temperature (Fig. 12). As a result, Cr
2
N begins to precipitate at the temperature at which the nitrogen content crosses the
solubility limits. More important, below about 1150 °C (2100 °F), additional nitrogen is absorbed from the sintering
atmosphere, leading to more Cr
2
N precipitation and chromium depletion along the grain boundaries. The net result is
inferior corrosion resistance due to grain-boundary corrosion.

Fig. 12 Solubility of nitrogen in austenitic stainless steel in equilibrium with gaseous nitrogen or Cr
2
N.
Source:
Ref 9
The rate of this detrimental nitrogen absorption increases with decreasing part density and with decreasing dew point. A
high dew point, however, leads to the problem of excessive oxidation. The basic relationship of this phenomenon is
shown in Fig. 13. The data in Fig. 13, which were developed for the bright annealing of stainless steel in dissociated NH
3

atmospheres, show the extent of nitrogen and oxygen absorption as a function of dew point. At high dew points (higher
than about -37 °C, or -35 °F, depending on part size), the rate of oxidation is severe enough to produce a dull surface. At
dew points of about -45 °C (-50 °F) or lower, nitrogen absorption increases so much that the corrosion resistance
deteriorates because of excessive Cr
2
N formation. Thus, optimum bright annealing of austenitic stainless steels must be
done within a narrow dew-point range. Although the authors (Ref 14) caution against applying these findings to sintered
stainless steels based on the unexplained higher nitrogen contents found for their parts sintered in dissociated NH

3
, it
should be noted that such higher nitrogen contents are expected on the basis of known solubility data for nitrogen in type
316L (Fig. 12) considering the differing methods of nitrogen analysis used.

Fig. 13 Safe operating parameters with r
espect to dew point can be developed for a specific set of operating
conditions and quality requirements. The safe zone here is for sintering in an atmosphere of 30% H
2
-70% N
2
at
1035 °C (1900 °F). Source: Ref 14
Chromium nitride sensitization may in some cases be limited to a very shallow surface depth of the part. With very slow
cooling, however, absorption and precipitation proceed toward the interior of the porous part. Figure 14 shows Cr
2
N
precipitates in the grain boundaries of parts that were sintered under conditions that produced nitrogen contents from 55
to 6650 ppm. Increasing nitrogen content correlates with increasing amounts of precipitation and increasing localized
corrosion (Fig. 14). Figure 15 shows the microstructure of a type 316L part that was sintered in dissociated NH
3
and
cooled very slowly. Slow cooling produced a lamellar structure of Cr
2
N and low-chromium austenite of very poor
corrosion resistance.

Fig. 14 Scanning electron micrographs of type 316L stainless steel. (a) Sintered 45 min in 100% H
2
at 1350 °C

(2460 °F); 66 ppm N. (b) Sintered 45 min in 75% H
2
at 1350 °C
(2460 °F); 3100 ppm N. (c) Sintered 45 min
in 25% H
2
at 1350 °C (2460 °F); 4300 ppm N. (d) Sintered 45 min in 25% H
2
at 1150 °C (2100 °F); 6650 ppm
N. The amount of intergranular precipitate increases with nitrogen content. Source: Ref 13

Fig. 15 Micrograph showing the lamellar structure of Cr
2
N and low-
chromium austenite in sintered type 316L
that was slowly cooled in dissociated NH
3
. Etched with Marble's reagent. 700×. Source: Ref 9
Corrosion resistance data for sintered types 304L and 316L in NaCl solutions and in 10% HNO
3
, reflecting the effect of
Cr
2
N precipitation, are shown in Fig. 8, 16, and 17. Figures 8 and 16 show that a higher sintering temperature, fast
cooling rates (75 °C/min, or 135 °F/min, versus 8 °C/min, or 14 °F/min), and the use of type 316L rather than type 304L
provide better corrosion resistance. That these measures are beneficial follows directly from the austenite-nitrogen phase
diagram (Fig. 12).


Fig. 16

Effect of composition, cooling rate, and sintering temperature on corrosion resistance of type 304L and
tin-modified type 304L P/M stainless steels (sintered density: 6.5 g/cm
3
; sintering atmosphere: dissociated
NH
3
) in 5% aqueous NaCl. B rating in
dicates <1% of specimen surface stained. Parenthetical values designate
sintering temperature and cooling rate. Source: Ref 19

Fig. 17 Weight loss of austenitic stainless steel in 10% aqueous HNO
3
as a function of absorbed nitrogen
content. Curve A: sintered in dissociated NH
3
at 1150 °C (2100 °F) with a dew point of -43 °C (-
45 °F).
Density: 5.10 to 5.20 g/cm
3
. Curve B: sintered in various atmospheres with different dew points.
Density: 5.2
to 5.8 g/cm
3
. Source: Ref 9
Figure 18 shows potentiodynamic corrosion curves for sintered type 316L in 10% HNO
3
. The corrosion current density in
the passive range increases and the corrosion potential decreases under conditions that promote Cr
2
N precipitation, that is,

lower sintering temperature, slower cooling rate, and high nitrogen concentration of the sintering atmosphere. Figure 19 is
similar to Fig. 18 except that internal rather than external cross sections were used. The significantly lower corrosion
currents of the internal surface confirm that Cr
2
N precipitation is most severe on the surface of a sintered part.

Fig. 18 Forward scan
potentiodynamic corrosion curves for external surfaces of three sintered type 316L
stainless steel samples in 10% HNO
3
at 25 °C (75 °F). Note the increasing corrosion currents in the 0- to 1-
V
range and the decreasing corrosion potential with nitrogen add
itions to the atmosphere, slow cooling, and lower
sintering temperatures. SCE, saturated calomel electrode. See also Fig. 19. Source: Ref 20

Fig. 19
Forward scan potentiodynamic corrosion curves of the internal microstructure (metallographic cross
section) for type 316L stainless steel samples sintered in 25% H
2
. Corrosion susceptibility in 10% HNO
3
at 25
°C (75 °F) increases with a lower sintering temperature and slow cooling. Cr
2
N precipitation is most severe on
the surface of a sintered part. See also Fig. 18. Source: Ref 20
Recently developed tin-containing grades of type 304L (Table 2) and 316L stainless steels have shown less sensitivity to
nitride precipitation and correspondingly improved corrosion resistance (Fig. 8 and 16). The beneficial effect of tin has
been confirmed in several studies (Ref 9, 10, 16, 19, 20, 21, 22) and has been attributed to an enrichment of the surfaces

of both the water-atomized powder and the sintered part with tin, presumably as a result of the low solubility of tin in
solid stainless steel (Ref 10). Tin may also form stable acid-resistant passive films in a crevice and may cause cathodic
surface poisoning, but its major beneficial effect is believed to lie in its formation of an effective barrier to nitrogen (and
possibly also to oxygen) diffusion. This reduces the rate at which nitrogen is absorbed on the surface of the sintered part
as it enters the cooling zone of the furnace. Auger composition depth profiles of regular and (1.5%) tin-containing type
316L parts sintered in dissociated NH
3
(Fig. 20) show that the presence of tin on the surface effectively suppresses
nitrogen absorption. In addition, on the basis of potentiodynamic polarization tests in 10% HNO
3
and 5 N H
2
SO
4
,
improvement in corrosion resistance has also been reported due to the presence of tin (Ref 20, 22).

Fig. 20 Auger composition depth profiles of P/M type 316L stainless steel parts sintered in dissociated NH
3
at
1175 °C (2150 °F). (a) Type 316 L. (b) Tin-modified type 316L. Source: Ref 9
The effect of oxygen on the corrosion resistance of sintered stainless steels is probably the most complex and least
understood variable for several reasons. First, commercial water-atomized compactible stainless steel powders have
typical oxygen contents of about 2000 ppm or more. Although much of this oxygen resides on the surfaces of individual
powder particles as oxidized silicon (Fig. 21a), the exact nature and distribution of the oxides depends on atomizing
conditions. Second, with typical industrial sintering practice, the reduction of these oxides remains incomplete and
depends on many process parameters. Lastly, as a sintered part enters the cooling zone of the furnace, certain elements
will oxidize upon reaching the temperature for the oxide-metal equilibrium of the high oxygen affinity elements (Fig. 22).
Thus, a sintered part still reflects the history of its powder-making process, compaction, and sintering. Figure 21(b) shows
the Auger composition depth profile of a type 316L part after sintering in hydrogen at 1260 °C (2300 °F). It is apparent

that much of the oxidized silicon present in the green part has become reduced and that severely depleted chromium has
been replenished.

Fig. 21 Auger composition depth profiles of P/M type 316L stainless steel. (a) Green part. (b) Sintered part


Fig. 22 Redox curves for chromium and silicon alone and in solution. Source: Ref 9

An empirical correlation between the saltwater corrosion resistance of sintered type 316L and the oxygen content of the
sintered parts suggests that sintering conditions resulting in lower oxygen contents provide better corrosion resistance
(Fig. 23). With excessive dew points (>-34 °C, or -30 °F), the oxygen content of a sintered part may increase
considerably. The microstructure (Fig. 24) of such a part shows a lack of particle bonding (compare with Fig. 10a for low
oxygen content), and its mechanical strength and corrosion resistance are both inferior.

Fig. 23 Effect of oxygen content on corrosion resistance of sintered type 316L and tin-
modified type 316L
(sintered density: 6.65 g/cm
3
; cooling rate: 75 °C/min, or 135 °F/min). Parenthetical values are sinteri
ng
temperature (°C), dewpoint (°C), and nitrogen content (ppm), respectively. Time indicates when 50% of
specimens showed first sign of corrosion in 5% aqueous NaCl. Source: Ref 10

Fig. 24 Microstructure of type 316L stainless steel sintered i
n a high dew point atmosphere. Oxygen content:
5100 ppm; sintered density: 7.5 g/cm
3
. Etched with Marble's reagent. 200×. Source: Ref 9
For optimum corrosion resistance, it appears that the following precautions are beneficial:
• Use of a powder with low oxygen content

• Sintering conditions that ensure a high degree of oxide removal
• Fast cooling through the high-temperature range after sintering
Cooling in a hydrogen atmosphere should be done with a water vapor content of less than 50 ppm (Ref 12). Cooling in a
nitrogen-containing atmosphere should be done with a dew point between about -37 and -45 °C (-35 and -50 °F) (Ref 14).
Effect of Sintered Density. Applications of sintered stainless steels cover a wide density spectrum. Low densities of
about 5 g/cm
3
may be typical of filters, but densities of 6.5 g/cm
3
or greater are typical of structural parts. It is therefore of
interest to know the effect of density on corrosion resistance. Corrosion studies of sintered austenitic stainless steels have
shown that the corrosion resistance improves significantly with increasing density in acidic environments, such as dilute
H
2
SO
4
, HCl, and HNO
3
. Figure 25 illustrates this behavior for three austenitic stainless steels (18Cr-11Ni to 18Cr-14Ni)
that were vacuum sintered 1 h at 1150 and 1250 °C (2100 and 2280 °F) and tested in boiling 40% HNO
3
.

Fig. 25
Relationship between sintered density and weight decrease of three austenitic stainless steels in 40%
HNO
3
solution. Source: Ref 23
For saline solutions, some investigations have found the effect of increasing density to be beneficial (Ref 20) while others
have found it to be detrimental (Ref 10, 16, 23). This lack of agreement is perhaps not surprising considering that

concentration changes in several of the critical variables, such as oxygen, carbon, and nitrogen, also depend on the density
of a part. It should be noted, however, that the positive relationship between density and corrosion resistance was derived
from short-term potentiodynamic polarization measurements (Ref 20), whereas the negative relationships were all derived
from longer-term salt immersion tests.
Table 6 summarizes recent results on the effect of density on the salt corrosion resistance immersion in 5% aqueous
NaCl) of vacuum-sintered type 316L parts. Unlike sintering in a reducing atmosphere, vacuum sintering does not lower
the oxygen content with decreasing density. Thus, an improvement in corrosion resistance with decreasing density, as
shown in Table 6, should not be attributed to a lower oxygen content, but is perhaps better explained in terms of reduced
crevice corrosion as a result of the improved circulation of the corrodent through large pores (Ref 9). The average pore
diameters of the parts pressed at the lower compacting pressures (Table 6), as measured by mercury porosimetry, were 60
to 80% larger than the size of pores of the high-density parts. The standard deviations of the pore size distributions were
similar and were around 2. Therefore, sintered stainless steel parts with densities from about 60 to 90% of theoretical have
average pore sizes from about 10 to 2 or 3 m that are likely to affect the circulation of the corrodent and thus its
resistance to crevice corrosion.




Table 6 Effect of density on corrosion performance of vacuum-sintered type 316L stainless steel
Compacting
pressure
Sintering
temperature
Corrosion rating
(b)
for four specimens, each immersed in 5%
aqueous NaCl, h
MPa

tons/in.

2


°C °F
Sintering

time,
min
Sintered

density,

g/cm
3

Median
pore size

of pore
volume
(a)


1 3 19

27

46

91

140

210

314

380

558

525

A

A

A A A B B B A B B B
A

A

A B B B B B B B B C
A

A

B B B B B B B B C C
276
20 1205

2200

45 5.67 9
A

A

B B B B B B B C C C
A

B

B B C C D . . . . . . . . . . . . . . .
A

B

. .
.
B B C C C D . . . . . . . . .
A

A

B B C C C C D . . . . . . . . .
552
40 1205

2200

45 6.53 5
A

A

B C C C C D . . . . . . . . . . . .
A

A

A A B B B B B B B B
A

A

A A B B B B B B B B
A

A

A A B B B B B B B B
276
20 1315

2400

45 5.86 8
A

A

A A B B B B B B B B
A

A

B B B B B C C C D . . .
A

A

B C C C C D . . . . . . . . . . . .
A

A

B C C C C D . . . . . . . . . . . .
552
40 1315

2400

45 6.57 5
A

A

A A B B B C C D . . . . . .

(a)
Determined by mercury porosimetry.
(b)
A, sample free from any corrosion; B, 1% of surface covered by stain; C, 1 to 25% of surface covered by stain with slight corrosion
product; D, >25% of surface covered by stain with heavy corrosion product.

Effect of Copper Additions to Type 304L. One study found that the corrosion resistance of copper-containing type
304L vacuum-sintered parts (1 h at 1200 °C, or 2190 °F; 88% dense) improved with increasing copper content (Ref 17).
Figure 26 shows the weight loss of the parts kept for 6 h in boiling 5% H
2
SO
4
. Higher nickel content is also beneficial.
Salt spray testing for 24 h with 5% NaCl solution resulted in almost no pitting. The effect of copper in P/M stainless steels
is said to be identical to that observed in cast stainless steels.

Fig. 26 Effect of nickel and copper additions on the corrosion rate of sintered austenitic sta
inless steel
compacts exposed to boiling H
2
SO
4
for 6 h. Relative sintered density is 88%. Source: Ref 17
Oxidation Resistance. Sintered stainless steels are not widely used for elevated-temperature service. Thus,
information on elevated-temperature oxidation resistance is scarce.
Figure 27 shows the weight gain in air at 700 °C (1290 °F) for type 310L stainless steel parts that were vacuum sintered 1
h at 1250 °C (2280 °F) as a function of sintered density (circular plates), mesh size of powder used, and sintering
temperature. This initial weight gain did not always show a parabolic course of oxidation. Within the density range
studied, oxidation increased almost exponentially with decreasing density. Silicon-modified (4.06% Si) type 310L
stainless steel showed weight gains that were less than 50% of those of regular type 310L. The increased oxidation of the

parts made from the finer powder fraction is due to their large internal surface area. Higher sintering temperature and
higher compacting pressure (higher densities) reduce surface porosity and specific pore surface area, thus lessening
internal oxidation through early pore closure. The maximum recommended operating temperature for sintered austenitic
stainless steels is 700 °C (1290 °F).

Fig. 27
Weight gain versus sintered density curves for materials prepared from powders of various particle
sizes. Parts were sintered 1 h at 1250 °C (2280 °F). Source: Ref 24
Higher-Alloyed Stainless Steels. Although the common stainless steel grades used in industry have maximum
chromium and nickel contents of 20 and 14%, respectively (Table 2), higher-alloyed stainless steels have been used in the
past to obtain improved corrosion resistance. Such steels are available from powder procedures. In one investigation, a
high-nickel/chromium/molybdenum austenitic stainless steel P/M material (SS-100) performed comparably to wrought
type 216, 316, or 317 in 16-h salt solution immersion tests (Ref 25).
Other Approaches to Improving the Corrosion Resistance of Sintered Stainless Steels. If the corrosion
resistance of sintered stainless steel parts remains inadequate after composition and process optimization, passivation and
coating treatments are sometimes used. Chemical and thermal passivation treatments for sintered type 316L, effective in
dilute H
2
SO
4
, are described in Ref 26. Chemical passivation with HNO
3
solutions similar to those applied to wrought
stainless steels is not suitable for every material. On the basis of rest potential measurements of sintered type 316L,
thermal passivation by heating the sintered parts for 20 to 30 min in air at temperatures of 400 to 500 °C (750 to 930 °F)
is recommended.
In another study, the corrosion resistance of vacuum-sintered type 304L (6.9 g/cm
3
) in 5% H
2

SO
4
was improved by
activating the parts in a mixture of 13 to 15% HNO
3
, 2% hydrofluoric acid (HF), and 0.3% hydrochloric acid (HCl),
followed by passivation for 30 min in 30% HNO
3
at 70 °C (160 °F) (Ref 27). After testing for 2 h in 5% H
2
SO
4
(Fig. 28),
the passivated specimens showed no weight loss, whereas the as-sintered specimens rapidly lost weight and turned the
solution green. In addition, Ref 28 describes a phosphate-base passivating treatment for sintered stainless steels that is
effective in acetic acid.

Fig. 28 Relationship between weight loss and corrosion time of vacuum-
sintered type 304L stainless steel in
5% H
2
SO
4

Improvement of the corrosion resistance of sintered stainless steel through the chemical vapor deposition (CVD) or
chromium is discussed in Ref 29. The chemical vapor deposition of chromium onto sintered stainless steel parts was
applied by pack cementation. Considerable infilling of the pores with chromium takes place; 50- m thick pores with
diameters of up to 50 m may become sealed. Immersion of coated and uncoated specimens in 5 and 10% H
2
SO

4

solutions for 168 h at room temperature showed significant attack of the uncoated specimens and no noticeable attack of
the coated specimens. Electrochemical testing in 5% H
2
SO
4
gave similar results, and a 3% salt spray test at room
temperature showed many local sites of corrosion for the uncoated specimen and no corrosion after 250 h for the coated
specimen. Sealing or coating of the pores of a sintered stainless steel part with an organic resin (Ref 25) or with water
glass is sometimes recommended, but performance data proving the effectiveness of this treatment are lacking.
Fully Dense P/M Stainless Steels
In the fully dense category of P/M stainless steels, parts made from water-atomized powders must be distinguished from
those made from inert gas atomized powders.
Water-atomized powders, because of their irregular particle shape, are cold compactible and permit the pressing of
complex parts, which, at temperatures approaching the melting point of the material, can be sintered to nearly full density.
However, water-atomized stainless steel powders typically have oxygen contents of 2000 ppm or more, and sintering to
full density usually does not reduce the oxygen content to the low level of the corresponding ingot material. The
commercial production of such parts is still in its infancy, and corrosion data are not yet available.
Inert gas or centrifugally atomized powders are spherical and noncompactible. They have low oxygen contents
(about 50 to 200 ppm) and are consolidated to full density by such process as hot isostatic pressing (HIP), hot forging, and
extrusion.
One company has manufactured seamless stainless steel tubes from gas-atomized powder since 1980. The P/M method is
said to offer a competitive alternative to conventional production methods due to:
• Efficient use of raw materials
• Low energy consumption
• Short total production time
• High flexibility (less material in process; short delivery times)
• The ability to make difficult compositions (Ref 30)
The process consist of cold isostatic compaction of the encapsulated nitrogen-atomized powder, followed by heating to

the extrusion temperature and hot extrusion. The capsule material is removed by decladding. Standard grades include
most of the common austenitic stainless steels as well as some special austenitic, ferritic-austenitic, and ferritic stainless
steels, together with nickel-base alloys.
In comparison to conventional material, the extruded P/M products possess a more homogeneous structure with reduced
microsegregation due to the rapid cooling of the powder particles. Also, the grain size is somewhat finer, slag inclusions
(particularly sulfides) are smaller, and the nitrogen content is somewhat higher (900 versus 500 ppm for wrought type
316).
Mechanical Properties and Corrosion Resistance. Attributed to the above differences are slightly higher yield
and tensile strength (Table 7) without a loss in elongation. Mechanical properties at elevated temperatures are practically
identical to conventionally produced materials. The impact toughness of the P/M material, although good, is lower than
that of conventional material when tested in the longitudinal direction. Creep strength is similar to that of conventional
material.
Table 7 Typical mechanical properties of cold-wo
rked and annealed stainless steel tubes extruded from
powders
Yield strength,

0.2% offset
Tensile
strength
Grade Type
(a)


Number of

samples
MPa ksi MPa

ksi

Elongation,

%
C 84 302 44 582 84 57
Type 304L

P/M 18 325 47 609 88 58
C 133 321 46 600 87 57
Type 304
P/M 72 350 51 660 96 55
C 90 319 46 604 88 53
Type 316L

P/M 128 336 49 632 92 52
C 134 306 44 584 85 54
Type 316
P/M 125 346 50 649 94 51
Type 904L

C 49 334 48 651 94 45
P/M 112 382 55 681 99 43
Source: Ref 30
(a)
C, conventional production; P/M, powder metallurgy.

No difference between P/M and conventional material has been found regarding the resistance to intergranular corrosion
according to practice C and practice E of ASTM A 262 (Ref 31). As shown in Fig. 29, the resistance to pitting attack, as
measured by the pitting corrosion breakthrough potential, is superior for several P/M grades compared to the
corresponding conventional grades. Table 8 gives the general and selective corrosion information from tests according to

ASTM A 262, practice C (Ref 31) for two austenitic P/M grades. The improved corrosion resistance of the P/M grades is
attributed to their lower segregation rate, their finer and more uniform distribution of inclusions, and their finer grain size.
Table 8 Huey test (ASTM A 262, practice C) corrosion data for two P/M extruded stainless steels

Corrosion rate, m/48 h

Selective attack, m

Grade Number of

samples
Average Specific Average

Specific
Type 725LN

14 0.57-0.69 1.5 max <50 100 max

Type 724L
14 1.48-1.79 3.3 max <30 200 max





Nominal composition, % Alloy
C(max)

Fe Cr Ni Mo

Others

303
0.10 rem

18 8.5 . . . . . .
304
0.05 rem

18.5

8.5 . . . . . .
329
0.05 rem

26 5 1.5 . . .
316
0.05 rem

17 11.5

2.2 . . .
44LK
0.03 rem

25 6 1.6 . . .
904L
0.02 rem

20 25 4.5 Cu

984LN

0.05 rem

20 33 2.2 Cu,N


Fig. 29 Comparison of pitting resistance of P/M and conventional stainless steels. Source: Ref 30

The enhancement of the corrosion resistance of stainless steel parts made from rapidly solidified powders has been
confirmed by several investigators. For example, the significantly superior oxidation resistance of type 303 stainless steel,
made by extrusion of rapidly solidified powder, was attributed to the elevated-temperature grain growth inhibiting effect
of uniformly dispersed manganese sulfide (MnS) particles (Ref 32). Figure 30 shows that this material maintains its good
corrosion performance in aqueous environments, and potentiodynamic polarization curves in 1 M H
2
SO
4
indicate that the
P/M material exhibits the lowest corrosion rate at the corrosion potential. Finally, although wrought type 303 was highly
susceptible to pitting, the P/M alloy showed no obvious pits on the surface and only a low pit density within the material.
The pits were related to the presence of sulfide stringers in the wrought material, from which it was concluded that P/M
steels with lower sulfur contents and with spherical sulfide morphology, such as type 304 and 316, might exhibit
improved pitting resistance.



Fig. 30
Potentiodynamic polarization curves for conventional type 303 and 304 stainless steels and for rapidly
solidified type 303 in deaerated 1 M H
2

SO
4
at 30 °C (85 °F). Source: Ref 33
Injection molding technology (see the article "Powder Injection Molding" in Powder Metal Technologies and
Applications, Volume 7 of the ASM Handbook) is currently used to manufacture small and nearly fully dense stainless
steel parts from the powders. However, information on the corrosion performance of such parts is unavailable.
P/M Superalloys
Development of P/M superalloys began in the 1960s with the search by the aerospace industry (and later the electric
power industry) for stronger high-temperature alloys in order to operate engines at higher temperatures and thus improve
fuel efficiency. Figure 31 illustrates the great advances achieved since the 1940s by the introduction of new processes and
alloys, such as vacuum melting, directional solidification of eutectics, development of alloys with high volume fractions
of ' phase, and P/M processing with and without oxide dispersions.

Fig. 31 Trends in alloy processing and development. Source: Ref 34
Initially, lower production costs were a major objective in exploring the P/M approach. Figure 32 illustrates the material
savings possible with two different P/M methods due to their near net shape capabilities. Later, specific advantages linked
on the P/M approach, such as the use of more complex and greater volume fractions of dispersoids, reduced segregation,
and improved workability, led to the development of stronger alloys and to the use of these alloys not only in turbine
disks but also in the higher-temperature turbine blades.

Fig. 32 Processing sequence in the production of jet engine compressor disks
Much effort is currently being directed toward reducing the cost of consolidating superalloy powders, particularly of
oxide dispersion strengthened (ODS) superalloys, through the development of suitable forging techniques (Ref 34).
Efforts are also underway to exploit the advantage of microcrystallinity and extended solid solutions of rapid
solidification technology.
Uses and State of Commercialization. P/M superalloys were first used in military engines in the mid-1970s. Table
9 summarizes the uses of P/M superalloys in terms of components, engine use, and reasons for using P/M technology.
Other uses of superalloys include nuclear reactors, heat exchangers, furnaces, sour gas well equipment, and other high-
temperature applications.

Table 9 Aerospace applications of P/M superalloys
Reasons for using P/M
technology
P/M superalloy Component Engine Aircraft/
manufacturer
Cost
reduction

Improved

properties

IN-100
Turbine disks, seals, spacers F-100 Pratt &
Whitney
X X
René 95
Turbine disks, cooling plate T-700 Helicopter/G.E. . . . . . .
René 95
Turbine disks, compressor shaft F-404 F-18 Fighter X . . .
René 95
Vane F-404 G.E. . . . . . .
René 95
High-pressure turbine blade retainer, disks, forward
outer seals
F-101 . . . X . . .

Astroloy
High-pressure turbine disks JT8D-17R
Turbofan
. . . X . . .
Merl 76
Turbine disks Turbofan . . . X X
Inconel MA-754
Turbine nozzle vane F-404 F-18 Fighter . . . X
Inconel MA-754
High-and low-pressure turbine vanes Selected
engines
. . . . . . X
Stellite 31
Turbine blade dampers TF 30-P100 USAF F-111F X . . .
Inconel MA-
6000E
Turbine blades TFE 731 . . . . . . X
Source: Ref 35
Manufacturing of P/M Superalloys. An important prerequisite for making P/M superalloys that possess reliable
dynamic properties is the use of clean powders. Years of intensive work were spent in identifying and controlling the
problems related to unclean powders. Today, argon and vacuum (also known as soluble gas process) atomization, as well
as atomization by the rotating electrode process, are known to be suitable for producing powders with the required low
oxygen content and low degree of contamination (details on these processes are available in the article "Atomization" in
Powder Metal Technologies and Applications, Volume 7 of the ASM Handbook). The so-called prior particle boundary
(PPB) problem, that is, the presence of carbides segregated at PPBs, was solved through the development of low-carbon
alloys. Special equipment is used for removing ceramic particles and particles containing entrapped argon. Some of these
problems are minimized or avoided in ODS alloys made by mechanical alloying. In mechanical alloying, elemental and
master alloy powders as well as refractory compounds are mechanically alloyed by high-energy milling (Ref 36, 37).
Two established powder consolidation techniques for P/M superalloys are hot isostatic pressing (HIP) and isothermal
forging. Figure 33 illustrates schematically the steps of the P/M processes in comparison to conventional processing. Both

P/M methods permit the manufacture of so-called near net shape parts with attendant improved material use and reduced
machining costs. Powder metallurgy forging exploits the improved forgeability deriving from the higher incipient melting
temperature and reduced grain size of P/M material. Hot compaction by extrusion leads to very fine grain size, improved
hot ductility, and superplasticity.

Fig. 33
Comparison of conventional processing and P/M processing for the fabrication of superalloy disks.
Source: Ref 35
Depending on the application of a superalloy part, the powder consolidation process can be controlled in order to yield
either a fine or a coarse grain size. Fine grain size is preferred for intermediate temperatures (up to about 700 °C, or 1290
°F) because of its higher strength and ductility at these temperatures. For high-temperature blade and vane applications,
however, a large grain size (ASTM 1 to 2) provides superior creep strength due to reduced grain-boundary sliding. Grain
coarsening of ODS alloys is achieved through special heat treatments after consolidation (Ref 34).
Compositions and Properties. Table 10 shows the compositions of the best known P/M superalloys. Many have the
same compositions as cast alloys but are manufactured similarly to wrought alloys. The important P/M superalloys IN-
100, René 95, and Astroloy were adapted to the P/M process by reducing their carbon content and by adding stable
carbide formers to eliminate the problem of PPB carbides. To facilitate HIP, alloy compositions were modified to increase
the temperature gap between the ' solvus (above which HIP has to be carried out for increasing grain size) and the
solidus temperature.
Table 10 Nominal compositions of several P/M superalloys
Composition, % Alloy
C Cr Mo

W Ta

Ti Nb

Co Al Hf

Zr B Ni Fe V Y

2
O
3


IN-100
0.07

12.5

3.2 . . .

. . .

4.3

. . .

18.5

5.0

. . .

0.04

0.02 rem . . . 0.75

. . .
René 95

0.07

13.0

3.5 3.5

. . .

2.5

3.5

8.0 3.5

. . .

0.05

0.01 rem . . . . . . . . .
MERL 76
0.02

12.4

3.2 . . .

. . .

4.3

1.4

18.5

5.0

0.4

0.06

0.02 rem . . . . . . . . .
AF 115
0.05

10.5

2.8 6.0

. . .

3.9

1.7

15.0

3.8

2.0

. . . . . . rem . . . . . . . . .
PA101
0.1 12.5

. . . 4.0

4.0

4.0

. . .

9.0 3.5

1.0

. . . . . . rem . . . . . . . . .
Low-carbon Astroloy

0.04

15.0

5.0 . . .

. . .

3.5

. . .

17.0

4.0

. . .

0.4 0.025

rem . . . . . . . . .
MA 754
0.05

20.0

. . . . . .

. . .

0.5

. . .

. . . 0.3

. . .

. . . . . . rem . . . . . . 0.6
MA 956
. . . 20.0

. . . . . .

. . .

0.5

. . .

. . . 4.5

. . .

. . . . . . . . . rem

. . . 0.5
MA 6000
0.05

15.0

2.0 4.0

2.0

2.5

. . .

. . . 4.5

. . .

0.15

0.01 rem . . . . . . 1.1
Oxidation. Nickel-, cobalt-, and iron-base superalloys use the selective oxidation of aluminum or chromium to develop
oxidation resistance (Ref 38). These alloys are therefore often referred to as aluminum oxide (Al
2
O
3
) or chromium oxide
(Cr
2
O
3
) formers, depending on the composition of the oxide scale that provides protection. Alloy composition, surface
conditions, gas environment, and cracking of the oxide scale affect the selective-oxidation process (Ref 38). Figure 34
shows the development of superalloys in terms of the progress achieved against high-temperature oxidation.