Volume 07 - Powder Metal Technologies and Applications Part 8 pdf

Fig. 16 Redox curves for chromium and silicon in 316L in H
2
at atmospheric pressure. Source: Ref 17

What happens after sintering, that is, during cooling, is even more important. Figure 16 shows two scenarios. For both
scenarios, the sintering temperature is 1200 °C (2192 °F). In the first scenario, the dew point of the sintering atmosphere
is -40 °C (-40 °F); in the second scenario, it is -60 °C (-76 °F). In the first scenario, the stainless steel part crosses the
SiO
2
/Si redox curve (Fig. 16) upon cooling at 1070 °C (1960 °F). At this temperature, the rate of oxidation of silicon is
quite rapid and, therefore, rapid cooling is necessary to prevent or minimize, the formation of silicon oxides on the
surface of the stainless steel part. Figure 17 shows a scanning electron microscopy (SEM) of such oxide precipitates for
316L. These precipitates do not cause the part to discolor, and they are visible only under a microscope. When tested in
aqueous FeCl
3
, in accordance with ASTM G 46 (20 °C), such parts exhibit inferior corrosion resistance due to pitting.
Higher-alloyed stainless steel powder such as SS-100 (20Cr-17Ni-0.8Si-5Mo) appear to be more immune to this type of
corrosion (Ref 10).

Fig. 17 Spheroidal silicon oxide particles formed on 316L part on cooling

With the second scenario in Fig. 16, the lower dew point of -60 °C (-76 °F) causes the parts to cross the SiO
2
/Si redox
curve in the cooling zone of the furnace at the much lower temperature of 890 °C (1632 °F). At that temperature, the
rate of silicon oxidation either is very slow, or any oxides formed at that temperature have little effect on the corrosion
resistance of the part. Figure 18 shows tentative critical cooling rate-dew point combinations as a function of dew point
for three hydrogen sintered austenitic stainless steels. Upper critical cooling temperatures, that is, the lowest high
temperatures where rapid cooling is to commence, are shown in Fig. 19.

Fig. 18
Effect of cooling rate and dew point upon corrosion resistance (5% aqueous NaCl) of hydrogen sintered
stainless steels. Dashed curves representing
maximum corrosion resistance are tentative. Corrosion resistances
shown in parentheses are percentages of maximum corrosion resistance for given grade and density.

Fig. 19 Upper critical cooling temperature and iso-corrosion resistance curves (5% aqueous NaCl) for H
2

sintered 316L (schematic)
Processing to the right of the cooling rate-dew point curves produces maximum corrosion resistance. Processing to the
left results in rapid deterioration of corrosion resistance as shown schematically in Fig. 19 for 316L. It is clear from these
relationships that for maximum corrosion, resistance, sintering in hydrogen requires very low dew points and/or rapid
cooling after sintering. Mechanical properties of hydrogen sintered stainless steels are given in the article "Powder
Metallurgy Stainless Steels" in this Volume.
Sintering in Vacuum. In the early years of commercial sintering of stainless steel parts, vacuum furnaces were said to
be good alternatives to other types of sintering because of their low consumption of gas. After years of experience, parts
producers learned that in addition to high initial capital cost, vacuum furnaces also were costly to maintain. Nevertheless,
it is clear from the previous section on sintering in hydrogen that with the typical furnaces (belt, pusher, and walking
beam) used presently in the industry, the number one property of stainless steel, superior corrosion resistance, is for most
presently used P/M stainless steels not attainable to a sufficient degree. Therefore, future P/M opportunities that require
excellent corrosion resistance cannot be realized until furnace manufacturers construct furnaces that are capable of lower
dew points and parts producers equip their furnaces with (already available) rapid cooling devices. This is where vacuum
furnaces are used. With a state of the art vacuum furnace, it is much easier to maintain a low dew point and to obtain
rapid cooling than it is with a typical atmosphere furnace. Nevertheless, certain precautions are necessary.
For maximum corrosion resistance of vacuum sintered stainless steel, surface depletion of chromium due to high-vapor
pressure and the presence of original surface oxides must be minimized. Sintering under a partial pressure of nitrogen or
argon of 1500 m of mercury effectively reduces chromium losses. Reference 9 shows that after high temperature
sintering (>1205 °C, or 2200 °F), where chromium losses are more severe, holding the parts prior to cooling for a short
time at a lower temperature, or by increasing the partial pressure of argon to 1 at the lower temperature, significantly

improves corrosion resistance. Both measures allow the parts to replenish (from the interior) surface chromium that was
lost at the high-sintering temperature.
In spite of the high-oxygen contents of water atomized stainless steel powders, vacuum sintered stainless steel parts
usually are bright. This is because some of the surface oxides, during sintering, diffuse into the interior.
In the absence of an external-reducing gas atmosphere, vacuum-sintered stainless steel parts have relatively low carbon
and oxygen contents due to the reaction between carbon and oxygen (or oxides) particularly at high-sintering
temperatures, to form carbon monoxide (Fig. 20) (Ref 18 and 19). There is evidence, however, that the typical amounts of
carbon present in a stainless steel part after delubrication are insufficient for removing most of the original oxide particles
present on the outer surfaces of a part. These unreduced original oxide particles give rise to pitting corrosion. Admixing
small amounts of graphite to the stainless steel powder, overall oxide reduction, particularly at the higher-sintering
temperatures, is greatly enhanced and is also sufficient to reduce surface oxides. Alternatively, a small partial pressure of
H
2
should accomplish the same.

Fig. 20
Oxygen versus carbon contents of vacuum and atmosphere sintered P/M austenitic stainless steels of
varying compositions
When making graphite additions to a stainless steel powder, it should be kept in mind that the carbon content of the
sintered part will increase. Thus, the optimum graphite addition is the maximum addition that produces no chromium
carbide precipitates in the cooling zone of the sintering furnace. It depends, among other factors, on the composition of
the stainless steel, the oxygen content of powder, the sintering temperature, and the cooling rate employed.
Vacuum-sintered stainless steels should be rapidly cooled in a non-oxidizing gas to prevent the formation of deleterious
surface oxides. Cooling in nitrogen will result in the formation and precipitation of chromium nitrides on the surfaces of
the parts. The attendant chromium depletion will cause the parts to have very low-corrosion resistance. Mechanical
properties of vacuum sintered stainless steels are shown in the article "Powder Metallurgy Stainless Steels" in this
Volume.
Sintering in H
2
-N

2
Gas Mixtures. Sintering at 1120 °C (2050 °F) in dissociated ammonia was the most widely used
method for sintering stainless steels in the 1960s and 1970s. Dissociated NH
3
was not only less expensive than H
2
, but it
also increased the strength of the sintered parts, although at some reduction in ductility, to levels comparable to wrought
stainless steels of the same composition, at densities of 85 to 90% of theoretical. The strengthening is the result of
nitrogen absorption during sintering. The amount of nitrogen absorbed follows known phase equilibria in accordance with
Sievert's law, that is, nitrogen absorption is proportional to the square root of the partial pressure of nitrogen in the
sintering atmosphere. (See the article "Corrosion-Resistant Powder Metallurgy Alloys.") The relationships for 304L are
shown in Fig. 21 (Ref 20), showing both the amount of nitrogen absorbed as a function of sintering temperature and
sintering atmosphere (dissociated ammonia and nitrogen) and the strength increase due to nitrogen absorption.

Fig. 21
(a) Effect of nitrogen content on ultimate tensile strength and elongation of 304L stainless steel. (b)
Effect of sintering temperature on amount of absorbed nitrogen for 304L. Source: Ref 20
The problems with nitrogen absorption and chromium nitride (Cr
2
N) precipitation during cooling (after sintering), and
sensitization and loss of corrosion resistance, when sintering is done in dissociated ammonia, are described in Ref 21.
They were not appreciated for many years, in part because corrosion resistance demands were modest and/or corrosion
resistance was not assessed. Later, with increasing demands for improved corrosion resistance, and as more quantitative
information on the effect of sintering in dissociated ammonia became available, recommendations were made to limit
nitrogen absorption to some 3000 ppm in austenitic stainless steels. While this limitation seemed to satisfy some
corrosion resistance requirements in an acidic environment, it was unsatisfactory for parts tested in aqueous NaCl. Even
with small amounts of Cr
2
N precipitates, rust spots would form in a short time. Figure 22 shows examples of Cr

2
N
precipitates in sintered austenitic stainless steels.

Fig. 22 Chromium nitride precipitates in 316L (a) sintered at 1150 °C (2100 °F) in dissociated NH
3
; 4500 ppm
N
2
; Cr
2
N precipitates along grain boundaries (1), and within grains (2). (b) sintered at 1120 °C (2050 °F) in
dissociated NH
3
and slowly cooled; 6500 ppm N
2
; Cr
2
N in lamel
lar form near surface (1) and as grain boundary
precipitates in the interior (2).
In an early study, Sands et al. (Ref 22) pointed out that 316L sintered in dissociated NH
3
required a cooling rate of 200
°C/min for preventing nitrogen absorption and precipitation of Cr
2
N. More recently, Frisk et al. (Ref 23) determined in a
laboratory study that sintering of 316L in dissociated NH
3
at 1250 °C (2280 °F) required cooling rates of >450 °C/min

(Fig. 23). The higher critical cooling rate of Frisk et al. can probably be ascribed to their much lower dew point (-100 °C
versus -40 to -60 °C for Sands), which allows for more rapid nitrogen absorption during cooling as illustrated in Fig. 24
(Ref 24) for the bright annealing of stainless steel strip. The deleterious reactions of increasing nitrogen absorption with
decreasing dew point, and of increasing oxidation with increasing dew point, leave a relatively narrow dew point window
for sintering in dissociated ammonia.

Fig. 23
Effect of cooling rate on presence of chromium nitrides in microstructure of 316L parts sintered at 1250
°C in dissociated NH
3
. Source: Ref 23

Fig. 24 Effect of dew point on nitrogen absorption and oxidation of 316L shim, disk, and ba
r stock annealed for
15 min at 1038 °C (1900 °F) in 30%H
2
-70%N
2
. It took 2.3; 2.8; and 4.7 min respectively to cool the three
materials from 1038 °C (1900 °F) to 538 °C (1000 °F). Source: Ref 24
These cooling rate requirements are even higher than those for sintering in hydrogen to prevent oxidation during cooling,
and the cooling-rate dew point relationship appears to be reversed. It is, therefore, not surprising that stainless steel parts
producers are increasingly shifting towards hydrogen sintering at the expense of sintering in dissociated NH
3
.
Sintering in an atmosphere of 10%N
2
-90%H
2
may be a more practicable compromise that greatly reduces the high-

cooling rate requirements of dissociated NH
3
to more manageable levels, while still benefiting substantially from the solid
solution strengthening obtainable with the lower nitrogen concentration. Good corrosion resistance for such conditions
were reported by Larsen (Ref 25) and Mathiesen (Ref 26).
The positive effect of nitrogen on corrosion resistance, as documented for wrought stainless steels, is expected to apply
equally to sintered stainless steels. However, this beneficial effect has not yet been well documented for. sintered stainless
steels, probably because of the overshadowing negative effect of excessive nitrogen absorption on the surface of parts
from the sintering atmosphere during cooling. Mechanical and other properties for 316L and 434L sintered in dissociated
NH
3
are given in the article "Powder Metallurgy Stainless Steels" in this Volume.

References cited in this section
9.
E. Klar and P.K. Samal, Optimization of Vacuum Sintering Parameters for Improved Corrosion Resistance
of P/M Stainless Steels, Advances in Powder Metallurgy and Particulate Materials,
Vol 7, Metal Powder
Industries Federation, 1994, p 239-251
10.

E. Klar and P.K. Samal, Effect of Density and Sintering Variables on the Corrosion Resistance of Austenitic
Stainless Steels, Advances in Powder Metallurgy & Particulate Materials,
Vol 3, Metal Powder Industries
Federation, 1995, p 11-3 to 11-17
17.

E. Maahn, S.K. Jensen, R.M. Larsen, and T. Mathiesen, Factors Affecting the Corrosion Resistance of
Sintered Stainless Steel, Advances in Powder Metallurgy & Particulate Materials, Vol 7, 1994, p 253-271
18.

C. Lall, Fundamentals of High Temperature Sintering: Application to Stainless Steels and Soft Magnetic
Alloys, Int. J. of Powder Metall., Vol 27 (No. 4), 1991, p 315-329
19.

E. Klar, M. Svilar, C. Lall, and H. Tews, Corro
sion Resistance of Austenitic Stainless Steels Sintered in
Commercial Furnaces, Advances in Powder Metallurgy & Particulate Materials,
Vol 5, Metal Powder
Industries Federation, 1992, p 411-426
20.

N. Dautzenberg, Paper No. 6.18, Eigenschaften von Sinters
tählen aus Wasserverdüsten Unlegierten und
Fertiglegierten Pulvern, 2nd European Symposium on Powder Metallurgy, 1968, EPMA, Vol II
21.

M.A. Pao and E. Klar, On the Corrosion Resistance of P/M Austenitic Stainless Steels,
Proceedings of the
International Powder Metallurgy Conference (Florence, Italy), Associazone Italiano di Metallurgia, 1982
22.

R.L. Sands, G.F. Bidmead, and D.A. Oliver, The Corrosion Resistance of Sintered Stainless Steels,
Modern
Developments in Powder Metallurgy, Vol 2, H.H. Hausner, Ed., Plenum Press, 1966, p 73-85
23.

K. Frisk, A. Johanson, and C. Lindberg, Nitrogen Pick up During Sintering of Stainless Steel,
Advances in
Powder Metallurgy & Particulate Materials, 1992, Vol 3, Metal Powder Industries Federation, p 167-179

24.

R.H.
Shay, T.L. Ellison, and K.R. Berger, Control of Nitrogen Absorption and Surface Oxidation of
Austenitic Stainless Steels in H-N Atmospheres, Progress in Powder Metallurgy,
Vol 39, H.S. Nayar, S.M.
Kaufman, and E.E. Meiners, Ed., Metal Powder Industries Federation, 1983, p 411-430
25.

M. Larsen, "Debindering and Sintering in Powder Metallurgy Processes," Ph.D. thesis, Technical
University of Denmark, 1994 (in Danish)
26.

T. Mathiesen, "Corrosion Properties of Sintered Stainless Steels," Ph.D. thesis, Techn
ical University of
Denmark, 1993 (in Danish)
Production Sintering Practices

Sintered Density
For many years, inferior corrosion resistance of sintered stainless steel parts has been and continues to be mistakenly
attributed to the presence of pores in accordance with the mechanism of crevice corrosion. That this hypothesis is
untenable follows from widely available evidence (Ref 7) that sintered stainless steel parts of identical composition and
similar pore volumes, pore sizes, and pore shapes, but sintered under varying conditions, may have corrosion resistances
(in 5% aqueous NaCl) that can vary by two orders of magnitude. Furthermore, a comparison of wrought and sintered
(85% of theoretical density) type 316L for susceptibility to crevice corrosion in 10% FeCl
3
, in accordance with ASTM G
48, showed that the wrought part was actually more severely attacked than the porous part (Ref 7, 27). More recent
studies have shown that most cases of inferior corrosion resistance of sintered stainless steels are due to incorrect
sintering, as previously described.

Also, for many years, there existed controversy as to the effect of sintered density upon corrosion resistance. Corrosion
testing in an acidic environment, such as dilute H
2
SO
4
, HCl, and HNO
3
, always showed that corrosion resistance
improved with increasing density. Testing in a neutral salt solution, however, showed the corrosion resistance to decrease
with increasing density (Ref 17, 28, 29, 30, 31), when corrosion testing was done in long-term immersion or salt spray
tests, whereas higher density was found to be beneficial in short-term potentiodynamic polarization tests (Ref 32). Maahn
and Mathiesen (Ref 8) attributed the failure to observe this important relationship between corrosion resistance and
density in short-term, potentiostatic anodic polarization tests to the lack of time for the time-consuming build up of
localized attack within the pores, in analogy to the mechanism of crevice corrosion. By using slow stepwise polarization,
the expected relationship, that is, a decrease of the stepwise initiation potential with increasing density (equivalent to
deteriorating corrosion resistance), was observed (Ref 33).
Recent studies, with austenitic stainless steels (Ref 10, 17, 31) showed that the corrosion resistance in 5% aqueous NaCl
(by immersion) can be reduced by up to two orders of magnitude due to the presence of porosity. The negative effect
appears at sintered densities of 6.7 g/cm
3
, reaches a minimum corrosion resistance between 6.9 and 7.2 g/cm
3
,
depending on pore morphology, and thereafter disappears at densities of 7.4 g/cm
3
(Fig. 25). To the left of the
minimum, corrosion resistance decreases with decreasing pore size due to increasing oxygen depletion and failure to
maintain the passive layer. To the right of the minimum, corrosion resistance improves as pores become closed off and
inaccessible to the surface. This type of corrosion can be reduced by impregnating the pores with a resin, by metallurgical
modification of the pore surfaces, or by the use of higher-alloyed stainless steels, particularly those containing higher

concentrations of molybdenum.

Fig. 25
Effect of density on corrosion resistance (B rating) of pressed, sintered, repressed, and annealed 317L
parts
Another approach to avoiding the problem of long-term corrosion in a neutral salt solution due to the presence of certain
size pores is to make use of the various forms of liquid-phase sintering (transient, persistent, and supersolidus) and to
achieve sintered densities >7.4 g/cm
3
. For austenitic stainless steels, silicon additions of several percent (Ref 34, 35) or
smaller amounts of boron (Ref 17), have been used. For ferritic stainless steels phosphorus additions (Ref 35) have been
used. The large shrinkage occurring during liquid-phase sintering is often accompanied by increasing loss of dimensional
stability. Sizing is usually employed to establish dimensional accuracy. Also, depending on the composition of the liquid-
phase sintering additive, secondary phases may be present after sintering, which can have a negative effect on corrosion
resistance (Ref 17).

References cited in this section
7. E. Klar, Corrosion of Powder Metallurgy Materials, Corrosion Metals Handbook,
9th ed., ASM
International, 1987, p 823-845
8. E. Maahn and T. Mathiesen, "Corrosion Properties of Sintered Stainless Steel," p
resented at UK Corrosion
91 (Manchester), 1991
10.

E. Klar and P.K. Samal, Effect of Density and Sintering Variables on the Corrosion Resistance of Austenitic
Stainless Steels, Advances in Powder Metallurgy & Particulate Materials, Vol 3, Metal Powder Ind
ustries
Federation, 1995, p 11-3 to 11-17
17.

E. Maahn, S.K. Jensen, R.M. Larsen, and T. Mathiesen, Factors Affecting the Corrosion Resistance of
Sintered Stainless Steel, Advances in Powder Metallurgy & Particulate Materials, Vol 7, 1994, p 253-271
27.

D. Ro and E. Klar, Corrosion Behavior of P/M Austenitic Stainless Steels,
Modern Developments in
Powder Metallurgy,
Vol 13, H.H. Hausner and P.W. Taubenblat, Ed., Metal Powder Industries Federation,
1980, p 247-287
28.

M. Svilar and H.D. Ambs, P/M Martensitic Stainless Steels: Processing and Properties,
Advances in Powder
Metallurgy, Vol 2, 1990, p 259-272
29.

S.K. Chatterjee, M.E. Warwick, and D.J. Maykuth, The Effect of Tin, Copper, Nickel, and Molybdenum on
the Mechanical Properties and Corrosion Resistance of Sintered Stainless Steel (AISI 304L),
Modern
Developments in Powder Metallurgy,
Vol 16, E.N. Aqua and C.I. Whitman, Ed., Metal Powder Industries
Federation, 1984, p 277-293
30.

F.M.F. Jones, The Effect of Processing Variables on the Properties of
Type 316L Powder Compacts,
Progress In Powder Metallurgy, Vol 30, Metal Powder Industries Federation, 1970, p 25-50
31.

R.M. Larsen, T. Mathiesen, and K.A Thorsen, The Effect of Porosity and Oxygen Content on the Corrosion
Resistance of Sintered 316L, Powder Metallurgy World Congress
(Paris), Vol III, Les Editions de
Physique, 1994, p 2093-2096
32.

G. Lei, R.M. German, and H.S. Nayar, Corrosion Control in Sintered Austenitic Stainless Steels,
Progress
in Powder Metallurgy, Vol 39, H.S. Nayar, S.M. Kaufma
n, and K.E. Meiners, Ed., Metal Powder Industries
Federation, 1984, p 391-410
33.

T. Mathiesen and E. Maahn, Corrosion Behavior of Sintered Stainless Steels in Chloride Containing
Environments, 12th Scandinavian Corrosion Congress (Helsinki), 1992, p 1-9
34.

W.F. Wang and Y.L. Su, Powder Metallurgy, Vol 29, 1986
35.

N.S. Mikkelsen and M. Jensen, "Sinterwerkstoff", European Patent Publication 0564778 A1, 1993
Production Sintering Practices
Sintering of High-Speed Steels and Tool Steels

Powder metallurgy processing offers several advantages to costly and highly alloyed tool steel materials. These
advantages include uniform and finer microstructure, improved grindability, improved cutting performance, and
capabilities of high-speed steels and tool steel alloys that cannot be made by conventional ingot metallurgy.
The use of conventional press-and-sinter technology offers the additional advantage of net shape or near-net shape
capability. Cutting tools, bearings, and wear parts are being produced commercially by fully dense sintering. Wear parts
also are produced by conventional P/M techniques to densities of 80 to 90%.

Current commercial fully dense sintering uses high green strength and compressible water-atomized tool steel powders
that are compacted in rigid dies using uniaxial pressing or in flexible rubber molds using cold isostatic pressing to make
green tools and parts. These parts are then sintered in a microprocessor-controlled vacuum furnace near the solidus
temperature of the alloy to virtually full density (at least 98% and frequently 99+% of theoretical).
The flexibility of this process allows production of pressed and sintered tool and die parts in various forms. Generally,
parts are pressed to 70 to 85% of theoretical density before sintering to full density. Pressing to lower green densities
tends to require longer times at temperature to obtain full density and results in coarser microstructures. Pressing at higher
pressures, which is required for increased green densities, results in increased tool wear and breakage. Parts with high
green density also may require extra care in sintering. The surface can sinter rapidly to high density and entrap gases from
the center of the part.
Production Sintering Practices

Sintering Mechanisms
Important contributions to sintering arise from diffusion and viscous flow. Diffusion rates increase with increasing
temperature because of the increased number of vacancies that promote diffusion of substantial alloying elements.
Sintering temperature must be held very close to the solidus temperature to attain full density in a reasonable time. A
sintering temperature 5.5 to 11 °C (10 to 20 °F) above the solidus reportedly forms a small amount of liquid, which
allows high-diffusion rates for enhanced sintering. These results are based on relatively rapid densification at high
temperatures. However, metallographically, there is no resemblance between a successfully sintered high-speed steel and
a typical liquid-phase sintered material, such as tungsten carbide or tungsten heavy metal. Additional research is required
to fully understand high-temperature sintering.
Factors Affecting Sintering. Sintering is a series of complex processes, of which densification is only one phase. As
green parts are heated to the sintering temperature, gases can be adsorbed (nitrogen, hydrogen, or oxygen) or evolved
(nitrogen, hydrogen, or carbon monoxide). Admixed carbon dissolves and homogenizes, while carbides dissolve and
grow. Grain growth also occurs as pores shrink and virtually disappear.
Sintering Time and Temperature. Increasing the sintering temperature decreases the amount of time required to
achieve full density. Higher temperatures also reduce the time between reaching full density and oversintering. These
relationships are shown schematically in Fig. 26. A sintering curve for M2 high-speed steel is shown in Fig. 27. Table 13
provides typical properties of M2 tool steel for various sintering temperatures.
Table 13 Sintering data for M2 tool steel

Testing was conducted on five lots. All compacts were pressed at 827 MPa (60 tsi) and sintered for 60 min.
Sintered properties:
At 1240 °C (2260 °F) At 1250 °C (2280 °F) At 1260 °C (2300 °F)
Total
carbon
(a)
, %

Added
graphite, %

Density, g/cm
3


Characteristics

Density, g/cm
3


Characteristics

Density, g/cm
3


Characteristics

Lot 1

0.77
0.0 7.19 MLP, FG 7.29 FG, SP 7.47 MSP, VFG
0.87
0.1 7.25 VFG, LP 8.03 FP, LP 8.04 FP, SP
0.97
0.2 7.48 VFG, LP 8.11 FP 8.09 FP
1.07
0.3 8.01 FG 8.09 FP, SP 8.07 E, LP
Lot 2
0.79
0.0 7.00 VFG, MP 7.25 FG, MP 7.77 MP
0.89
0.1 7.18 VFG, MP 7.76 MP 8.08 FP
0.99
0.2 7.58 MP, VFG 8.08 FP 8.06 FP, LG
1.09
0.3 8.07 VFP 8.00 LG, SP 8.05 E, LP, FP
Lot 3
0.86
0.0 7.17 VFG, MP 7.47 MP 8.09 FP
0.96
0.1 7.29 VFG, MP 8.08 MLP 8.08 FP, LP
1.06
0.2 7.84 VFG, MP 8.09 SP 8.06 E, LP, LG
1.16
0.3 8.07 FG, VFP 8.07 SP, E 8.07 E, LG
Lot 4
0.87
0.0 7.22 VFG, MP 7.87 MP 8.09 FP, LP
0.97

0.1 7.47 FG, MP 8.08 FP, SP 8.07 SP
1.07
0.2 8.07 FP 8.06 FP, SP 8.06 E, MSP
1.17
0.3 8.08 FP 8.06 E, SP 8.05 E, MSP
Lot 5
1.00
0.0 7.82 VFG 8.09 FP, SP 8.06 E, FP, LG
1.10
0.1 8.08 FG 8.07 FP, SP 8.05 E, EP, LG
1.20
0.2 8.07 VFP 8.06 LG, E 8.05 E, FP, LG
1.30
0.3 8.06 LG 8.06 LG, E 8.04 E, FP, LP
M, many; V, very; FG, fine grain; LG, large grain; SP, small pores; LP, large pores; FP, few pores; MP, many pores; E, eutectic.
(a)
wt% C in powder plus wt% admixed graphite; does not account for carbon loss during sintering due to
deoxidization


Fig. 26 Relationship among sintered density, sintering time, and sintering temperature

Fig. 27 Sintering curves for compacted M2 high-speed steel. Sintering time for all materials is 1 h.

Effect of Carbon Content. Carbon content has a significant effect on sintering temperature. Increasing carbon content
reduces the sintering temperature by lowering the solidus temperature.
Carbon can be controlled closely by blending graphite or lampblack into the powder before pressing. Carbon dissolves
rapidly, and its high diffusivity enables rapid homogenization above 980 °C (1800 °F). However, blending of more than
0.15 to 0.2% C can cause several detrimental effects, such as nonuniform distribution of carbon, variable response to
sintering, variable part-to-part density, and part distortion.

More than 0.15% C can be admixed with 0.05 to 0.30% blending oil, such as mineral oil. However, blending oils may
create powder flow difficulties and can lead to excessive outgassing during vacuum sintering.
Figure 28 shows the effect of carbon content on microstructure for M2 high-speed tool steel. The amount of mixed carbon
critically affects the microstructure. For a given sintering cycle, insufficient carbon produces incomplete sintering, and
excessive carbon produces oversintering. Table 13 gives the effect of carbon additions for five heats of M2 tool steel and
three sintering temperatures.

Fig. 28 Effect of increasing carbon content on the sintered microstructure of M2. All samples were sintered i
n
the same run. Carbon was increased by blending in graphite. Note the microstructure changes from being
undersintered (round pores) for 0.0 and 0.1% C to being oversintered (sharp angular pores, eutectic) for 0.2
and 0.3% C. 250×
Effect of Oxygen Content. Oxygen affects the sintering of tool steels by reducing carbon content during sintering.
Oxygen from the annealed powder reacts with carbon during sintering to form carbon monoxide. Approximately 0.01% C
is lost in the atmosphere for every 0.01% of oxygen present in the annealed powder. This reaction reduces the oxygen to
less than 200 ppm during vacuum sintering.
Effect of Silicon Content. Silicon aids sintering by suppressing the melting point of tool steels. A carbon equivalent
(CE) for estimating sintering temperature can be calculated as:


Admixed silicon does not homogenize rapidly because of low diffusivity; consequently, it usually causes incipient
melting.
Effect of Boron Content. Boron forms a low-melting-point eutectic with steels. Only small amounts can be added
without changing carbide morphology. Boron can be added to powder as an alcoholic solution of boron oxide, followed
by drying at 95 °C (200 °F). This results in an even coating of the particle with boron oxide. The treated powder thus
sinters uniformly.
Effect of Carbide Formers. The amount of carbon required for sintering and heat treating increases with the
concentration of tungsten, molybdenum, and vanadium. Tungsten, molybdenum, and vanadium combine with carbon to
form carbides that deplete the matrix of carbon. Vanadium is especially potent, because the vanadium-rich MC carbide is
very stable and resists dissolution at high-sintering temperatures.

The stability of the vanadium carbide effectively retards grain growth. Generally, increasing vanadium content (when
properly balanced with carbon) facilitates sintering. High-speed steels containing 3% or more vanadium (M3 type 2 and
T15) can be sintered to full density with microstructures finer than wrought.
In high-speed steels, chromium content does not affect sintering of alloys due to the diminished stability of chromium
carbide compared to molybdenum, tungsten, and/or vanadium carbides. Chromium carbide is formed only in annealed
material. Chromium content affects the sintering of alloys that contain only minor amounts of vanadium, molybdenum, or
tungsten, such as the high-chromium die steels.
Production Sintering Practices

Vacuum Sintering to Full Density
The prime object of vacuum sintering to full density is to uniformly expose all parts to sufficient temperature and time.
This is typically done to produce the minimum required density and to prevent overheating and/or excessive carbide
growth.
Sintering Cycles. Figure 29 shows typical sintering cycles. These sintering cycles consist of a lower temperature hold,
followed by one or more high-temperature holds to densify compacts.

Fig. 29 Sintering cycles for fully dense sintering of high-speed steels.(a) British Patent 1,562,788.
(b) U.S.
Patent 4,063,940
Deoxidation. Heating rates to the deoxidation temperatures are not critical. Parts should be heated at the maximum
practical heating rate to minimize furnace time, provided the temperature of the load does not greatly surpass the
deoxidation soak temperature. Batch-type vacuum furnaces can be backfilled with inert gas or hydrogen to >46 kPa (350
torr) to improve the heating rate and temperature uniformity of the load.
The primary function of the deoxidation step is to react oxygen with carbon to form carbon monoxide, which is removed
by the vacuum pumps. Oxygen, nitrogen, and other gases that may be present in the powder compact must be removed
before sintering closes the interconnected porosity to the surface. Once the porosity is no longer interconnected, gas
evolution results in blistering.
The deoxidation soak or hold also provides an opportunity for load temperature to become uniform. Temperature in any
part of the load should not vary more than ±9 °C (±16 °F) before heating to the sintering soak. The load may be soaked at
1040 °C (1900 °F) for several hours to improve temperature uniformity without seriously affecting the microstructure.

Heating to sintering temperature is slow and closely controlled to ensure temperature uniformity. Typical ramp
rates are 0.5 to 5.5 °C/min (1 to 10 °F/min).
Sintering Soak. Sintering treatments may consist of an isothermal soak or a very slow ramp (typically 3 to 33 °C/h, or
5 to 60 °F/h). Temperature uniformity is essential for successful sintering. Temperature range should not vary by more
than 8 °C (15 °F) at any location within the load. Table 14 gives typical sintering temperatures for two-soak sintering
cycles for several alloys.
Table 14 Typical sintering temperatures and compositions for several high-speed steels

Sintering
temperature
(a)


Composition, % Alloy
°C °F C Cr

Mo

W V Co

Relative
sinterability
(b)


M2
1245

2270

0.85

4.2

5.0 6.3 1.9

. . .

3
M2 (high carbon)

1240

2260

1.00

4.2

5.0 6.3 1.9

. . .

3
M3 type 2
1255

2290

1.20

4.1

5.0 6.0 3.0

. . .

2
M4
1260

2300

1.32

4.5

4.5 5.5 4.0

. . .

3
M35
1225

2240

1.15

4.2

5.1 6.4 2.0

5.0

4
M42
1220

2230

1.10

3.8

9.5 1.5 1.2

8.0

5
(a)
Approximate sintering temperature for powders annealed at atmospheric pressure and
pressed at 830 MPa (60 tsi).
(b)
1 represents the easiest, 5 is the most difficult.

Cooling. Generally, cooling from the sintering temperature is performed as rapidly as possible by backfilling the furnace
with inert gas and using forced convection cooling (also known as fan cooling or gas quenching) to minimize furnace
time. Typically, the hardnesses of gas-quenched fully dense tool and high-speed steels are 50 to 60 HRC. These
hardnesses are substantially lower than the austenitized and quenched hardnesses of 55 to 65 HRC because of excessive

dissolution of carbides, which results in excessive retained austenite.
Nitriding. Tool steels and high-speed steels can be alloyed with nitrogen by maintaining a nitrogen partial pressure after
deoxidation, but before sintering. Significant alloying does occur with partial pressure at >133 Pa (1 torr). Increasing the
partial pressure increases the nitrogen content. Depending on the alloy, nitrogen contents from 4000 ppm (M2) to over
8000 ppm (T15) can be produced by backfilling to atmospheric pressure.
Nitrided cases can be produced by introducing a suitable nitrogen partial pressure after sintering has closed off the
interconnected porosity. Use of nitrogen as the quenching gas during forced convection cooling does not result in any
significant nitriding.
Production Sintering Practices

Heat Treatment
Heat treatment of sintered high-speed steels is similar to the heat treatment of wrought counterparts. In both cases, heat
treatment consists of hardening an annealed structure, followed by tempering to achieve desired properties. The finer
microstructure of P/M high-speed steels may require slightly lower temperatures than wrought components to optimize
performance.
Annealing. Sintered parts should be annealed before austenitizing. This treatment provides grain refinement and
transforms large amounts of retained austenite. Sintering temperatures and times are higher and longer than optimum
austenitizing heat treatment times. Consequently, parts quenched from the sintering temperature contain large amounts of
retained austenite, which lower material properties. A suitable annealing cycle for sintered parts is heating to 900 °C
(1650 °F) for 4 h and cooling at a rate of 50 °C/h (90 °F/h) to 500 °C (930 °F), followed by rapid cooling to ambient
temperature.
Austenitizing. The austenitizing temperature of high-speed steels is influenced by the exact composition of the alloy
(particularly carbon), as well as the carbide size. Fine carbides dissolve more rapidly than coarse carbides; consequently,
parts with fine carbides should be heat treated at a lower austenitizing temperature. Table 15 gives hardness values,
hardening and tempering temperatures, and the average Snyder-Graaf intercept grain size for several high-speed steels.
Table 15 Recommended conditions for salt bath hardening of sintered high-speed steels

Requires hardness

Hardening

temperature
(a)


Tempering
temperatures
(b)


Alloy
HRC DPH °C °F °C °F
Intercept

grain
size
63-64

790-815 1170

2138

570 1058 . . .
64-65

815-840 1180

2156

560 1040 . . .
M2

65-66

840-870 1200

2192

550 1022 . . .
64-65

815-840 1170

2138

570 1058 15
65-66

840-870 1180

2156

570 1058 15
66-67

870-905 1200

2192

560 1040 15
M3 type 2

67-68

905-940 1220

2228

550 1022 14
63-64

790-815 1180

2156

570 1058 10.5
64-65

815-840 1200

2192

560 1040 8
M4
65-66

840-870 1200

2192

540 1004 8
63-64

790-815 1180

2156

560 1040 17
64-65

815-840 1200

2192

580 1076 17
65-66

840-870 1200

2192

570 1058 17
M15
66-67

870-905 1210

2210

560 1040 15
63-64

790-815 1180

2156

580 1076 12
64-65

811-840 1180

2156

560 1040 12
M35
65-66

840-870 1200

2192

560 1040 12
65-66

840-870 1170

2138

570 1058 17
66-67

870-905 1180

2156

565 1049 17
67-68

905-940 1200

2192

550 1022 16
T15
68-69

940-980 1220

2228

520 968 13.5
66-67

870-905 1160

2120

570 1058 15
67-68

905-940 1180

2156

570 1058 15
T42
68-69

940-980 1210

2210

560 1040 14
(a)
All samples were hardened for 150 s soak time.
(b)
All samples were triple tempered (× 3 × 1 h)

The Snyder-Graaf method for determining intercept grain size is based on an actual count of the grains. In the as-
quenched condition, the grain boundaries of high-alloy tool steels are clearly revealed by deep etching in nital. The test
method is conducted on a metallograph, with the structure shown on a ground-glass screen at a magnification of 1000×. A
127 mm (5 in.) line drawn on the ground glass represents a length of 127 m (0.005 in.) on the sample. The number of
grains crossed or touched by this 127 mm (5 in.) line is counted; the average of ten readings at random points on the
sample gives the intercept grain size.
Tempering temperatures of P/M parts are similar to those used for wrought materials. Depending on property
requirements, temperatures range from 540 to 595 °C (1000 to 1100 °F). Figure 30 shows tempering curves for M2 and
T15. Heat treated properties of sintered M2, M35, and T15 are given in Table 16.
Table 16 Typical mechanical properties of commercially sintered M2, M35, and T15 high-speed steels

Grade Property
M2 M35 T15
Density, g/cm

3

8.05-8.2 8.05-8.2 8.15-8.3
Ultimate tensile strength
(a)
, MPa (ksi)

750-800 770-820 770-830
Elongation
(a)
, %
12-14 6-9 3-6
750-2000

770-2000

770-2000

Ultimate tensile strength
(b)
, MPa (ksi)

(108-290)

(112-290)

(112-290)

(a)
Fully annealed.

(b)
Depending on heat treatment


Fig. 30 Tempering curves for high-speed steels at varying hardening temperatures. (a) Mean ha
rdening
response of sintered M2. (b) Mean hardening response of sintered T15
Production Sintering Practices

Sintered Microstructures
Fully dense sintered tool steel and high-speed steel metallurgy and microstructures are similar in most respects to wrought
counterparts. Lower significant differences in carbide size and uniformity of carbide distribution may have significant
consequences. Sintered T15 is capable of finer carbides and grain size than the wrought alloy. Sintered M2 exhibits
carbide and grain sizes that generally are not as fine as sintered T15 and are frequently coarser than wrought M2. T15
contains a large amount of vanadium-rich carbides (MC) that are very stable and do not coarsen readily. These carbides
also inhibit grain growth during sintering.
Sintered D2 microstructure exhibits somewhat finer carbides than wrought D2. Wrought D2, like wrought T15, frequently
has a segregated or "banded" structure in which large carbides group together. Improved properties of P/M tool steels are
attributed to their finer and more uniform microstructure.
As-sintered microstructures consist of untempered martensite within prior austenite grain boundaries and large
amounts of retained austenite. Etching in nital reveals the presence of prior austenitic grain boundaries during sintering.
Etching also can reveal gray or white areas in the matrix. Gray areas are reported to be caused by the very fine carbide
dispersions that result from marginal or insufficient vacuum quenching. White areas indicate locations in which
precipitation has not occurred.
Figure 31 shows undersintered, correctly sintered, and oversintered microstructures. Undersintered microstructures (Fig.
31a) have fine or very fine carbides and grains in addition to porosity. Porosity is rounded (without sharp ends) and can
be irregular in shape when sintered density is low. Correctly sintered microstructures have little or no porosity, no eutectic
structure or evidence of melting, uniformly dispersed carbides, and uniform grain size, as shown in Fig. 31(b).

Fig. 31 Photomicrographs of as-sintered T15 high-

speed steel. (a) Undersintered structure. 190×. (b) Correctly
sintered structure. 190×. (c) Oversintered structure. 190×
The oversintered microstructure has large carbides and grains. Porosity with sharp ends, usually at grain triple points, is
evident. A eutectic structure from localized melting and a continuous carbide network around grains may also be present
in oversintered microstructures (Fig. 31c). Oversintered parts have slightly lower sintered densities ( 0.2 to 0.5 g/cm
3
)
than correctly sintered parts because of the formation of porosity at the grain triple points.
Contamination. Metallic contamination can be readily detected in sintered tool steel microstructures. Contamination
appears as carbide-free areas with porosity and/or carbide networks. Such defects occur when lower melting alloy
particles or alloy particles that react with carbon from the tool steel powder to form low-melting alloys are present in
higher melting tool steel alloys. Iron, stainless steel, and low-alloy steel powders form such defects in M2, T15, and other
high-speed steels. These defects also occur when particles of T15 are present in M2 high-speed steel.
Tool steel powder producers have significantly reduced typical levels of contamination by using powder cleaning,
compaction, and sintering equipment only for tool steel powders and by keeping powder-producing environments
meticulously clean. Tool steel powder users must also clean powder hoppers and blenders to maintain high-quality P/M
part production.
Production Sintering Practices

Atmospheric Pressure Sintering
Although vacuum is the preferred mode of sintering, wear parts made of tool steels and high-speed steels can be sintered
to conventional P/M densities or to full density at atmospheric pressure and in atmospheres with dew points below -40 °C
(-40 °F). Atmosphere-sintered parts have higher oxygen and nitrogen levels than vacuum-sintered parts.
Nitrogen-based, dissociated ammonia, and hydrogen atmospheres are viable alternatives to vacuum. Atmosphere
composition has little effect, however, on sintered density. Parts sintered in pure nitrogen or nitrogen-base atmospheres
are nitrided. Increasing the nitrogen content decreases the sintered transverse rupture strength. Consequently, as-sintered
transverse rupture strength is lowest for parts sintered in nitrogen and highest for parts sintered in hydrogen. Sintered
high-speed steels and tool steels should be double tempered to maximize transverse rupture strength.

Production Sintering Practices
Sintering of Copper-Base Alloys
Alain Marcotte, United States Bronze Powders, Inc.

Copper-base P/M materials rank second only to iron-base parts in terms of commercial applicability. As with other P/M
materials, the final properties and related performance of copper-base parts depend on successful sintering techniques.
This section reviews sintering practices for Cu-Sn, Cu-Zn-Pb, and Cu-Ni-Zn alloys.
Production Sintering Practices

Sintering of Bronze
Sintered bronze can be produced either from mixtures of copper powder and tin powder or from a prealloyed tin bronze
powder. The nominal composition of 90Cu-10Sn can be complemented with other constituents such as dry organic
lubricants, graphite, lead, and iron, depending on the specified grade.
Premix/Diffusion Alloyed Bronzes
Premix or partially diffused bronzes are used extensively in the manufacture of porous, self-lubricating bushings and
bearings and for more complex structures requiring superior bearing and mechanical strength. Self-lubricating bushings
and bearings are produced at nominal densities (oil impregnated) of 5.8 to 7.2 g/cm
3
, with oil contents ranging from 24 to
11 vol%, respectively. Corresponding radial crush values (K strength constant) are approximately 69 MPa (10 ksi) of the
lower density, increasing to 228 MPa (33 ksi) at the highest nominal density.
The basic manufacturing procedure consists of compacting the powder shapes to the appropriate green density and
sintering to achieve a homogeneous metallurgical alpha bronze structure, followed by oil impregnation. A sizing
operation completes the process to ensure dimensional precision and general surface integrity.
Dimensional Change. Several methods are used in the industry to produce bearings and structural parts from sintered
90Cu-10Sn bronze. For practical purposes, either pure copper and pure tin or prealloyed bronze and tin powder can be
mixed. There are advantages and disadvantages to both methods. Mixed powders possess relatively good pressing
properties, but there is always a risk of segregation. Premix powder containing prealloyed bronze (e.g., 94Cu/6Sn + 4Sn)

have less liquid phase during sintering, which accounts for a comparatively lower sintered strength at similar green
density. However, "partially prealloyed" bronze minimizes the risk of segregation and still maintains acceptable sintered
strength.
Partially diffused bronze is possible when premix bronze is presintered at a temperature range from 400 to 750 °C (750 to
1380 °F), such that a metallurgical bond between tin and copper powder is created. The sintered cake is then crushed,
ground, and screened. This treatment minimizes segregation and gives similar sintered properties to premix bronzes.
To include all dimensional patterns with related absolute dimensional change values for each commercially available
premix system is prohibitive. A variety of premixed powders are available with specific sintered dimensional patterns to
satisfy customer design and tooling needs. Despite the dimensional magnitude of the particular premix being sintered,
compositions of this type exhibit a common sintered dimensional pattern. Absolute sintered dimensional characteristics
typically are unique to a specific source of copper and tin powders. For example, sintered dimensional consistency can be
obtained by blending two or more copper-base powders that exhibit different growth characteristics and/or by use of tin
powders that also exhibit different growth characteristics.
Generally, copper-tin blends composed of relatively coarse powder sinter to higher growth values than a blend composed
of finer powders. After powder blends have been tested and adjusted to provide an approximation of target dimensions,
final adjustments are made during production sintering to obtain dimensional precision. For discussion purposes, Fig. 32
shows the relationship between dimensional change, sintering time, and green density of a premix bronze and a partially
diffused bronze, using the same source of powders and a similar particle size distribution (see Ref 36).

Fig. 32 Effect of properties on dimensi
onal change of bronze. Sintered at 820 °C (1500 °F), under dissociated
ammonia atmosphere.
Factors affecting the ultimate, or peak, dimensional values include physical characteristics of the constituents and
compacted density. Control of sintered dimensions in premix systems is achieved by manipulating sintering time and/or
temperature.
Sintering Time and Temperature. Typical sintering furnace temperatures for bronze range from 815 to 860 °C
(1500 to 1580 °F); total sintering time within the hot zone ranges, from 15 to 30 min, depending on the furnace
temperature selected, required dimensional change, and most importantly, the presence of an optimum alpha grain
structure. Figure 33 shows the typical microstructure of a premix bronze part.

Fig. 33 Typical microstructure of sintered premix bronze. 90-10 bronze, etched in K
2
Cr
2
O
7
solution. 175×

Sintering atmospheres should be protective and reducing to facilitate sintering. Reduction of the copper oxides that may
surround each copper powder particle and reduction of tin oxide formation allow for increased diffusion rates.
Consequently, faster sintering rates and more homogeneous structures can be obtained.
Prealloyed Bronzes
Sintered bronze alloys are rather uncommon in powdered metal usage. This is primarily due to their relatively high cost
compared to low-alloy steels. However, sintered bronze properties can be advantageous for non-magnetic applications
that require very good corrosion resistance, good mechanical strength, and excellent ductility.
Prealloyed 80Cu-9Sn-2Zn bronze powders with a select lubricant are intended for the fabrication of high-density P/M
structural components. Unlike many elemental copper-tin premixes, the sintering of prealloyed bronze results in the
attainment of high sintered densities (85 to 90% of theoretical ), that provide correspondingly high strengths and
hardnesses.
Sintering properties. Prealloyed bronze powders are relatively easy to work due to the excellent ductility obtained on
sintering. For example, an 89Cu-9Sn-2Zn bronze pressed at 414 MPa (30 tsi) to 85% of theoretical density, exhibits
>30% elongation thus allowing for a substantial degree of cold working. Compared with sintered brass, bronze powders
reach higher yield strength and hardness levels when compacted under similar conditions. For example, pressing at 550
MPa (40 tsi) results in hardness of 45 HRB for an 89Cu-9Sn-2Zn-2Fe bronze and of 90 HR for a 70Cu-30Zn brass,
respectively (after 30 min sintering at 840 °C (1550 °F) under dissociated ammonia).
The addition of 2% of a select grade iron in he composition improves the sintered structure, which has a more uniform
grain size and results in comparatively higher yield strength and hardness. Figure 34 shows the microstructure of a
prealloyed bronze. Sintering temperature and atmospheres of prealloyed bronzes are similar to those of bronze premix.
Figure 35 shows the dimensional change and transverse rupture strength of prealloyed bronzes as a function of green
density.

Fig. 34 Typical microstructure of sintered prealloyed 89Cu-9Sn-2Zn-2Fe bronze, etched in K
2
Cr
2
O
7
solution.
350×

Fig. 35 Properties of prealloyed bronzes. Sintered 30 min at 840 °C (1550 °F), under DA


Reference cited in this section
36.

E. Peissker, Modern Developments in P/M, Vol 7, Metal Powder Industries Federation, 1974, p 597-613

Production Sintering Practices

Sintering of Brass and Nickel Silvers
Powders of brasses and nickel silvers are prealloyed, single-phase (alpha) powders that, on sintering, yield moderate
mechanical strength, excellent ductility, and good corrosion resistance. Parts can be subsequently burnished to improve
surface finish. The various alloy compositions produced also provide suitable color, or shade, selection for applications
that require a high degree of surface finish and appearance.
During alloy preparation, lead can be added to improve machinability of the sintered forms. Typical machining operations
include drilling, tapping, turning, threading, and grinding. Excellent sintered ductility also facilitates secondary
operations, such as sizing, cold densification, swaging, and staking. By using multiple pressing and sintering operations,
the yield strength and hardness of the P/M structure may approach those of its wrought alloy counterpart (Ref 37).
Standard prealloyed brass and nickel silver powder compositions are controlled to conform to existing materials

standards. This conformance precludes the additions of foreign metallic constituents, such as tin and iron, that affect
sintered mechanical and dimensional characteristics. In spite of the various compositions comprising these standard alloy
powders, most exhibit similar characteristics. After blending with lubricant, powders densify approximately 10% more
than their as-atomized apparent density. For example, an as-atomized powder with an apparent density of 3.0 g/cm
3

usually blends to a density of 3.3 g/cm
3
minimum with the addition of dry lubricant. Compressibilities are excellent, as
lubricated powders compact to 85% of wrought counterpart densities at 414 MPa (60 ksi). Compression ratios of
lubricated powders range from 2.0-to-1 to 2.2-to-1.
Effect of Lubricant. Blending of powders with dry organic lubricants is normally accomplished in a double cone-type
blending unit. To minimize the inclusion of large lubricant agglomerates and other undesirable particles, the material is
passed through a 40 mesh sieve, or screening is recommended prior to blending. The primary lubricant employed with
brasses and nickel silvers is lithium stearate.
Lithium stearate provides an apparent scavenging or cleansing effect that enhances the sinterability of these powders. It
may also result in spotty or speckled superficial stains on sintered surfaces. These phenomena affect the appearance of the
components, but are not detrimental to mechanical properties. To minimize staining, lithium stearate additions of less than
0.5 wt% are recommended, and zinc stearate can be added to provide the additional required lubricity.
Different types of lubricant have a marked effect on the physical and mechanical properties of nonferrous prealloyed
powder (Ref 37). To illustrate the effects on mechanical properties, Fig. 36 shows data for three frequently used lubricants
in the P/M industry lithium stearate, zinc stearate, and stearic acid. The beneficial effect of 1 wt% lithium stearate and
the deleterious effect of 1 wt% stearic acid on mechanical properties are shown.

Fig. 36 Effects of lubricants and sintering time at temperature on tens
ile properties, sintered density, and
dimensional change of brass compacts
Compacting of lubricated powders is performed with standard types of compacting presses employing steel or carbide
dies and punches. Excellent compressibility and good green strength permit compacting to 75% of theoretical density at
pressures as low as 207 MPa (30 ksi).

Although powders are normally free of gangue and other inclusive, abrasive material that may cause tool wear, their
relative softness does cause tooling difficulties. If powders are not adequately lubricated, fines within tool clearances gall
the die wall and adjacent punch areas, thereby requiring tool removal and cleaning. The amount of lubricant added to the
powder should be proportional to the total surface area of the die assembly requiring lubrication during forming and
ejection. Low-profile, minimum die wall contact parts may only require 0.5 wt% added lubricant, whereas a high die wall
contact part having core rods for holes or internal cavities may require 1.0 wt% lubricant.
Generally, there are no restrictions on the compacted configuration of these powders. In tool design, particularly with
regard to die fill, consideration must be given to the relatively higher apparent densities of lubricated brass and nickel
silver powder. Typical apparent densities for lubricated powders range from 3.3 to 3.6 g/cm
3
.
Sintering of brasses and nickel silvers typically is not difficult; however, basic sintering practices do differ from those
employed with other common alloy systems, such as elemental copper-tin blends and iron powder blends. These
differences include sintering temperature, time at sintering temperature, and atmosphere protection.
Sintering temperatures for standard brasses range from 760 to 925 °C (1400 to 1700 °F). Temperature selection
depends on the brass alloy being sintered and the mechanical properties desired after sintering. Lower brasses with higher
zinc contents and lower melting points are sintered at the lower temperature. Generally, a starting temperature of 100 °C
(180 °F) below the solidus temperature (as determined from any copper-zinc binary alloy constitutional diagram) is
suitable.
Nickel silver can be sintered at 870 to 980 °C (1600 to 1800 °F). Currently, only one base alloy is used for the
manufacture of P/M structural parts; it has nominal composition of 64Cu-18Ni-18Zn. The leaded alloy composition
contains 1.5% Pb. Sintering characteristics are similar to those of the brasses; therefore, responses to sintering parameters
that affect dimensional and mechanical properties of brass are equally applicable to nickel silver.
Sintered Properties. Dimensional and mechanical properties of brasses and nickel silvers are primarily affected by
compact density and the amount of time at temperature, as well as the sintering temperature itself. As mentioned
previously, other elements that affect dimensional and mechanical properties usually are not added to powders. However,
sintered properties, especially dimensional change, can be effectively controlled by manipulation of sintering time at the
appropriate temperature. Each alloy exhibits unique dimensional characteristics a 90Cu-10Zn brass compacted at 414
MPa (30 tsi) and sintered for 30 min at 870 °C (1600 °F) may shrink 0.5%, while a 70Cu-30Zn brass similarly treated
may shrink 2.5%.

Figures 37 and 38 show typical property relationships that can be controlled through manipulation of time at temperature.
The leaded 80Cu-20Zn brass shown in Fig. 37 and 38 is commonly used for structural parts fabrication. The density of
7.6 g/cm
3
is "average" for compacting lubricated prealloyed powders containing 0.375% lithium stearate and 0.375% zinc
stearate at 414 MPa (30 tsi). As shown, close dimensional control can be obtained with a minimum reduction in
mechanical properties after 15 min at temperature. Ductility is increased for subsequent forming operations, such as
sizing, cold repressing for densification, or coining, by increasing sintering time.

Fig. 37 Transverse strength of 80Cu-20Zn brasses. Sintered in hot zone at 870 °C (1600 °F) in DA


Fig. 38 Dimensional change of 80Cu-20Zn brasses. Sintered in hot zone at 870 °C (1600 °F) in DA

Non-leaded machinable brasses have been recently introduced on the market to reduce or eliminate lead. It has
been shown (Ref 38) that by replacing lead with select alloy additions, brass parts still maintain similar mechanical and
physical properties, while having improved machinability. Figures 37 and 38 illustrate the properties of non-leaded brass,
leaded brass, and a non-leaded brass containing select alloy additions. Figure 39 also shows the machinability for those
different types of 80Cu-20Zn brasses.

Fig. 39 Machinability of 80Cu-20Zn brasses. Sintered in hot zone at 870 °C (1600 °F) in DA