Paper No. 2007-01-0145
SINTERING OF POWDER PREMIXES – A BRIEF OVERVIEW
Kalathur S.Narasimhan and Frederick J.Semel
Hoeganaes Corporation
1001 Taylors lane
Cinnaminson,NJ 08077
Abstract
Advances in the understanding of the sintering of powder premixes have
contributed significantly to the growth of the ferrous powder metallurgy
industry. This includes sintering both in the solid state and in the presence
of a liquid phase. In this article, the sintering of iron powder premixes
containing: 1) graphite; 2) nickel and graphite; 3) copper and graphite;
4) Phosphorus as ferrophosphorus; and, 5) boron as ferroboron are
discussed. The evolution of microstructure and mechanical properties are
discussed as well.
Paper No. 2007-01-0145
INTRODUCTION
The growth of ferrous powder metallurgy (P/M) over the past five decades has been
outstanding as this technology is proving itself as an alternative lower cost process to
machining, casting, stamping, forging, and other similar metal working technologies.
The growth has been spurred by continuing technological advancements in powder
making and processing, alloy development, and parts production methods.
The markets for the powder metallurgical parts are varied. Automotive applications have
so far dominated the industry growth. However, non-automotive applications are
becoming increasingly important. The latter applications mainly include hand tools,
household and lawn appliances and industrial motor controls and hydraulics
The growth of P/M is dependent on the ability of technological advancements to deliver
products with higher densities than are currently available. Figure 1 shows how parts
distribution by density has progressed and the anticipated use of higher density parts in
the future. The indicated density increases to date are mainly due to improvements in
powder making and compaction processing. The future advances are likely to come
mostly from our understanding of the sintering process.
P/M parts are characterized by density, composition, and microstructure. For a given
part, these three parameters are optimized.
Sintering and additional heat treatments of powder premixes generate the
microstructures that are needed to meet the required performance. Generally, there are
four types of powder premixes that are used in the P/M industry, (2). These are: 1) iron
powder mixed with alloying ingredients such as C, Cu, Ni and P; 2) iron pre-alloyed with
Mo or with Mo and Ni and mixed with C; 3) partially alloyed iron with Mo, Ni and Cu and
mixed with carbon; and, 4) hybrid systems that use various combinations of 1, 2, and 3.
All of these premixes contain a lubricant and in many cases also contain a binder to
prevent compositional variations due to demixing and dusting. Prior to sintering and any
additional heat treatments, the premixes are compacted to near net shape. Typically,
the compaction process determines the final density to within one or two percent and
the sintering process provide the balance.
Paper No. 2007-01-0145
P/M Part Distribution by Density
2005
7.8
6.2
6.2-6.6
7.2-7.6
6.6-6.9
1979
7.2-7.6
1989
6.9-7.2
7.8
6.2
7.2-7.6
7.8
7
6.2
6.9-7.2
6.2-6.6
6.9-7.2
6.2-6.6
6.6-6.9
6.6-6.9
Figure 1: History projection of density advances with time.
In this article, we will discuss the sintering, microstructure, and mechanical properties of
the four most common alloy systems including: 1) iron-carbon; 2) iron-carbon-nickel; 3)
iron-carbon-copper; and, 4) iron-phosphorus. The iron-boron system as potential future
alloy will also be discussed.
In general, when powder premixes are sintered, sintering is by solid-state processes or
a combination thereof with liquid phase processes. In both cases, atom movements
plays a key role. The basic mechanisms include vapor transport and surface, volume
and grain boundary diffusion. The vapor transport and surface diffusion mechanisms
effect metallurgical bonding and pore rounding but not densification. The volume and
grain boundary diffusion mechanisms effect all three: bonding, pore rounding and
densification. Typically, in ferrous PM, all four processes occur to some extent. Since
vapor transport and diffusion are thermally activated processes and iron is subject to
oxidation, sintering necessarily involves high temperatures and protective atmospheres.
In the case of iron base premixes, this usually means temperatures in the range of 1100
to 1300 oC, (2000 to 2350 oF) and nitrogen/hydrogen atmospheres ranging anywhere
from 5% to 100% hydrogen. Hold times are typically from 15 to 30 minutes but on
occasion may be as long as an hour or more. Excellent reviews of the sintering process
and of the various sintering mechanisms and their effects are available in the open
literature, (3, 4).
Fe-C and Fe-C-Ni Systems:
The response of Fe-C and Fe-C-Ni premixes to sintering in terms of the microstructures
and mechanical properties that are obtained are very similar. A typical Fe-C
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microstructure based on an iron powder premix containing 0.5% graphite as the carbon
source is shown below in Figure 2. Such a composition is normally designated as
F-0005, (F = Iron, 5 = 0.5% C).
F- 0005
Figure 2: Microstructures of sintered iron with 0.5% carbon. Sinter bonding of the
powder particles and two phases, pearlite and ferrite, are visible.
The micrographs each show two phases. The light-etching phase is ferrite, essentially,
pure iron or if nickel is present, a solid solution of nickel in iron. The dark etching phase
is pearlite, a lamellar structure of the ferrite and iron carbide, Fe3C, known as cementite.
The relative amounts of the two phases are largely dependent on the carbon content
and to a lesser extent on the process conditions. Their morphologies, on the other hand,
are almost entirely dependent on the process. Typically, higher carbons and faster
cooling rates subsequent to sintering lead to higher pearlite contents, decreased
lamellae spacing and finer grains of both phases. The micrographs also clearly show a
good ‘degree of sinter’ in terms of pore rounding and sinter bonding of the iron particles.
The porosity seen here is typical of most ferrous P/M materials. More generally, the
numbers and shapes of the pores are a function of the as-compacted density and the
sintering conditions, especially the temperature.
The microstructures of the sinterings are important, of course, because they largely
determine the resultant mechanical properties. In general, increases in the pearlite
content and decreases in the lamellae spacing and grain size increase strength but
decrease ductility while increases in the ferrite content increase ductility but decrease
strength. Increases in the density and degree of sinter increase both strength and
ductility. The mechanical properties of an iron powder premix with 0.8% graphite, (i.e. F0008), are shown in Table 1, (5).
Paper No. 2007-01-0145
Table 1: Mechanical Properties of F-0008 carbon steel sintered at 1120°C
Density
UTS
YS
Elongation
Impact
Hardness
(g/cm³)
(MPa)
(MPa)
In 25.4 mm
Energy (J)
(HRB)
6.2
240
210
<0.5
4
50
6.6
290
240
0.5 to 1.0
5
60
7.0
390
260
1.0
7
70
The microstructure of an iron powder premix containing 0.5% graphite and 2%
elemental nickel, (i.e. FN-0205), and otherwise submitted to the same processing as the
earlier F-0005 premix is shown below in the micrographs of Figure 3. Notice, in
particular, the effect of the nickel in increasing the pearlite and decreasing the ferrite
contents relative to those of the F-0005 microstructure. The nickel does this in part by
decreasing the carbon content of the pearlite and in part by altering the response of the
phases to the process conditions.
FN-0205
Figure 3: Microstructure of sintered iron with 0.5% graphite and 2% nickel.
Nickel also increases the diffusivity of the iron and consequently has noticeable effects
on the densification that occurs during the process as well as on the pore morphology
and the degree of sinter of the resultant microstructure. Although the latter effects are
not as qualitatively obvious as its effect on the pearlite content, they are easily
demonstrated by quantitative means.
Paper No. 2007-01-0145
FN-0208
Figure 4: Microstructure of sintered iron with 0.8% graphite and 2% nickel. Note the
nickel-rich regions.
Another interesting effect of the nickel is illustrated by the micrographs of Figure 4.
These show the microstructure of an iron powder premix containing 0.8% graphite and
2% nickel, (i.e. FN-0208). As the carbon content of steel increases to 0.8%, the pearlite
content should increase to 100% and correspondingly, the ferrite content should
decrease to 0%. However, the micrographs in this case clearly show the existence of
light etching colonies, which are superficially indistinguishable from the ferrite of the
earlier F-0008 microstructure. As it turns out, these are so called nickel-rich areas, i.e.,
the remnants of the elemental nickel due to its incomplete solution during the sintering
process. These areas are also present at lower carbon contents as well but, of course,
are not so obvious as here. In any case, since the evident effect of the nickel in these
lower carbon steels is to increase the pearlite content, the mechanical property
expectation is that it will correspondingly increase strength and decease ductility.
However, as it happens, the anticipated strength increases are observed but the
ductility decreases are not. Instead, the ductility either remains the same or increases
slightly. This effect is generally attributed in part to the improved degree of sinter due to
the nickel and in part to the presence of these nickel rich areas.
The mechanical properties of the most common of the Fe-C-Ni and Fe-C-Cu premix
compositions are shown in Figures 5 through 9. The sintering and microstructure of the
Fe-C-Cu alloys will be discussed shortly. The data in the figures indicate the effects at
two carbon contents, 0.5% and 0.8%, and of the 2% alloy level in each case. The FN0205 and FN-0208 designations correspond to the nickel bearing compositions and the
FC-0205 and FC-0208 designations likewise correspond to the copper compositions.
Paper No. 2007-01-0145
The figures are all based on the Standard 35 data compilations of the Metal Powder
Industry Federation and are presented to reflect the impact of sintered density on
properties, (6).
Properties of the Nickel and Copper Steels For Optimized
Microstructures
Figure 5: Variation of the ultimate tensile strength versus sintered density.
Figure 6: Variation of the compressive yield strength with sintered density.
Paper No. 2007-01-0145
Properties of the Nickel and Copper Steels Continued
Figure 7: Variation of apparent hardness with sintered density.
Figure 8: Variation of impact energy with sintered density -unnotched specimens.
Figure 9: Rotating bending fatigue strength versus sintered density.
Paper No. 2007-01-0145
Heat Treated Properties
As will be evident, these nickel and copper premix compositions generate excellent
mechanical properties, which typically increase with increasing density. If the application
requires even higher properties, then the parts can be submitted to a hardening
treatment. This typically involves reheating to a temperature in the range of 850 to 950 oC
followed by rapid cooling as by quenching in oil and a subsequent low temperature stress
relief or so-called tempering treatment. The tensile properties of the nickel and copper
steels resulting from this heat treatment are shown below in Figures 10 and 11. A typical
quenched and tempered microstructure is shown in Figure 12. The dark etching phase in
this micrograph is so-called tempered martensite and the light-etching phase is again an
instance of the nickel rich areas that were previously discussed.
Properties of the Nickel and Copper Steels in the Heat Treated
Condition
Figure 10: Variation of the ultimate tensile strength versus sintered density.
Figure 11: Variation of the compressive yield strength with sintered density.
Paper No. 2007-01-0145
Quenched and Tempered FN0208
Figure 12: Microstructure of heat-treated nickel steel showing tempered martensite
And nickel-rich areas.
Fe-C-Cu System
The Fe-Cu-C system is widely used in the P/M industry whenever hardness, higher
strength, and moderate wear resistance is required. As a review of the earlier
mechanical property compilations in Figures 8 through 13 will confirm, the copper steels
exhibit higher strengths and hardnesses than the nickel steels at similar copper addition
levels and otherwise the same carbon contents.
Premixes of the Fe-C-Cu system are typically made with elemental copper additions.
Since copper melts at 1084 oC and these premixes are normally sintered at
temperatures in the range of 1100 to 1300 oC, their sintering behavior differs from those
of the Fe-C and Fe-C-Ni systems by virtue of the presence of a liquid phase.
As a general matter, sintering in the presence of a liquid phase can have profound
effects. However, the effects may be positive or negative and depend on several
factors. In general, the three main factors concern both the liquid and the solid phases
and include: 1) their surface tension properties; 2) their relative solubilities; and 3) their
volume fractions.
In the case of the Fe-C-Cu premixes, neither the solubilities nor the surface tension
properties that characterize them are very favorable. In addition, there are certain
Paper No. 2007-01-0145
adverse effects having to do with the appearance of the liquid and the behavior of the
copper subsequently. Consequently, the contribution of the copper to the overall
sinterability of these premixes is essentially negative. However, as was indicated in the
earlier mechanical property compilations, the nominal copper addition is typically 2%.
Thus, the volume fraction of the liquid that forms is fairly low and its effects, although
generally negative, are also reasonably moderate. The details are as follows.
The effects of a liquid phase during sintering depend primarily on its ability to penetrate
the interparticle boundaries and secondarily on its ability to dissolve and transport the
solid phase. If the liquid is able to penetrate the boundaries and is present in sufficient
quantity to spread throughout the compact, then significant densification will occur in a
relatively short period of time after the liquid first appears due to the effects of its
capillary action in rearranging the particles in a more compact form. Thereafter,
provided the solid phase is reasonably soluble in the liquid and the liquid is conversely
not too soluble in the solid, the densification will continue by a process that is similar to
but generally faster than the grain boundary diffusion mechanism in the solid state.
In order that the liquid penetrate the grain boundaries sufficiently to effect significant
densification, the dihedral angle that it makes with the solid at the point of the
boundaries must be low. Dihedral angles of 5o or less are ideal but higher angles up to
about 20o are known to be suitable. In the case of the binary Fe-Cu system, the dihedral
angle is in the neighborhood of 10o or less. However, it is also very sensitive to the
presence of carbon and increases sharply as the carbon content increases. Thus, for
the premix compositions under discussion, (i.e. 0.5 to 0.8%C), the angle is in the
neighborhood of 35 to 45o. As a result, the liquid penetrates the boundaries but only at
their intersections, (e.g. at so-called triple points), and consequently, not sufficiently to
promote much densification, (7, 8). Moreover, even if this were not the case, there is
another difficulty in this system. As discussed below, the relative solubilities of the solid
and liquid phases are also unfavorable.
At the carbon and copper contents of interest, the phase relations indicated by the
binary Fe-Cu diagram provide a reasonable approximation of the relations in the ternary
Fe-C-Cu system. This diagram is shown overleaf in Figure 13, (9).
After the rearrangement step or in lieu of it, the liquid phase promotes sintering by
providing a medium of rapid transport for the solid. To do so, the solid must dissolve in
the liquid and the greater its solubility, the better. On the other hand, as already
indicated, the liquid should not be too soluble in the solid. The reason is that its
dissolution will diminish its presence and hence, its sintering effect. Unfortunately, in the
case of the Fe-Cu and Fe-C-Cu systems, the solubilities of the two phases are precisely
the opposite of what they should be. As indicated by the phase relations in Figure 13,
the copper only dissolves up to about 4.5% iron at all but the highest temperatures
under discussion whereas the iron dissolves upwards of 9% copper. As a result, the
Paper No. 2007-01-0145
Figure 13: The Fe-Cu binary diagram.
copper's capacity as a medium of transport for the iron is very limited and its existence
as a liquid is likewise limited. In fact, as to the latter, the copper starts to dissolve in the
iron even before it melts and then dissolves even faster once it does melt so that the
actual liquid presence during the sintering of these compositions is short lived at best.
Nevertheless, if the foregoing were the only effects of the copper, then the overall effect
of the presence of the liquid phase would be innocuous or, perhaps, even somewhat
beneficial. However, both the melting of the copper and its subsequent dissolution are
accompanied by growth effects, which more than offset any beneficial effects that might
otherwise accrue. The growth associated with the melting of the copper is thought to
arise as follows. During the compaction step, some of the copper particles are
completely encapsulated by the iron and some end up forming part of the internal
surfaces of the pores. When the copper melts, it undergoes a substantial volume
expansion. The copper particles that are part of the pore surfaces simply expand into
the pores. However, the particles that were initially encapsulated force the adjacent
material to expand as well and so produce the indicated growth effect. The growth
associated with the dissolution of the copper is generally attributed to a difference in the
diffusion characteristics of the copper and the iron. In general, copper diffuses faster in
iron than visa versa. Thus, as it goes into solution, it causes growth by adding its
volume to the iron while simultaneously leaving an almost equivalent volume of porosity
in its wake. The latter effect is frequently manifest as small isolated pores in the iron
matrix. For example, a number of such pores are evident in the high magnification
micrograph of Figure 14 which otherwise shows the typical microstructure of a sintered
FC-0208 composition.
Paper No. 2007-01-0145
Figure 14: Microstructure of a Fe-Cu-C sintered part. Notice the small isolated pores
in the micrograph on the right which are possibly due to copper solution
during sintering.
Fe-P System
Iron-phosphorus compositions find application in parts that require a combination of
high yield strength and better than average ductility and also in parts that require
improved magnetic properties. The increased strengths and superior magnetic
performance of these compositions are directly traceable to the solution hardening
effects of the phosphorus. The associated distortion of the iron lattice both strengthens
it and increases it’s resistivity. The ductility improvements are likewise traceable to the
phosphorus but in this case, primarily to its effects in improving the sintering response
of the iron.
The yield strength and elongation properties of a series of iron-phosphorus
compositions as sintered to different densities are shown overleaf in Figure 15. As a
review of this figure will confirm, the yield strength is a monotonic increasing function of
the phosphorus content but the elongation first decreases then increases and finally
decreases again with increasing phosphorus. The yield strength variation is thought
Paper No. 2007-01-0145
Figure 15: Yield strength and elongation values of Fe-P premixes as a function of
sintered density and phosphorus content.
to be mostly attributable to the comparatively large solution hardening effect of the
phosphorus. For example, according to findings based on microalloying theory, its effect
in increasing both the yield and tensile strength of iron is upwards of 20 times that of
copper, (10). In contrast, the more complex variation of the elongation property with
increasing phosphorus is apparently due to a combination of three effects including: 1)
enhanced sintering; 2) solution hardening; and, 3) at high phosphorus contents, the
precipitation of Fe3P as an embrittleling grain boundary phase. Thus, as shown below in
Figure 16, the Fe-P binary diagram is relevant to explain the elongation variation of
these mixes as well as their sintering response, (11).
Figure 16: Fe-P binary diagram.
The diagram shows that the phosphorus closes the γ loop and stabilizes the α phase.
Since the iron diffusivity of the α is about an order of magnitude greater than that of the
Paper No. 2007-01-0145
γ, the general effect of stabilizing this phase is to increase the sinterability of the
resulting alloys. The insert on the right hand side of the diagram indicates that the knee
of the loop occurs at about 1100 oC and that the γ phase is stable at this temperature for
contents up to about 0.3% P. Further increases, however, destabilize it in favor of the α
which eventually becomes the dominant phase at contents in excess of about 0.65% P.
Thus, the greatest sintering response is typically observed at this content with higher
phosphorus contents having little additional effect.
The diagram also indicates the existence of a binary eutectic at 1050 oC and 10.5 % P.
Since the solidus limit on the iron rich side here occurs a 2.8% P, the existence of this
eutectic would ordinarily have little bearing on the behavior of the low alloy compositions
that are typical of most P/M applications. However, phosphorus-containing premixes are
normally made with the intermetallic compound Fe3P as the phosphorus source. As will
be evident from the diagram, the Fe3P determines the solidus limit on the phosphorus
rich side of the eutectic. Thus, in spite of their relatively low alloy contents, the premixes
benefit from the presence of a transient liquid phase, which appears during the early
stages of the sintering process. Since its appearance is, at best, very brief, the evident
effect of the phase is to even out the phosphorus distribution along the interparticle
boundaries of the compact and thereby create a more continuous intergranular network
of the rapid sintering α phase than would otherwise be the case. In other words, its
unlikely that the enhanced sintering response that is observed with these alloys is due
to actual liquid phase sintering per se but rather to an effect of the transient presence of
the liquid in promoting improved conditions for subsequent rapid sintering in the solid
state.
Lastly, the phase diagram also indicates Fe3P precipitation during cooling subsequent to
sintering for alloys with phosphorus contents in excess of about 0.3%. The precipitates
are dispersed in the α matrix and at grain boundaries. At phosphorus contents in excess
of about 0.6%, the latter are sufficient to form essentially continuous films and embrittle
the grain boundaries. Thus, as indicated in the earlier Figure 15, the ductility of these
alloys tends to peak with increasing phosphorus at about 0.6% and decrease thereafter.
Since parts of the phosphorus alloys may undergo substantial shrinkage during
sintering, they are sometimes alloyed with Cu and/or C to return them to near die size.
The Cu does this by virtue of the growth that accompanies its dissolution during the
sintering process whereas the C does it by stabilizing the γ phase and reducing the
amount of α phase sintering that occurs. Since the Cu and C additions also improve the
strength of the resulting parts, several material systems containing one or both of these
alloys in addition to phosphorus have been developed accordingly.
Fe-B System
Boron is capable of producing extraordinary effects in ferrous-based alloys. However, in
P/M, these effects largely remain as untapped potentials since they are accompanied by
equally extraordinary process control and/ or metallurgical problems. For example, in
as-cast and wrought steels, as little as 0.0015% B has been shown to double and, in
Paper No. 2007-01-0145
many cases, more than double the hardenability of virtually any ferrous base alloy, (12).
However, boron’s affinity for oxygen and, to a lesser extent, nitrogen and the associated
inconsistencies which attend their perennial presence in iron and iron base powders
have so far completely prevented its use as a practical hardenability agent in P/M
applications.
Similarly, Jandeska has developed an extremely effective abrasion resistant alloy for
automotive applications, which takes advantage of boron’s ability to combine with iron
and carbon to form massive hard carbides in an otherwise relatively ductile pearlitic
matrix, (13). But, here again, the necessity to sinter in pure hydrogen and to use
expensive oxygen and nitrogen getters such as titanium or aluminum as well as the
unfortunate further complication of having to sinter the alloy in a narrow temperature
window to form the carbides and control dimensions has largely mitigated against its
widespread usage.
In P/M terms, probably the most spectacular attribute of boron is the potential it offers to
achieve high sintered densities, (14). For instance, an addition of as little as 0.15% is
sufficient to attain near full density in most compositions. Of course, once again, these
increases are accompanied by the aforementioned process problems relating to
oxygen, nitrogen and dimensional control. However, given the fact that the indicated
density potential has been known for twenty years or more, if these were the only
problems, they would have almost certainly been solved by now. Unfortunately, they are
not the only problems. In particular, as explained below, it turns out that in spite of the
density improvements, there are microstructual changes attending the use of the boron
that prevent all but marginal improvements in the resulting mechanical properties.
The present findings are based on an investigation that included studies of three alloy
systems: Fe-B; Fe-B-C; and Fe-B-X where X was any one or a combination of carbon
and one or two of the most common of the first and second series transition elements.
Most of the latter work was with iron based prealloyed Mo as in Ast 85 HP or prealloyed
iron with Mo and Cr as in various experimental compositions that were made for
Högänas AB in Sweden during their development of Astaloy 3CrM.
The phase relations in the Fe-B system are shown below in Figure 17, (15). According
to the indications of this figure, the principal phases at all temperatures on the iron rich
side of the diagram are Fe and the intermetallic compound, Fe2B or so-called iron
boride. At high temperatures, the phase relations in the Fe-B-C and Fe-B- X systems
containing C are similar. They include the iron base phase with whatever carbon is
present as a dissolved solute and Fe2B. At room temperature, the phases are
marginally different. They include the iron base phase in the essentially carbon free
condition and either the so-called borocarbide phase, Fe23(C,B)6 which, of course, is
also an intermetallic compound or for carbon lean compositions, a mixture of the
borocarbide with Fe2B, (16). The borocarbide phase is the product of a ternary reaction
that starts at 965o C or somewhat higher and is essentially complete at 800o C or
slightly below. In the case of the Fe-B-X system in general, the indications of
observations made in this laboratory are that the phase relations are qualitatively similar
to those of the Fe-B system. Accordingly, on the iron rich side of the diagram, the
principal phases at all temperatures appear to be Fe with most of the X as a dissolved
Paper No. 2007-01-0145
solute and a simple substitution intermetallic that appears at high temperature as the
alloy boride, (Fe,X)2B and at lower temperatures as nearly pure X2B.
Figure 17: Fe-B phase diagram.
The binary Fe-B alloy and, owing to the similarity in phase relations, each of the ternary
alloys mentioned are classic liquid phase sintering systems. Contrary to the indications
of the foregoing figure, the eutectic liquid in the binary system appears to form at about
1160o C and anywhere from 5 to 15o C lower in the ternary systems. In all three
systems, it’s dihedral angle is of the order of 10o or less, so it both wets the iron matrix
and easily penetrates the original interparticle boundaries and apparently with more
difficulty or, at least, more slowly, the grain boundaries. In accordance with the
indications of the figure, the boron is virtually insoluble in the solid phase, so unlike the
transient nature of the liquid in the case of the Fe-Cu-C system, the liquid presence in
this case is permanent or so-called, persistent . Since, as also indicated in the diagram,
the liquid’s iron content is upwards of 80% atomic, (i.e. approaching 97% by weight), its
sintering potential in terms of ability to transfer iron is substantial. Of course, the actual
amount of sintering that takes place is also dependent on how much liquid is actually
present, (i.e. on its volume fraction). This, in turn, is determined mainly by the boron
content and the temperature. As a general matter, increasing the boron content at a
constant temperature or increasing the temperature at a constant boron content will
increase the amount of liquid and thus the amount of sintering that occurs. For example,
the effect of increasing the temperature at a constant boron content is shown below in
Figure 18 for sintering in the Fe-B-(C, Cr, Mo) system.
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Figure 18: Variation of density as a function of sintering temperature.
The mechanical properties of these same sinterings are shown below in Figure 19. As a
review of this figure will confirm, the indicated properties are unexpectedly poor in view
of the high densities that were achieved. In fact, as shown by the data presented in the
earlier Figure 10, equivalent or better properties are obtainable by conventional
sintering of a simple FC-0208 composition. Unfortunately, the present poor response in
terms of properties is typical of all three of the subject alloy systems and is in every
case a consequence of the disastrous effects of the room temperature remnants of the
liquid phase. More specifically, the boron latent liquid that initially penetrated the
interparticle and grain boundaries and was responsible for the observed densification is
likewise responsible for the poor properties because during cooling it subsequently
remains in the boundaries and precipitates embrittling intermetallic borides and/or borocarbides. The only possibility to improve matters with these alloys appears to be by a
spheroidizing anneal. However, preliminary trials with this objective suggested that the
process is likely to be costly and the potential improvements too small to justify the
effort.
Paper No. 2007-01-0145
Figure 19: Tensile properties and ductility properties of the sinterings of
1.5Cr – 0.5 Mo Pre-alloy with 0.5%C, 0.4%B.
To conclude, the use of boron to improve hardenability and abrasion resistance appears
to have potential for future applications. Although processing difficulties tend to mitigate
against its present widespread usage in these regards, its only a matter of time before
the general advance of technology provides the process control remedies that are
needed to overcome the problems. On the other hand, it seems unlikely that boron will
play a significant role as a liquid phase sintering additive in the foreseeable future.
Unless a practical expedient can be found to alter the morphology of the boron
containing phase from what it presently is in the as-sintered condition, the liquid phase
compositions will continue to exhibit properties that simply do not justify the effort and
expense that are required to produce them. Unfortunately, based on present theory, the
possibility that such an expedient will turn up any time soon appears minimal.
All the systems discussed thus far use pure iron powder as a base. However
prealloyed materials and diffusion-alloyed materials are also used replacing pure iron.
The driver for these is to increase tensile strength, hardenability and fatigue
performance. Number of alloy systems have been developed and reported at various
MPIF conferences in the past years. Of particular interest are two alloy systems namely,
a nickel-molybdenum- manganese prealloyed material(Ancorsteel737SH) and a
chromium-silicon-molybdenum (Ancorsteel4300) material. Both of these systems were
developed to reduce cost of processing powder metallurgy parts.
Table (2) below shows the effect of copper and graphite additions to Ancorsteel 737SH
on the mechanical properties:
Paper No. 2007-01-0145
Table 2:
Ancorsteel 737SH Mixes with 0.75w/oAcrawax C
Sintered at 2080F with full VariCool, Tempered at 400F for 1 hour.
Pressure
(tsi)
Pressure
(MPa)
Green
Density
(g/cm³)
737SH + 0.7 w/o Gr
30
415
6.55
40
550
6.85
50
690
7.03
737SH + 1.0 w/o Cu + 0.7 w/o Gr
30
415
6.58
40
550
6.87
50
690
7.06
737SH + 2.0 w/o Cu + 0.7 w/o Gr
30
415
6.59
40
550
6.88
50
690
7.04
(%)
Sintered
Density
(g/cm³)
0.13
0.16
0.2
6.49
6.79
6.98
0.19
0.22
0.25
0.14
0.15
0.19
6.51
6.80
6.99
0.13
0.17
0.19
6.51
6.78
6.96
Green Exp.
Dimensional
Hardness
Change
(%)
(HRC)
YS
3
YS
TS
3
TS
Elong.
(10 psi)
(MPa)
(10 psi)
(MPa)
(%)
22
32
37
84
88
114
579
607
786
88
97
118
607
669
814
0.7
0.6
0.7
0.22
0.26
0.30
26
32
37
103
114
134
710
786
924
106
119
151
731
821
1041
0.9
1.0
1.3
0.29
0.38
0.41
26
31
35
101
124
141
697
855
972
102
124
146
703
855
1007
0.8
0.9
1.0
Microstructure developed in a 1Cu ,0.7% Gr containing material is shown below. It is
fully martensitic in as sintered and rapidly cooled condition.
Figure 20: Microstructure of a Sample Produced from Premix #1-5 (1 w/o
Copper – 0.7 w/o Graphite). Etched with 1% Nital / 4% Picral. Original
Magnification 500X.
In the case of Ancorsteel 4300,chromium is introduced as an alloying element
along with silicon without adversely affecting the compressibility .Further this material
can be sintered at conventional sintering temperature of 1120C to 1150C.Sintering at
higher temperature than this can result in superior properties. A comparison is made of
this material with a diffusion-alloyed material and a hybrid material of
0.85%molybdenum prealloyed material premixed with 2%nickel and 0.6% graphite.
Ancorsteel 4300 contains lower alloying elements than FD0405 and does not require
heat treatment as is necessary with FLN2 4405.
Paper No. 2007-01-0145
Table 3: Summary of mechanical properties for the three alloys with
conventional cooling.
Table 4: Nominal compositions of the alloys.
APPLICATIONS
P/M parts manufactured today utilize several of the material systems described above.
Listed below are some of the applications. Any specific application is dependent on
design and material performance to meet the design needs. For the same application,
there could be more than one material system and process to match the performance
characteristics. The exact material selection is made by discussions with the powder
producers, parts producers, and the end user. The list below shows typical applications,
and is by no means exhaustive.
Paper No. 2007-01-0145
Table 5: Application of sintered powder mixtures.
Material System
Application
Fe
Magnetic clutch housings
Lock-up clutch collar
Magnetic applications
Fe-Cu-C
Powder forged connecting rod
Crankshaft belt drive
Camshaft pulley
Valve seat inserts
Clutch adjustment ring
Converter turbine sleeve
Oil pump gears
Fe-Cu-C (double pressed-double sintered)
Synchronizer hubs
Fe-Ni-C
Crankshaft belt drive
Camshaft sprocket assembly
Engine balance shaft components
Water pump pulley
Output shaft hub
Fe-Cu-C (heat-treated)
Parking gears
Fe-P
Crankshaft sprocket
Synchronizer hubs
Synchronizer interlock sleeve
Magnetic applications
Fe-Cu-P-C
Bearing cap
CONCLUSIONS
Iron powder premixes containing carbon and either or both of nickel and copper sinter
mainly by surface diffusion which produces metallurgical bonding but has no effect on
dimensions. Thus, they provide the opportunity to manufacture parts to net shape with a
wide range of mechanical properties. The use of phosphorus in the form of Fe3P
provides strength and/or magnetic properties via solution hardening as well as good
ductility as a result of densification due to the combined effects of a transient liquid
phase and alpha phase sintering. Boron systems have potential for future applications
requiring either optimum hardenability or abrasion resistance but probably not
densification by liquid phase sintering.
Growth of P/M will continue over the next several years. Sintering, a key technique, will
be needed to reach the densities and properties required to optimize this growth.
However, persistent liquid phase sintering as exemplified by the present boron
compositions is not likely to be one of these techniques. Nevertheless, there are three
other liquid phase processes that have yet to be thoroughly explored including in order
Paper No. 2007-01-0145
of greatest probability of near term success: 1) infiltration; 2) alloy design for transient
liquid phase sintering; and, 3) supersolidus liquid phase sintering. In addition, there are
also the possibilities of using either persistent or transient alpha phase sintering as well.
Unfortunately, none of these processes is as easy or straightforward as simple
compaction. However, compaction is limited by the pore free density of the premix,
which in some cases may be quite low. Moreover, although its possible to increase the
pore free density by reducing the lubricant content and/or otherwise re-designing the
premix, its easily shown that even under the best of circumstances, compaction is
incapable of achieving the density values that are physically possible by sintering. Thus,
to realize its maximum potential, the P/M industry will inevitably need to take advantage
of this key technique.
ACKNOWLEDGEMENTS
Tom Murphy and his group provided microstructures provided in this report.
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1.
Source – Automotive Magazines – Data Compiled by Eric Boreczky (Private
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Paper No. 2007-01-0145
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