Iron Base Infiltration For High Density
F. J. Semel
Hoeganaes Corporation, Cinnaminson, NJ 08077
ABSTRACT
A new P/M process for making parts potentially having equivalent or better properties than the
ductile cast irons and densities up to 7.55 g/cm3 is described. The process is in the early stages
of development and is based on an essentially pioneer technology that does not depend
significantly either on high pressure compaction or high temperature sintering. The process
appears to offer considerable potential for economic production of large parts as the associated
green densities and process temperatures that are required to implement it are typically below
6.8 g/cm3 and 1175 oC respectively. What is currently known of the process and underlying
technology is presented in detail.
INTRODUCTION
The mechanical properties of ferrous based PM parts are density limited. In general, the higher
the density at any given alloy content, the higher the resultant properties. Consequently, in order
to increase mechanical properties without resort to increased alloy content as well as to
increase applicability with minimal increase in cost, the major thrust of R&D in ferrous PM in the
last quarter century or so has been to increase density. In general, there are only two ways to
do this: compaction and sintering. Of the two, compaction is by far the simpler and has
consequently received the most attention.
The maximum achievable density by compaction is dependent on the admix composition of
interest and in the case of the most common compositions of industrial importance is currently
about 7.40 g/cm3, (i.e. using traditional molding grade powders). Assuming no densification in
sintering, inevitable weight losses during the process due to lubricant burn-off and incipient
deoxidation of the base powder by admixed graphite will reduce this to about 7.35 g/cm3. Of
course, if there is densification during sintering, then higher densities are possible. However,
owing to the technical and economic limitations of the sintering process, the maximum
additional density that can reasonably be achieved in practice is only about 0.10 g/cm3. Thus,
the current overall maximum achievable density in a single press and sinter process is about
7.45 g/cm3.
Exclusive of powder forging which is far too expensive to be competitive in the context of the
average PM part, there is only one way to achieve yet higher densities and simultaneously
maintain the flexibility to do so without significant compositional compromise. This, of course, is
the double press and sinter process in which the part is initially compacted, submitted to what is
basically a combination lubricant burn-off and sub-critical anneal at a low temperature, then
compacted a second time and finally sintered at a high temperature. The extra processing, of
course, adds costs but has nevertheless been found to be sufficiently economical to be
competitive, especially for parts that would otherwise require several machining steps to
produce from wrought or cast stock. Unfortunately, the additional densification afforded by the
process is only about 0.15 to 0.20 g/cm3 at best. Moreover, it decreases sharply with increase in
the density of the first compaction step and at a first compaction density of 7.40 g/cm3 is only
about 0.05 to 0.10 g/cm3. So, given the inevitable density decreases due to weight losses during
sintering, the potential in terms of the maximum achievable density is typically no greater than
about 7.50 g/cm3.
Copper infiltration provides yet another means to achieve high densities in ferrous based parts.
Here, however, the expense of the infiltrant adds even more to the costs than in the double
press and sinter process. In addition, although the achievable mechanical properties are quite
good, they are comparatively limited by the fact that copper is necessarily the dominant alloy
addition and results in a composite microstructure in which it is essentially the properties limiting
‘soft phase’.
Iron Base infiltration
A potential improvement to using copper in the infiltration process would be to use an iron base
alloy as the infiltrant. However, a survey of the PM literature employing a recently published
data base which includes upwards of 4500 citations worldwide as reported over the last fifty
years, (1), indicated only six titles on the subject, (2-7). These articles were all published before
1985 and were apparently the output of independent research in three laboratories, one in the
Soviet Union and two in Europe. The articles reported studies on three different alloy systems.
One system may be described qualitatively as iron-boron with various first and second transition
series metals as ternary additions. The second system was a near eutectic iron-carbonphosphorous alloy. The third system employed copper coated particles as a barrier to carbon
diffusion and actual cast iron as the infiltrant.
Unfortunately, none of these studies was able to make a sufficiently convincing case for the
process to generate continued interest either then or at any time in the intervening interval.
Possibly, this was because it was evident from the aggregate of experience which they
contained that iron base infiltration is a relatively complex process.
In any case, starting in late 2001, research efforts in this laboratory were directed to conducting
a project with the objective to develop the iron base infiltration process as a practical matter.
The purpose of the present report is to document the findings of this effort to date.
EXPERIMENTAL PROCEDURE
In view of the general lack of success of the earlier referenced research as well as the evident
complexities which were indicated, it was decided to take a relatively basic approach and
develop the process essentially as a pioneer technology.
General Alloy and Process Design Considerations
There are five basic parameters to be considered in designing an infiltration process as follows:
1.
2.
3.
4.
5.
Alloy system;
Equilibrium phase relations;
Base compact density;
Infiltrant weight; and,
Process conditions.
The considerations relating to each of these parameters that apply specifically to iron base
infiltration are detailed below. Significantly, iron base infiltration can not be practiced as a
general matter without knowledge of these details and very little of what is presented here that
is specific to it appears elsewhere in the open literature including the aforementioned early
articles on the subject.
Alloy System
The alloy system is perhaps the most important consideration in terms of the mechanical
properties that are likely to result from the process. For example, based on extensive studies
that were done in this laboratory on the liquid phase sintering of iron-boron systems, boron as
the eutectic forming alloy would not be considered to be a good choice as a candidate for iron
base infiltration, (8). Even at near full density and with the added benefit of a precipitation
hardening effect from the addition of molybdenum to form molybdenum borides, the iron-boron
system failed to produce tensile properties that were appreciably better or, in many cases, even
as good as those observed in any number of well known PM compositions at moderate
densities. Not surprisingly, a review of the properties reported in the aforementioned articles that
teach infiltration in iron-boron systems only served to confirm this view. Instead, a better choice
of a candidate system would be the simplest possible alloy that is known to produce good
properties. Thus, in this study, only steel and/or cast iron systems, (i.e. well known Fe-C base
alloys), were considered. In the studies reported here, these included compositions in the Fe-C
and Fe-C-Si systems.
Equilibrium Phase Relations
The equilibrium phase relations are important because they essentially indicate the permissible
infiltrant and base compact compositions to be used as well as specifying their respective
melting points. These data along with other information that the phase relations provide are
needed to design the infiltration process. In this study, “ThermoCalc”, a commercially available
thermodynamics program that calculates the properties of alloys from the known properties of
their components, was used to estimate the equilibrium phase relations of each of the systems
of interest, (9). As an example, the equilibrium phase relations of the binary Fe-C system in the
vicinity of the eutectic composition are shown overleaf in Figure 1.
Based on a cursory review of this figure, it may seem that there are any number of compositions
in the system to choose from in selecting the infiltrant and base compact compositions. In fact,
however, the actual choice of compositions is very limited. For example, suppose that the
liquidus and solidus compositions on the isotherm at 1200 oC were selected as the infiltrant and
base compact compositions and that the process temperature was correspondingly set at
o
1225 C. As it turns out, careful consideration of the phase relations indicated in the figure will
show that both of these compositions as well as the selected process temperature are
problematical.
Figure 1 - Fe-C Binary Diagram
Consider the infiltrant composition first. According to the indications of the figure, it will first start
to melt at the eutectic temperature, (i.e. at 1153 oC), dividing as it does so into liquid and solid
phases of the eutectic liquidus and solidus carbon contents. Based on the lever rule and the
compositional values indicated in the figure, the residual solid phase at this point will constitute
about 20% by weight of the original infiltrant. The indicated difficulty, as confirmed by
experiment, is that after the liquid phase forms or perhaps coincident with its formation, it
infiltrates the base compact leaving this solid phase behind which because of its low carbon
content will now not melt completely at the process temperature, (i.e. 1225o). In fact, according
to the phase relations indicated in the figure, it melts over a range of temperatures with its final
melting point being about 1400 oC. Thus, the implication is that in order to effect complete
infiltration at a reasonable temperature, the selected infiltrant composition must be near or equal
to the eutectic liquidus value.
Now consider the same base compact composition and process temperature but, as suggested
by the foregoing, that they are combined with an infiltrant of the eutectic liquidus composition. In
this case, when the system attains the eutectic temperature, the infiltrant will melt completely
and, as previously, start to infiltrate the base compact. However, here again, it may not
complete the process. The difficulty in this instance is the tendency for diffusional solidification
which will adversely affect the infiltration rate and may even prevent the process altogether. As
explained below, diffusional solidification is directly the result of differences in the base compact
and the equilibrium solidus compositions.
As will be recalled, the base compact composition was taken as the solidus value at 1200 oC.
According to the indication of Figure 1, the solidus carbon content at this temperature is
generally lower than the carbon contents needed for equilibrium with liquids at all lower
temperatures including, in particular, that of the solidus at the eutectic temperature.
Consequently, as the eutectic liquid enters the base compact, the system will attempt to effect
equilibrium between the two by transferring carbon from the liquid to the solid. In accordance
with the phase relations, this transfer will be accompanied by a partial freezing of the liquid
coincident with the depletion in its carbon content vis-à-vis the required liquidus value. The
extent to which the liquid solidifies will depend on the magnitude of the carbon differences
involved. Subsequently, since the temperature at this point is continuing to increase towards the
process temperature, the same units will ordinarily re-melt coincident with further temperature
increase and the associated decrease in the carbon requirement of the liquidus. Thus, the
partial freezing process described will not necessarily stop the infiltration process but since it
occurs along the plane of the liquid front as the liquid advances into the compact, it can be
expected to impede its progress. Based on studies of the effect in the Fe-C system, the critical
parameter appeared to be the magnitude of the difference in the carbon content of the base
compact and that of the eutectic solidus. The indications were that if this difference is much
greater than about 0.10%, the infiltration process is either stopped completely or is sufficiently
slowed to prevent it from going to completion in a reasonable time. Thus, the choice of the base
compact composition is linked to that of the infiltrant and must be decided accordingly.
The choice of the process temperature is also limited. For instance, if it is set too low, then the
units of the infiltrant which freeze during infiltration due to the difference in the carbon content of
the base compact and the equilibrium solidus value may not re-melt sufficiently to facilitate
complete infiltration. Or if it is set too high, it may lead to excessive liquid phase formation. The
potential drawbacks of too high a liquid phase content, (e.g. > ~25%), are that it may promote
microstructural coarsening which is detrimental to mechanical properties or, in a worse case
scenario, lead to slumping or other undesirable shape changes, (10).
Finally, the eutectic in the Fe-C system is a three phase equilibrium. As it turns out, the
eutectics in each of the other alloy systems of interest are also three phase equilibriums.
Consequently, although their equilibrium phase relations are generally more complicated, they
are similar in the sense that precisely the same infiltration process design considerations apply
as regards the selection of the critical process parameters to be used in these systems.
Base Compact Density
The density of the base compact along with its weight determines the volume of the pores to be
filled by the infiltrant. The density also determines the so-called open or interconnected porosity
which is a measure of the fraction of the pores that are accessible to the surface of the compact.
In general, the open porosity is a decreasing function of density but fortunately the function is
non-linear and the greatest rate of decrease occurs at high densities, typically in excess of 90%
of the theoretical maximum or so-called pore free density, (11). Thus, it’s essential in order to
optimize the density potential of the infiltration process to limit the density of the base compact
to a value that is equal to or less than about 90% of the pore free value.
In Fe-C alloys, the pore free density is largely dependent on the carbon content. In accordance
with the earlier discussion of its equilibrium phase relations, the base compact carbon content
that is indicated for use in the Fe-C system is about 2%. As will be seen, the equilibrium phase
relations of each of the other alloy systems of interest all indicated lower but essentially similar
values. Thus, it was decided to use a common base compact density in the initial studies. With
graphite as the carbon source in the base compact, it turns out that the 2% value mentioned
corresponds to a pore free density of 7.49 g/cm3. Hence, the maximum base compact density
was set at 6.8 g/cm3 which is slightly above 90% of this value.
Infiltrant Weight
The infiltrant weight to achieve maximum density is given by the product of the density of the
infiltrant and the pore volume of the base compact at the infiltration temperature. Although the
infiltrant density can be estimated with reasonable accuracy, the pore volume parameter that is
needed in this relation can not. This is primarily because the base compact is subject to
unpredictable volume changes due principally to graphite solution and to sintering during
heating in advance of infiltration. As a consequence, the infiltrant weight can not be determined
without resort to experiment. However, it can be approximated on the basis of the ambient
temperature values of the indicated parameters. Owing to the graphite solution effect, this will
inevitably underestimate the full weight value, typically by 15 to 20%. Nevertheless, it provides a
consistent starting point to determining the required weight that is generally applicable to parts
of different density, geometry, and weight. The indicated calculation procedure is set out in
detail in the Appendix.
Process Conditions
Once the infiltrant and base compact compositions are selected, it remains to prepare the
associated premixes and choose the process temperature, the time at temperature and the
furnace atmosphere.
Premix Preparation – Preliminary studies indicated that it is absolutely essential to minimize
gross variations due to segregation by de-mixing in both the infiltrant and the base compact as
well as to prevent significant carbon losses due either to dusting and/or decarburization during
processing, especially, in advance of infiltration. To exemplify, the indications were that graphite
segregation in the infiltrant caused uneven and incomplete melting which often led to localized
erosion of the infiltrated surface and incomplete infiltration. Similarly, segregation in the base
compact typically caused random defects due to local melting in its un-infiltrated surfaces and
appeared to contribute to localized erosion of the infiltrated surface as well. Carbon losses in
either case, but especially in the infiltrant, normally led to incomplete infiltration. Consequently, it
is absolutely essential to prepare the premixes of both the infiltrant and the base compact as
binder-treated compositions as well as to add additional graphite to each to offset the losses
due to carbon reduction of the residual oxides of the base powder. As described below, special
precautions were also taken to prevent decarburization by oxygen impurities in the furnace
atmosphere.
Process Temperature – The choice of the process temperature is based on two considerations.
In accordance with the earlier discussion, it must be high enough to insure that units of the
infiltrant which freeze during infiltration will re-melt yet low enough to prevent excessive liquid
phase formation once infiltration is complete. Values which satisfy these criteria must be
determined from the applicable phase relations and the compositional details of the system.
Time at Temperature - This particular parameter remains to be studied in detail. Based on the
experimental evidence to date, infiltration per se appears to occur in a relatively short span of
time, perhaps, in as little as 5 minutes or less in the Fe-C and Fe-C-Si systems but is known to
take much longer in other systems. In addition, there have also been indications that it is
important to approach the infiltration temperature slowly so as to effect a uniform temperature in
the base compact. Moreover, when infiltration is complete, the system naturally enters a liquid
phase sintering stage that is applicable both to consolidate the early sinter bonds of the base
compact as well as to eliminate any residual porosity that may yet be present. On the other
hand, once these effects are realized, the potential for microstructural coarsening suggests an
imperative to minimize the time. Consequently, the times at temperature that have generally
characterized the studies to this point have seldom exceeded 30 minutes and have on occasion
been as short as 15 or 20 minutes.
Furnace Atmosphere - Here again, this parameter has yet to be investigated in detail. In studies
to date, either hydrogen or synthetic dissociated ammonia, (i.e. 75% H2 and 25% N2 by volume),
which happen to be the atmospheres of common usage in this laboratory have been used. It
remains to study nitrogen based atmospheres as these provide significant economies relative to
the hydrogen based ones and consequently tend to be the P/M industry’s atmospheres of
choice. Present concerns as regards the nitrogen atmospheres are their relatively poor heat
transfer characteristics and most importantly, their as yet unknown effects on the wetting and
spreading properties of the infiltrant. If either of these prove to be problematic, it will be
necessary to revert to the more costly atmospheres.
In any event, independently of base chemistry, the dew point and carbon potential of the
atmosphere otherwise appear to be the most important parameters affecting the outcome of the
process. Control of each of these is necessary to prevent decarburization in advance of
infiltration. This is especially true of the infiltrant since decarburization will normally manifest as
incomplete infiltration. Decarburization after infiltration, of course, is also undesirable but is not
as critical to the outcome of the process.
In any case, to prevent decarburization by the atmosphere, the carbon potential in the furnace
must be of the order of the carbon potential of graphite. This is because until the infiltrant and
base compacts attain the eutectic temperature much of the carbon they contain is present as
graphite. In fact, in the case of the infiltrant, most of the carbon is present as graphite. In heating
to the eutectic temperature, the balance of the carbon in both compositions dissolves in the iron
and is correspondingly less susceptible to oxidation. For example, oxidation of carbon in
solution can often be prevented by use of an atmosphere of sufficiently low dew point or one
having both an high hydrogen potential and a low dew point. However, this is not the case with
carbon as graphite which, at the temperatures typical of P/M processing, will reduce hydrogen
from water vapor regardless how little water is present in the atmosphere or how high the
hydrogen potential opposing the reaction.
Of course, an atmosphere of low oxygen content and hence of a low dew point is nevertheless
essential to minimize the potential for graphite oxidation regardless of what other precautions
are taken. However, the only possibility to prevent graphite oxidation altogether is to increase
the carbon potential by introducing a carbon containing compound, usually an hydrocarbon,
which thermodynamics indicates(,) has a greater susceptibility to oxidation than graphite.
Methane and propane, for example, are two such compounds. Both spontaneously decompose
to their constituent elements at high temperature and in thermodynamic terms are therefore
more susceptible to oxidation than graphite. Of course, the amount of either that is needed in a
particular case will depend in part on the oxygen purity of the base atmosphere and in part on
the ‘oxygen tightness’ of the furnace. Since the latter typically varies from furnace to furnace
according to design and maintenance, the actual percentage additions that are needed are
generally not quantifiable. As a practical matter, they are ordinarily determined empirically by
trial and error. Moreover, as those familiar with the art will agree, even if they were quantifiable,
its highly likely that they would be determined empirically anyhow.
Another method to prevent graphite oxidation and the one that was used exclusively in the
present studies is to enclose the parts in a graphite gettered box such as a ceramic sintering
tray with a close fitting cover and process them accordingly. Of course, this method is only
applicable on a small scale. However, it is especially practical in a laboratory environment
where owing to the necessity to share equipment, it’s inconvenient or inappropriate to add an
hydrocarbon to the furnace atmosphere.
Materials, Procedures and Equipment Specific to the Present Study
The base powder used in the studies was a standard Hoeganaes Ancorsteel 1000B powder.
As mentioned, infiltrant and base compact compositions in the Fe-C and Fe-C-Si systems were
investigated. These compositions were in all cases prepared as binder treated admixtures of the
aforementioned base powder with two or more of the several admix ingredients as listed below.
The binder treatment processing was generally in accord with the standard Hoeganaes
ANCORBOND process, (12). The infiltrant and base compact mix sizes were typically 200 and
1000 grams respectively. The admix ingredients mentioned included graphite, lubricant, and
SiC. The graphite was Asbury grade 3203, a natural flake type graphite with a minimum carbon
content of 95% and an average particle size of less than 10 micrometers. The lubricant was
Acrawax C, the standard PM grade of the Lonza Division of IMS Company. The SiC was SaintGobain Ceramics Company grade F-600, a commercially pure SiC nominally containing 70%
silicon and 30% carbon and having an average particle size under 15 micrometers.
The base compacts of the study were compacted to various densities that were typically equal
to or less than 6.8 g/cm3. The compacts were in all cases in the form of standard Transverse
Rupture Strength specimens, (ASTM 528), but to a nominal constant weight of 35 grams
throughout, (i.e. to a nominal heights in the range of 12.5 to 14 mm). The infiltrant slugs were
compacted in the same form using a standard pressure throughout of 550 MPa. Their weights
were typically in the range of 3 to 5 grams varying in accordance with the result of the Infiltrant
Weight calculation as set out in the earlier referenced appendix and other considerations as will
be discussed.
The specimens were infiltrated in an high temperature Hayes pusher type furnace. As
discussed above, the specimens were processed in graphite gettered boxes in either a
commercially pure hydrogen or synthetic dissociated ammonia atmosphere. Process
temperatures varied according to the aims of the particular trial but were typically in the range of
o
o
1155 to 1195 C, (2110 to 2185 F). Process times varied likewise but again were typically of
the order of ½ hour at temperature.
RESULTS AND DISCUSSION
The studies that are reported here started with a limited investigation of the Fe-C system and
continued to a fairly detailed investigation the Fe-C-Si system. The early work in the Fe-C
system was basically in the nature of defining studies with the objectives to delineate the
important process variables and to assess the potential of the process in terms of the maximum
achievable density. The subsequent research into the effects of silicon as an alloy addition had
the objective to effect essential microstructural improvements relative to the Fe-C system with
the ultimate aim being to produce the best possible mechanical properties. In the course of
these studies, a peculiar dimensional change effect was observed that had been overlooked in
the early studies. The effect was in the nature of a dimensional non-uniformity and large enough
to negate the net shape advantage of the process. In what follows, the major findings of the
studies of the two alloy systems are reported along with a description of the aforementioned
effect and an abstract of the research that since been conducted with regard to it. The general
aim of the report is to present as accurate a picture of the present state of development of the
technology as is reasonably possible.
The Fe-C System
The studies of the Fe-C system essentially had the objective to establish the important process
variables but, in fact, were never intended or expected to be conclusive since it was anticipated
that alloy additions other than carbon would eventually be necessary to achieve reasonable
properties. Moreover, as may be evident, many of the findings of these studies formed the basis
of the Alloy and Process Design section of the Experimental Procedure and have consequently
already been indicated. Three important findings, however, which have yet to be presented
include: 1) the maximum achievable density of the process; 2) early indications as to the
existence of densification mechanisms other than simple pore filling; and, 3) the microstructures
and properties of the resulting parts.
The theoretical maximum or pore free density of an Fe-C alloy is dependent on its carbon
content, the microstructural constituent which the carbon precipitates and the density and
content of the Fe phase which composes the balance of the microstructure. If the carbon
containing precipitate is assumed to be cementite, (i.e. Fe3C), which is typically the case in PM
processing, then it is easily shown that the pore free density of the alloy, ρFe-C, can be calculated
as a function of the carbon content, %C, in accordance with the following expression:
1)
1/ρFe-C = 1/ρFe + 0.1495%C[1/ρcementite - 1/ρFe].
where ρFe and ρcementite are the pore free densities of the constituent phases and the numerical
constant on the right is 1/100 the quotient of the molecular weights of the Fe3C and C, (i.e.
179.56 and 12.01 respectively).
According to this result, the maximum achievable density of an infiltrated part will depend on the
final carbon content of the part. However, since the final carbon content is also dependent on
the infiltrant weight and since, as earlier indicated, the infiltrant weight to full density can not be
determined without recourse to experiment, the final carbon content and hence the maximum
density to be expected in the part is essentially unpredictable. Consequently, experiments with
the objective to determine the maximum achievable density by infiltration must be designed and
interpreted accordingly.
In a series of trials involving increasing infiltrant weights and hence increasing final carbon
contents, the details of the last trial in the series were as follows. The base compacts contained
nominally 2% carbon which is just under the eutectic solidus content of 2.03% and were
pressed to 6.8 g/cm3 and weighed 35 grams. The infiltrant contained 4.34% carbon which is
equal to the eutectic liquidus value and weighed 4.5 grams. As a matter of interest, the latter
value is ~0.8 grams greater than the infiltrant weight in accordance with the appended
o
o
calculation procedure. Infiltration was at 1180 C, (nominally 2150 F), for 1/2 hour at
temperature in synthetic DA. The results of the trial are shown below in Table 1.
Table 1 - Infiltrated Properties of an Fe-C Alloy at an Average Carbon Content of 2.27%
Specimen
Density
Dim. Chg. vs Green
Number
g/cm3
%
7.63
-0.54
1
7.59
-0.37
2
Average
7.61
-0.46
Taking the density of a commercially pure iron, (e.g. Ancorsteel 1000B), as 7.85 g/cm3 and that
of cementite as 7.4 g/cm3, the pore free density according to Eq. 1 and the indicated carbon
content is 7.69 g/cm3. In comparison, the observed average density of 7.61 g/cm3 is just under
99% of this value. The implication is that if the infiltrant weight had been greater by about 1% of
the final total infiltrated weight, (e.g. by ~0.4 grams), it would have been sufficient to fill the
remaining pores and effect infiltrated densities that approached the theoretical limit. However,
there is also evidence in the data to suggest that simple pore filling is not all that is involved in
the process. Based on the dimensional change values, it’s apparent that sintering also made a
significant contribution to the observed densification. Thus, while the results clearly show that
the infiltration process is capable of producing densities that approach the pore free value, it’s
equally clear that the underlying mechanism is not a simple volume displacement process but
includes densification by solid state and, very probably, liquid phase sintering as well.
Figure 2 overleaf is a micrograph of a typical Fe-C alloy in the as-infiltrated condition. The
relative density in this case is just under 98%. Apart from the pores, the evident microstructural
features shown in the figure include a predominantly pearlitic matrix in an essentially continuous
network of hyper-eutectoid grain boundary carbides. Owing to the presence of the grain
boundary carbides, the mechanical properties of the alloy were not expected to be much better
than those of a standard low density press and sinter composition of similar pearlite content and
were consequently not determined. It was likewise evident that it would be necessary to find
suitable ways to modify the structure and, in particular, to disrupt or, better yet, eliminate the
grain boundary carbides if iron base infiltration was to survive as a practical matter.
The best known metallurgical methods to modify a microstructure include alloying, heat
treatment and hot working. Of the three, alloying is intrinsically the most economic but is also
the least reliable since, in general, there is no way as yet to predict what the effects of an alloy
addition will be on the precipitation behavior of a particular phase. Fortunately, however, the
phase in this instance, (i.e. Fe3C), is metastable relative to graphite and it is well known in the
cast iron industry how to modify it by graphitization.
The graphitizing elements in order of decreasing graphitizing power reportedly include: tin,
phosphorus, silicon, aluminum, copper and nickel, (13). Of these, tin, phosphorus and copper
were eliminated from further immediate consideration because their ternary phase relations are
complicated by the presence of low temperature eutectic or peritectic reactions. Aluminum, of
course, was eliminated because of its high affinity for oxygen. Remaining, therefore, were
silicon and nickel. As it happens, both have ternary phase relations that are similar to those of
the Fe-C system. Since silicon was listed as being the more powerful graphitizer of the two, it
was decided to work with it first.
Figure 2 - Fe-C Composition Infiltrated to ~7.52 g/cm3
- Nital/Picral @ 200X
The Fe-C-Si System
Virtually all trials in this phase of the study were at a silicon content of 1.05%, (i.e. 1.5%
admixed SiC). According to the Thermocalc program, the corresponding eutectic liquidus and
solidus carbon contents are 4.01% and 1.87% respectively and the eutectic temperature is
nominally the same as in the Fe-C system, (i.e. 1153 oC ≈ 2107 oF). Results typical of the
several studies that were done in this system are presented in Table 2.
Table 2 - Infiltrated Properties of an Fe-C-Si Alloy at an Average Silicon Content of 1.05%
Density
Dim. Chg. vs Green
Specimen
Number
%
g/cm3
7.53
0.48
1
7.56
0.63
2
Average
7.54
0.56
The base compacts corresponding to the data in the table contained 1.75% carbon which is
0.12% below the eutectic solidus value. They were pressed to 6.7 g/cm3 and weighed 35 grams.
The infiltrant nominally contained 4.05% carbon which is just above the eutectic liquidus value
and weighed 5.25 grams. The latter value is ~1 gram greater than the infiltrant weight in
o
o
accordance with the appended calculation procedure. Infiltration was at 1163 C, (2125 F), for
1/2 hour at temperature in synthetic DA.
The findings in this instance can not be properly interpreted without reference to the
microstructure. For example, while the density values are lower than earlier, it turns out that this
is essentially a graphitization effect and contrary to being inferior, they are, on a relative density
basis, actually slightly better than earlier. The microstructure is shown below in Figure 3.
Figure 3 - Fe-C-Si Composition Infiltrated to 7.54 g/cm3 - Nital/Picral @ 200X
A cursory comparison of the structural details in this figure with those of the earlier Figure 2 will
show that the silicon addition had a profound effect. Amazingly, it produced an essentially
ductile cast iron structure, (14). The eutectoid or near eutectoid pearlitic matrix that was seen in
the earlier Fe-C alloy remains but the grain boundary networks of hyper-eutectoid carbides have
virtually all been replaced by a random dispersion of graphite precipitates. The graphite
precipitates that are most evident in the figure are of the so-called ‘bull’s-eye’ variety. This type
occurs chiefly in ductile or nodular cast irons and consists essentially of a spheroidal graphite
nodule within an encapsulating annular sphere of ferrite. Less evident but also present in this
micrograph and, more generally, in the numerous others that have been examined in this study
are so-called vermicular or compacted graphite precipitates as well as occasional flake type
precipitates. The latter morphologies occur chiefly in the so-called compacted and gray cast
irons.
Commensurate with the change in microstructure, the precipitation of graphite also changes the
pore free density of the alloy. Since the density of graphite is lower than that of the carbide, the
general effect of increasing degrees of graphitization is to decrease the pore free density. The
magnitude of the effect as determined on the basis of the pore free densities of the constituent
phases is shown below in Table 3. The total carbon value in the table is nominally the same as
that of the subject composition.
Table 3 – Effects of Graphitization on Infiltrated Pore Free Density and Microstructure
Pore Free
Microstructure
Composition
Total Carbon
Graphite
Fe3C
Density
Graphite
Fe3C
Fe
3
%
%
%
g/cm
Volume Fractions
2.05
0
30.7
7.71
0.0%
31.9%
68.1%
2.05
25
23.0
7.65
1.7%
23.8%
74.5%
2.05
50
15.3
7.59
3.4%
15.7%
80.9%
2.05
66
10.3
7.57
4.4%
10.6%
85.0%
2.05
75
7.7
7.54
5.0%
7.8%
87.2%
2.05
100
0.0
7.48
6.6%
0.0%
93.4%
The microstructure in the present case approximates to complete or near complete
graphitization of the hyper-eutectoid carbon content of the alloy. Thus, assuming cooling typical
of normal PM processing, the corresponding pore free density of the alloy may be estimated on
the basis of the equilibrium eutectoid carbon content of the alloy and the findings in the above
table as follows. According to the indications of the Thermocalc program, the eutectoid carbon
content of an Fe-C-Si alloy at 1% Si is 0.69%. This is equivalent to just under 34% of the
indicated total carbon value. Hence, the degree of graphitization to effect complete graphitic
precipitation of the hyper-eutectoid carbon and produce the indicated microstructure would be
just over 66%. Thus, as indicated in the highlighted row of data in the table, the corresponding
pore free density is about 7.57 g/cm3.
As a matter of interest, since the carbon contents of casting alloys are necessarily higher than
those typical of the infiltration process, their as-cast densities are correspondingly lower. For
example, the pore free density of a fully pearlitic ductile cast iron at a typical carbon content of
3.8% is about 7.25 g/cm3. The implication is that parts with ductile iron structures that are
produced by infiltration have the potential to exhibit better mechanical properties than their ascast counterparts, including particularly, higher elastic modulus values which have been shown
to be especially sensitive to the carbon content, (15).
Now, returning to the infiltrated properties in Table 2, it will be evident that the observed density
values approached the pore free density and, on a relative basis, are therefore comparable to
the earlier results in the Fe-C system. On the other hand, in contrast with the earlier indications
of significant densification by sintering in addition to infiltration, the present dimensional change
values are positive and, of course, give no indication of a sintering contribution. Presumably, the
relative increase in these values is another effect of the observed graphitization.
The Potential To Supplement Infiltration By Liquid Phase Sintering
Once infiltration is complete, the average compositions resulting from the infiltrant and base
compact compositions comprise supersolidus liquid phase systems. Thus, it’s potentially
possible to combine iron base infiltration with liquid phase sintering to obtain results that are not
possible by either process alone. For example, if the infiltrant weight is initially adjusted so that
residual porosity exists after infiltration, then additional densification is possible by liquid phase
sintering. The advantage in combining the two processes lies in the potential to manipulate the
final dimensional change of the resultant parts and still achieve full density.
The general process changes that are needed to implement the indicated combination of the
two processes include: 1) reducing the infiltrant weight below the weight that would effect full
density; and, 2) employing either a two step process involving infiltration at one temperature and
liquid phase sintering at a higher temperature or simply increasing either or both the infiltration
temperature and the time at temperature. Otherwise, as will be seen, the specific changes that
are required in a particular case must be determined experimentally.
Results typical of the several studies that were conducted to examine the effects of combining
the two processes are shown in Table 4. The compositional and geometric details of the base
compacts and infiltrant that were used in this particular study were precisely the same as those
that were used to generate the data in Table 3 of the earlier study. However, in this case, three
different infiltrant weights were employed. The highest weight at 5.25 grams corresponded to
the weight to effect full density without benefit of significant liquid phase sintering as established
in the earlier study. The two remaining weights nominally corresponded to consecutive 15%
decrements of this value. Infiltration was at 1163 oC, (2125 oF) for 15 minutes at temperature
followed by liquid phase sintering at 1182 oC, (2160 oF), again for 15 minutes at temperature.
Shown also in the table are the associated total carbon contents and the liquid phase contents
at the higher temperature. As will be explained, in addition to the infiltration weight and the
process conditions, these parameters also affected the outcome of the trial.
Table 4 – Effects of Infiltration at 1163 oC Followed by Liquid Phase Sintering at 1182 oC
Liquid Phase
Infiltrated
Dim. Chg. vs
Infiltrant Weight Total Carbon
Content at 1182 oC
Density
Green
3
grams
%
%
grams/cm
%
2.05
16.6
7.54
0.59
5.25
2.01
14.6
7.57
0.13
4.50
1.97
12.7
7.43
- 0.01
3.75
According to these findings, the dimensional change decreased with decrease in the infiltrant
weight and did so without significant adverse effect to the final density, especially at the
intermediate weight. Thus, the data generally confirmed the expected greater contribution of
sintering to the outcome of the processing. The slightly higher final density at the intermediate
infiltrant weight and the lower density at the lowest weight are each thought to be attributable to
the decrease in the total carbon which the data show accompanied the weight changes. As to
the first effect, as previously indicated, the pore free density of these alloys increases with
decrease in the total carbon and, as it turns out, the slight increase in the density at the
intermediate weight that is indicated here is just accounted for by the accompanying decrease in
the total carbon value. In the case of the low density value at the lowest infiltrant weight, the
connection to the total carbon is less direct. Evidently, the amount of sintering that occurred in
this case was not sufficient to eliminate all of the residual porosity that was created by use of the
low infiltrant weight. As a general matter, the densification that occurs in liquid phase sintering
varies directly as the liquid phase content, (9). Thus, the low final density in this case is
apparently attributable to the accompanying decrease in the liquid phase content as shown in
the data. However, as may already be evident, the liquid phase content is determined by the
phase relations and, at a given temperature and base alloy content, is entirely a function of the
associated carbon content in accordance with the well known lever rule.
The Distortion Effect
The normal procedure in this laboratory for determining the volume and/or dimensional change
of a TRS specimen is to measure its dimensions at roughly the centers of the conjugate
transverse faces. In effect, this method involves the implicit assumption that the specimen is
dimensionally uniform. However, in view of the novelty of the infiltration process, it was decided
at the start of the studies of the present alloy system to make more careful determinations and,
in particular, to include dimensional uniformity checks in addition to the usual measurements. As
it turned out, the very first checks of this property showed the existence of a type of dimensional
non-uniformity that may be unique to iron base infiltration and which subsequently came to be
called the ‘distortion effect’.
As a general matter, the distortion effect is characterized by density gradients in the through
thickness direction of the infiltrated compact that are manifest as differences in its lateral
dimensions. The greatest variations always occur in and just under the infiltrated surface to a
depth of a few millimeters but smaller variations may occur elsewhere as well. As a result, the
magnitude of the effect is measured simply as the difference in the lengths of the infiltrated and
opposing uninfiltrated surfaces. Typically, the effect is large enough that if not otherwise
mitigated, the resultant parts will require a machining step before they can be put into service in
all but the least demanding applications. For example, the distortion values typical of the
specimens of the present studies ranged from 0.1 to 0.4 mm, (0.004 to 0.015 ins.), with
occasional higher values up to a maximum of ~0.5 mm, (0.02 ins.).
Subsequent to the discovery of the distortion effect, a very substantial research effort was made
to mitigate it but regrettably, with only partial success. In view of the complexities that the effect
was found to involve, a detailed review of this research is beyond the scope of the present
effort. The current plan, however, is to report it later this year at the European PM conference.
In the interim, the essential findings of the several studies that were done in this connection
were as follows.
It was determined that the effect has two general causes. The primary cause is liquid
penetration and separation of the sinter bonds of the particles in and just under the surface of
the base compact followed by lateral expansion of the affected elements under the influence of
the surface tension forces that act on the uninfiltrated liquid. The secondary cause is incomplete
graphitization of the hypereutectoid carbon content of the compact. Distortion due to the liquid
penetration mechanism is always observed and is normally fairly substantial in magnitude. In
comparison, distortion due to incomplete graphitization only occurs intermittently and is
generally of a smaller magnitude.
Analysis of the ‘liquid penetration’ mechanism suggested that it was susceptible to two different
alloying strategies. One was to prevent the operation of the mechanism altogether and the other
was to impede it until infiltration was complete. Of the two, the second strategy appeared to be
the simpler and was subsequently tested. However, owing primarily to the failure of the
materials systems that were selected to implement it to behave as expected, all such efforts
were unsuccessful. In contrast, the incomplete graphitization mechanism was found to be
amendable principally, to the alloy content of the graphitizing element and to a lesser degree, to
processing as well.
The major conclusion of these studies was that while distortion due to the liquid penetration
mechanism may yet prove to be preventable by alloying, a near term solution by such methods
appears to be unlikely. Moreover, since this type distortion is entirely manifest in and just under
the surface of the part, die design may offer a viable alternative means to mitigate it. Thus, it
makes sense to continue the further development of the technology, on the one hand, with due
regard for the existence of this effect but on the other, without significant additional delay to deal
with it as a research issue.
SUMMARY AND CONCLUSIONS
The development to date of the iron base infiltration process as a pioneer technology was
presented. Initially, the five basic elements that are needed to design an infiltration process and
the general choices that were made in each case were discussed. The five basic elements
mentioned included the alloy systems, the relevant equilibrium phase relations, the maximum
base compact density, the infiltrant weight to effect full density, and the selection of the
applicable process conditions.
The alloy systems studied were generally limited to the simplest possible steel and/or cast iron
compositions known to produce good properties. The investigation started with defining studies
in the Fe-C system and advanced to alloys in the Fe-C-Si system. As confirmed by experiment,
the equilibrium phase relations of the Fe-C system indicated that the applicable infiltrant and
base compact compositions were limited to the eutectic liquidus and solidus or near hyposolidus compositions. Significant compositional deviations in either case led to incomplete
infiltration due apparently to diffusional solidification. The selection of the infiltration temperature
was likewise limited. In the case that the base compact composition was in the hypo-solidus
range, the temperature had to be set high enough to reverse diffusional solidification at lower
temperatures in order to assure complete infiltration. On the other hand, it was also necessary
not to set the infiltration temperature too high to avoid the adverse effects of too much liquid
phase formation subsequent to infiltration. To ensure that the total pore content of the base
compact was readily accessible to the infiltrant, the maximum green density in the study was set
at 90% of the corresponding pore free value, (i.e. at 6.8 g/cm3). In actual practice, most trials
were based on a base compact density of 6.7 g/cm3. In contrast, the infiltrant weight to full
density was indeterminate without recourse to experiment. However, a first order estimate which
predicted the weight to within 80% of the experimentally indicated value proved helpful to
conduct the trials in a systematic manner. Finally, other than the infiltration temperature which
was decided in accordance with the phase relations and the infiltration time which was set
somewhat arbitrarily at about 1/2 hour at temperature in most cases, the balance of the process
conditions were largely dictated by the laboratory environment and equipment. The specimens
were processed in a pusher type furnace under cover of hydrogen or a synthetic dissociated
ammonia atmosphere. In lieu of the possibility to control the carbon potential of the furnace
atmospherically, graphite gettered and covered trays were used to prevent decarburization in
advance of infiltration, especially of the infiltrant composition.
Other than confirming the indications of the phase relations with regard to the permissible
compositions and helping to define the applicable temperature range of the process, the initial
studies of the Fe-C system also showed that the maximum achievable density approached the
pore free density and that in addition to simple pore filling, some of the densification was due to
sintering and very probably, to liquid phase sintering subsequent to infiltration. Metallographic
examinations of the resulting parts, however, showed a predominantly white cast iron structure
and in view of the well known fact that such structures are inherently brittle, it was evident that
either alloying or additional processing would eventually be necessary to effect good properties.
Based on the success of the Cast Iron Industry to modify such structures compositionally by the
use of graphitizing elements, it was decided to try the alloying approach. Of the several alloying
elements that are known to be effective in this regard, silicon and nickel were selected for study
primarily because of the similarity of the phase relations of the associated ternary systems to
those of the Fe-C system.
Initial trials with alloys in the Fe-C-Si system confirmed their general applicability to the iron
base infiltration process as well as demonstrating the effectiveness of the silicon to modify the
resulting microstructure. Infiltrated densities were typically in the neighborhood of 7.55 g/cm3
and metallographic examinations showed the presence of a predominantly nodular graphite or
so-called ductile cast iron structure. The slightly lower densities of these alloys relative to the
earlier Fe-C alloys were explained on basis of the observed graphitization of the hyper-eutectoid
carbides, (i.e. Fe3C), and the lower density of graphite compared to the carbide.
The alloys in the Fe-C-Si system were also used to demonstrate the potential to supplement
infiltration by liquid phase sintering with the aim being to manipulate the final dimensional
change of the part. It was shown that comparable densities with decreasing dimensional change
values could be achieved by decreasing the infiltrant weight below the full density weight and
adding a liquid phase sintering step to the process. In the example study, this involved
increasing the process temperature by about 20 oC, (36 oF), and holding for a short additional
time after the infiltration step. Alternatively, simply increasing the infiltration time has also been
found to be effective in many cases.
Finally, a dimensional change anomaly of a sufficiently pernicious nature as to negate the net
shape advantage of the process was described. Although presentation of the very substantial
efforts that were subsequently made relative to it was generally beyond the scope of the report,
the key findings of this research were briefly outlined. It was indicated that this so-called
distortion effect had two causes, a primary one and an intermittently occurring secondary one.
Analysis of the underlying mechanism of the primary cause indicated that it was amendable to a
particular alloying strategy. However, owing to an unfortunate choice of the materials system
selected to implement this strategy, the subsequent efforts to mitigate the associated distortion
failed. In contrast, efforts to mitigate the distortion due to the secondary cause yielded to both
alloying and processing. Based on the difficulties with the primary cause, it was concluded that
a near term metallurgical solution of the distortion problem was unlikely and consequently, it
was appropriate to continue the development of the technology without further delay to deal with
it now.
APPENDIX
First Order Estimate of the Infiltrant Weight to Full Density
The infiltrant weight to achieve maximum density is, as earlier indicated, given by the product of
the density of the infiltrant and the pore volume of the base compact at the infiltration
temperature. Although the infiltrant density can be estimated with reasonable accuracy, the pore
volume parameter that is needed in this relation can not. This is primarily because the base
compact is subject to unpredictable volume changes due to the α to γ transformation, graphite
dissolution and sintering in advance of infiltration. As a consequence, the infiltrant weight can
only be approximated. One way to do this with minimal recourse to experiment is to calculate it
on the basis of the ambient temperature values of the indicated parameters; the rationale being
that if the weight of material so calculated is sufficient to fill the volume in this case, it will be a
reasonable approximation of the actual case as well. The calculation may be generalized as
follows. The infiltrant weight is given by the product of the ambient temperature values of the
infiltrant density ρI and the pore volume of the base compact VP as follows,
1)
WI = ρI VP
The value ρI of the infiltrant density in this case must be calculated as a pore free value from the
infiltrant composition. If fIFe, fIC and fIA are the fractional weights of the iron, graphite and alloy
contents of the infiltrant and ρFe, ρC and ρA are their pore free densities, then:
ρI = 1/(fIFe / ρFe + fIC / ρC + fIA / ρA)
2)
The pore volume parameter VP is given by the product of the volume of the base compact VB
and its volume fraction of porosity, fpores including the volume fraction that is occupied by the
admixed organics. If WB is the weight of the base compact and ρG is its green density, then
VB = WB / ρG. The calculation of the pore fraction fpores is more complicated. Without going into
the details, the value of this parameter can be shown to be given by,
fpores = 1 - ρG (fBFe / ρFe + fBC / ρC + fBA / ρA)
3)
where fBFe, fBC, fBA are the fractional weights of the iron, graphite and alloy contents of the
compact and ρFe, ρC and ρA are as previously defined. As a matter of interest, the fractional
weight of the organics content of the compact cancels out in the operations leading to this
expression.
As an example, consider the infiltrant weight required to achieve maximum density in an Fe-C
alloy. Assume that the infiltrant is an admixture of graphite and iron and that its composition is
selected from Figure 1 in accordance with the earlier considerations as the eutectic value, (i.e.
as 4.34% C and the balance iron). Assume further that the base compact is a transverse rupture
bar that weighs 35 grams and, again in accordance with the earlier considerations, has a green
density of 6.8 g/cm3, a carbon content of 1.93% which reference to Figure 1 will show is 0.1%
lower than the eutectic solidus value, an organics content of 0.8% and the balance iron. Then,
given that the pore free densities of the graphite and the iron are 2.32 and 7.85 g/cm3, the
infiltrant density ρI will be found to be 7.114 g/cm3. In the case of the base compact, the volume
VB will be found to be 5.147 cm3, the pore fraction fpores to be 0.101 and the resulting pore
volume VP to be 0.519 cm3. Then, as is easily confirmed, when this value is combined with the
earlier infiltrant density value of 7.114, it yields an infiltrant weight of 3.7 grams. In the actual
case, the infiltrant weight to full density was found to be ~ 4.5 grams.
ACKNOWLEDGMENTS
Special thanks are due to the Ben Franklin Technology Partners of Pennsylvania for funding a
part of this research and to Messrs. W. B. Bentcliff, G. Golin and T. Murphy of the Hoeganaes
Laboratory for their help in obtaining the data and figures used in preparing the manuscript.
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