A New Process For Making High Density Parts
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 process referred to in the title of this paper is Iron Base Infiltration. It is new in the sense
that the P/M industry does not currently make parts by this process. Otherwise, a survey of the
open literature will show the existence of at least six earlier articles on the subject, all of which
were published before 1985, (1-6).
The findings of the initial efforts to develop the process as a viable parts making technology
were reported earlier this year at the 2004 International Conference on Powder Metallurgy and
Particulate Materials, (7). The purpose of the present report is to expand on an important
segment of the work that was only described in summary form in the earlier report. This had to
do with a dimensional change anomaly called the ‘distortion effect’ which, as far as presently
known, is unique to iron base infiltration. A brief review of certain of the earlier findings is
crucial to an understanding of the work that will be reported.
Summary of Relevant Earlier Development Efforts
In spite of the existence of earlier research on iron base infiltration, the general approach that
was taken in the present effort was to develop the process as a pioneer technology. Initially, the
five basic elements that are needed to design an infiltration process were considered. These
included the alloy system, 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 hypo-solidus
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. Simultaneously, it had also to be set low
enough 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 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.
In addition to 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, (i.e. 7.71 g/cm3), and that in addition to simple pore filling, some of the
densification was due both to solid state sintering in advance of infiltration and liquid phase
sintering thereafter. 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, (8). 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, (9). The slightly lower densities of these alloys relative to
the earlier Fe-C alloys was explained on basis of the observed graphitization of the hypereutectoid carbides, (i.e. Fe3C), and the lower density of graphite compared to the carbide.
The Distortion Effect
In view of the novelty of the iron base infiltration process, it was decided at an early stage of the
studies to include dimensional uniformity checks in addition to the usual part measurements that
are typically made in P/M research. As it turned out, the very first checks of this property
showed the existence of a type of dimensional non-uniformity that, as mentioned, 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 a result of density gradations in the infiltrated
compact that are manifest as a disparity in the lateral dimensions of the infiltrated and opposing
uninfiltrated surfaces. The greatest variations always occur immediately under the infiltrated
surface to a depth of a few millimeters but may occur elsewhere as well. As a consequence, 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.
Experimental findings, as later presented, indicated that there are two different causes that
contribute to the effect: a primary one; and, an intermittently occurring secondary one. The
primary cause appears to be liquid penetration and dissolution during the early stages of
infiltration of the sinter bonds existing between the particles in and just below the surface of the
base compact followed by lateral expansion of the affected elements under the influence of the
surface tension forces which act on the as yet uninfiltrated liquid. The secondary cause is
incomplete graphitization of the hypereutectoid carbon which is typically limited to the lower
regions of the infiltrated compact. This effect appears to occur intermittently and can only be
verified by metallography. When it is observed, it is usually manifest as a fairly abrupt change
from a relatively low density ductile cast iron structure, normally in the top and middle portions
of the compact, to a high density white cast iron structure in the bottom portions.
Subsequent work very quickly showed that the distortion due to incomplete graphitization was
amendable primarily to the alloy content of the graphitizing element and to a lesser degree, to
processing, including particularly cooling after infiltration. Thus, the distortion that remained and
the type that is the subject of the balance of the report was that due to the primary or ‘liquid
penetration’ mechanism.
General Alloying Approaches to Mitigating Primary Type Distortion
Based on the liquid penetration mechanism, two general strategies ostensibly existed to prevent
or, at least, to limit the resulting distortion. One was to manipulate the surface tension forces to
minimize the lateral spreading tendency of the infiltrant and the other, to prevent the penetration
and dissolution of the sinter bonds of the base compact until the infiltration step is complete. Of
the two, only the second is actually viable. As it turns out, the lateral spreading tendency of the
liquid under the action of surface tension forces can be shown to be directly proportional to the
cosine of the contact angle. The implication being that a large contact angle, (e.g. ~ 90o), is
needed to alleviate the distortion. However, the obvious difficulty with this idea is that the
infiltration process itself is critically dependent on the infiltrant’s ability to spread which, of
course, requires a low contact angle, (10).
In essence then, there was really only the one potential strategy; that being to prevent the liquid
from dissolving the sinter bonds of the base compact. Once again, however, there were two
possibilities to implement this strategy. One was to prevent dissolution of the sinter bonds
altogether by increasing the dihedral angle between the liquid and the solid. The other was to
forestall their separation until infiltration is complete by creating a sufficiently large
compositional imbalance between the infiltrant and the base compact to effect limited diffusional
solidification of the liquid as it first enters the compact. Upon careful consideration, it turned out
that neither of these possibilities was without difficulties. However, as outlined below, of the
two, the obstacles associated with the first appeared to be especially problematical and it was
decided to limit the subsequent studies accordingly.
There were three evident difficulties with the idea to increase the dihedral angle. They are as
follows:1) the dihedral and contact angles are not independent entities and it’s easily shown that
an increase in one is almost certain to lead to an increase in the other, (11); 2) the dihedral angle
and the interfacial energies on which it depends are relatively difficult properties to measure and
there are as yet neither sufficient data, especially as regards the compositions of present interest,
nor a reliable predictive theory on which to base efforts to change it; and, 3) increasing the
dihedral angle sufficiently to prevent separation of the sinter bonds during infiltration will almost
certainly limit the potential for significant densification by liquid phase sintering after
infiltration, (12).
In comparison, there was only one apparent drawback to implementing the second possibility
involving limited diffusional solidification of the infiltrant during the early stages of the process.
This was that the underlying theory of the compositional differences that are needed to
implement this strategy requires that the individual compositions each be homogeneous at the
moment of infiltration. If they are not homogeneous, the outcome of the differences between the
two, vis-à-vis the desired solidification, is impossible to predict. The obvious practical result of
this requirement was that the presence of alloying elements other than graphite in the associated
compositions must either be based on pre-alloyed powders to start or, in the case of admixes,
submitted to a homogenizing step in advance of infiltration. Since, of the two, the use of prealloyed powders appeared to offer potential economies in comparison with the second, it became
the preferred approach and has thus far been the subject of most of the work to date.
Diffusional Solidification Design To Mitigate The Distortion Effect
Silicon’s well known affinity for oxygen combined with a preliminary study that demonstrated
that as little as 1.8% pre-alloyed nickel was sufficient to effect graphitization in infiltrated
compacts led to the initial abandonment of the earlier Fe-C-Si system in favor of the Fe-C-Ni
system. In comparison with nickel containing pre-alloys, it was anticipated that the silicon alloys
would be difficult, if not impossible, to water atomize and anneal to the low oxygen contents that
appeared to be necessary to implement the technology.
The design of the alloys in this case was essentially based on the equilibrium phase relations of
the Fe-C-Ni system as indicated by the “Thermo-Calc” program, (13), and certain a’priori
considerations as outlined below. Combined with the above finding as to the nickel content
required to effect graphitization, they led eventually to the 2% and 3% Ni containing pre-alloys
as set out below in Table 2 of the Experimental Procedure section. According to the scheme that
was devised in connection with these powders, the 2% nickel containing alloy was the base
powder of the Base Compact composition while the 3% Ni alloy was the base powder of the
Infiltrant composition. The respective carbon contents in each case were taken as the equilibrium
eutectic solidus and liquidus values of the 3% Ni alloy, (e.g.1.91% and 4.19% C respectively).
Thus, in effect, the imbalance in the resulting compositions was due entirely to the difference in
their respective nickel contents.
The logic underlying the idea to create a compositional imbalance between the infiltrant and the
base compact that results in partial solidification during infiltration is that if the liquid solidifies
upon contact with the solid phase of the compact, it can not at the same time penetrate and
dissolve its sinter bonds. Given the requirement that the composition of the infiltrant must be that
of the eutectic liquidus, the only way to create the indicated imbalance is to select a base
compact composition that is on the lean side of the eutectic solidus. In ternary and higher order
alloys, there are several ways to accomplish this. For example, in the general case, the imbalance
can be based on any one of the minor alloy additions or on any combination of two or more of
them. Thus, in the present case, either the carbon or the nickel or both could have been selected
to create the required imbalance. However, as outlined below, consideration of the potential
kinetic consequences of the selection led to the decision to base it solely on the nickel content.
Once the infiltrant enters the base compact, the resulting two phase system will immediately act
to equalize any compositional differences that exist, first by partial solidification of the liquid as
already indicated and thereafter by diffusional processes which occur in both phases. Thus,
depending on the kinetics of these processes, the system may equalize before infiltration is
complete. If this happens, the sinter bonds of the base compact will lose the protection afforded
by the compositional imbalance and the compact will again become susceptible to disruption by
the liquid and consequent lateral spreading of the affected elements. Hence, in choosing the alloy
addition on which to base the imbalance, it was essential to select the one that would have the
greatest effect in retarding these processes. As between carbon which is an interstitial solute in
iron and diffuses rapidly and nickel which is a substitutional solute and diffuses slowly, the
obvious choice was the nickel.
Then given this and setting the nickel content of the Base Compact composition at 2% to ensure
graphitization of the hypereutectoid carbon, it remained at this point to select the nickel content
of the Infiltrant composition. In deciding this, particular consideration was given to the potential
consequences of the partial solidification effect on the kinetics of the infiltration process. The
early studies of the Fe-C system had shown that solidification during infiltration generally
slowed the process and could stop infiltration altogether if the extent to which it occurred was
too great. Thus, it was essential to limit the solidification potential of the Infiltrant composition
accordingly.
Unfortunately, it was impossible to predict the extent of solidification needed to produce a
favorable result. The earlier findings of the Fe-C system were helpful in this regard but were
certainly not decisive. For example, they indicated that the maximum allowable partial
solidification was about 35%. However, since the corresponding compositional differences
underlying this value were based on carbon which, of course, diffuses faster than nickel, it could
not be expected to give a reliable indication of what to expect of differences based on nickel.
In the general case, the extent to which the infiltrant will solidify during infiltration is determined
in accordance with the equilibrium phase relations by the magnitude of the compositional
differences between the liquid and the solid phases, the average composition after infiltration and
the infiltration temperature. Its value for any fixed set of compositional values is a maximum at
the eutectic temperature which, of course, is the lowest possible infiltration temperature and
otherwise decreases with increasing temperature up to some maximum temperature which in
addition to the compositional values mentioned is a function of the phase relations.
In the present case, the data in Table 1 below show the extent of solidification versus the nickel
content of the infiltrant for various nickel contents up to 4%. The values in each instance are
based on the equilibrium phase relations and otherwise assume that the infiltrant weight is 10.5%
of the final infiltrated weight. Shown also in the table are the eutectic liquidus and solidus carbon
contents and the value of the maximum infiltration temperature in each case. The latter data
define the corresponding carbon contents of the Infiltrant and Base Compact compositions and
the temperature values at which the potential for solidification of the infiltrant diminishes to
zero.
Table 1 - Maximum Diffusional Solidification Versus the Nickel Content of the Infiltrant
Max.
Eutectic Carbon Contents
Max.
Silicon Content
Infiltration
Diffusional
of the Infiltrant
Liquidus
Solidus
Temperature
Solidification
o
%
%
%
%
C
(oF)
2.5
8.4
4.21
1.93
1157 (2115)
3.0
15.9
4.19
1.91
1160 (2120)
3.5
24.5
4.16
1.89
1164 (2127)
4.0
32.7
1.14
1.87
1167 (2133)
Based on the maximum permissible partial solidification indicated by the Fe-C system, these
findings suggested that the nickel content of the Infiltrant could be set as high as 4%. However,
as already indicated, it was decided instead to set it at 3%. Other than the obvious economic
implications, there were two technical considerations underlying this choice as follows. First, the
solid state diffusivity of nickel is about four orders of magnitude smaller than that of carbon,
(14). Thus, it was very possible that it was neither permissible to set the solidification potential
of the infiltrant as high as in the earlier case nor for that matter actually necessary to do so in
order to preclude the equalization of the system before infiltration was complete. Second, in view
of the earlier experiences with infiltrants based on admixed silicon, it seemed highly likely that if
the need arose, additional nickel, up to perhaps as much as 1%, could be admixed without
serious detriment to the homogeneity of the infiltrant.
EXPERIMENTAL PROCEDURE
Laboratory scale quantities of the indicated 2% and 3% nickel containing powders were water
atomized in accordance with standard procedures. Their respective chemical and screen analyses
after hydrogen annealing at 930 oC were as listed below in Table 2.
Table 2 - Chemical and Screen Analyses of the Ni Pre-alloys
Chemical Analysis
Screen Analysis
Std US Screen
2% Nickel
2% Nickel
3% Nickel
Opening in
Powder
Powder
Powder
ums
Element
0.011
0.006
0
%C
+ 250
0.011
0.009
4.0
%S
- 250 / +180
0.103
0.104
5.4
%O
- 180 / +140
<.0002
0.0004
15.0
%N
- 140 / +106
1.94
2.99
24.0
%Ni
- 106 / +75
0.11
0.08
13.9
%Mn
- 75 / +60
0.05
0.03
20.7
%Cu
- 60 / +45
0.08
0.08
17.0
%Cr
-45
3% Nickel
Powder
0
0
0
0
0
0
1.3
98.7
The Infiltrant and Base Compact compositions based on these powders were made to the 3% Ni
eutectic liquidus and solidus carbon contents as indicated in the earlier Table 1. The carbon was
admixed in the form of Asbury Grade 3203 graphite having a minimum carbon content of 95%
and an average particle size of less than 10 micrometers. Sufficient additional graphite was
added in each case to allow for the inevitable carbon losses to the oxygen of the base powders.
The Infiltrant powder was made without lubricant. The Base Compact powder was lubricated
with 0.7% Acrawax C, the standard PM grade of the Lonza Division of IMS Company. The
infiltrant and base compact mix sizes were typically 200 and 1000 grams respectively. Both were
prepared as binder treated admixtures. The binder treatment processing was generally in accord
with the standard Hoeganaes ANCORBOND process, (15).
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 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 paper detailing the initial studies, (7).
The specimens were processed in an high temperature Hayes pusher type furnace. To prevent
decarburization, especially of the infiltrant composition in advance of complete graphite solution
by melting at the eutectic temperature, the specimens were processed in graphite gettered boxes
in a commercially pure synthetic dissociated ammonia atmosphere. Process temperatures varied
according to the aims of the particular trial but were typically in the range of 1160 to 1195 oC,
(2120 to 2185 oF). Process times varied likewise but again were typically of the order of ½ hour
at temperature.
RESULTS AND DISCUSSION
As it turned out, the necessity to use a higher nickel content in the infiltrant did not arise.
Contrary to expectation, the 3% Ni infiltrant would not infiltrate the compacts completely unless
the process temperature was well in excess of the maximum temperature for which the phase
relations indicated the possibility of partial solidification. For example, according to the data in
the table, the latter temperature is 1160 oC, (2120 oF), whereas the minimum temperature
required to effect complete infiltration was found to be substantially higher at about 1180 oC,
(2155 oF).
Subsequent to this finding, a fairly extensive series of trials was conducted in an effort to lower
the infiltration temperature sufficiently to test the efficacy of the diffusional solidification
concept. These included studies at lower base compact densities down to 6.0 g/cm3 and the use
of small additions of various surface active additives to the infiltrant in an effort to improve its
wetting characteristics. As it turned out, both remedies helped but disappointingly did not
succeed in lowering the applicable infiltration temperature to much below 1165 oC, (2130 oF).
Interestingly, similar difficulties were not encountered in the earlier studies of the Fe-C and
Fe-C-Si systems. In both cases, infiltration occurred without difficulty at temperatures as low as
1157 oC, (2115 oF).
In any event, of the numerous studies that were done in this connection, the following one is of
interest both to document infiltration in the Fe-C-Ni system as well as to further elucidate the
distortion effect and more particularly, its apparent causes. To facilitate infiltration, the Infiltrant
in this case was modified with 0.15% Si in the form of admixed SiC. Otherwise, the Infiltrant
and Base Compact compositions were the same as earlier indicated. The base compact density
and weight were nominally 6.7 g/cm3 and 35 grams and the infiltrant weight was 4 grams.
Infiltration was in synthetic DA at 1168 oC, (2135 oF) for 10 minutes followed by a liquid phase
sintering step at 1182 oC, (2160 oF) for 20 minutes. The results are shown below in Table 7.
Table 7 - Infiltrated Properties of an Fe-C-Ni Alloy at an Average Nickel Content of 2.05%
Distortion Value
Density Dim. Chg. vs Grn.
Specimen
Top to Bottom
Middle to Bottom
Number
3
g/cm
%
mm
(ins)
mm
(ins)
7.60
-0.46
0.16
(0.006)
0.04
(0.0015)
1
7.61
-0.61
0.20
(0.008)
0.06
(0.0024)
2
Average
7.61
-0.53
0.18
(0.007)
0.05
(0.002)
According to these data, the average infiltrated density was 7.61 g/cm3 and its easily shown on
the basis of the negative dimensional change values that sintering contributed at least as much as
0.10 g/cm3 to this result. Assuming complete graphitization of the hypereutectoid carbon, the
expected maximum density of this alloy would be about 7.54 g/cm3. Thus, in view of the higher
observed value, it was evident that graphitization in this case was probably incomplete.
The top to bottom distortion values shown in the table are typical of the values that were
generally seen in connection with the distortion effect both in this alloy system and in earlier
admixed Fe-C-Si alloys. In this particular case, metallographic examinations, as set out below,
showed that the values were attributable in part to lateral spreading by sinter bond separation
during infiltration, (i.e. to the liquid penetration mechanism), and in part to the aforementioned
incomplete graphitization. The middle to bottom distortion values shown in the table give an idea
of how much of the distortion is attributable to each cause. As shown below, they are primarily
indicative of the incomplete graphitization effect and the fact that it was localized to the bottom
of the compacts. In contrast, if graphitization had been complete, these values would have
approximated to zero since the lateral spreading due to the liquid penetration mechanism is
always limited to the top portions of the compact. Thus, in effect, they indicated that about 25%
of the overall distortion in this case was due to incomplete graphitization and hence, that the
balance, or about 75%, was due to the liquid penetration mechanism.
Perhaps, the most graphic indication of the causes of the distortion effect is provided by simple
visual inspection of the mounted cross section of an infiltrated specimen in the etched condition.
A macrograph of such a section is shown in Figure 1a.
A review of this figure will show that the etch has delineated three distinctly different areas of
the section. Starting at the top of the figure which corresponds to the infiltrated surface, they may
be roughly described in order of appearance as: 1) a relatively thin upper band of dark etching
material that traverses the section just under the infiltrated surface; 2) a broad middle band of
light etching material that likewise traverses the section over most of its depth below the
infiltrated surface except near the lower left edge of the specimen; and, 3) a lower band of
moderately dark etching material that joins the middle band at the aforementioned lower left
edge and otherwise traverses the section at the bottom of the specimen. As suggested by the
different etching characteristics, metallographic examination of the specimen confirmed that
each of the three areas exhibited a different microstructure. Of the three, the microstructure just
under the infiltrated surface was the most distinctive. A micrograph typifying this structure is
presented in Figure 1b.
(a)
(b)
Figure 1 - a) Macrograph of the mounted cross section of an infiltrated specimen in the polished
and nital/picral etched condition at 10x; b) Micrograph of an area just under the infiltrated
surface of the same specimen at 100x.
A review of Figure 1b will show that it is characterized by the presence of structural motifs that
exhibit a relatively complex hyper-eutectoid carbide/graphite morphology and otherwise
diminish in content with depth below the infiltrated surface, (i.e. the upper surface in the figure).
In fact, in this particular specimen, they disappeared completely from the microstructure at
depths that were within two fields of the surface at this magnification, (i.e. at depths of less than
~0.20 mm). More generally, both these motifs and their disappearance just a short distance
below the infiltrated surface are typical features of the distortion effect. The motifs per se are
thought to be the remnants of the areas of most severe infiltrant penetration of the sinter bonds
and grain boundaries of the base compact and thus to provide direct evidence of the liquid
penetration mechanism.
Exclusive of the few remaining instances of these motifs, the structure in the lower half of this
micrograph is typical of the structure in the middle band of light etching material of Figure 1a
and otherwise very like the ductile iron microstructure that was seen in the earlier studies of the
Fe-C-Si system. Similarly, the structure corresponding to the lower band of moderately dark
etching material in Figure 1a is typical of a white cast iron, (i.e. essentially un-graphitized) and
very like the microstructure seen in the earlier studies of the Fe-C system.
SUMMARY AND CONCLUSIONS
Recent efforts to develop iron base infiltration as a viable parts making capability revealed a
dimensional change anomaly of a sufficiently pernicious nature as to negate the net shape
advantage of the process. The purpose of the report was to describe this so-called distortion
effect, the apparent strategies available to mitigate it, and the results to date of the research
conducted in regard to it. It was indicated that the effect has two causes, a primary one and an
intermittently occurring secondary one. The primary cause was indicated to be the result of
liquid penetration and separation of the sinter bonds in and just below the surface of the base
compact followed by lateral expansion of the affected material under the influence of the surface
tension forces that act on the uninfiltrated liquid. The secondary cause was indicated to be the
result of incomplete graphitization of the hyper-eutectoid carbon and was subsequently found to
be amendable to alloying and processing. Analysis of the primary cause suggested two possible
strategies to mitigate it. One was to prevent liquid penetration of the sinter bonds by increasing
the dihedral angle between the infiltrant and the solid phase of the base compact and the other
was to forestall sinter bond separation by use of compositional differences designed to effect
diffusional solidification during infiltration. Of the two, the second appeared to be the simpler
and the subsequent efforts to implement it using alloys in the Fe-C-Ni system were described.
Regrettably, due to the unexpected poor infiltratability of the Ni containing infiltrant, these
efforts were unsuccessful. The balance of the report was devoted to documenting the distortion
effect in the subject alloy system and to presenting evidence to further elucidate its causes.
ACKNOWLEDGMENTS
Special thanks are due 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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J. R. Davis, Ed. “Cast Irons”, ASM International, Materials Park, Ohio, 1996, pp 3-15.
Ibid, pp 67-68.
R. M. German, “Liquid Phase Sintering”, Plenum Press, New York, 1985, pp 157-162.
Ibid, pp 15-21.
R. M. German, “Sintering Theory and Practice”, John Wiley & Sons, Inc., New York,
1996, pp 111 & 155.
“Thermo-Calc” Graphical User Interface Version: 2.1.2.0, Foundation for Computational
Thermodynamics, Stockholm, Sweden, 1995.
J. F. Elliott, M. Gleiser, and V. Ramakrishna, “Thermochemistry For Steelmaking”, Vol II,
Thermodynamic and Transport Properties, Addison - Wesley Inc, Reading, MA, 1963,
pp 687-697.
F. J. Semel and S. H. Luk, “Iron-Based Powder Mixtures Containing Binder-Lubricant”,
U. S. Patent 5,298,055.