EFFECT OF COPPER AND NICKEL ALLOYING ADDITIONS ON THE TENSILE AND FATIGUE
BEHAVIOR OF SINTERED STEELS
N. Chawla, D. Babic, J.J. Williams, and S.J. Polasik
Department of Chemical and Materials Engineering
Arizona State University
P.O. Box 876006
Tempe, AZ 85287-6006
M. Marucci and K.S. Narasimhan
Hoeganaes Corp.
1001 Taylors Lane
Cinnaminson, NJ 08077
ABSTRACT
The influence of Cu and Ni on tensile and fatigue behavior of Fe-0.85Mo prealloy with 0.6% graphite
addition was investigated. The transient liquid phase formed by the presence of Cu resulted in more
rounded pores in the Fe-Ni-Cu alloy compared to Fe-Ni. While the total porosity was similar in both
alloys, a larger fraction of secondary pores was present in the Fe-Ni-Cu. The addition of Cu resulted in an
increase in proportional limit stress, ultimate tensile strength, and fatigue strength over the Cu-free alloy.
The higher fatigue resistance was attributed to the enhanced solid solution strengthening provided by Cu
in Fe, and the higher proportional limit in the Fe-Ni-Cu alloy, caused by the more rounded nature of the
pores, and stronger Fe-Ni-Cu matrix surrounding the pores. The proportional limit stress appears to be a
good indicator of fatigue strength, since it quantifies the onset of localized plasticity in these materials.
INTRODUCTION
Sintered ferrous alloys are frequently alloyed with Cu or Ni for enhanced sinterability and mechanical
properties [1-5]. The two alloying elements play very different roles during the sintering process,
particularly because Cu has a much lower melting point than Ni. Thus, at the typical temperatures used to
sinter ferrous alloys, Cu forms a liquid phase while Ni remains in the solid state. The implication of Cu
being in liquid phase is that when the Cu particles melt, they leave their original positions, creating very
small pores in their wake [6]. This type of porosity is characterized as secondary porosity, because the
pores are much smaller than the larger primary pores associated with the sintering of ferrous alloys.
Another effect attributed to Cu liquid phase is that it forms a semi-permeable liquid layer around the Fe
particles, which improves iron particle bonding but also reduces carbon diffusion into iron [5]. Since Ni
remains in the solid state during sintering, partial diffusion of Ni into Fe contributes to a heterogeneous
microstructure, consisting of nickel-rich ferrite as well as nickel-rich areas predominantly at the periphery
of the pores [4].
In this study we have attempted to separate the effects of Cu and Ni on the tensile and fatigue properties
of Fe-0.85Mo prealloy steels. Two alloys were studied, both containing Ni, but only one containing Cu.
Microstructure characterization, including pore size and pore shape characterization, was performed to
better understand the tensile and fatigue behavior of the alloys. It will be shown that the combination of
Cu and Ni provides better tensile and fatigue properties than Ni alone. The effect of Cu on tensile and
fatigue behavior will be discussed in terms of: (a) Cu as a strengthening agent in the Fe particles and (b)
Cu liquid phase effect on pore morphology, size, and distribution.
EXPERIMENTAL PROCEDURE
Two Fe-0.85Mo alloys were investigated in this study, both containing approximately the same amount of
Ni and graphite, Table 1. The Fe-Ni-Cu alloy contained 1.5 % of copper, while the Fe-Ni alloy did not
contain any copper.
Table 1. Composition of powder metallurgy alloys examined in this study.
Alloy
Fe-Ni-Cu
Fe-Ni
Mo
0.85
0.85
Composition (wt. %)
Cu
Ni
C
1.50
1.75
0.60
0.00
2.00
0.60
Fe
Balance
Balance
The powders were blended and binder treated using a proprietary process developed by Hoeganaes Corp.
[7,8]. All powders were pressed into rectangular blanks, to a green density of 7.0 g/cm3, and sintered at
1120ºC for 30 minutes in a 90% N2–10% H2 atmosphere. In order to characterize the pore structure, both
the pore size and pore shape distributions were quantified. Analysis of primary pores (taken as pores with
areas above 25 µm2) was taken from digital image analysis of optical micrographs of the samples’ crosssection. The smaller, secondary pores (taken as pores with areas smaller than 25 µm2) were analyzed by
scanning electron microscopy (SEM). Pore size was estimated by measuring the pore area fraction, while
pore shape was characterized by a pore shape factor, F [5]:
F=
4πA
P2
where A and P are the measured pore area and measured pore perimeter, respectively. In the analysis of
the pore shape factor a value of 1 denotes a perfectly circular pore, and values that approach zero denote
increasingly irregular pores.
Tensile and fatigue testing were conducted on cylindrical axial specimens machined from the sintered
rectangular blanks. The specimens had to an overall length of 75 mm, gage length of 20 mm, and gage
diameter of 4.75 mm. Specimens were hand polished, using diamond paste, to a 1 µm finish. Tensile tests
were conducted at a constant strain rate of 10-3/s, while axial fatigue testing was conducted at an R ratio
(σmin/σmax) of –1, and a sinusoidal frequency of 40 Hz. All mechanical testing was conducted on a
precision aligned servohydraulic load frame. SEM and energy dispersive spectroscopy (EDS) were
performed for fracture surface analysis and compositional analysis, respectively.
RESULTS AND DISCUSSION
The sintered density of both alloys was about 7.0 g/cm3 and the dimensional change after sintering was
also similar, Table 2. Microstructural analysis revealed that porosity in both alloys to be about 12%,
which was consistent with the density of the alloy relative to the pore-free density. While the total
porosity was similar, optical micrographs of both alloys, Fig. 1, revealed a larger fraction of secondary
pores in the Cu-containing alloy. This was to be expected, as mentioned in the introduction, since the Cu
formed a transient liquid phase during sintering and left the secondary pores in its wake.
Table 2. Sintered properties of Fe-0.85Mo alloys.
Alloy
Fe-Ni-Cu
Fe-Ni
(a)
Green
Density
(g/cm3)
7.05
6.97
Sintered
Density
(g/cm3)
7.01
6.98
Dimensional
Change
(%)
-0.06
0.02
Pore area
fraction
(%)
12.0
12.1
(b)
Figure 1. Optical micrographs of as-sintered microstructure: (a) Fe-Ni-Cu and (b) Fe-Ni alloy.
Note the higher degree of secondary porosity in the Cu-containing alloy.
The absence of a large amount of dimensional change in the Cu-containing alloy can be explained by the
competing mechanisms of (a) transient liquid phase formation and diffusion of this phase into Fe particles
[6] and (b) formation of secondary pores formed by the liquid phase [9]. The absence of liquid phase in
the Fe-Ni alloy resulted in an almost negligible dimensional change. Scanning electron microscopy of the
polished surfaces further indicated that Ni did not diffuse completely into iron and was present
predominantly at the periphery of the primary pores in both alloys, Fig. 2.
Figure 2. Backscattered scanning electron micrograph showing nickel-rich areas from incomplete
diffusion of Ni (bright regions in the micrograph), predominantly around primary pores.
The pore size and distribution analysis quantified primary pores (> 25 µm2) and secondary pores (< 25
µm2) separately. A comparison of the primary pore size and pore shape distribution for both alloys is
shown in Fig. 3. The Fe-Ni-Cu alloy had a lower frequency of pores between 25-75 µm, but a larger
frequency of pores above 75 µm. Thus, from the pore size distribution, the primary pores in the Fe-Ni-Cu
alloy appeared to be larger. Pore shape analysis, Fig. 3(b), indicated that these pores were also slightly
rounder than those in the Fe-Ni alloy. Analysis of secondary pores, which required microstructure
characterization using SEM, is shown in Fig. 4. The pore size was normalized, due to the proprietary
nature of the Cu particle size used in these alloys. The analysis indicated that a larger fraction of
secondary pores was present in the Fe-Ni-Cu alloy, and these pores were also more rounded than those in
the Fe-Ni alloy. These pores may be attributed to the voids left behind by the original Cu particles [9].
While the pore shape analysis was useful, it did not reflect the dependence of pore shape on pore size.
Analysis of pore shape for the distribution of pore sizes, Fig. 5, indicated that for a given pore size the
pores in the Fe-Ni-Cu alloy were more rounded than that in the Fe-Ni alloy.
25
70
Fe-Ni-Cu
Fe-Ni
60
Fe-Ni-Cu
Fe-Ni
50
Frequency (%)
Frequency (%)
20
40
30
15
10
20
5
10
0
0
75 150 225 300 375 450 525 600 675 750 825 900 975
2
Pore Size (µm )
0
0.2
0.4
0.6
0.8
2
Pore Shape Factor (F=4πA/P )
2
1
Figure 3. Pore size and pore shape distribution of primary pores (larger than 25µm ), revealing
larger and slightly rounder pores in Cu-containing alloy.
80
40
Fe-Ni-Cu
Fe-Ni
70
30
Frequency (%)
Frequency (%)
60
50
40
30
25
20
15
20
10
10
5
0
Fe-Ni-Cu
Fe-Ni
35
0.01
0.1
0.3
0.4
0.5
0.6
Normalized pore size
0.8
0.9
1
0
0
0.2
0.4
0.6
2
Pore Shape Factor (F=4πA/P )
2
0.8
Figure 4. Pore size and pore shape distribution of secondary pores (smaller than 25µm ), showing
higher density of secondary pores and rounder pores in Cu-containing alloy.
1
1
2
Pore Shape Factor (F=4 πA/P )
Fe-Ni-Cu
Fe-Ni
0.8
0.6
0.4
0.2
0
0
0.2
0.4
0.6
0.8
Normalized Pore Size
1
Figure 5. Pore shape versus pore size relationship in both alloys, indicating that pores in the Cu-containing
alloy were more rounded.
The tensile behavior of the alloys was profoundly influenced by the presence of Cu, Table 3. The Cucontaining alloy exhibited higher proportional limit stress (the stress defined as the onset of plasticity in
the material, taken here as 1% deviation in linearity of the stress-strain curve) and higher ultimate tensile
strength. The Young’s moduli of the alloys were similar, since both alloys contained the same relative
porosity. The higher tensile strength in Fe-Ni-Cu can be attributed to the presence of Cu in solid solution,
which is effective in increasing the resistance to dislocation motion and increasing the work-hardening
rate of the alloy. The higher proportional limit was also affected by the presence of Cu in solid solution,
but may have been more influenced by the more rounded nature of pores in the Fe-Ni-Cu. The stress
concentration at a spherical pore is clearly less severe than that of a very irregular pore. Fractographic
analysis of tensile fracture surface of both alloys was characteristic of most sintered ferrous powder
metallurgy alloys, Fig. 7, showing evidence of localized ductile rupture of sintered bonds.
Alloy
Fe-Ni-Cu
Fe-Ni
Table 3. Tensile behavior of Fe-Ni-Cu and Fe-Cu alloys.
Young’s
Proportional Limit
Ultimate Tensile
Modulus
Stress
Strength
(GPa)
(MPa)
(MPa)
135
221
797
131
190
570
Strain-to- Failure
(%)
1.91
2.12
800
Fe-Ni-Cu
700
Stress (MPa)
600
Fe-Ni
500
400
300
200
100
0
0
0.5
1
1.5
Strain (%)
2
2.5
Figure 6. Tensile behavior of Fe-Ni-Cu and Fe-Ni alloys. Note the higher strength and work
hardening rate in the Fe-Ni-Cu alloy.
Figure 7. Tensile fracture surface, characteristic of both alloys, shown localized ruptured at sintered bonds.
The fatigue behavior for both alloys was also quite different. The stress versus cycles behavior for both
alloys, Fig. 8, showed that the addition of Cu significantly increased fatigue resistance. In order to
compare the fatigue strength (taken as fatigue runout at 107 cycles) to tensile behavior, we examined the
ratio of fatigue strength to tensile strength, σfat/σuts, and the ratio of fatigue strength to proportional limit
stress, σfat/σpl, Table 4. Due to the porosity in these materials, the fatigue ratio of σfat/σuts was significantly
lower than for fully dense Cu-containing steels, where the ratio is between 0.5-0.6 [10]. While the σfat/σuts
ratio was around 0.25 for both alloys, the σfat/σpl ratio was around 0.80 for both alloys. This can be
explained by the fact the proportional limit stress is a very good indicator of the onset of plasticity, which
in sintered materials takes place at the tips of the pores. In fatigue loading, this localized damage is more
likely to be controlling fatigue life, than large-scale plasticity that takes place close to the ultimate tensile
strength. Thus, the proportional limit stress may be a much better indicator of the fatigue strength than
ultimate tensile strength. It should be noted that the method by which proportional limit is measured is
limited by the accuracy of the technique and resolution of the stress-strain data acquisition, so more
effective measurement techniques of the proportional limit stress may yield even better correlation to
fatigue strength.
400
Fe-Ni-Cu
Fe-Ni
Stress Amplitude (MPa)
350
300
250
200
150
100 3
10
4
10
5
6
10
10
Cycles to Failure, N
7
10
8
10
f
Figure 8. Stress versus cycles behavior of Fe-Ni-Cu and Fe-Ni alloys. The addition of Cu
enhances the fatigue resistance of the alloy.
Table 4. Fatigue endurance limit and its ratio with ultimate tensile strength and proportional limit stress.
Alloy
Fatigue Strength
(MPa)
Fe-Ni-Cu
170
Proportional
Limit Stress
(MPa)
221
Fe-Ni
160
190
Ultimate Tensile
Strength
(MPa)
797
570
σfat/σpl
σfat/σuts
0.77
0.21
0.84
0.28
Fractography indicated that fatigue crack initiation took place at large pores, particularly at the surface of
the specimen, Fig. 9. Porosity has long been attributed as the preferred site for fatigue crack initiation at
or near the surface in sintered materials [1-4, 11-16]. Polasik et al. [4] have also shown that in the high
cycle fatigue regime, most of the fatigue life is spent in crack nucleation rather than crack propagation.
Thus, while the Cu-containing alloy had a larger density of small secondary pores, which would be
conducive to easier crack propagation since the mean free path between pores decreases, this did not
significantly deter the fatigue strength. Rather, since the primary pore distribution in both alloys was not
significantly different, the higher strength and fatigue resistance of the Fe-Ni-Cu matrix around the pores
was more significant in contributing to the higher fatigue strength.
Figure 9. Surface pore associated with fatigue crack nucleation site.
The fatigue fracture surface also exhibited localized fatigue striations. Unlike striations in wrought
materials, however, the striations in porous sintered materials seemed to be highly localized. This may be
due to the multiple sites of favorable orientation and size, with respect to the loading axis, that underwent
blunting and sharpening and give the step-like appearance of the striations. Similar localized fatigue
striations were also observed by Rodzinak and Slesar [17].
Figure 10. Localized fatigue striations observed in both alloys.
CONCLUSIONS
The mechanical behavior of binder-treated Fe-0.85Mo prealloy with Cu, Ni, and graphite additions was
investigated. In particular, the effect of Cu on tensile and fatigue behavior was studied, yielding the
following conclusions:
1. The transient liquid phase formed by the presence of Cu resulted in more rounded pores in the Fe-NiCu alloy compared to Fe-Ni. While the total porosity was similar in both alloys, a larger fraction of
secondary pores was present in the Fe-Ni-Cu.
2. The addition of Cu resulted in an increase in proportional limit stress and ultimate tensile strength
over the Cu-free alloy. Copper diffusion into iron seems to contribute to enhanced resistance to
dislocation motion resulting in an increase in work hardening rate over the Fe-Ni alloy. Both alloys
exhibited similar fracture mechanisms, consisting primarily of ductile rupture at sintered particle
bonds.
3. Fe-Ni-Cu alloy had a higher fatigue resistance than the Fe-Ni alloy. The higher fatigue resistance can
be attributed to the enhanced solid solution strengthening and more rounded pores provided by Cu in
Fe, particularly since the amount of primary porosity in both materials was similar. The proportional
limit stress appears to be a good indicator of fatigue strength, since it quantifies the onset of localized
plasticity in these materials.
ACKNOWLEDGMENTS
The authors would like to thank Hoeganaes Corp. for financial support of this work.
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