FATIGUE CRACK GROWTH OF Fe-0.85Mo-2Ni-0.6C STEELS
WITH A HETEROGENEOUS MICROSTRUCTURE
G. Piotrowski, X. Deng, and N. Chawla
Department of Chemical and Materials Engineering
Arizona State University
Ira A. Fulton School of Engineering
Tempe, AZ 85287-6006
K.S. Narasimhan and M. Marucci
Hoeganaes Corporation
Cinnaminson, NJ 08077
ABSTRACT
Powder metallurgy processing of steel alloys typically results in a material with heterogeneous
microstructure and residual porosity. The fatigue crack growth behavior of these materials is strongly
affected by the nature of porosity and heterogeneous microstructure. Notched fatigue specimens were
prepared from a Fe-0.85Mo prealloy mixed and binder-treated with 2%Ni and 0.6%C. The alloys were
tested at three different densities: 6.98 g/cm3, 7.36 g/cm3, and 7.53 g/cm3. The microstructure at each
density was characterized to determine the porosity, microconstituents, and phase fractions. Fatigue
testing was performed at various R-ratios, ranging from –2 to 0.8. Increasing porosity and increasing Rratio resulted in a decrease in ∆Kth. In situ observation of crack growth showed that the cracks propagated
through Ni-rich regions. It appears that pearlite regions, and, to some extent bainite regions, however,
contributed to toughening and crack deflection.
INTRODUCTION
Steel products made by press and sinter powder metallurgy (P/M) processing are replacing products made
by more conventional procedures, such as casting and forging, due to the cost savings associated with
near-net shape processing [1-4]. Unfortunately, P/M parts usually have some degree of residual porosity
after sintering, which can adversely affect the mechanical properties of these materials [4-8]. The nature
of porosity is determined by several processing variables, including the type and amount of alloying
additions, powder size distribution, green density, sintering temperature, and sintering time [1,5].
In general, an increase in porosity is associated with more irregular pores, and a greater fraction of
interconnected pores. Increased pore interconnectivity and pore clustering leads to increased strain
localization and damage, thus reducing the strength and ductility of the steel [5,9]. Porosity has also been
shown to significantly influence the fatigue response of sintered P/M steels [4-8,10-13]. Danniger et al.
[5] found that pores and pore clusters act as sites of crack initiation, while Polasik et al. [8] have shown
that small cracks that nucleate at pores coalesce to form large fatigue cracks leading to fracture. A
comprehensive understanding of the effects of porosity and microstructure on the fatigue crack growth
behavior of sintered steels is required. In this study, we have examined the effect of porosity and
microstructure on fatigue crack growth behavior of a Fe-0.85Mo-2Ni-0.6C at three different sintered
densities: 6.98 g/cm3, 7.36 g/cm3, and 7.53 g/cm3. Quantitative characterization of the heterogeneous P/M
microstructure was performed to determine the degree of porosity and nature of microstructure at all
densities. With increasing density the degree of pore interconnectivity decreased, although density did not
significantly affect the fraction of the microconstituents in the microstructure. Fatigue crack growth
experiments were performed at constant R-ratio, ranging from -2 to 0.8. It will be shown that increasing
porosity and R-ratio decrease the fatigue crack growth resistance of the material, and that the
heterogeneous nature of the microstructure also influences the fatigue crack growth behavior.
MATERIALS AND EXPERIMENTAL PROCEDURE
An Fe-0.85Mo prealloy powder was blended and binder treated with 2 wt.% Ni and 0.6wt.% graphite
[14,15]. All powders were pressed into rectangular blanks and sintered at 1120ºC for 30 minutes in a 90%
N2–10% H2 atmosphere. The samples were pressed and sintered to obtain three different sintered
densities: 6.98 g/cm3, 7.36 g/cm3, and 7.53 g/cm3. The samples with 7.53 g/cm3 density were made by a
double-press/double-sintering process. Porosity was determined by image analysis of several
representative micrographs using both optical and scanning electron microscopy (SEM). The micrographs
were segmented into black and white images, and the porosity determined by an automated procedure.
Quantitative characterization of the phase fractions of the microstructure was performed by both image
analysis of the optical and SEM images and Vickers hardness tests. The phase area fractions of each
density were determined using a segmentation procedure with various colors, whereby each phase was
assigned a specific color. The segmented images were then analyzed based on color to determine the
phase fractions. Hardness measurements were conducted using a Vickers hardness indenter with a 50 g
applied load. A minimum of 20 measurements were conducted in each phase.
Fatigue tests were performed on a servo-hydraulic load frame equipped with a Questar (Questar Corp.,
New Hope, PA) traveling microscope. The traveling microscope was used for in situ measurement and
observation of fatigue crack growth. This was particularly important because it allowed visualization of
the interactions between the crack and the microconstituents in the microstructure. All fatigue tests were
performed using a single edge notch axial fatigue configuration [16-18]. The samples were machined by
electrodischarge machining (EDM) to the following dimensions: height of 45 mm, width of 11.5 mm,
thickness of 7.48 mm. The edge notch was machined by EDM to a length of 4.5 mm. Fatigue tests were
performed at constant R-ratio, ranging from 0.8 to -2, in load control at a frequency of 30 Hz. A
decreasing ∆K procedure was used until ∆Kth was reached. At this point, the ∆K was raised to a value in
the Paris law regime, and the ∆K increased to obtain the higher ∆K behavior. Samples were pre-cracked
to a length of about 600 µm. The typical growth increment for a given ∆K was 200 µm, such that the
increment encompassed several phases in the heterogeneous microstructure and was sufficiently larger
than the plastic zone from the previous ∆K increment.
RESULTS AND DISCUSSION
Microstructure Characterization
The porosity at each density was determined from several optical micrographs using image analysis,
Table 1. At the lowest density, 6.98 g/cm3, the measured porosity was 9.5%, while at the highest density
7.53 g/cm3, the porosity was about 3.9%. The values obtained from the image analysis technique
Sintered Density
(g/cm3)
6.98
7.34
7.53
Table 1. Porosity Measurements versus Density
Porosity from Sintered Density
Porosity from Image Analysis
(%)
(%)
10.3
9.5 ± 0.8
4.5
4.5 ± 0.6
3.2
3.9 ± 0.2
correlated well with those obtained from calculation of the sintered density from the pore-free density The
microstructure at the three densities showed noticeable differences in porosity, Figure 1. At the highest
porosity, larger, more irregular, and interconnected pores were observed. With decreasing porosity, the
overall pore size was smaller, and the pore shape was more regular.
(a)
(b)
50 µm
(c)
50 µm
50 µm
Figure 1. Microstructure of Fe-0.85Mo-2Ni-0.6C steels at: (a) 6.98 g/cm3, (b) 7.34 g/cm3, and (c) 7.53 g/cm3. At
higher porosity the pores are larger, more irregular, and more interconnected.
Etched microstructures showed the heterogeneous nature of the microconstituents, in addition to the
porosity, Fig. 2. These included coarse pearlite, fine pearlite, Ni-rich regions (likely Ni-rich austenite,
surrounding the pores), and bainite, which appears at the periphery at the Ni-rich regions. Ni-rich regions
likely formed as a result of incomplete diffusion of the elemental Ni into the surrounding matrix upon
sintering.
Coarse
Pearlite
Quantitative characterization of the phases and phase
fractions in the microstructures, at all three densities,
Ni-rich
was also conducted, Fig. 3. The phase fractions were
Bainite
determined from both optical and SEM micrographs.
Each phase was segmented in the micrograph and
indexed by a specific color. A similar analysis has
been conducted by Komenda et al. [19] to
Fine
characterize a heterogeneous steel microstructure.
Pearlite
The analysis showed that the phase fractions in the
Pore
microstructure were relatively uniform at all
densities. As expected, the microstructure consisted
primarily of coarse pearlite, with smaller amounts of
25 µm
fine pearlite and Ni-rich phase, and minor amounts
Figure 2. Characteristic heterogeneous
of bainite. The bainite microconstituent was
Microstructure of p/m steel consisting of porosity,
observed primarily at the periphery of the Ni-rich
coarse and fine pearlite, bainite, and Ni-rich phase
regions. It is interesting to note that the results from
(likely austenite).
optical micrographs overestimated the amount of
coarse pearlite and Ni-rich regions, at the expense of the amount of fine pearlite. This is understandable,
since the fine pearlite is more difficult to resolve in the optical microscope. Thus, the true phase fraction
measurements were obtained from the higher magnification SEM images.
Vickers hardness of the several microconstituents
showed that the bainite regions were the hardest,
followed by fine pearlite, coarse pearlite, and Ni-rich
areas, Fig. 4. The relatively low hardness of the Nirich phase suggests that it may be retained austenite
[20]. The hardness of each phase, as expected, did not
vary significantly with density.
Phase Area Fraction (%)
100
SEM
Optical
80
Total Pearlite
60
Coarse Pearlite
40
Fine Pearlite
Fatigue Behavior
Ni-rich
20
Bainite
0
6.9
7
7.1
7.2
7.3
7.4
7.5
Density (g/cm3)
7.6
7.7
Vickers Hardness Number (VHN)
Figure 3. Phase and microconstituent area fractions
by optical and scanning electron microscopy
(SEM). Optical microscopy underestimates the
fraction of fine pearlite.
700
600
Bainite
500
400
300
Fine Pearlite
200
Coarse Pearlite
Ni-rich
100
Fatigue crack growth experiments at the three
densities showed that porosity had a strong effect on
the crack growth rate (da/dN) and the threshold stress
intensity factor, ∆Kth. Figure 5 shows the da/dN
versus ∆K behavior for all densities, at various Rratios. ∆Kth at the three densities varied between 14.716.4 MPa.m1/2 at R=-2, to 2.9-4.3 MPa.m1/2 at R =
0.8. These values are in the range of those found by
others in similar Fe-Mo sintered steels [4,21-23]. The
slope in the steady state or Paris law regime, m, was
also measured [24]. Increasing R-ratio resulted in an
increase in m from about 2 at low R-ratio to around
10 at R-ratio of 0.8. This increase in slope is
indicative of a much higher crack growth rate, for a
given ∆K, due to increasing Kmin. We now compare
the fatigue crack growth behavior, at all three
porosity levels for three R-ratios: -1, 0.1, and 0.8, Fig.
6. At all R-ratios, decreasing porosity resulted in a
higher ∆Kth. This is particularly clear at the highest Rratio of 0.8.
The fatigue crack growth data, at all R-ratios, was
analyzed using the two- parameter approach proposed
Density (g/cm3)
by Vasudevan and Sadananda [26,27]. In the twoFigure 4. Vickers hardness measurements of
parameter approach there exist two critical parameters
individual phases. As expected, density does not
that dictate when fatigue crack propagation will take
affect the hardness of the individual phases.
place, at a given crack growth rate, Fig. 7. The two
critical components are the cyclic component ∆K and the static component Kmax. The critical values for
crack propagation then, shown in Fig. 7, are the asymptotic values in the curve, ∆K* and K*max. By
measuring the values ∆K* and K*max over a range of da/dn values, “fatigue trajectory” plots can be
generated. These provide a means to elucidate changes in material behavior with increasing crack growth
rate. The two-parameter approach has also been used to categorize material behavior into different
classes, Fig. 8 [26]. Materials which are dominated by environmental effects exhibit the behavior
described by Class IV. Materials which are dominated by the static component, Kmax*, will show a
significant monotonic component to fatigue damage, similar to that shown by the Class II material. The
so-called “ideal material,” described as Class III behavior in Fig. 8, exhibits a perfect “L-shaped”
behavior.
6.9
7
7.1
7.2
7.3
7.4
7.5
7.6
6.98 g/cm3
10
10
10
(a)
-6
R = 0.1
R = 0.3 m = 5 R = -1
m=4
m=5
R = -2
m =2
R = 0.8
m = 10
-7
-8
-9
m
-10
10
10
da/dN (m/cycle)
da/dN (m/cycle)
10
7.34 g/cm3
10
10
10
(b)
-6
R = 0.8
m = 10
R = 0.3 R = 0.1 R = -1
m=6 m=4 m=3
-7
-8
-9
m
-10
10
-11
10
-11
2
4
6
8 10
10
30
2
4
∆K (MPa-m1/2)
6
8 10
30
∆K (MPa-m1/2)
da/dN (m/cycle)
7.53 g/cm3
10
10
10
10
(c)
-6
R = 0.1 R = -1
R = 0.8 R = 0.3 m = 5 m = 5
m = 10 m = 5
-7
-8
R = -2
m=3
-9
-10
Figure 5 da/dn versus ∆K for: (a) 6.98 g/cm3 (b) 7.34 g/cm3,
and (c) 7.53 g/cm3. Porosity has a strong influence on the
fatigue behavior of the P/M steel. Increased porosity and Rratio decrease the ∆KTH values of the steel. The slope of the
steady-state region of curves was also found to increase with
R-ratio, indicative of an increased monotonic contribution to
fatigue.
m
10
-11
10
2
4
6
8 10
30
1/2
∆K (MPa-m )
R = 0.1
-6
da/dN (m/cycle)
10
R = 0.8
R = -1
∆K
-7
Fatigue Crack
Growth Region
10
-8
10
-9
∆K*
10
6.98 g/cm
-10
10
3
3
7.36 g/cm
7.53 g/cm3
-11
10
2
4
6
8 10
30
∆K (MPa-m1/2)
Figure 6. Effect R-ratio on fatigue crack
growth, as a function of density. Increasing
porosity and R-ratio resulted in a decrease in
fatigue crack growth resistance.
K*max
Kmax
Figure 7. Schematic of two parameter
approach to fatigue (after Vasudevan and
Sadananda [27]). Two critical parameters are
required for crack growth, a cyclic
component, ∆K, and a static component,
Kmax.
∆K
Environment effect
We begin by examining the
∆K versus R-ratio behavior.
V
V
Strong Weak
The curves are relatively
I
linear, although it can been
I
seen that with increasing
density, a higher ∆K is
II
required
to achieve a given
II
crack growth rate. Also, at
III
III
least up to an R-ratio of 0.8,
IV
there does not seem to be a
IV
plateau in ∆K, as described
by the Class III behavior.
R-ratio
Kmax
Rather, is likely that large
Figure 8. Classification of different types of materialsby the two parameter
degree of plasticity is
approach. Class I materials exhibit limited sensitivity to environmental effects.
present at the crack tip at
Class III material is termed the “ideal material” behavior, while Class II
very high R-ratio, that
materials are influence by large plasticity at the crack tip.
makes the P/M materials fall
into the Class II category. The changes in material behavior with increasing R-ratio and with increasing
crack growth rate are more easily observed in a series of plots of ∆K versus Kmax, Fig. 9. A series of
curves are shown at each of the crack growth rates, shown in Fig. 10. As the degree of porosity increases,
the monotonic contribution to the fatigue crack growth behavior increases. This can be seen graphically as
the deviation from the horizontal asymptote in the ∆K-Kmax plots. The behavior can be rationalized by the
fact that with increasing R-ratio, the localization of strain and plasticity at the crack tip are enhanced in a
material with a higher degree of porosity. Thus, a lower threshold is required for crack growth to take
place with a larger degree of porosity at the crack tip. This behavior is exacerbated with increasing crack
growth rate, as shown by the increased deviation of the “L-shaped” curves with increasing crack growth
rate. This behavior is more predominant at the lower densities, because of the enhanced effect of porosity.
At the highest density, the material behavior is closer to the classical “ideal material.” To quantify the
degree of monotonic superposition, a comparison between the experimental data and a theoretical pure
fatigue line was made. The condition for “pure fatigue” is that ∆K* equal K*max, Fig. 11(a). Thus, any
devitation from this 45o line, must be due to environmental, crack closure, or static effects. As the line
moves toward the horizontal axis, this signifies an increased contribution from monotonic loading, since
Kmax is the dominating component to fatigue [25]. Our data show that at the two lower densities the
behavior is close to parallel to the horizontal axis, indicating significant static effects associated with
strain localization between pores, Fig. 11(b). At the highest density, however, the line is close to parallel
to the pure fatigue line. This indicates that the static contribution with increasing R-ratio is not as
predominant. Rather, the off-set between the two parallel lines may be attributed to two components: (a)
fatigue crack closure at low R-ratio and (b) static effect at higher R-ratio.
∆K
6.98 g/cm3
50
7.34 g/cm3
50
1e-7
30
1e-8
20
1e-9
10
K*max
∆K
∆K
-2
-1.5
-1
-0.5
R-ratio
0
0.5
1e-7
30
1e-8
20
1e-9
10
TH
*
0
-2.5
40
1
0
-2.5
K
∆K
40
30
20
max
10
∆K
-1.5
-1
-0.5
R-ratio
0
1e-8
1e-9
*
K
*
TH
-2
(c)
1e-7
∆ K [MPa-m1/2]
40
(b)
∆ K [MPa-m1/2]
∆ K [MPa-m
1/2
]
(a)
7.53 g/cm3
50
0.5
1
∆K
max
TH
*
0
-2.5
∆K
-2
-1.5
-1
-0.5
R-ratio
Figure 9. ∆K versus R for: (a) 6.98 g/cm3, (b) 7.34 g/cm3, and (c) 7.53 g/cm3.
0
0.5
1
*
6.98 g/cm3
10-7
m/cycle
R = -1 Pure Fatigue
R=0
R = 0.1
30
10-8
R = 0.3
10-9
20
10-10
10
R = -1 Pure Fatigue
R=0
(b)
40
∆K (MPa-m1/2)
∆K (MPa-m1/2)
R = -2
(a)
40
7.34 g/cm3
R = 0.1
10-7
m/cycle
30
R = 0.3
10-8
20
-9
10
-10
10
10
R = 0.8
0
R = 0.8
0
0
10
20
30
40
Kmax (MPa-m1/2)
0
10
20
30
Kmax (MPa-m1/2)
40
7.53 g/cm3
(c)
∆K (MPa-m1/2)
40
10-7
m/cycle
R = -2
R = -1 Pure Fatigue
R=0
10-8
R = 0.1
30
-9
10
20
R = 0.3
10-10
Figure 10. Trajectory maps of ∆K versus Kmax at various
da/dN: (a) 6.98 g/cm3, (b) 7.34 g/cm3, and (c) 7.53 g/cm3.
Note that for highest porosity, the horizontal asymptote is
diminished due to the significant contribution of plasticity at
the crack tip (from Kmax).
10
R = 0.8
0
0
10
20
30
Kmax (MPa-m1/2)
40
Experimental Results
(a)
Pure Fatigue
Line
Increasing
Crack Growth Rate
∆K*
Monotonic
Contribution
(b)
Pure Fatigue
15
11
7.55 g/cm3
∆K*
7
7.36 g/cm3
6.98 g/cm3
3
Kmax*
3
7
11
Kmax*
15
Figure 11. Quantifying the Degree of Monotonic Superposition: (a) monotonic contribution quantified by
measuring the deviation of the experimental data from the “pure fatigue” line, and (b) as porosity increases, the
degree of monotonic superposition also increases. This is indicated by the increasing deviation from the pure
fatigue line.
In this section we describe fractographic analysis after fatigue crack growth. Since the crack path
appeared quite tortuous, due to the hetrogeneous microstructure of these steels, we have also quantified
the degree of crack deflection on the measured ∆K, per the model of Suresh [28]. The crack appears to be
highly dependent on the phase(s) at the crack tip, Fig. 12. For the Ni-rich regions, cracks tend to
propagate in a linear fashion, suggesting that the Ni-rich regions offer little to no resistance to crack
propagation. This is further by the Vickers hardness data, which showed the Ni-rich phase to be very soft,
indicating that it might be Ni-rich austenite phases [20]. For the pearlite regions, cracks tend to be highly
deflected, with some evidence of the Fe3C particles in the ferrite matrix bridging the crack, Fig. 13. Pores
inside the Ni-rich regions can also act as nucleation sites for secondary cracks. It has been demonstrated
that these cracks originating at pores ahead of the crack tip, often join the main crack [8,13], Fig. 14.
Cracks propagating through the coarse pearlite, fine pearlite, and bainite all show large increases in the
degree of fatigue resistance due the crack deflection, Fig. 14. Crack arrest is often present, and further
induces crack deflection through branching. The role of the Ni-rich areas in fatigue is still not quite clear.
Quantitative measurements of the crack growth rate in the different microstructural constitutents are being
conducted, and will shed some light on the role of Ni-rich areas.
(a)
(b)
20 µm
20 µm
Figure 12. Fatigue Crack Behavior through heterogeneous microstructure: (a) Ni rich, and (b) Pearlite. The crack
propagates through the Ni-rich areas, but it tortuous and deflected through pearlite due to the Fe3C needles.
(a)
(b)
300
Counts
Fe3C
Fe
Fe3C
250
200
150
C
100
Fe
50
Fe
0
0
1 µm
2
4
6
8
Energy (keV)
Figure 13. Crack Bridging due to Fe3C: (a) Fe3C needles pulled out of the ferrite matrix during fatigue, and (b)
EDS analysis showing the composition corresponding to Fe3C.
(a)
(b)
50µm
(c)
50µm
(d)
50µm
50µm
Figure 14. In situ observation of cracking during fatigue: (a) Initial, (b) 10,000 cycles, (c) 16,000 cycles, and (d)
22,000 cycles.
The crack deflections induced by the heterogeneous microstructure increased the fatigue resistance of the
steel. The increase in the fatigue resistance of
the steels can be quantified by applying the
crack deflection model by Suresh [28]. The
model assumes that for a given crack length
there exist linear portions, S, and deflected
portions, D, Fig. 15. The angle between the
Figure 15. Suresh’s model [28] for quantifying crack
linear portion and the deflected portion is given
deflections. The crack consists of linear regions, S, and
deflected regions, D. The angle between S and D is θ.
by θ. By measuring S, D and θ, the true ∆K,
corrected for crack deflection, can by calculated
from the following equation:
∆K I,app


2 θ 
 D cos  2  + S 
 

≈ ∆k eff 
D+S






−1
[1]
The values of ∆KI,app calculated from the use of the Suresh deflection model represent increases in fatigue
resistance due to crack deflection. The corrected ∆K values are shown in Fig. 16. Note that the curves
shift to the right, indicating some contribution from deflection, although this contribution is quite small.
Also, the contribution from deflection does not appear to be noticeably different between the three
densities. This may be due to the fact that all three densities exhibit relatively equivalent heterogeneous
microstructures.
6.98 g/cm3
10
-7
10
-8
R = 0.8
10 -6
-2
10 -9
10 -10
10 -7
100
10
∆K (MPa-m1/2)
7.53 g/cm3
(c)
10 -6
0.3 0.1
-1
R = 0.8
0.3 0.1 -1
-7
10 -8
10 -9
10 -11
1
10
∆K (MPa-m1/2)
100
-2
R = 0.8
10 -8
10
(b)
10 -10
10 -11
1
da/dN (m/cycle)
0.3 0.1 -1
(a)
da/dN (m/cycle)
da/dN (m/cycle)
10 -6
7.34 g/cm3
Figure 16. Increase in Fatigue Resistance Due to
Deflection: (a) 6.98 g/cm3, (b) 7.34 g/cm3, and (c) 7.53
g/cm3. The extent of deflection is similar for all the
three densities.
10 -9
10 -10
10 -11
1
100
10
1/2
∆K (MPa-m )
CONCLUSIONS
In this study the effect of porosity and heterogeneous microstructure on the fatigue crack growth behavior
of sintered steels was systematically studied. The following conclusions can be made based on our results:
•
The heterogeneous microstructure plays an important role in fatigue crack behavior, in particular,
porosity significantly influences fatigue crack growth.
•
Increasing porosity and R-ratio resulted in a decrease in ∆KTH values. Increased R-ratio and
porosity also increased the monotonic contribution during fatigue due to strain localization
between pores at the crack tip.
•
The fatigue crack path is tortuous and highly dependent on microstructure. Areas of pearlite cause
crack arrest, deflection and branching, while Ni-rich and porous areas do not appear to offer
resistance to crack growth.
•
Crack deflections increase the fatigue resistance of the material for all three densities, although
the contribution from crack deflection is quite small.
ACKNOWLDEGMENTS
The authors gratefully acknowledge Hoeganaes Corp. for financial support of this research.
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