FATIGUE CRACK GROWTH OF PREALLOY Fe-0.85Mo-2Ni-0.6C STEELS
WITH A HOMOGENEOUS MICROSTRUCTURE
X. Deng, G.B. Piotrowski, and N. Chawla
Department of Chemical and Materials Engineering
Ira A. Fulton School of Engineering
Arizona State University
Tempe, AZ 85287-6006
K.S. Narasimhan and M. Marucci
Hoeganaes Corporation
1001 Taylors Lane
Cinnaminson, NJ 08077
ABSTRACT
The fatigue crack growth behavior of powder metallurgy steels (P/M steels) is strongly affected by the
nature of porosity and microstructure of steel matrix. Our previous work has focused on a premix P/M steel
prepared from Fe-0.85Mo prealloy mixed and binder-treated with 2%Ni and 0.6% graphite. In this study,
we have studied the fatigue crack growth behavior of a prealloy steel of similar composition. Use of the
prealloy powder resulted in more homogenous microstructure than the premix steel. The alloys were tested
at three different densities: 7.0 g/cm3, 7.3 g/cm3, and 7.5 g/cm3. Microstructure characterization was
conducted by optical and scanning electron microscopy (SEM). Fatigue testing was performed at various
R-ratios, ranging from –2 to 0.8. Increasing porosity and increasing R-ratio resulted in a decrease in ∆Kth.
The degree of crack closure was measured for both premix and prealloy steels at different R-ratios, and is
discussed.
INTRODUCTION
The use of powder metallurgy (P/M) steels is expanding due to the cost savings associated with near-net
shape processing [1-4]. The porosity of P/M steels is always an important concern due to its negative effect
on the mechanical behavior of these materials [4-8]. Pores are locations of high stress concentration, which
will introduce localized yielding and cracking [5,8,9]. The pore can also reduce the load bearing area and,
thus, reduce mechanical performance [4]. Besides porosity, the microstructure of the steel matrix has a
significant effect on mechanical behavior. We have found that the rate of crack propagation behavior is
highly dependent on the phases in the microstructure [10]. Thus, a comprehensive understanding of the
effects of porosity and microstructure on the fatigue crack growth behavior of sintered steels is required.
In our previous study, we examined the effect of porosity and microstructure on fatigue crack growth
behavior of a premix Fe-0.85Mo-2Ni-0.6C, The steel consisted of a heterogeneous microstructure of
pearlite, Ni-rich areas, and bainite at three different sintered densities [10]. The crack grew fastest in the Ni
rich area while it propagated significantly more slowly in fine pearlite and bainite. In the current study, we
investigated the fatigue crack growth behavior of a prealloy steel with similar composition. This steel had a
homogeneous microstructure of pearlite, although the distribution of pores was different from that of the
premix steel. The crack growth behavior was studied at three different sintered densities: 7.0 g/cm3, 7.3
g/cm3, and 7.5 g/cm3. Quantitative characterization of the microstructure was performed to determine the
degree of porosity and nature of the microstructural constituents at all densities. Fatigue crack growth
experiments were performed at a constant R-ratio, ranging from -2 to 0.8. It will be shown that increasing
porosity and R-ratio decreased the fatigue crack growth resistance of the material. Crack closure was
measured and compared for premix and prealloy P/M steels.
MATERIALS AND EXPERIMENTAL PROCEDURE
The prealloy P/M steel consisted of Fe-0.85Mo-2Ni prealloyed powder binder-treated with 0.6% graphite.
The prealloyed samples were pressed and sintered to obtain three different sintered densities: 7.0 g/cm3, 7.3
g/cm3, and 7.5 g/cm3, which are similar to the premixed densities in our previous study [10]. The 7.5 g/cm3
prealloy samples were made by a double-press/double sintering process. All samples were pressed into
rectangular blanks and sintered at 1120ºC for 30 minutes in a 90% N2–10% H2 atmosphere. The
compacting pressure needed to achieve a given density in the prealloy samples was higher than the premix
samples, due to the higher incompressibility of the Fe-Mo-Ni prealloy powder [11].
Porosity was determined by image analysis of several representative micrographs, 350 µm by 450 µm in
size, using scanning electron microscopy (SEM). The micrographs were segmented into black and white
images, and the porosity determined using image analysis software (ImageJ, Bethesda, MD).
Fatigue tests were performed on a servo-hydraulic load frame equipped with a traveling microscope
(Questar Corp., New Hope, PA). 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 [12-14]. The samples were machined by
electrodischarge machining (EDM) to the following dimensions: height of 45 mm, width of 11.5 mm,
thickness of 7.5 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. The testing
procedure is the same as the one for premix P/M steel [10]. A clip-on crack opening displacement gauge
was used to measure load-displacement data for calculating the closure load (Pcl), following ASTM E 647
guidelines. Kcl was then calculated using Pcl and the effective stress intensity, ∆Keff, was then calculated
(∆Keff = Kmax – Kcl).
RESULTS AND DISCUSSION
Microstructure Characterization
1.
Porosity
The porosity at each density was determined from several SEM micrographs using image analysis, and is
summarized in Table 1. At the lowest nominal density, 7.0 g/cm3, the measured porosity was 7.2%, while
at the highest density 7.5 g/cm3, the porosity was about 2.9%. The values obtained from the image analysis
technique correlated well with those obtained from calculations of 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.
(c)
(b)
(a)
50 µm
50 µm
50 µm
Figure 1. Microstructure of Prealloy Fe-0.85Mo-2Ni-0.6C steels at: (a) 6.98 g/cm3, (b) 7.33 g/cm3, and
(c) 7.48 g/cm3. At higher porosity the pores are larger, more irregular, and more interconnected.
Table I. Fraction of porosity versus density in prealloy Fe-0.85Mo-2Ni-0.6C steels
Sintered Density
Porosity from Sintered Density
Porosity from Image Analysis
(g/cm3)
(%)
(%)
7.0
10.3
7.2 ± 0.4
7.3
4.5
4.9 ± 0.7
7.5
3.8
2.9 ± 1.2
Etched microstructures showed the homogeneous nature of the microconstituents, in addition to the
porosity, Figure 2. The prealloy P/M steel was mainly composed of pearlite. The premix steel [10], on the
other hand, consisted of a heterogeneous microstructure of coarse pearlite, fine pearlite, Ni-rich regions
(likely Ni-rich austenite, surrounding the pores), and bainite.
2.
Pore size, shape, interpore spacing, and pore
clustering
Comparison of the pore size, shape, and clustering
shows fundamental differences in the pore size, shape,
Pearlite
interpore spacing, and pore clustering attributes of
premix and prealloy P/M steels. Pore size was measured
in a manner similar to the measurement of the fraction
of porosity. Several representative SEM images for each
density were segmented into black and white images
and analyzed using image analysis software. This
allowed measurement of pore size based on the area of
pixels of each pore. As a result, “pore size” is actually a
measure of the pore area, µm2. For the lowest density,
7.0 g/cm3 for both premix and prealloy P/M steels, the
Figure 2. Homogeneous microstructure of
distribution of pore areas is very broad, with a
prealloy P/M steel consisting of porosity and
maximum pore area of 800-1000 µm2, Figs. 3 and 4. As
pearlite.
a comparison, for the highest density, premix 7.5 g/cm3,
3
prealloy 7.5 g/cm , the distribution of pore areas is much narrower, with the maximum pore area decreased
to 400-500 µm2.
Figures 3 and 4 show that increased density decreases the fraction of larger pores. The compacting pressure
required to press a P/M sample to a higher density is much greater than the pressure required for lower
densities. As a result, the higher compacting pressure collapses most of the large pores present into clusters
of much smaller, and as will be shown, rounder pores. The effect of the increased compacting pressure will
also be evident upon examination of the interpore spacing.
3
50
(a)
3
Premix 7.5 g/cm , Pore Area Fractions
3
Premix 7.0 g/cm , Pore Area Fractions
Premix 7.4 g/cm , Pore Area Fractions
50
(b)
50
40
40
40
30
30
30
20
20
20
10
10
10
0
0
0
25 125 225 325 425 525 625 725 825 925
2
Pore Area (µm )
(c)
25 125 225 325 425 525 625 725 825 925
2
Pore Area (µm )
25 125 225 325 425 525 625 725 825 925
2
Pore Area (µm )
Figure 3. Premix Pore Areas versus Density; (a) 7.0 g/cm3, (b) 7.4 g/cm3, and (c) 7.5 g/cm3.
Increased density resulted in a decrease in the pore area and pore area distribution. In general, pores
are larger and less susceptible to changes to density in the premix samples than the prealloy samples.
3
50
(a)
3
3
Prealloy 7.0 g/cm , Pore Area Fractions
Prealloy 7.5 g/cm , Pore Area Fractions
Prealloy 7.3 g/cm , Pore Area Fractions
50
(b)
50
40
40
40
30
30
30
20
20
20
10
10
10
0
0
0
25 125 225 325 425 525 625 725 825 925
2
Pore Area (µm )
(c)
25 125 225 325 425 525 625 725 825 925
2
Pore Area (µm )
25 125 225 325 425 525 625 725 825 925
2
Pore Area (µm )
Figure 4. Prealloy Pore Areas versus Density; (a) 7.0 g/cm3 (b) 7.3 g/cm3 and (c) 7.5 g/cm3.
Increased density resulted in a decrease in the pore area and pore area distribution. Pores
tended to be smaller in area with decreased porosity and in the prealloy samples.
The fundamental physical differences between premix and prealloy powder cause a decrease in the
compressibility of the prealloy powder, [4, 15]. As a consequence, prealloy powders require a greater
compacting pressure to achieve densities similar to their premix counterparts. Comparing the pore area
distributions for the two alloying methods, for all three densities, the prealloy samples exhibit a narrower
pore area distribution and exhibit a greater number of smaller pores. The effect of changes in density have
a similar effect on the pore area distributions for the two alloying methods, in that increased density
decreases pore size. Changes in density, however, appear to have a greater impact on the pore area
distributions of the premix samples.
To further understand the impact of density and alloying method on the pore characteristics, it is necessary
to examine the pore morphology. With the image analysis software, the pore morphology was defined
using the shape of the pixilated pore perimeter and pore area, such that the pore shape, F, is determined
from the equation:
F = 4πA/P2
where A is the pore area, and P is the pore perimeter. An F value of one denotes a perfectly round pore,
whereas F values approaching zero denote increasingly irregular pores.
(1)
Analysis of the pore morphology was performed in conjunction with the pore size measurements on the
representative SEM micrographs for each density and type. Increased density resulted in an increase in the
circularity of the pores present in all samples, Figures 5 and 6. At the highest porosity, in the 7.0 g/cm3
samples for both premix and prealloy, the pore shape distribution is very broad, with a slightly greater
fraction of irregular pores than circular pores. At the lowest porosity, in the premix 7.5 g/cm3 and prealloy
7.5 g/cm3 samples, the shape distribution is skewed towards the right, representing an increase in
circularity.
3
3
Premix 7.0 g/cm , Pore Shape
20
20
(a)
3
Premix 7.4 g/cm , Pore Shape
20
(b)
(c)
15
15
15
10
10
10
5
5
5
0
0
0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1
Pore Shape, F
Premix 7.5 g/cm , Pore Shape
0
0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1
Pore Shape, F
0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1
Pore Shape, F
Figure 5. Premix Pore Shape versus Density; (a) 7.0 g/cm3, (b) 7.4 g/cm3, and (c) 7.5 g/cm3.
Decreased density resulted in an increase in the pore irregularity and pore shape distribution. For
all densities, premix pores tend to be less circular than prealloy pores.
3
20
Prealloy 7.0 g/cm , Pore Shape
3
3
20
(a)
Prealloy 7.3 g/cm , Pore Shape
20
(b)
(c)
15
15
15
10
10
10
5
5
5
0
0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1
Pore Shape, F
0
0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1
Pore Shape, F
Prealloy 7.5 g/cm , Pore Shape
0
0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1
Pore Shape, F
Figure 6. Prealloy Pore Shape versus Density; (a) 7.0 g/cm3 (b) 7.3 g/cm3 and (c) 7.5 g/cm3.
Decreased density resulted in an increase in the pore irregularity and pore shape distribution. Pores
tended to be more circular with decreased porosity and in the prealloy samples.
Comparing the two alloying methods for the three densities, the prealloy samples tend to possess a greater
fraction of near circular pores. As the porosity decreases, the difference in the fraction of circular pores
between the two alloying methods increases significantly. For the lowest density, 7.0 g/cm3 samples for the
two methods, the number of near circular pores with a F-value greater than 0.7 is approximately 13.6% for
premix samples and approximately 24.3% for the prealloy samples. Comparing these values with values
from the highest density, premix 7.5 g/cm3 and prealloy 7.5 g/cm3, the difference is greater, with values of
approximately 34.7% and 52.8%, respectively. The increase in the fraction of circular pores at the lower
porosities is, again, indicative of the higher compacting pressure required to produce prealloy P/M parts at
higher densities. As was briefly discussed previously, the creation of small, circular pores is a product of
the collapse of larger pores and of the decrease in the distance between previous powder particles during
pressing and later sintering. The creation of smaller, regular pores corresponds to a decrease in the
interpore spacing.
Interpore spacing, λ, is a measure of the distance between near neighbor pores, and can be calculated by
using a Dirichlet tessellations, Figure 7, [16]. In Dirichlet tessellations, pores in the microstructure occupy
the center of tessellated cells, where each cell wall is equidistant between two adjacent pores. For this
analysis, the interpore spacing is defined as the shortest edge-to-edge distance between adjacent pores that
share a cell wall in the tessellated micrograph [17].
Figure 7. Dirichlet Tessellations for Interpore
Spacing; The interpore spacing was determined
using Dirichlet tessellations. Pores occupy the
center of each cell in the tessellation, and the
interpore spacing is calculated as the shortest
edge-to-edge distance between adjacent pores.
The interpore spacing was determined from the
SEM micrographs used to measure pore area and
shape. Image analysis software was used to perform
the watershed procedure needed to create the
Dirichlet tessellations and subsequent near neighbor
measurements. For the different densities,
decreasing the amount of inherent porosity causes a
decrease in the interpore spacing, Figure 8 and
Figure 9. The decrease in λ is consistent with the
decrease in pore size resulting from the higher
compacting pressure. As was briefly discussed
previously, the creation of small, circular pores is a
product of the collapse of larger pores and of the
decrease in the distance between powder particles
during pressing and later sintering. The creation the
smaller, regular pores corresponds to a decrease in
the interpore spacing.
Comparing the two alloying methods, the premix
samples tend to have a much broader distribution of
interpore spacings, with a smaller fraction of pores within 40 µm or less. For the premix samples, changes
in density have a greater effect on the interpore spacing. At the lowest density, 7.0 g/cm3, approximately
31.17% of the pores lie within 20 µm or less. As density increases, a greater fraction of pores fall within
this range, such that at the highest density, 7.5 g/cm3, 45.36% of pores are within 20 µm or less. Changes
in density have less effect on the prealloy samples, such that the change from 7.0 g/cm3 to 7.5 g/cm3 results
in approximately in a change from 43.97% to 53.21% of pores within 20 µm or less.
Measurement of the variance in interpore spacing provides a relative idea of the amount of pore clustering
present in the microstructure, given that samples with more variance in λ values will have a greater degree
of pore clustering than λ values with less variance. Measurement of the coefficient of variance (COV)
provides an accurate and convenient method of quantifying pore clustering. The COV value is a measure
of the calculated interpore spacing standard deviation, σd, divided by the average interpore spacing, λavg;
COV = σd/λavg
(2)
COV values approaching zero signify an increase in ordering of the pore distributions, as the COV value
increases, so does the amount of pore clustering [9, 17].
Figure 10 shows that pore clustering increases with density for both premix and prealloy P/M steels. At
similar densities, prealloy always has higher pore clustering than premix. Since pore clustering can lead to
the early crack initiation and can also accelerate crack propagation, it may have a significant effect on the
fatigue crack growth behavior of P/M steels.
3
Premix 7.0 g/cm , Interpore Spacing
60
(a)
50
3
3
Premix 7.4 g/cm , Interpore Spacing
60
(b)
50
40
40
30
30
30
20
20
20
10
10
10
20
100 180 260 340 420
Interpore Spacing, (µm)
500
0
20
(c)
50
40
0
Premix 7.5 g/cm , Interpore Spacing
60
100 180 260 340 420
Interpore Spacing, (µm)
0
500
20
100 180 260 340 420
Interpore Spacing, (µm)
500
Figure 8. Premix Interpore Spacing versus Density; (a) 7.0 g/cm3, (b) 7.4 g/cm3, and (c)
7.5 g/cm3. Decreasing porosity resulted in a decrease in the interpore spacing. This is
most likely due to the collapse of larger pores with increased compacting pressure.
3
60
Prealloy 7.0 g/cm , Interpore Spacing
(a)
50
3
60
Prealloy 7.3 g/cm , Interpore Spacing
(b)
50
3
60
40
40
30
30
30
20
20
20
10
10
10
20
100 180 260 340 420
Interpore Spacing, (µm)
500
0
20
(c)
50
40
0
Prealloy 7.5 g/cm , Interpore Spacing
100 180 260 340 420
Interpore Spacing, (µm)
500
0
20
100 180 260 340 420
Interpore Spacing, (µm)
500
Figure 9. Prealloy Interpore Spacing versus Density; (a) 7.0 g/cm3, (b) 7.3 g/cm3, and (c)
7.5 g/cm3. Decreasing porosity resulted in a decrease in the interpore spacing. The prealloy
samples tend to have tighter pore spacing and less variation in spacing distance than do the
premix samples, due to the higher compacting pressure required.
0.8
Fatigue Behavior
0.7
1.
da/dN vs. ∆K Curves
0.6
Fatigue crack growth experiments on prealloy P/M
steels at the three densities showed that porosity had a
strong effect on the crack growth rate (da/dN) and the
Premix
0.4
threshold stress intensity factor, ∆Kth. Figure 11 shows
Prealloy
the da/dN versus ∆K behavior for all densities, at
0.3
various R-ratios. ∆Kth at the three densities varied
6.8
7
7.2
7.4
7.6
7.8
3
between 14.4-15.6 MPa·m1/2 at R = -2, to 3.3-4.9
Density (g/cm )
MPa·m1/2 at R = 0.8. The slope in the steady state or
Figure 10. Relationship between COV and
Paris law regime, m, was also measured. Increasing Rdensity of premix and prealloy P/M steels.
ratio resulted in an increase in m from about 4 at low
R-ratio to around 9 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 of premix
[10] and prealloy P/M steels, at all three porosity levels for three R-ratios: -1, 0.1, and 0.8, Figure 12.
Comparison of the fatigue crack growth behavior for the premix and prealloy samples at three R-ratios, -1,
0.1, and 0.8, shows that in both cases decreased porosity resulted in an increase in ∆Kth. This trend is
particularly evident at the 0.8 R-ratio, where Kmax is dominant. At higher, positive, R-ratios, the ∆Kth
behavior of the two steel types was similar. As R-ratio decreases to negative values, the premix steel
samples exhibited somewhat improved ∆Kth compared to the prealloy steel. This improvement is marginal
0.5
at higher densities, but increased with increased porosity. The increase in the premix ∆Kth values at
negative R-ratio could be a result of decreased pore clustering in the premix steel, as shown in Figure 10.
As a result, the premix steel had an increased resistance to cracking, especially in compression and
experienced less damage at lower growth rates.
da/dN (m/cycle)
-7
10
7.0 g/cm3
R = 0.8
R = 0.1 R = -1 R = -2
m = 9 R = 0.3
m=5 m=5 m=4
m=6
da/dN (m/cycle)
(a)
-6
10
-8
10
-9
10
-10
10
m
-11
10
2
4
6
8 10
30
10
-6
10
-7
10
-8
10
-9
10
-10
10
-11
7.3 g/cm3
(b)
R = 0.8 R = 0.3 R = 0.1 R = -1 R = -2
m=9 m=6 m=5 m=6m=4
m
2
4
1/2
da/dN (m/cycle)
∆K (MPa-m )
10
-6
10
-7
10
-8
10
-9
10
-10
10
-11
R = 0.8 R = 0.3 R = 0.1R = -1R = -2
m=9 m=6 m=5 m=5 m=3
m
10
4
6 8 10
1/2
∆K (MPa-m )
R = 0.8 R = 0.1
Figure 11. da/dN versus ∆K curves for preally P/M
steels: (a) 7.0 g/cm3 (b) 7.3 g/cm3, and (c) 7.5 g/cm3.
Porosity has a strong influence on the fatigue behavior
of the P/M steel. Increased porosity and R-ratio
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.
30
(a)
-6
10
-6
10
-7
10
10
-8
10
-8
10
-9
10
-9
-10
10
-11
2
4
6 8 10
1/2
∆K (MPa-m )
(b)
R = -1
-7
10
30
7.5 g/cm3
(c)
2
6 8 10
1/2
∆K (MPa-m )
7.0 g/cm
3
7.3 g/cm
3
7.5 g/cm
3
30
10
-10
10
-11
2
R = 0.1
R = 0.8
4
6
R = -1
8 10
7.0 g/cm
3
7.4 g/cm
3
7.5 g/cm
3
30
1/2
∆K (MPa-m )
Figure 12. Effect R-ratio on fatigue crack growth, as a function of density. (a) prealloy and (b) premix
P/M steel. Increasing porosity and R-ratio resulted in a decrease in fatigue crack growth resistance.
2.
Fatigue Crack Closure
Kmax
Based on the driving forces for crack
propagation there are two main models, that
∆Keff = Kmax - Kcl
have been proposed to explain fatigue crack
growth. One is the classical crack closure
K
∆K = Kmax - Kmin
model [19] and the other is a more novel
Kcl
two-parameter (Kmax - ∆K) model [20,21].
The concept of crack closure was first
Crack remains closed
proposed by Elber [22], who observed that
Kmin
fatigue cracks could be closed, even at a far
field tensile load, due to plasticity at the
wake of the crack tip. Once in the crack
Number of Cycles (N)
wake, plastic deformation partially fills the
crack mouth causing premature contact of
Figure 13. Schematic showing decrease in effective
the cracked surfaces. When the crack
stress intensity, ∆Keff, due to premature closure of the
surfaces are in contact (although the far
crack.
field applied stress is tensile), it is assumed
that no damage occurs at the crack tip. Kcl is the stress intensity factor corresponding to the load at which
the crack surfaces come in contact. Thus the stress intensity factor, which drives fatigue crack propagation
(∆Keff), is less than the applied stress intensity factor, ∆K (Kmax – Kmin). In other words, ∆Keff is equal to
Kmax – Kcl as shown in Figure 13. This closure mechanism has been used to explain the effects of load ratio
on fatigue crack growth rates as well as the influence of load history [23], and is termed plasticity induced
closure. Following Elber’s work, additional mechanisms of crack closure were also identified. These
include closure due to oxides and asperities wedging the crack mouth, roughness of the crack surfaces,
transformation toughening ahead of the crack tip, and the presence of viscous fluid between the crack
surfaces [24-29]. For P/M steels, little work has been done on the crack closure effect.
To incorporate closure data into the current analysis on both premix and prealloy P/M steels, fatigue crack
closure measurements were made using crack opening displacement (COD) gauge. Load-displacement
data was acquired to measure the crack closure load. It was found that for P/M steels there was noticeable
deviation from linearity in the load-displacement curve at threshold and near threshold region. The
deviation from linearity in the load-displacement curve is the proof that the crack closes prematurely [22].
Crack closure load, Pcl, was calculated from the unloading portion of the load-displacement curve
following ASTM E 647 guidelines to calculate effective stress intensity factor. The plot of crack growth
rate (da/dN) versus both nominal (∆K) and effective stress intensity (∆Keff) for both prealloy and premix
P/M steels are shown in Figure 14 and Figure 15 respectively. Both premix and prealloy P/M steels show
similar da/dN-∆Keff behavior. From Figure 14 and 15, it can be seen there is a significant different between
∆K and ∆Keff for low R ratios, especially for R < 0, demonstrating there is considerable crack closure.
When closure is subtracted, the curves collapse into a single curve for R ≤ 0.3 for both premix as well as
prealloy P/M steels. When R > 0.3, there is no significant closure effect since Kmin > Kcl.
The similarity in fatigue threshold values between premix and prealloy steels can be explained as follows.
In our previous work we showed that, in premix steels, crack growth was much faster through Ni-rich areas
than the rest of the heterogeneous microstructure. In the current work, we have shown that the degree of
pore clustering in prealloyed material is higher than that of premix steels. Thus, although the prealloy
material has a homogeneous microstructure, consisting primarily of tough pearlite, the cracks grow through
prior particle boundaries and link the more clustered pores. Thus, the bulk of the pearlite in the prealloy
microstructure is avoided by the fatigue crack and results in a ∆Kth that is similar to that of the premix
alloy.
10
-6
10
-7
10
-8
10
-9
(a) 7.0 g/cm3
R = 0.8
10
10
R = 0.3
-10
∆K
R = 0.1
10
-7
10
-8
10
-9
(b) 7.3 g/cm3
R = 0.8 R = 0.3 R = -1
10
-10
10
-11
∆K
6 8 10
1/2
∆K, ∆K (MPa-m )
2
30
4
-6
10
-7
10
-8
10
-9
10
-10
10
-11
R = 0.3 R = 0.1 R = -1
R = 0.8
Figure 14. Crack growth rate as a function of both
nominal and effective stress intensity factor for
prealloy P/M steel, (a) 7.0 g/cm3, (b) 7.3 g/cm3,
and (c) 7.5 g/cm3.
∆K
∆K
4
6
∆K, ∆K
8 10
eff
1/2
eff
30
(MPa-m )
(a) 7.0 g/cm3 R = 0.3 R = 0.1
10
-6
10
-7
10
-7
10
-8
10
-8
10
-9
10
-9
-6
R = 0.8
10
-10
10
-11
R = -1
∆K
∆K
2
4
6
8 10
1/2
∆K, ∆K (MPa-m )
-6
10
-7
10
-8
10
-9
R=0.3
10
-10
10
-11
R=0.8
R=0.1 R=-1
∆K
∆K
30
2
4
6 8 10
1/2
∆K, ∆K (MPa-m )
eff
30
eff
(c) 7.5 g/cm3
R = 0.1
R = -1
R = 0.8
10
-10
10
-11
Figure 15. Crack growth rate as a function of both
nominal and effective stress intensity factor for
premix P/M steel, (a) 7.0 g/cm3, (b) 7.4 g/cm3, and
(c) 7.5 g/cm3.
∆K
∆K
R = 0.3
3.
(b) 7.4 g/cm3
eff
eff
10
30
(c) 7.5 g/cm3
2
10
6 8 10
1/2
∆K, ∆K (MPa-m )
eff
eff
eff
10
∆K
R = 0.1
eff
-11
4
-6
R = -1
∆K
2
10
2
4
6 8 10
1/2
∆K, ∆K (MPa-m )
eff
30
eff
Fractographic analysis
Figure 16 shows the typical fatigue crack propagation in prealloy P/M steels. In our previous study on
premix P/M steels, fatigue crack growth behavior is quite different in different phases [10]. For the Ni-rich
regions, cracks tend to propagate in a linear fashion, suggesting that the Ni-rich regions offer little
resistance to crack propagation. For the pearlite regions, cracks tend to be highly deflected, with some
evidence of the Fe3C plates in the ferrite matrix bridging the crack. For prealloy steels, the crack
propagates mainly inside pearlite. Figure 16 shows significant crack defection and branching partially due
to the Fe3C plate deflection and bridging, and partially due to the pore clustering. The high degree of
wedging and branching of crack path leads to the significant mismatch in the crack surfaces while
unloading, which is attributed to the considerable crack closure.
(b)
(a)
50µm
Figure 16. Typical side view of the fatigue cracks for prealloy PM steel, (a) crack wedging, and
(b) crack branching.
Figure 17 shows the fractograph of prealloy P/M steels after fatigue test. Two important features were
observed at different fatigue crack growth region. In the threshold regime corresponding to low stress
intensity, the fatigue striation was dominant. In this regime, the delaminating of Fe3C plate happens
significantly, Figure 17 (b). In Paris law regime, the plastic dimple becomes another important feature
besides fatigue striation, Figure 17 (c) and (d), indicating the increasing monotonic contribution in this
regime.
CONCLUSIONS
In this study the fatigue crack behavior of prealloy Fe-0.85Mo-2Ni-0.6C P/M steels was systematically
studied. The following conclusions can be made based on our results:
•
Porosity, along with pore shape, pore clustering significantly influences fatigue crack growth of
prealloy P/M steels. Compared with premix P/M steel, prealloy steels have higher fraction of
smaller or circular pores, narrower distribution of pore size and interpore spacing, and higher pore
clustering.
•
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.
•
There is significant crack closure for both premix and prealloy P/M steels at R ratio less than 0.3.
The significant crack wedging and branching is the main reason for crack closure of P/M steels.
The effective stress intensity factor is the actual driving for crack growth under low R ratios.
There is no significant crack closure effect for high R ratios more than 0.3.
•
For prealloy P/M steels, the fatigue crack path is tortuous. The pearlite causes crack arrest,
deflection and branching. Fatigue striation happens dominantly in threshold regime and plastic
dimple becomes another important feature in Paris law regime.
(a)
(b)
1µm
50µm
(c)
(d)
50µm
5µm
Figure 17. Fatigue fracture surface of prealloy PM steel at (a) threshold regime, (b) higher
magnification of threshold regime, (c) Paris law regime, and (d) higher magnification of Paris law
regime. Fatigue striations are the dominant feature in threshold regime, while dimples are the
significant features in the Paris law regime.
ACKNOWLDEGMENTS
The authors gratefully acknowledge Hoeganaes Corp. for financial support of this research.
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