Introduction
Thanks to its castability and mechanical properties, A319 alloy
belonging to commercial cast Al-Si alloys, is widely used in automotive
components such as for cylinder head of engine where the alloy should
undergo Low Cycle Fatigue (LCF) loading due to the engine startup and
shutdown1–8. A319 alloy is a hypoeutectic Al–Si
alloy with two main solidification stages: (1) formation of aluminum
rich (α-Al) dendrites, (2) development of two-phase (Al)–Si
eutectic3. However, the presence of additional
alloying elements such as Mg and Cu, as well as of impurities such as
Fe, leads to more complex solidification reactions and thus results in
complex multiphase microstructures that have an influence on the macro
mechanical properties3,8.
Since a decade, the traditional casting process in automotive industry,
i.e. the Die Casting (DC) process, is being replaced by the Lost Foam
Casting (LFC) process in order to optimize geometry, reduce cost and
control consumption 9,10. However, the A319 alloy
casted by LFC has a much shorter LCF life than that by DC due to a
slower cooling rate in LFC which results in numerous large pores and a
coarser microstructure 11. Understanding the
microstructures in A319 alloy casted by LFC and especially their
influence on the LCF damage mechanisms is crucial for the casting
process improvement and the strength and life prediction.
Some studies12–19 have been performed to understand
the damage mechanisms of A319 alloy in past years. Large pores above a
critical size are assumed to be the most likely crack initiation
sites12,17; while hard inclusions, i.e. eutectic Si,
iron intermetallics and copper containing phases are generally
responsible for crack propagation but also for crack initiation in the
absence of large pores13–19. However, most of the
studies were performed in High Cycle Fatigue (HCF) regime. Only a few
studies focus on the damage mechanisms of Al-Si alloys with large
casting defects in LCF regime11,20. Moreover, even if
post-failure analysis allows identifying defect at initiation site, it
does not allow establishing the relation between crack growth and
microstructure. This analysis is made even more difficult in LCF as
failure often results from multicracking rather than from a single
crack18. Recently, 3D in-situ analysis with X-ray
Computed Tomography (X-ray CT) has been used to study the damage
mechanisms of A319 alloy in LCF by the authors21,22.
It allows following crack initiation and propagation in the bulk of the
studied material and highlights the influence of the strain
localizations due to the pores on crack initiation. However, the
understanding of the damage mechanisms during the propagation stage is
quite limited due to the non-real-time method as only a few scans could
be recorded during the LCF tests and the damage evolutions during two
scans could not be observed.
In order to further study the LCF damage micromechanisms in LFC A319
alloy, an experimental protocol, that not only allows following crack
initiation and propagation on surface in real-time but also measuring
surface in-plane displacement and strain field and linking the surface
damage evolutions with the defects below surface, has been set up. At
first, the microstructure of the studied material was characterized by
using X-ray CT and Scanning Electron Microscope (SEM) together with
Energy Dispersive Spectroscopy (EDS). Laboratory Computed Tomography
(Lab-CT) was used to characterize pores in the bulk of specimen for the
selection of the region where the following in-situ observation will be
performed, as well as for the comprehensive analysis after failure. LCF
tests were performed with surface in-situ observations to follow crack
initiation and propagation in real-time. Post-mortem analysis using
Optical Microscope (OM) and SEM completed the in-situ surface analysis.
Digital Image Correlation (DIC) method was used to measure the
mechanical full-fields during LCF tests. An etching method, which gives
a natural texture to the aluminium dendrites by revealing segregation of
Si, was developed to make DIC feasible to an acceptable resolution
without adding a speckle pattern that will also mask the microstructure.
The relation between microstructure, crack initiation and propagation,
and strain field was established by using the developed experimental
protocol to allow a better understanding of the LCF damage mechanisms in
LFC A319 alloy.
Material and
experimental methods
Material
The material used in this study was extracted from an A319 (7.85 wt.%
Si, 0.30 wt.% Fe, 3.05 wt.% Cu, 0.19 wt.% Mn, 0.28 wt.% Mg)
prototype cylinder head, which was manufactured using LFC process, from
the most critical area, i.e. the fire deck22 (Fig.
1a).
2D characterizations were first performed using SEM together with EDS in
order to identify the different phases. The microstructural constituents
in the studied material are pores and hard inclusions including eutectic
Si, iron intermetallics and copper containing phases (Fig. 2a). Iron
intermetallics involve α and β phases: α(AlFeMnSi) phase usually
presents a ”Chinese script” morphology (in Fig. 2a-b, the iron
intermetallics marked by an arrow is an α phase), while β (AlFeSi) phase
almost always presents a needle-like morphology (Fig. 2c). The copper
containing phases involve Al2Cu and AlCuMgSi phases the
latter being always connected with the former (Fig. 2b-c). Both copper
containing phases have a very complex microstructure while the size of
AlCuMgSi phases is usually much smaller (less than 1 µm) than that of
other hard inclusions. The copper containing phases may connect with the
Al-Si eutectic phase (Fig. 2a) or with iron intermetallics (α and β)
(Fig. 2b-c).
3D characterizations were performed using X-ray CT in order to
understand the complex microstructural morphologies of the constituents.
The reader can refer to Ref. 21 for the 3D
characterizations of pores that have larger sizes than hard inclusions.
This paper herein focuses on the 3D characterizations of hard inclusions
while the reader can refer to Ref. 3 for the
experimental set-up details.
Fig. 3a (resp. b) gives the 3D rendering of iron intermetallics (resp.
copper containing phases) in the whole studied volume while Fig. 3c
gives the 3D rendering of Si in a subvolume of 500 × 500 × 500
µm3 to allow a better understanding of its morphology.
Both Si and iron intermetallics (including α and β phases that are
connected in volume) form complex and highly interconnected networks and
extend in the whole characterized volume, i.e. the maximum Feret
diameter 21 of individual segmented objects reaches
the maximum size of the characterized volume; the example of iron
intermetallics is shown in Fig. 3a. The copper containing phases, shown
in Fig. 3b, are considered as a whole as AlCuMgSi phase cannot be
distinguished from the connected Al2Cu phase due to its
small size compared to the current resolution. They also form a complex
microstructure but are less extended than Si and iron intermetallics
with a maximum Feret diameter of 0.98 mm that is smaller than the
maximum size of the characterized volume, i.e. 2.92 mm. The volume
fractions of the different microstructural constituents estimated under
the current used resolution are: pores≈1.0% 21,
Si≈9.7%, iron intermetallics≈7.4%, copper containing phases ≈1.3%3.
The Feret diameter, which is used to characterize the size of pores, is
no longer suitable to characterize the size of hard inclusions due to
their interconnected and extended morphologies. Thus a granulometric
analysis 23, which measures the thickness distribution
of a given phase, was introduced to characterize the size of hard
inclusions. Fig. 4 compares the size distribution, i.e. Feret diameter
for pores 21 and local thickness for hard inclusions,
of the microstructure constituents; the volumetric frequency is defined
as the volume of phase having the considered size divided by the total
volume of phase. Pores are much larger than hard inclusions. The pores
having Feret diameter between 400 µm and 1200 µm represent 98% of the
total volume of pores. In the same time, the iron intermetallics (resp.
copper containing phases) having a local thickness between 6 µm and 18
µm represent more than 82% of the total volume of iron intermetallics
(resp. copper containing phases) and the Si having a local thickness
between 3 µm and 12 µm represent more than 99% of the total volume of
Si. From a size point of view, the pores can be considered as the most
critical defects in the volume.
Specimens
Two standard notched plate specimens (Fig. 1), named as “Specimen 1 and
2”, were used for LCF tests with surface in-situ observations. Fig. 1c
summarizes the procedure of extraction, preparation of specimens and
selection of region of interest (ROI) where in-situ observations will be
performed.
In order to ensure a high enough resolution for surface in-situ
observations, the field of view is restricted. Thus, it is important to
identify the area where cracks are most likely to initiate on surface in
order to select this area as the ROI. Cracks are assumed to be more
inclined to nucleate at large pores, especially at surface large
pores12,14,17,24,25.
Lab-CT allows pores in the bulk of specimen being revealed in 3D
nondestructively 21 and was thus used to select ROI:
the selected areas have a cluster of subsurface pores or pores near the
surface. Such a characterization of pores will also help to establish
the relations between surface observations and pores in the bulk for
comprehensive analysis after failure.
The key parameters of Lab-CT used for the characterizations of pores in
the bulk of “Specimen 1 and 2” are listed in Table 1. The used
resolutions (voxel size ~ 6 µm) allow the large pores in
the whole critical volume, i.e. notched area, to be revealed.
Grinding was performed on four sides and each corner of specimens using
abrasive papers of grades up to 4000 grits in succession to ensure that
all the surfaces are smooth and thus prevent the generation of stress
concentration. A fine polishing was performed using polishing cloths and
diamond suspensions up to 1/4 μm for surface observations. The selected
ROI was then recorded using OM to give an overview of microstructural
constituents’ morphology before loading.
Testing
procedure
The tensile-tensile fatigue tests were performed at room temperature
with an Instron 8501 servohydraulic testing machine (Fig. 5b). The
cyclic loading was controlled in displacement at a speed of 0.15 mm/s.
During the test, the applied load was recorded and an extensometer,
which was installed across the notched area of the specimen, was used to
measure the macroscopic deformation (Fig. 5c). Surface in-situ
observations were performed at microstructural scale on surface, using
Questar long distance microscope equipped with a JAi 500 CCD camera. It
allows the surface to be followed in real-time during tests. The surface
on which optical observation was performed is called “flat surface”
hereafter by opposition with the notch root called “notched surface”.
The Questar was mounted on a translation stage that allows displacement
in three directions. The resolution, i.e. the pixel size, was adjusted
with additional lens and by changing the distance between the specimen
and the camera. The field of view, i.e. the diagonal length of one image
of size 2456 × 2050 pixels2, is about 1.2 mm with the
used resolution, i.e. 1 pixel ~ 0.38 μm. A coaxial
lighting was used for surface observations, and its intensity was
adjusted to give an appropriate texture.
The tests were periodically interrupted with the specimen held under
load when obvious damage evolutions were observed on surface, otherwise
the tests were also interrupted after a few cycles (less than
10n+1 /3 cycles for the tests performed between
10n cycles and 10n+1 cycles) even
without evidence of damage evolutions . Once the tests were interrupted,
six images (3×2, named as ‘Zone 0’, ‘Zone 1’ … ‘Zone 5’ in Fig.
5c) were taken in order to cover the selected ROI (about 2.7×1.5
mm2). An appropriate overlap between adjacent images
is necessary for the further stitching of the images to a larger one,
which was performed using automatic Image Stitching
plugins26 in ImageJ. The basic loading conditions of
each specimen are listed in Table 2. Post mortem analyses were performed
on the fracture surface by using OM and SEM in order to identify the
possible crack initiation sites in fatigue and to study the fracture
morphologies.
2.4 Surface strain
field measurement
2.4.1 Development
of etching speckle pattern
The surface full-field measurement was performed using DIC technique and
the accuracy and the spatial resolution of the measurement directly
depend on the presence of numerous and finely dispersed markers on the
specimen surface 27. This random texture is usually
obtained through the application of a paint speckle pattern. The obvious
drawback is that the speckle also masks the microstructure underneath
and thus prevents the study of the relationship between microstructural
features and strain heterogeneities.
Etching on polished surfaces may be an alternative to paint speckle as
it can reveal the different phases hence increasing the image texture.
Tint etching consists in the formation of a film whose thickness is
different from one constituent to the other. Coloration is due to
interference effects (under white light) in the deposited film so that a
difference of film thickness will result in different colors. A change
in etch duration will also result in a different film thickness and will
thus give rise to a totally different set of colors. A Weck’s reagent -
100 mL water, 4 g KMnO4 and once dissolved 1 g NaOH –
was used for etching with a 14 s duration. This method reveals the
segregation (coring) in the dendrites when applied to cast aluminum
alloys 28. Its main advantage is that it does not
attack intermetallic phases and precipitates 29 so
that it is not expected to interfere with the damage mechanisms to be
studied. Indeed, no pit or grooves were observed under SEM. The quality
of the so obtained natural speckle can be assessed by comparing the
gray-level distributions in the same area of the specimen surface before
and after tint etching (Fig. 6a). Before etching, the histogram shows
two narrow peaks, i.e. one small for the hard inclusions and one tall
peak for the aluminum, as the image is almost binary. After etching, the
image dynamic shows a wider range of gray-levels within the Aluminum
matrix where etching has revealed Si segregation and a broad peak for
hard inclusions within the interdendritic space. The large standard
deviation of the gray-level distribution (Std. Dev. in Fig. 6a) of the
specimen surface, as compared to that before etching, indicates a richer
and more random speckle pattern whose suitability for image registration
was tested.
2.4.2 Validation of
the developed DIC measurement method in tensile test
In order to validate the ability of the developed DIC measurement method
by using etching of the surface to reveal strain heterogeneity at the
microstructure scale, this method was primarily applied in an in-situ
tensile test performed on another Al-Si alloy (A356 alloy: 6.5-7.5 wt.%
Si, <0.3 wt.% Fe, 0.25-0.4 wt.% Mg). The natural tracers
revealed using the etching method for A319 and A356 alloys are similar
(Al matrix (majority) and hard inclusions)3. The
tensile test was performed at a displacement rate of 1 µm/s using a
servo-hydraulic testing machine (Fig. 6b). A small size dog-bone-shaped
specimen, i.e. minimum cross-section of the flat specimen is 1.43×1.34
mm2, was used and a shallow notch was introduced to
have a slight stress concentration in the central part of the specimen
and force the damage events to occur in this area, i.e. ROI. The test
was periodically interrupted with the specimen held under load and the
in-situ surface observation system, i.e. Questar with camera, was used
to capture images in four adjacent zones of the ROI with a pixel size in
the image of about 0.338 µm in order to cover the whole specimen
cross-section in the notch area. Stitching of the four adjacent images
acquired at each load step was performed using the plugin
MosaicJ30 in Image J software that allows rigid
registration of images. DIC was performed using the publicly available
computer program Elastix31 for intensity-based image
registration on the stitched images.
The measurement accuracy32,33 of DIC was evaluated
through the comparison between a prescribed displacement and the
displacement measured from DIC 34. One reference image
was taken on the specimen surface, then a given displacement (10 µm in
the present case) was imposed on the Questar translation stage in a
direction parallel to the specimen surface and the moving image was
taken at this position. The uncertainty is estimated from the standard
deviation of the displacement field measured by DIC between the two
images, i.e. reference image and moving image35–37.
The estimated uncertainty of DIC measurement in this study shows that
the natural speckle pattern provided by color etching can be considered
good enough to yield an uncertainty that remains low, i.e. 0.08 µm (0.23
pixels), at a size of element small enough to result in a high spatial
resolution, i.e. 10.72 µm (32 pixels), in the measured field. Besides
the average strain computed from the strain fields derived from the
measured displacement fields allows to draw a stress-strain curve
(represented with symbols in Fig. 6b) that is consistent with the curve
obtained on larger specimens with appropriate extensometer measurement
(represented with a plain line in Fig. 6b). Thus, the experimental
protocol developed in this study to perform DIC at the scale of the
microstructure can be used to understand damage mechanisms at a fine
spatial scale in Al-Si cast alloys.
In Fig. 6c, strain heterogeneity is observed in the ROI of the specimen
and was compared with SEM observations at a magnification large enough
(500) to detect microcracks (circled in white in the lower image in Fig.
6c) that were not visible in the optical image. In fact, strain
localization appeared in the eutectic Al-Si constituent at a stress
level of 220 MPa (average strain of about 0.44 %) and increased at a
stress level of 237 MPa (average strain of about 0.71 %). The
comparison of the microcracks (white circles in Fig. 6c) and the
debonding (red circles) of eutectic Si particles observed after the test
with the strain field at the last loading step before failure shows (1)
that most of the strain localization can be ascribed to Si particles and
(2) that Si damage has occurred at areas where a strain localization was
indeed visible at the last loading step of the tensile test. To the
authors’ knowledge, this type of comparison between DIC and fracture in
such materials at such a scale is a new result.
Thus, the developed DIC measurement method established the relation
between the strain heterogeneities and the damage events successfully in
tensile tests. The monitoring conditions (camera, pixel size) are indeed
similar to those of the fatigue test even if the size of the specimen
and the loading conditions are different as larger specimens are
preferred for fatigue test in order to study the propagation behavior.
This method will then be used in the following fatigue tests.
2.4.3 DIC
measurement for fatigue test
The polished specimens in 2.2 were etched by using the developed method
in 2.4.1 on the selected ROI surface prior to the fatigue tests. The
digital images acquired at the minimum loading of 1stcycle was used as the reference for DIC measurements and the images
acquired at the maximum loading of other cycles were used as moving
images. A 32 pixels element size (12.0 µm), which corresponds to an
acceptable uncertainty (0.09 µm), was selected.
Results
3.1 Surface in-situ
observation
The basic results of the fatigue tests performed on Specimens 1 and 2
are listed in Table 2. Surface cracks were observed to initiate only in
the selected ROI for both specimens. Thus, the selection of ROI using
Lab-CT before fatigue tests was successful.
3.1.1 Crack
initiation
In Specimen 1 (Fig. 7), no obvious cluster of pores was observed near
the flat surface in the notched cross-section using Lab-CT. However, a
surface pore located in the notch root (marked by arrows in Fig. 7a) is
less than 0.1 mm from the flat surface. The pore’s volume is about 0.012
mm3, and its max. Feret diameter is about 0.38 mm.
Although it is not very large compared with standard large pores in this
material 21, this pore is supposed to be the most
probable crack initiation site on the observed surface through the
comprehensive evaluation based on the influence of pores’ size and
location12,21,25: (a) it is quite near the flat
surface and the notch root contrary to other larger pores in bulk, (b)
it is larger than other pores located on the observed surface. Thus, the
flat surface area in front of this pore was selected as the ROI. In
fact, Specimen 1 failed from a single crack (Fig. 7b) that nucleated at
Si phase on the flat surface right in front of this pore. In order to
check the influence of this pore on crack initiation, fracture analysis
was performed.
One half of the broken specimen was observed using a single-lens reflex
camera (Fig. 7d). Zone A is more likely to be crack initiation and
propagation areas instead of final fracture area as it is more flat than
other areas in fracture surface. River patterns 38, of
which the direction represents the direction of the crack propagation,
were indeed observed around the above mentioned pore in Fig. 7f with a
higher resolution. Thus, the crack initiation, which was observed on the
flat surface, could be ascribed to this pore that is surface-breaking in
the notch area but is some distance below the flat surface where in-situ
observations were done.
SEM observations were also performed on the whole fracture surface. The
area, i.e. Zone A in Fig. 7d, where the river patterns were observed
corresponds to the area where crack initiated and propagated. The other
areas without river patterns are the fast fracture area in the final
stage of fatigue test. This is consistent with the in-situ surface
observation as cracks were only observed to grow on surface above this
area.
Several pores, marked by arrows in Fig. 7e, were found in the crack
propagation area, i.e. Zone A. They present evidence of river patterns
so that it is difficult to identify the first initiation site inside the
specimen. There are several possible scenarios: (1) One crack initiated
from one of these pores, and then grew through other pores towards the
free surface. (2) Cracks initiated from several pores at the same time
or at different cycles, and then coalesced during the propagation. (3)
Cracks initiated from hard inclusions in bulk that are not connected
with pores under the influence of the stress concentration induced by
pores and then propagated through the pores.
By using surface in-situ observations and post-mortem analysis, one
cannot confirm which case is the true scenario for the crack initiation
and propagation in bulk. However, the influence of pores below the
observed flat surface on crack initiation observed on surface is
certain.
Fig. 8 presents a comprehensive analysis of Specimen 2 including 3D
rendering of pores in the bulk of specimen (Fig. 8a), in-situ surface
observations (Fig.8b, e), OM surface observation (Fig. 8d) and fracture
analysis (Fig. 8c, f and g). In Specimen 2, cracks initiated on surface
from five different sites identified as A, B, C, D, and E in Fig. 8b.
Cracks initiated from sites A and B (Fig. 8e) before the maximum load of
the 1st cycle was reached, i.e. below an applied
stress of 80MPa. The opening of these two cracks became larger with a
further increase of the stress up to the maximum value of the
1st cycle (in the top right of Fig. 8e). Thus, the
maximum loading in the 1st cycle was stopped at about
98 MPa, which is lower than for Specimen 1.
SEM post-mortem analysis showed that these two crack initiation sites on
surface were not only at hard inclusions but also in Al matrix and that
they are located just above Pore A (Figs. 8a-c). Thus, the cracks were
likely to have initiated from Pore A and then propagated to the flat
surface. Fracture analysis in Fig. 8f shows that the distance between
initiation sites A/B and Pore A is less than 0.1 mm. The cracks
initiated on surface at sites A/B during the tensile stage of the
1st cycle, thus the damage mechanisms could not be
ascribed to fatigue but to the large stress concentration caused by the
large Pore A in subsurface. This crack corresponds to brittle fracture
zones with cleavage of Si (Fig. 8f).
Cracks were observed to initiate at sites C and D (Fig. 8b) suddenly
between the 15 860th cycle and the
15 890th cycle. These two initiation sites were across
hard inclusions and Al matrix and cracks opened very quickly. Pore B is
below the initiation site C (Figs. 8a-c) at a depth of about 0.8 mm as
shown in Table 3. Although the river patterns marked in Fig. 8g revealed
the direction of cracks originated from the cleavage of a Si particle in
a local area, one cannot obtain more information about crack propagation
in a larger area. However, the ‘ridge’ line (Fig. 8g), which formed
during the coalescence of cracks from two different
planes39, was observed between initiation site C and
Pore A (Fig. 8c). This implies that crack initiation C may be due to
coalescence of two cracks.
Cracks were observed in-situ to initiate simultaneously at site C and at
site D where post-mortem analysis confirms the presence of Si phase.
Site D is not present on fracture surface as it is covered by the
fracture surface caused by Crack 1 branch in Fig. 8d. No subsurface pore
was found around initiation site D and no crack was observed to
propagate to this area from other pores. Thus, hard inclusions may have
acted as initiation sites at or around site D as hard inclusions are
indeed observed to act as crack initiation site due to strain
localizations in previous study by using 3D in-situ
analysis21,22.
Crack 2, which is marked in Fig. 8d, was observed to initiate from site
E (Fig 8b) near Pore C after 19 350 cycles. The influence of surface
Pore C on crack initiation at site E is obvious.
3.1.2 Crack
propagation
The cumulative crack lengths of Specimens 1 and 2 were measured using
ImageJ software on in-situ captured images for the quantitative analysis
of crack propagation (Fig. 9a). The measured crack length corresponds to
the distance between the initiation site at the notch root and the
farthest crack tip. Crack observed on surface of Specimen 2 is much
longer than that of Specimen 1 due to numerous large near surface pores
in Specimen 2 (Figs. 7-8).
In Specimen 1, crack initiating at Si particle after 7 000 cycles (Fig.
7b) propagated along hard inclusions, including Si phase, iron
intermetallics (both α and β phases), Al2Cu and AlCuMgSi
phases, until 65 000cycles where the crack tip was arrested by hard
inclusions (Fig. 7c). Meanwhile, another crack initiated at hard
inclusions at a 2nd initiation site below the arrested
crack tip (Fig. 7c) after 85 000 cycles. Inside the specimen, this
secondary crack may be a branch of the main crack as it is close to the
arrested principal crack tip (Fig. 7c) or it may also be an independent
nucleation site. However, one cannot confirm it using fracture analysis
as the 2nd initiation site is covered in the final
fracture surface. The principal crack was arrested until final failure.
When the crack from the 2nd initiation site (Fig. 7c)
connected with this principal crack after 90 000 cycles, failure
occurred quickly.
After nucleation at sites A and B in Specimen 2 at 1stcycle, cracks were arrested for more than 2 000 cycles
(step
1 in Figs. 8d and 9a) by hard inclusions perpendicular to each crack
tips. Once the hard inclusions had failed, the two cracks grew again and
connected on surface. Both cracks (Fig. 8e) coalesced due to crack
propagation below the surface under cyclic loadings as river patterns
were observed between Pore A and these two initiation sites (Fig. 8f).
Then, this coalesced crack was observed to propagate along hard
inclusions (step 2 and 3 in Figs. 8d and 9a). After 15 860 cycles, one
crack, which has nucleated from Pore A, propagated towards site C on
surface, and another crack initiated at site D on surface. They joined
with each other (step 4 in Figs. 8d and 9a) and also joined with the
cracks that nucleated from site A/B through cracks growth along hard
inclusions after 18 000 cycles. Transcrystalline fracture40 was observed in the fracture surface below site D
(Fig. 8g). It may be caused by monotonic loading when cracks coalesced
between sites C and D. However, one cannot confirm it just from fracture
analysis. This large coalesced crack, i.e. Crack 1 in Fig. 8d,
propagated until Crack tip 1where the crack was arrested by hard
inclusions. Then, Crack 2 propagated along hard inclusions towards the
arrested Crack 1 (Fig. 8d). In the same time, a branch crack propagated
in the direction perpendicular to the loading direction and along hard
inclusions. When Crack 2 coalesced with the branch of crack 1 after
20 400 cycles, fracture occurred in the following 30 cycles.
3.2 DIC
measurements
DIC measurements are then applied to measure the surface in-plane
displacement and strain field for the comparisons with the observations
shown in section 3.1.
For Specimen 1, the cracks initiated in Zone 0 after 7 000 cycles and
the propagation remained in Zone 0 (Fig. 7c) even in the last in-situ
observation step (about 105 cycles) before final
failure; no cracks could be observed in the other five zones. On the
contrary for Specimen 2, the cracks initiated in Zones 0 and 5 during
the tensile stage of 1st cycle and they were in-situ
observed to grow in all the six zones before final failure.
DIC measurements were performed between the reference image at the
minimum load of the 1st cycle and the images at the
maximum load of different cycles. The average longitudinal strain
deduced from DIC at the maximum loading of different cycles in Zones 0
and 2 of Specimen 1 and that in the whole ROI of Specimen 2 are shown in
Fig. 10. The macroscopic average strain measured by extensometer is also
shown for comparison.
For both Specimens 1 and 2, the average longitudinal strain, deduced
from DIC, in notched area zones was a little larger than the macroscopic
average strain measured by the extensometer due to the strain
localization in this area. This localization is more significant around
the notch root; indeed, the average strain measured by DIC in Zone 0 is
generally higher than that in Zone 2.
It should be mentioned that the measured strain localizations in the
field from DIC may arise by the possible out-of-plane displacement that
induce errors for DIC measurements as 2D-DIC method (instead of 3D-DIC
method) is applied in the present study. In addition, the following two
situations should also be considered: (1) Before crack initiation, they
represent the real localizations in the case of the nonexistence of
out-of-plane displacement; (2) After crack initiation, they indicate the
discontinuities appeared at the crack. For Specimen 1, DIC measurements
are focused on two zones: Zone 0 where crack initiation and propagation
were observed and Zone 2 where no crack was observed. Cracks were
observed to initiate in Zone 0 due to the pores below the surface in an
area where a larger deformation was indeed measured. During the fatigue
test, the gap (Fig. 10) between the measured strain in Zone 0 and that
in Zone 2 grew larger after crack initiation in Zone 0 after 7000
cycles: the measured strain in Zone 2 is basically below 0.2% while
that in Zone 0 increases from less than 0.3% to more than 0.5%. Thus,
as mentioned above, the large deformations measured by DIC for Zone 0 of
Specimen 1 from 7000 cycles in Fig. 10 were mainly induced by the
discontinuities appeared at the crack. The similar situation was also
observed for Specimen 2. On the contrary, the average strain in Zone 2
of Specimen 1, where there is no crack, keeps relatively constant during
the test.
Several examples of strain fields measured by DIC for both Specimens 1
and 2 are shown in Figs. 11 and 12. The deformed image, i.e. the image
taken at the corresponding cycle, was put behind the measured field to
allow its comparison with surface microstructure. In fact, as the
measured fields are related to the reference image instead of the
deformed images, a little difference may exist between the translucent
field and the microstructure behind.
In Fig. 11, the large deformations at the cracks (see red areas in Fig.
11) are caused by opening of cracks. In the areas marked with circles in
(a) and (b), strain localizations (>3% for (a), while
>6% for (b)) occurred while there was still no crack.
Surface crack passed through the area marked in (a) in the following
cycles, and a 2nd crack initiated in the area marked
in (b). Thus, strain localizations were measured in these areas prior to
the observations of cracks on surface. Besides, the strain measured by
DIC in the areas marked with circles in (a) and (b) may also be
overestimated as out-of-plane deformation may have occurred in these
areas where cracks appeared in the following cycles.
One example of strain fields in the whole ROI of Specimen 1 is shown in
Fig. 11c. Besides the Zone 0 (the lower left corner area in the stitched
image), no obvious strain localization was observed until the last step
before failure (after 90 000 cycles). The DIC performed in the other
five zones highlights the role of strain localizations in crack
nucleation from another perspective: no strain localizations no crack
nucleation. Besides, the strain localization (marked by a circle in Fig.
11c) prior to the appearance of crack occurred where the final fracture
took place.
The relations between strain localization and crack were also observed
for Specimen 2 in Fig. 12. The crack tips after 18 410 cycles are marked
by arrows, and the cracks initiated from sites C and D have not
connected yet. The strain localizations measured in the area (marked
with a white circle around Initiation site E) where a secondary crack
initiated around the surface pore after 19 350 cycles when the main
crack was arrested by hard inclusions around crack tip, indicates the
influence of strain localizations on crack initiation. Meanwhile, the
strain localizations measured in the area (marked with a white circle
between crack tips) where two separate cracks connected on surface in
the following cycles are mainly ascribed to the coupling influence of
the two separate crack tips.
4. Discussion
4.1 Crack
initiation
Pores are assumed to be the most probable crack initiation sites in
HCF12,14,17,24,25and in LCF as demonstrated in a
previous study in the laboratory 21,22.They were thus
characterized in the bulk of specimens using Lab-CT to identify possible
crack initiation areas, i.e. ROI. The successful selection of ROIs for
both tested specimens confirms the role of pores for crack initiation.
Even if the pores are below the flat surfaces for surface in-situ
observations, their influence on crack initiation on surface are further
confirmed through the comprehensive analysis using surface in-situ
observations, fracture analysis by SEM and OM and 3D characterizations
of pores by Lab-CT (Figs. 7 and 8).
Thus, pore is undoubtedly the most important microstructural feature for
crack initiation in the studied material under LCF loading. This is
firstly ascribed to the large pores formed in the studied material due
to LFC process 21: the minimum size of pores around
which cracks initiated on surface in the tested specimens is 0.4 mm,
which is much larger than the critical sizes in the
literature12,17.
The location of pore is also demonstrated to be an important factor on
crack initiation: Cracks initiated around a surface pore at the notch
root in Specimen 1 (Fig. 7) instead of other larger pores far from
surface; In Specimen 2 (Fig. 8), more cracks initiated from the near
surface Pore A and the surface Pore C than from the larger subsurface
Pore B (Table 3).
The comparison between the two tested specimens implies that the number
of large pores has an influence on crack initiation and fatigue life. In
Specimen 2, which has more large pores near surface than Specimen 1,
cracks initiated on surface in the tensile stage of
1st cycle while cracks initiated on surface of
Specimen 1 after 7 000 cycles. Specimen 1 has a longer fatigue life than
Specimen 2 while the former has a higher loading. Thus a large number of
surface or near surface large pores accelerates crack initiation and
also reduces fatigue life 41.
A previous study in the laboratory 21,22 shows that
the strain localizations around large pores are responsible for crack
initiation. In addition to large pores, hard inclusions may have also
acted as individual crack initiation site in the studied material, such
as 2nd crack initiation site for Specimen 1 (Fig. 7c)
and Initiation site D for Specimen 2 (Fig. 8b). However, these
initiation sites are not the initial nucleation sites. Cracks nucleated
at these sites with the propagation of main crack on surface. The most
likely reason is that the local strain levels in these sites are
increased by the propagation of main crack and thus result in failure of
hard inclusions. DIC measurements indeed show the strain localizations
at these sites. For example, strain localization is observed at
2nd Crack initiation site before crack nucleation in
Fig. 7b. Thus, crack initiation can be ascribed to strain localizations
induced by: (1) large pore; (2) the propagation of a formerly nucleated
crack; (3) their coupling influence. Fig. 12 demonstrates the third
case: with the coupling influence of the propagated main crack and the
large Pore C, strain localizations are observed at Initiation site E at
18 410 c. before crack nucleation after 19 350c.
4.2 Crack
propagation
Once cracks nucleated on surface, they were observed to propagate along
hard inclusions under cyclic loading. The influence of the direction of
hard inclusions on crack propagation was observed for both specimens.
When the direction of the crack propagation was along the direction of
the longest axis of hard inclusions at the surface, the crack grew
quickly. However, when the direction of crack propagation was
perpendicular to the direction of hard inclusions, the crack may be
arrested 42. For Specimen 1, cracks arrest at hard
inclusions was observed at 16/25/45/65 K cycles marked in Fig. 7c; the
corresponding points were also marked in the crack propagation curve in
Fig. 9a. More crack arrests were observed for Specimen 2 as shown in
Fig. 8d. The crack grew slowly when its tip encountered hard inclusions
having their long axis perpendicular to the crack propagation direction.
The crack tip needed many cycles to pass these barriers by two ways
(Fig. 7c and Fig. 8d): (a) Change its propagation direction, i.e. grow
along these barriers. Once these barriers have been bypassed, the crack
propagated along a new route perpendicular to the loading direction. (b)
Break these barriers. Once the arrested crack has passed these barriers,
it grew very fast.
When cracks were arrested by the barriers, some similar phenomena were
observed in the area (Fig. 9b) around the crack tip and along the
loading direction:
- For Specimen 1: (a) When crack tip was arrested at 45 K cycles in Fig.
7c, Crack F1-a, which is just below the arrested crack tip, initiated
at Si particle between two surface pores, which are actually one pore
that is connected in subsurface as shown in the Lab-CT observation
(Fig. 7a). (b) When crack tip was arrested at 65 K cycles, Crack F1-b,
which is just below the arrested crack tip, initiated at copper
containing phases. In the same time, crack initiated at hard
inclusions (iron intermetallics, Al2Cu phase and
AlCuMgSi phase) at 2nd initiation site, which is
also below the arrested crack tip.
- For Specimen 2: When the main crack was arrested by hard inclusions at
crack tip 1 (Fig. 8d), Crack 2 initiated on surface around Pore C
after 19 350 cycles. Although strain localizations are easily
generated around such large pores, no crack was observed before the
crack was arrested at crack tip 1.
The above four secondary cracks, i.e. Crack F1-a, Crack F1-b,
2nd initiation site for Specimen 1 and Crack 2 for
Specimen 2, have a common feature: when the principal crack was
arrested, they initiated from hard inclusions that are located in the
same vertical direction as the main crack tip (Fig. 9b), i.e. the line
through the principal crack tip and along a direction perpendicular to
the crack propagation.
The failure of hard inclusions needs enough applied stress43. Thus, the above observations imply that stress
concentrations may occur in the same vertical direction as the crack
tip, i.e. the area marked in Fig. 9b, when the crack was arrested.
Actually, such strain localizations are confirmed by DIC measurements,
for example, 2nd Crack initiation site for Specimen 1
in Fig. 11b and Crack 2 for Specimen 2 (Initiation site E) in Fig. 12.
As discussed in 4.1, it should be noted that initiation site E is a
little away from the vertical direction of the arrested crack tip due to
the combined effect of the stress concentrations caused by arrested
crack tip and Pore C.
Conclusions
In order to study the LCF damage micromechanisms in LFC A319 alloy, an
experimental protocol has been set up to establish the relations between
microstructure, crack initiation and propagation and strain fields. The
studied material was characterized thoroughly by using X-ray CT and SEM
and Lab-CT was used to characterize pores in the bulk for the selection
of ROI and to link the relations between surface observations and
subsurface pores. LCF tests were performed with surface in-situ
observations to follow crack initiation and propagation and DIC method
was used to measure the mechanical fields. Post-mortem analysis was
performed using OM and SEM.
- An
etching method, which gives a natural texture to the aluminium
dendrites by revealing segregation of Si, was developed to make DIC
feasible to an acceptable resolution without adding a speckle pattern
that will also mask the microstructure. The relations between cracks
and strain localizations are observed for the validation in the
tensile test and the application in LCF tests.
- Large pores, especially for surface or near surface pores, play the
most critical role for crack initiation, A large number of surface or
near surface large pores accelerates crack initiation and also reduces
fatigue life.
- Crack initiation is ascribed to strain localizations induced by: (a)
large pore; (b) the propagation of an already nucleated crack; (c)
their coupling influence.
- Cracks propagated along hard inclusions while the orientation of hard
inclusions has an influence on crack propagation: (a) When it is the
same as the direction of crack propagation, the crack is prone to
propagate along hard inclusions; (b) When it is perpendicular to the
direction of crack propagation, the crack is arrested, then the
arrested crack spends more cycles to continue to propagate by changing
the direction of propagation to the direction along hard inclusions or
by breaking hard inclusions to pass them.
Acknowledgements
This research work was funded by INDiANA-ANR project (grant
ANR-12-RMNP-0011)
and
PSA
Peugeot Citroën. The authors would like to thank TOMCAT beamline
at
Swiss Light Source (SLS) for providing SR-CT beamtime, and the China
Scholarship Council (CSC) for funding the PhD thesis of Long WANG.
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