FRACTURE AND FATIGUE BEHAVIOR OF CEMENTED CARBIDES

FRACTURE AND FATIGUE BEHAVIOR OF CEMENTED CARBIDES: 3-D FIB
TOMOGRAPHY OF CRACK-MICROSTRUCTURE INTERACTIONS
Jose María Tarragó, Emilio Jiménez-Piqué, Miquel Turón, Lorenzo Rivero and Luis Llanes
CIEFMA, Universitat Politècnica de Catalunya
Barcelona, Spain, 08028
Ihsan Al-Dawery, Ludwig Schneider
Sandvik Hyperion,
Coventry, UK, CV4 0XG
ABSTRACT
The fracture and fatigue phenomena in WC-Co cemented carbides (hardmetals) have been subjects
extensively investigated in the last thirty years. From these studies it is well established that metallic
binder phase plays a key role as the toughening and fatigue susceptible agent in those materials, as its
effective ductility is critical for defining crack growth resistance and cyclic induced degradation.
However, experimental proof of involved toughening and fatigue micromechanisms has usually been
presented on the basis of post-failure fractographic examination. In this work, the fracture and fatigue
behavior of a WC-Co is investigated and a 3D characterization of crack-microstructure interaction during
stable crack growth in hardmetals is carried out in order to gain a better understanding of the failure
processes in cemented carbides under the application of monotonic and cyclic loads. In doing so,
FIB/FESEM, 3D tomography and imaging reconstruction are combined with systematic mechanical and
indentation protocols for assessing crack extension behavior of cemented carbides. Experimental findings
clearly highlight existing differences regarding failure mechanisms operative under monotonic and cyclic
loads, and provide new and interesting insights for understanding them.
INTRODUCTION
Cemented carbides, also called hardmetals, are a group of powder metallurgical liquid phase sintered
composite materials with outstanding wear resistance and hardness consisting of brittle refractory carbides
of the transition metals (e.g. WC, TiC, TaC) embedded in a metallic matrix [1]. The success of this
ceramic-metal composite resides in the combination of the hardness and wear resistance of the ceramic
particles with the toughness of the metallic phase. It makes cemented carbides leader materials in several
engineering and tooling applications, such as metal cutting, mining, rock drilling, metal forming,
structural components and wear parts [2].
The exceptional toughness level exhibited by cemented carbides is related to the highly effective
toughening through constrained deformation of the ductile phase (e.g. Refs. [3–6]). Ductile-phase (cobalt
ligaments) reinforcement of a brittle matrix (tungsten carbide network) is a notable example of toughening
mechanisms which act in the crack wake such to screen the crack tip from the far-field driving force [7].
From a fracture viewpoint, such toughening has proven to be a successful microstructural design strategy
in brittle-like materials because it implies the existence of a crack growth resistance-curve (R-curve)
behavior which imparts damage tolerance; and thus, improved in-service reliability, to the corresponding
structural components [8]. Thus, the first step of crack propagation in cemented carbides is the formation
of a continuous crack in the carbide phase which subsequently advances continuously in the brittle phase
254
and circumvents the ductile regions leaving bridging ligaments across the crack faces [5]. In this way, a
multiligament zone of the order 2-4 ligaments is developed in the direction of crack propagation [4,5]. As
the crack propagates, the binder ligaments elongate and microcavities are formed inside the ligament to
uphold volume constancy. Using finite elements calculations it was demonstrated that binder sites exposed
to a high local plastic strain and stress triaxiality are preferred zones for microvoid nucleation [9]. First
void will be followed by others along the plane linking the adjacent cracks in the carbide phase and finally
the ligament fractures by void growth and coalescence. Fracture along the carbide-binder interface
proceeds in a similar manner to that in the binder, i.e. by the nucleation, growth and coalescence of
microcavities. However, shallow and closely spaced microvoids are evidenced when the crack runs
parallel and close to the binder/carbide interface. Fischmeister et al. [9] attributed the formation of such
finer dimple structure to the fulfillment of high triaxiality conditions in the binder zone close to the
carbide interface.
Cemented carbides, despite the plasticity developed by the binder ligaments, exhibit a brittle-like nature
behavior. This fracture mode, such as in other brittle materials, is governed by unstable propagation of
pre-existing flaws which may be inherent to the processing or induced in the forming process or in service
conditions [10,11]. Thus, strength and reliability of hardmetals is dependent on the size, geometry and
distribution of pre-existing flaws. Following this idea, Linear Elastic Fracture Mechanics (LEFM) is
recognized as the suitable theoretical basis to describe and rationalize the fracture behavior of cemented
carbides [12,13]. On the other hand, a statistical tool is required to rationalize the rupture strength of this
brittle material, as it is dependent on the size and geometry of the flaws. For that, Weibull statistics are
commonly employed to evaluate not only the rupture strength, but also the reliability of the material.
Meanwhile, as it is also the case for other brittle-like composite systems (e.g. ceramic- and intermetallicbased materials) where crack-tip shielding mechanisms prevail [14], the susceptibility of hardmetals to be
mechanically degraded under cyclic loading is well established (e.g. Refs. [15–19]). Schleinkofer and coworkers documented a strong strength degradation of cemented carbides under cyclic loads that
predominantly occur in the ductile binder phase [17,18,20] and established that subcritical crack growth is
the controlling stage for fatigue failure in cemented carbides over crack nucleation [17,20,21]. On the
other hand, taking into account this second assumption and that fracture rupture behavior of cemented
carbides has been extensively rationalized within the LEFM framework, Torres and co-workers proposed
the fatigue crack growth threshold as the effective toughness under cyclic loading [19]. Experimental
validation for such an approach was then presented for a series of WC-Co hardmetal grades [22].
Moreover, such results pointed out a strong microstructural influence on the fatigue sensitivity of
hardmetals, depending on the compromising role played by the metallic binder as both toughening and
fatigue susceptible agent [23]. These investigations also point out an extreme Paris-Erdogan law like
dependence of fatigue crack growth (FCG) kinetics similar to the behavior reported for structural ceramics
and intermetallics. This binder fatigue degradation, also observed in the experimental applied stress –
fatigue life (S-N curves) data published by Sailer et al. [24], is rationalized on the basis of fatigue-induced
accumulation of fcc to hcp phase transformation within the Co binder [17]. However, most of the
experimental support validating the above referred failure scenarios are either indirect, through clear
evidences of different fractographic features associated with fracture and fatigue (e.g. Refs. [4,17,23]) or
rather limited (but direct), by means of transmission electron microscopy of slices (i.e. very local and
small areas) in regions around crack tips (e.g. Refs. [17,21,25]).
Over the past few years, new advanced characterization techniques have been successfully implemented
for microstructure, surface modification and deformation/damage characterization in various areas of
materials research. Among them, Focused Ion Beam (FIB) technique has shown to be extremely versatile
for overcoming many of the experimental limitations/difficulties ascribed to more conventional ones, such
as scanning and/or transmission electron microscopy, atomic force microscopy, and X-ray diffraction,
among others. Fruitful examples of its implementation for quantifying microstructure as well as evaluating
255
tribological and mechanical phenomena in cemented carbides are the works by Cairney et al. [26], Beste
et al. [27], Mingard et al. [28] and Gant et al. [29]. It is the purpose of this study the fracture and fatigue
behavior of a WC-Co cemented carbide and to use FIB technique for bringing new insights into the
toughening and fatigue micromechanisms operative in hardmetals when subjected to monotonic and cyclic
loading conditions, respectively. In doing so, the main focus will be to document crack-microstructure
interaction at regions very close to the tip of cracks arrested after stable growth, and discuss them on the
basis of existing knowledge on fracture and fatigue phenomena for these materials.
EXPERIMENTAL ASPECTS
The studied material is a medium-grained WC-Co cemented carbide supplied by Sandvik Hyperion. The
key microstructural parameters: binder content (%wt), mean grain size (dWC), carbide contiguity (CWC), and
binder mean free path (λbinder) of the studied material are listed in Table I. Mean grain size was measured
following the linear intercept method using a JEOL-7001F field emission scanning electron microscopy
(FE-SEM) micrographs. Carbide contiguity and binder mean free path were deduced from best-fit
equations, attained after compilation and analysis of data published in the literature (e.g. [3,5]), on the
basis of empirical relationships given by Roebuck and Almond [30] but extending them to include carbide
size influence [12,13].
Table I: Microstructural parameters for the studied material
Wt.% binder dWC (µm)
CWC
λbinder (µm)
11% Co
1.12 ± 0.71 0.38 ± 0.11 0.42 ± 0.29
Mechanical characterization includes hardness (HV30), flexural strength (σr), fracture toughness (KIc) and
fatigue crack growth (FCG) parameters. Hardness was measured using a Vickers diamond pyramidal
indenter and applying a load of 294N. In all the others cases, testing was conducted using a four-point
bending fully articulated test jig with inner and outer spans of 20 and 40 mm respectively. Flexural
strength tests were performed on an Instron 8511 servohydraulic machine at room temperature. At least 15
specimens of 45x4x3 mm dimensions were tested per grade. The surface which was later subjected to the
maximum tensile loads was polished to mirror-like finish and the edges were chamfered to reduce their
effect as stress raisers. Results were analyzed using Weibull statistics. Fracture toughness and FCG
parameters were determined using 45x10x5 mm single edge pre-cracked notch beam (SEPNB) specimens
with a notch length-to-specimen width ratio of 0.3. Compressive cyclic loads were induced in the notched
beams to nucleate a sharp crack and details may be found elsewhere [31]. The sides of SEPNB specimens
were polished to follow stable crack growth using a high-resolution confocal microscope. Fracture
toughness was determined by testing SEPNB specimens to failure at stress-intensity factor load rates of
about 2 MPa√m/s. FCG behavior was assessed for two different load ratio (R) values, 0.1 and 0.5, using a
RUMUL resonant testing machine at load frequencies around 150 Hz.
The effective evaluation of crack-microstructure interactions requires the suitability of a procedure for
introducing sharp precracks into specimens free of residual stresses. From this viewpoint, sharp
indentation is suggested as a simple and practical precracking technique, as far as hardmetals are brittle
enough (e.g. Ref. [32]). Tomography characterization to be implemented in this study involves automated
serial sectioning using FIB and imaging of each milled surface with field-emission scanning electron
microscopy (FE-SEM). However, precise and detailed imaging of regions close to the tips of cracks, just
as they develop after unloading the indenter may become a difficult task. Main reason behind it is the
intrinsic indentation residual stress field. Thus, as material is removed through FIB milling, effective
stress state at the crack tip changes and the latter gets displaced. Such experimental inconvenience may be
overcome by relief of indentation residual stresses before FIB. This was approached in this investigation
256
by controlled crack extension under cyclic loads. It permitted to attain much longer cracks without any
indentation residual stresses acting at their tips.
One of the purposes of this investigation is to document crack-microstructure interaction during stable
crack growth under both monotonic (toughness) and cyclic (fatigue) loads. It then requires the arresting of
cracks after stable growth under each of the referred loading conditions. Indeed, the above residual stress
relief procedure could be taken as final step for assessing fatigue micromechanisms. However, to attain
similar cracks under monotonic loading demands one further step: testing to failure of specimens
containing multiple indentations (e.g. Ref. [33]). Strength-indentation tests were conducted on specimens
with 4x3x45 mm dimensions under four-point bending (outer and inner spans of 40 and 20 mm,
respectively). Four controlled flaw patterns were produced 2 mm apart in the center of the prospective
tensile surface of each flexure specimen by applying a load of 490N using a Vickers diamond pyramid.
Care was taken to orient one set of the corresponding Palmqvist cracks of each indentation flaw to be
parallel to the cross section of the specimen where the prospective rupture would occur. Indentation
residual stresses for all the controlled flaws were relieved by subjecting the specimens to tensile cyclic
bending (load ratio of 0.1 and frequency of 10 Hz) in order to induce stable crack growth. Finally, fracture
of the cracked specimens was induced in flexure, under either monotonic or cyclic loads. All the
specimens ruptured at one of the controlled flaws. On the other hand, stably grown and arrested cracks
remained for the other three indentations. FIB serial sectioning was then carried out in regions close to the
tips of these surviving cracks. A schematic of the process is shown in Figure 1.
2 mm
4 mm
20 mm
3 mm
40 mm
Figure 1. Schematic representation of the process followed to generate stable cracks for the subsequent
sectioning by means of the FIB/FE-SEM technique.
In order to document crack-microstructure interaction during stable crack extension, a combined FESEM/FIB (Zeiss Neon 40) was used. Before ion milling, a thin protective platinum layer was deposited on
small areas of interest, corresponding to regions close to tips of arrested cracks, along surface crack paths.
U-shaped trench with one cross-sectional surface perpendicular to crack path and to the specimen surface
were produced by FIB (Figure 2a). Subsequently series of crack-microstructure interactions were
obtained by periodical removing the material within U-shaped crater parallel with the cross-sectional
surface by FIB using automated software (Figure 2b). A 15x15x10 um3 volume was ion milled and about
600 images were obtained.
The 3D reconstruction of the obtained images was carried on using the commercial Avizo software. In the
first place, the images were aligned by subsequently applying image processing techniques in order to
correct, equalize and differentiate the grey level of the three phases (WC, Co and crack). The next step
consisted in segmenting the three phases to finally reconstruct the 3D volume.
257
(a)
5 μm
(b)
Figure 2. FESEM images of the trenches generated by FIB [(a) previously to sequential milling and (b)
once sequential milling has been completed] around the regions of interest.
RESULTS AND DISCUSSION
Hardness, flexural strength and fracture toughness
The hardness, flexural strength, Weibull modulus and fracture toughness of the investigated material are
listed in Table 2. The flexural strength dispersion evidenced is fairly small for; and accordingly, the
corresponding Weibull analysis yields high values, indicative of high reliability from a structural
viewpoint (Figure 3a). Fractographic examination reveals that the critical defects are abnormal grains or
carbide agglomerates (e.g. Figure 3b) with an equivalent diameters (2acr) of about 15-30 μm. It is in
agreement with values estimated from a direct implementation of the basic LEFM equation relating
strength, toughness and critical flaw size (see Table II), by considering defects as surface semicircular
flaws. This supports the use of LEFM for rationalizing the fracture behavior of the cemented carbide
studied here.
Table II: Hardness, strength and fracture mechanics parameters for the WC-Co hardmetal investigated.
HV30 (GPa) Flexural strength (MPa) Weibull modulus KIc (MPa√m) 2acr (μm)
12.8 ± 0.2
3101 ± 102
36
13.9 ± 0.3
31 ± 2
(a)
(b)
Figure 3. (a) Fracture Weibull distribution for the studied materials and (b) carbides agglomerate acting as
a critical flaw for fracture.
258
Fatigue behavior
FCG rates are plotted against the range and the maximum applied stress intensity factor, ∆K (Figure 4a)
and Kmax (Figure 4b) respectively, for the two load ratios studied. As it has been previously reported for
WC-Co hardmetals [19,23], the WC-Co grade under consideration exhibits: (i) a large-power dependence
of FCG rates on ∆K, and (ii) subcritical crack growth at ∆K values much lower than fracture toughness.
Also, a very pronounced load ratio effect is observed in the dependence of FCG on ∆K. However, as
observed for other brittle-like materials, R effects are largely reduced when plotting FCG against Kmax.
This is an indication of predominance of static over cyclic failure modes [16,23].
(a)
(b)
Figure 4. da/dN behavior vs. (a) ∆K, and (b) Kmax, for each load ratio studied.
The FCG – fatigue life relationship proposed and validated by Torres and co-workers for WC-Co
cemented carbides [19] has been extended for the studied cemented carbide. In doing so, a classical
approach on the basis of fatigue limit, within an infinite life framework, and FCG threshold (Kth) is
implemented by defining the latter as the effective toughness under fatigue for a given critical flaw size.
Thus, fatigue limit (σf) is deduced from the stress intensity factor threshold of a small non-propagating
crack emanating from a defect of critical size, 2acr, according to a relationship of type (1):
σf = Y ∗
Kth
√acr
(1)
Hence, the fundamental LEFM correlation among strength, stress intensity factor and defect size applies
also for natural defects too in cemented carbides. This assertion may be done considering that: (i) size of
the critical natural flaws are larger than the microstructural unit; (ii) plasticity is confined to process zone
ahead the crack tip; and (iii) process zone governing fracture (multiligament zone behind the crack tip)
extends over a relatively short distance (about five ligaments) [19,34]. Thus, fatigue limit values can be
estimated from the relation given by the expression (2) under the assumption that flaws controlling
strength have the same size, geometry and distribution under monotonic and cyclic loading.
K
σf = σr ∗ Kth
Ic
(2)
Attempting to validate the estimated fatigue limit, an experimental study was conducted using 15 samples
and following an up-and-down load (stair-case) fatigue test (Figure 5a). Predicted and experimentally
determined fatigue limits are listed in Table III. FCG threshold – fatigue limit correlation is validated by
the excellent agreement attained between them. Furthermore, the fractographic examination conducted on
failed specimens reveals that size, geometry and nature of the initial critical defects are similar under both
259
monotonic and cyclic loading conditions (Figure 5b). However, under the application of cyclic loads the
defects grow until they reach a critical size and the sample fractures in unstably.
Table III: Fatigue crack growth threshold and predicted and experimentally determined fatigue limit
values in terms of maximum applied stress.
FCG threshold (MPa√m) Predicted fatigue limit (MPa) Experimental fatigue limit (MPa)
7.6 ± 0.2
1696 ± 80
1632 ± 130
Initial critical defect
Defect growth under the application of cyclic loads
(a)
(b)
Figure 5. (a) Up-and-down fatigue test used to determine mean fatigue limit for the WC-Ni cemented
carbide studied and (b) example of a critical defect growth under the applications of cyclic loads.
Crack-microstructure interactions
After failure, the fracture surfaces of tested specimens were examined using FE-SEM. Clear differences
are evidenced when comparing fractographic aspects corresponding to stable (Figure 6a) and unstable
crack growth (Figure 6b). While in the former “step-like” fatigue damage features are discerned within
the binder, in the latter the metallic binder exhibit well-defined dimples, suggesting a pure ductile fracture
mechanism. Fracture under cyclic loading in the cobalt binder appears to follow a faceted,
crystallographic fracture mode, as can be appreciated by the sharp angular facets localized within broken
binder regions. With the purpose of documenting and understanding the crack growth processes under the
application of monotonic and cyclic loads, a tomographic reconstruction of the process zone at the crack
tip has been carried on by means of the FIB/FE-SEM technique.
(a)
(b)
Figure 6. Scanning electron micrographs corresponding to (a) stable crack growth under cyclic loads
(R=0.1) and (b) unstable crack growth under monotonic loading for the WC-Co hardmetal investigated.
Fatigue facets (a) and ductile dimples (b) are neatly discerned in the metallic constitutive phase.
260
Stable crack growth under Monotonic Loading
Figure 7 shows crack-microstructure details along the path of an arrested crack, after stable growth under
monotonic loading. From these preliminary findings, many of the conclusions drawn in the reference
papers by Sigl, Exner and Fischmeister [4,5] are validated: (1) continuous crack encloses cobalt regions
(ligaments) which elongate during crack opening by void formation with plastic deformation localized to
the bridges between the voids; (2) ligaments finally fail by void coalescence; (3) mean depth of the plastic
zone is always smaller than the mean free path in the binder; and (4) partition of crack path between
binder and carbide does not correspond to the volume fraction but favors greatly the ductile binder.
Additionally, local blunting effects as crack front ends within the binder (as invoked by McVeigh and Liu
[35], although at higher length-scale) and the role of carbide corners as stress risers are clearly discerned.
As a conclusion, the multiligament zone may be unambiguously established as main foundation for
rationalizing the crack growth resistance (R-curve) behavior and exceptional fracture toughness of
cemented carbides. Within this framework, description and understanding of microstructural effects on
fracture toughness, R-curve characteristics, strength variability and damage tolerance for hardmetals, are
validated too. In Figure 8 the 3D reconstruction of the microstructure containing the crack propagated
under monotonic loading is shown. The light blue corresponds to the WC carbides, the dark blue to the
Cobalt binder and the red to the crack path. An examination of the reconstruction permits a clear
visualization and understanding of the fracture process (i.e. microvoids nucleation and growth). A clear
example of the process is presented in Figure 9.
Figure 7. Stable crack growth under monotonic loading: FESEM micrographs showing crackmicrostructure interactions as imaged on serial sections obtained by means of FESEM/FIB tomography.
Binder ligaments are clearly discerned.
Figure 8. 3D reconstruction of the process zone for stable crack growth under monotonic loading. Light
blue corresponds to the WC skeleton; dark blue to the cobalt phase and the crack is represented in red.
261
Figure 9. Microvoids formation and growth within a Cobalt ligament.
Unstable crack growth under Cyclic loading
Crack-microstructure interaction during stable crack growth under cyclic loading is shown in Figure 10.
Different from the fracture (monotonic loading) scenario, fatigue micromechanisms in cemented carbides
are less defined and understood. Taking this into consideration, images attained through FIB “slice and
view” sheds further light for proposing and/or validating specific failure micromechanisms: (1) subcritical
fatigue crack growth is more predominantly located in the ductile binder phase than under monotonic
loading [9]; (2) fatigue crack extension is discerned to follow crystallographic-like paths (steps); and (3)
crack arrest phenomena within binder ligaments seems to be rather usual, particularly in regions far from
microstructure-related stress risers. Although the crystallographic nature of the referred steps is evident,
their intrinsic origin as related to shear bands, stacking faults and/or twins resulting from fcc to hcp phase
transformation is not clear. The fact that step-like markings in fatigue fracture surfaces have been
observed not only in Co-base hardmetals but also in Ni-base grades [36], points out that fatigue
susceptibility of cemented carbides goes beyond deformation mode changes induced by the referred fcc to
hcp transformation, well-established in plain WC-Co grades [37]. As in the case of monotonic loading, the
3D reconstructed images of the whole volume and the individual phases is presented in Figure 11. It is
interesting to note that under cyclic loading the crack propagates through the ductile metallic phase and
bridging ligaments are not formed. Figure 12 shows the process of fracture of a binder pool under the
application of cyclic loads.
Figure 10. Stable crack growth under cyclic loading: FESEM micrographs showing crack-microstructure
interactions as imaged on serial sections obtained by means of FESEM/FIB tomography.
262
Figure 11. 3D reconstruction of the process zone for stable crack growth under monotonic loading. Light
blue corresponds to the WC skeleton; dark blue to the cobalt phase and the crack is represented in red.
Figure 12. Crack propagation within a binder pool under the application of cyclic loads.
CONCLUSIONS
The fracture and fatigue behavior of a WC-Co cemented carbide has been investigated. A detailed study of
crack-microstructure interaction during stable crack growth under the application of monotonic and cyclic
loads has also been carried on by complementary use of systematic mechanical and indentation testing
protocols and FIB/FE-SEM tomography. From the attained results the following conclusions may be
drawn:
1. The fracture behavior of the studied WC-Co cemented carbide has been successfully rationalized
using LEFM, on the basis of a reliable assessment of fracture toughness as well as of the size and
geometry of critical flaws.
2. As previously reported for cemented carbides, WC-Co studied grade exhibit a large-power
dependence of FCG rates on ∆K, subcritical crack growth at ∆K values lower than the fracture
toughness, and very pronounced load ratio effects in FCG - ∆K curves. However, R effects are
largely reduced when plotting FCG against Kmax. This indicates the predominance of static over
dynamic failure modes.
3. The fatigue behavior (“infinite fatigue life”) can be rationalized by a fatigue mechanics approach.
This statement is valid considering the initiation of subcritical crack growth as the control
parameter of fatigue failure, i.e. Kth as the effective toughness under cyclic loads.
4. The key role played by the metallic binder phase as the toughening and fatigue susceptible agent
in cemented carbides is completely validated. Unequivocal proof of the multiligament zone as
main foundation for understanding toughness and R-curve behavior in cemented carbides is
provided. Indeed, fatigue susceptibility of hardmetals is then associated with operativeness of
such toughening mechanism under cyclic loading. In this case, stable crack growth follows a
rather crystallographic path through the metallic binder, quite distinct from the typical ductile
failure mode found under monotonic loading.
263
AWKNOLEDGEMENTS
This work was financially supported by the Spanish Ministerio de Economía y Competitividad (Grant
MAT2012-34602). Additionally, J.M. Tarragó acknowledges the Ph.D. scholarship received from the
collaborative Industry-University program between Sandvik Hyperion and Universitat Politècnica de
Catalunya.
BIBLIOGRAPHY
[1]
H. E. Exner, “Physical and chemical nature of cemented carbides,” Int. Met. Rev., 1979, vol. 24,
pp. 149-173.
[2]
L. J. Prakash, “Developments in tungsten carbide-cobalt cemented carbides,” Int. Powder Metall.
Directory, 2008, pp. 131-148.
[3]
R. K. Viswanadham, T. S. Sun, E. F. Drake, and J. Peck, “Quantitative fractography of WC-Co
cermets by Auger spectroscopy,” J. Mat. Sci., 1981, vol. 16, pp. 1029-1038.
[4]
L. S. Sigl and H. E. Exner, “Experimental study of the mechanics of fracture in WC-Co alloys,”
Metall. Mat. Trans. A, 1987, vol. 18, pp. 1299-1308.
[5]
L. S. Sigl and H. F. Fischmeister, “On the fracture toughness of cemented carbides,” Acta Metall.
Mater., 1988, vol. 36, no. 4, pp. 887-897.
[6]
P. A. Mataga, “Deformation of crack-bridging ductile reinforcements in toughened brittle
materials,” Acta Mater., 1989, vol. 37, no. 12, pp. 3349-3359.
[7]
R. O. Ritchie, “Mechanisms of fatigue crack propagation in metals , ceramics and composites: role
of crack tip shielding,” Mater. Sci. Eng. A, 1988, vol. 103, pp. 15-28.
[8]
A. G. Evans, “Perspective on the Development of high-toughness ceramics,” J. Am. Ceram. Soc.,
1990, vol. 73, no. 2, pp. 187-206.
[9]
H. F. Fischmeister, S. Schmauder, and L. S. Sigl, “Finite element modeling of crack propagation in
WC- Co hard metals,” Mat. Sci. Eng. A, 1988, vol. 105/106, pp. 305-311.
[10]
H. Suzuki and K. Hayashi, “The strength of WC-Co cemented carbide in relation to structural
defects,” Trans. Japan Ins. Met., vol. 16, pp. 353-360, 1975.
[11]
B. Casas, Y. Torres, and L. Llanes, “Fracture and fatigue behavior of electrical-discharge
machined cemented carbides,” Int. J. Refract. Met. Hard Mater., 2006, vol. 24, no. 1-2, pp. 162167.
[12]
Y. Torres, “Comportamiento a fractura y fatiga de carburos cementados WC-Co,” 2002, Ph.D.
Thesis, Universitat Politècnica de Catalunya, Catalunya, Spain.
[13]
D. Coureaux, “Comportamiento mecánico de carburos cementados WC-Co: Influencia de la
microestructura en la resistencia a la fractura, la sensibilidad a la fatiga y la tolerancia al daño
264
inducido bajo solicitaciones de contacto,” 2012, Ph.D. Thesis, Universitat Politècnica de
Catalunya, Catalunya, Spain.
[14] R. O. Ritchie, “Mechanisms of fatigue-crack propagation in ductile and brittle solids,” Inter. J. of
Fract., 1999, vol. 100, pp. 55-83.
[15]
W. Dawihl, “Die wissenschaftlichen und technischen Grudlagen der Pulvermetallurgic und ihrer
Anwendungsbereiche,” Stahl u. Eisen, 1941, vol. 61, pp. 909-919.
[16]
P. R. Fry and G. G. Garret, “Fatigue crack growth behaviour of tungsten carbide-cobalt
hardmetals,” J. Mater. Sci., 1988, vol. 23, pp. 2325-2338.
[17]
U. Schleinkofer, H. G. Sockel, K. Gorting, and W. Heinrich, “Microstructural processes during
subcritical crack growth in hard metals and cermets under cyclic loads,” Mater. Sci. Eng. A, 1996,
vol. 209, pp. 103-110.
[18]
U. Schleinkofer, H. G. Sockel, K. Gorting, and W. Heinrich, “Fatigue of hard metals and cermets,”
Mater. Sci. Eng. A, 1996, vol. 209, pp. 313-317.
[19]
Y. Torres, M. Anglada, and L. Llanes, “Fatigue mechanics of WC-Co cemented carbides,” Int. J.
Refract. Met. Hard Mater., 2001, vol. 19, pp. 341-348.
[20]
U. Schleinkofer, H. G. Sockel, K. Görting, and W. Heinrich, “Fatigue of hard metals and cermets new results and a better understanding,” Int. J. Refract. Met. Hard Mater., 1997, vol. 15, no. 1-3,
pp. 103-112.
[21]
S. Kursawe, P. Pott, H. G. Sockel, W. Heinrich, and M. Wolf, “On the influence of binder content
and binder composition on the mechanical properties of hardmetals,” Int. J. Refract. Met. Hard
Mater., 2001, vol. 19, pp. 335-340.
[22]
Y. Torres, M. Anglada, and L. Llanes, “Fatigue limit - fatigue crack growth threshold correlation
for hardmetals: influence of microstructure,” Proceedings of the 8th International Fatigue
Congress, compiled by A.F. Blom, EMAS, 2002.
[23]
L. Llanes, Y. Torres, and M. Anglada, “On the fatigue crack growth behavior of WC–Co cemented
carbides: kinetics description, microstructural effects and fatigue sensitivity,” Acta Mater., 2002,
vol. 50, pp. 2381-2393.
[24]
T. Sailer, M. Herr, H.-G. Sockel, R. Schulte, H. Feld, and L. J. Prakash, “Microstructure and
mechanical properties of ultrafine-grained hardmetals,” Int. J. Refract. Met. Hard Mater., 2001,
vol. 19, pp. 553-559.
[25]
G. Erling, S. Kursawe, S. Luyck, and H. G. Sockel, “Stable and unstable fracture surface features
in WC-Co,” J. Mat. Sci. Lett., 2000, vol. 19, pp. 437-438.
[26]
J. M. Cairney, P. R. Munroe, and J. H. Schneibel, “Examination of fracture surfaces using focused
ion beam milling,” Acta Metall., 2000, vol. 42, pp. 473-478.
265
[27]
U. Beste, E. Coronel, and S. Jacobson, “Wear induced material modifications of cemented carbide
rock drill buttons,” Int. J. Refract. Met. Hard Mater., 2006, vol. 24, pp. 168-176.
[28]
K. P. Mingard et al., “3D Imaging of structures in bulk and surface modified WC-Co hardmetals,”
Proceedings EuroPM2012, Basel, Switzerland, 2012, vol.2, pp. 155-160.
[29]
A. J. Gant, M. G. Gee, D. D. Gohil, H. G. Jones, and L. P. Orkney, “Use of FIB/SEM to assess the
tribo-corrosion of WC/Co hardmetals in model single point abrasion experiments,” Tribol. Int.,
2013, vol. 68, pp. 56-66.
[30]
B. Roebuck and E. A. Almond, “Deformation and fracture processes and the physical metallurgy
of WC-Co hardmetals,” Inter. Mater. Reviews, 1988, vol. 33, no. 2, pp. 90-110.
[31]
Y. Torres, D. Casellas, M. Anglada, and L. Llanes, “Fracture toughness evaluation of hardmetals:
influence of testing procedure,” Int. J. Refract. Met. Hard Mater., 2001, vol. 19, pp. 27-34.
[32]
D. K. Shetty, I. G. Wright, P. N. Mincer, and A. H. Clauer, “Indentation fracture of WC-Co
cermets,” J. Mat. Sci., 1985, vol. 20, pp. 1863-1882.
[33]
R. F. Krause, “Rising Fracture Toughness from the Bending Strength of Indented Alumina
Beams,” J. Am. Ceram. Soc., 1988, vol. 71, pp. 338-343.
[34]
Y. Torres, R. Bermejo, L. Llanes, and M. Anglada, “Influence of notch radius and R-curve
behaviour on the fracture toughness evaluation of WC–Co cemented carbides,” Eng. Fract. Mech.,
2008, vol. 75, no. 15, pp. 4422-4430.
[35]
C. McVeigh and W. K. Liu, “Multiresolution modeling of ductile reinforced brittle composites,” J.
Mech. Phys. Solids, 2009, vol. 57, no. 2, pp. 244-267.
[36]
J. M. Tarragó, C. Ferrari, B. Reig, D. Coureaux, L. Schneider, and L. Llanes, “Fatigue behavior of
a WC-Ni Cemented Carbide,” Proceedings 18th Plansee Seminar, compiled by Plansee Group,
Reutte, Austria, 2013.
[37]
V. K. Sarin and T. Johannesson, “On the Deformation of WC-Co Cemented Carbides,” Met. Sci.,
vol. 9, 1975, pp. 472-476.
266