Fracture and fatigue of rock bit cemented carbides: Mechanics and mechanisms of crack growth resistance under monotonic and cyclic loading

Fracture and fatigue of rock bit cemented carbides: Mechanics and mechanisms of crack growth resistance under monotonic and cyclic loading

Int. Journal of Refractory Metals and Hard Materials 45 (2014) 179–188 Contents lists available at ScienceDirect Int. Journal of Refractory Metals a...

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Int. Journal of Refractory Metals and Hard Materials 45 (2014) 179–188

Contents lists available at ScienceDirect

Int. Journal of Refractory Metals and Hard Materials journal homepage: www.elsevier.com/locate/IJRMHM

Fracture and fatigue of rock bit cemented carbides: Mechanics and mechanisms of crack growth resistance under monotonic and cyclic loading Y. Torres a,1, J.M. Tarrago a, D. Coureaux a, E. Tarrés a,b, B. Roebuck c, P. Chan b, M. James b, B. Liang d, M. Tillman b, R.K. Viswanadham e, K.P. Mingard c, A. Mestra a,f, L. Llanes a,f,⁎ a

CIEFMA — Universitat Politècnica de Catalunya, Barcelona 08028, Spain Sandvik Hard Materials, Coventry CV4 0XG, UK c NPL, Teddington, London TW11 0LW, UK d Smith International Inc., Houston, TX 77205, USA e Sandvik Rotary Tools, Shenandoah, TX 77385, USA f CRnE — Universitat Politècnica de Catalunya, Barcelona 08028, Spain b

a r t i c l e

i n f o

Article history: Received 18 February 2014 Accepted 14 April 2014 Available online 24 April 2014 Keywords: Fatigue Fracture mechanics Crack growth resistance Rock bit hardmetals

a b s t r a c t In an attempt to improve the material selection, design and reliability of rock bit WC–Co cemented carbides (hardmetals), an extensive and detailed study is conducted with the main goal of characterizing the fracture and fatigue crack growth (FCG) behavior of four hardmetal grades. Work includes basic microstructural and mechanical characterization of the materials, assessment of fracture toughness and FCG kinetics. It is found that rock bit cemented carbides exhibit relatively high fracture toughness values (between 17 and 20 MPa√m) in direct association with their specific microstructural characteristics, i.e. medium/coarse carbide grain size and medium cobalt content. The influence of microstructure on the measured crack growth mechanics under monotonic loading may be accounted by considering the effective operation of ductile ligament bridging and crack deflection as the prominent toughening mechanisms. Regarding FCG behavior, it is observed to exhibit a significant Kmax influence. Furthermore, relative increments in toughness are maintained, in terms of crack growth threshold, under cyclic loading. As a consequence, fatigue sensitivity for rock bit cemented carbides is found to be lower than that extrapolated from data reported for fine-grained grades. Crack growth resistance under cyclic loading for the hardmetals studied may be understood on the basis that prevalent toughening mechanisms (ductile ligament bridging and crack deflection) show distinct susceptibility to fatigue degradation and are thus critical in determining fatigue sensitivity. © 2014 Elsevier Ltd. All rights reserved.

Introduction The unique combination of hardness, toughness and wear resistance exhibited by WC–Co cemented carbides (hardmetals) has made them a preeminent material choice for extremely demanding applications, such as metal cutting/forming tools or mining bits, where improved and consistent performance together with high reliability are required (e.g. Ref. [1]). The remarkable mechanical properties of hardmetals result from a two-fold effectiveness associated with its composite character. On the one hand in terms of composite nature: combination of two completely different phases (hard, brittle carbides and a soft, ductile metallic binder) with optimal interface properties, as given by ⁎ Corresponding author at: CIEFMA — Universitat Politècnica de Catalunya, Barcelona 08028, Spain. Tel.: +34 934011083; fax: +34 934016706. E-mail address: [email protected] (L. Llanes). 1 Current address: Universidad de Sevilla, Sevilla 41092, Spain.

http://dx.doi.org/10.1016/j.ijrmhm.2014.04.010 0263-4368/© 2014 Elsevier Ltd. All rights reserved.

a very low interfacial energy, nearly perfect wetting and very good adhesion in the solid state for the WC and cobalt couple [2]. On the other hand as related to composite assemblage: two interpenetrating-phase networks where toughening through constrained deformation of the ductile phase is highly effective (e.g. Refs. [3–7]). 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 [8]. 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 (Rcurve) behavior which imparts damage tolerance, and thus, improved in-service reliability to the corresponding structural components [9]. Meanwhile, as it is also the case for other brittle-like composite systems (e.g. ceramic- and intermetallic-based materials) where crack-tip shielding mechanisms prevail [10], the susceptibility of hardmetals to be mechanically degraded under cyclic loading is

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well established (e.g. Refs. [11–18]). In this regard, existing data also points out that, as it is also evidenced for other structural materials (e.g. Refs. [19,20]), improvements through microstructure control on fracture (strength and/or toughness) do not directly translate into similar beneficial effects on the fatigue response of WC–Co cemented carbides. This is clearly discerned from the systematic studies on these materials by Sockel's and Llanes's groups where fatigue sensitivity, in terms of fatigue strength (slope of Wöhler plots) [17] and fatigue crack growth resistance (threshold/toughness ratio) [18] respectively, is found to rise as the metallic binder content increases. In both cases the observed behavior has been rationalized on the basis of the intrinsic susceptibility of the ductile phase to premature fatigue failure; the latter being more pronounced as the content and effective ductility (less constraining) of the metallic binder are also higher. The above issues are particularly relevant in applications where service conditions may place a critical demand on crack growth resistance (as compared to hardness), under both monotonic (fracture toughness) and cyclic (fatigue behavior) loading, as it is the case of rock bits used in mining and oil-field operations. For these heavy duty applications, the use of conventional coarse-grained grades (with low to medium binder contents) is well-established [21–23]. Unfortunately, literature data combining fracture and fatigue characteristics for WC–Co cemented carbides are mainly concentrated on relatively finer-grained grades (e.g. Refs. [17,18]), with only a few studies dealing with microstructures containing carbides of sizes above 2 μm [13,24–26]. In this regard, the investigation conducted by Fry and Garret on a large number of hardmetal grades covering wide ranges of binder content and carbide grain size should be highlighted [13,24]. Using a constant-stress intensity factor double torsion test specimen geometry, they reported an enhancement on fatigue crack growth resistance with increasing binder mean free path. In general, the discerned changes correlated relatively well with corresponding variations in fracture toughness, although microstructural influence became less relevant as stress intensity range decreases towards the threshold regime. Such findings are in satisfactory agreement with those presented by Knee and Plumbridge [25] for five microstructurally different cemented carbides. Indeed, these authors even pointed out the existence of a threshold value (below which fatigue cracks were non-propagating) independent of microstructure or toughness. Nevertheless, a detailed study on either microstructural coarsening effects on fatigue sensitivity issues, in terms of threshold/toughness ratio for instance, or the mechanics describing the local mechanisms operative during stable propagating cracks and their correlation with key microstructure features were not addressed. The objective of this article is to document and discuss the fracture and fatigue crack growth (FCG) resistance of four different rock bit (two mining and two oil-field) hardmetal grades, i.e. WC–Co cemented carbides with microstructures coarser than those previously studied by the authors in Ref. [18]. Particular attention is focused on identifying pertinent micromechanisms associated with crack growth resistance under both monotonic (toughening) and cyclic (fatigue degradation) loading. Experimental fracture toughness and FCG data are then utilized to propose and validate the mechanics behind crack growth resistance for medium- and coarse-grained hardmetals. From a technological perspective, the information gathered and analyzed in this investigation is clearly useful for improving material selection, design and reliability of rock bit cemented carbides. Microstructural characteristics of the hardmetal grades studied Sets of samples corresponding to four experimental hardmetal grades were investigated. Nomenclature, composition and key microstructural parameters for each grade are listed in Table 1. The microstructural description includes two-phase microstructural features,

Table 1 Microstructural characteristics, hardness and fracture toughness for the materials investigated. Grade

%wtCo

dWC (μm)

CWC

λCo (μm)

HV30 (GPa)

A B C D

9.5 10.0 14.0 12.0

2.21 1.96 1.92 4.11

0.32 0.33 0.27 0.19

0.64 0.61 0.80 1.29

11.2 11.3 10.6 9.6

± ± ± ±

KIc (MPam1/2) 0.6 0.6 0.7 0.6

17.1 17.9 17.7 20.4

± ± ± ±

0.8 0.8 0.6 0.5

i.e. carbide contiguity (CWC) and binder mean free path thickness (λCo), because they permit an effective rationalization of the combined influence of independently varied single-phase ones, i.e. carbide grain size (dWC) and cobalt content (%wtCo). Carbide grain size was measured using the linear intercept technique. In doing so, scanning electron microscopy (SEM) images were used and mean values were determined from over 500 intercepts. On the other hand, binder content values are given as supplied by the manufacturer (Sandvik Hard Materials, Coventry, UK). Values for the two-phase microstructural parameters (λCo and CWC) were estimated from best-fit equations, attained after compilation and analysis of data published in a large number of literature studies, on the basis of empirical relationships given by Roebuck and Almond [27] but extending them to include carbide size influence [28,29]. Hardness for the materials investigated is also included in Table 1. It was measured using a 30 kgf (294 N) Vickers diamond pyramidal indentation. Ten indentations were made per hardmetal grade. Combining the data presented in Table 1, a slight decreasing trend of HV30 with rising λCo may be inferred, particularly regarding D grade as compared to the other three materials. However, relative differences are small in concordance with the quite narrow range for binder mean free path measured. Microstructural coarsening effects on the fracture toughness of cemented carbides Fracture toughness assessment Fracture toughness was determined using single edge notched bend (SENB) specimens of dimensions 10 × 5 × 45 mm, with a notch lengthto-specimen width ratio, a/W, of 0.4. Notches were induced in the material by electrical discharge machining, followed by notch tip sharpening using a razor blade impregnated with diamond paste. The resulting notch tip curvature was between 10 and 20 μm. A precrack was introduced through application of cyclic compressive loads under reverse cyclic bending [30], and the loading conditions used during the precracking procedure included: sine waveform, load ratio (R) of 10, testing frequency of 15 Hz and maximum compressive stress in the range of 600–1000 MPa. Stable crack growth was observed during cyclic compression, with cracks arrested after extension between 80 and 100 μm from the notch tip. Fig. 1 shows a SEM micrograph of a precrack induced from a notch tip (and further propagated under tensile fatigue, as will be explained later). It is observed that generated cracks are extremely fine and sharp. In general, crack paths are transgranular either through the binder phase, although close to the carbide–binder interface, or through the carbide grains. However, the relatively coarse aspect of the microstructure points out the prevalence of a crack–microstructure interaction relatively uncommon in finergrained hardmetals, i.e. crack deflection (e.g. Figs. 2 and 3). Such phenomenon may be described as pronounced, in terms of both effective crack path perturbation and the frequency with which they occur. Hence, it should be considered as another relevant toughening mechanism for these hardmetal grades, besides the extrinsic one associated with plastic stretching of ductile ligaments, and both of them increase crack growth resistance of the composite with respect to the intrinsic one for the ceramic (WC) phase. Within the context of this study, the operation of this additional toughening mechanism is not only

Y. Torres et al. / Int. Journal of Refractory Metals and Hard Materials 45 (2014) 179–188

10 μm

20 μm Fig. 1. D grade — SEM micrograph showing a sharp crack induced from the notch from reverse cyclic bending and its corresponding interaction with microstructure.

important from the viewpoint of the enhanced crack growth resistance that it may provide (through the reduction of the effective applied stress intensity factor) but also from the one of their sensitivity to degradation or inhibition under cyclic loading, a subject to be treated in the Microstructural coarsening effects on the FCG behavior of cemented carbides section of this paper. Based on previous experience of precracking-induced residual stress effects on the toughness evaluation of hardmetals [31], precracks were further propagated in most of the tested samples before fracture toughness measurement. Such subcritical crack growth was conducted under far-field cyclic tensile loads (R value of 0.1) in four-point bending and a test frequency ranging from 0.5 to 10 Hz. Crack extension behavior under constant applied load ranges was monitored in situ using a high resolution (±5 μm) telescope. Finally, fracture toughness was determined by testing the precracked SENB samples to failure under constant loading rate values, between 200 and 400 N/s. Stress intensity factors given in the literature [32] were used in the fracture mechanics evaluation. Fracture toughness values for each hardmetal grade, determined from at least three tests per set, are given in Table 1. Relevant differences are only observed for the coarsest hardmetal grade D, this material exhibiting the highest fracture toughness level. Once the notched and precracked samples were tested to failure, an extensive SEM examination was conducted on the fracture surfaces. As expected, dimple ductile rupture in the metallic binder interdispersed with trans- and inter-granular fracture of the carbides were the relevant fractographic features discerned in all the cases (e.g. Fig. 4).

5 μm Fig. 2. D grade — optical micrograph indicating the magnitude and frequency of crack deflection phenomena during stable crack propagation.

181

Fig. 3. D grade — SEM micrograph detailing crack deflection mechanisms, both within a coarse carbide and through the subsequent binder ligament, during stable crack propagation.

Fracture toughness–microstructure correlation Microstructural effects on the fracture toughness of cemented carbides may be rationalized through an analysis involving both modeling of the toughening behavior, directly related to operative toughening mechanisms, and key two-phase microstructural parameters such as λCo (or CWC). Based on literature reports (e.g. Refs. [3–7]), the most relevant toughening mechanism to WC–Co hardmetals is shielding due to ductile Co ligament bridging behind the crack tip. It should also be applicable for the cemented carbides here studied as indicated by the sharply defined ductile dimpled fracture of the metallic binder (e.g. Fig. 4). In this case, toughness enhancement is mostly given by the energy expanded in the constrained plastic stretching of the binder ligaments and increases with crack extension up to a maximum steady-state level (KIc), which corresponds to the bridging length when the ligament zone is fully developed. Considering that the bridging zone size for hardmetals is small (about 5dWC, i.e. ~ 10–20 μm) compared to the crack length and specimen dimensions used here (several millimeters), an effective saturation toughness may be estimated from [33,34]: vffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi  ffi u λ  u D Eσ y A f Co u 2 t   K Ic ¼ K t þ 1−ν 2

ð1Þ

2 μm Fig. 4. C grade — SEM micrograph showing fractographic aspects associated with unstable failure in SENB specimens tested to rupture.

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where Kt is the critical crack-tip stress intensity factor required for crack initiation; E is the Young's modulus and ν is the Poisson ratio of the composite; σy is the binder flow strength determined from a Hall–Petch type relation (σy = σo + k y[λCo] − 1/2 [3,6], with σo ~ 480 MPa [35] and ky = 1.55 MPa√m [6]); Af is the area fraction of the metal binder intercepted by the crack plane; and D* is a “work of rupture” function that depends on the constraint (and thus on the strength of the binder–carbide interface) and constitutive properties of the binder. The validation of this analysis may be evaluated, for fracture toughness, microstructure and fractographic data reported in the literature for different sets of cemented carbides [3,5, 6,18,28] together with the one here determined, by comparing estimated and experimental values. Hence, it is found that a satisfactory agreement between predicted and experimental fracture toughness values is achieved by using linear-regression best fitting values of 6.9 MPa√m and 3.2 for Kt and D*, respectively (Fig. 5a, solid line). In doing so, values of E Co ~ 207 GPa, E WC ~ 703 GPa and νWC–Co ~ 0.21 were used. Additionally, where experimental data on Af was not available, it was estimated from the empirical relationships proposed by Sigl and Exner for describing the area fractions of fracture

a

paths corresponding to fracture through the binder phase in WC–Co cemented carbides [5]. A Kt value of 6.9 MPa√m is slightly higher than that experimentally determined by using Hertzian indentation for the intrinsic WC-matrix toughness (~ 3–6 MPa√m) [36] and close to the one extrapolated to zero volume fraction of cobalt of the fracture toughness data of several WC–Co alloys [3]. Within the analytical approach followed, this should be expected because at least another toughening mechanism (besides the one based on reinforcement by ductile ligaments) such as crack deflection should be incorporated into Kt, and the resulting effect added to the baseline value given by the intrinsic WC-matrix toughness. On the other hand, the D* value of 3.2 assessed for a high strength metal dispersion system such as WC–Co may be considered as intermediate (range of work of rupture values reported for different ductile-phase toughened systems is between 0.75 and 8 [37–39]) and should be related to the compromising effect on ductile reinforcement toughening of the strong WC–Co interface together with the high degree of work hardening in the cobalt binder [39]. Based on the physical meaning described above for Kt and D*, a further analysis of the experimental data was attempted as a function of microstructure. This was done because, as shown in Fig. 5a, it is discerned that a data fitting seems to be improved by using a potential-like function (dotted line). Hence, Fig. 5b shows the same data as in Fig. 5a, but separated in subgroups of hardmetal grades defined by three different carbide grain sizes: fine, medium and coarse. Under these conditions, and according to the same analysis procedure followed in Fig. 5a, tangency at the polynomial fitting would define specific Kt and D* values for each of the microstructural subgroups. The corresponding best fitting values are given in Table 2. From the results attained, it is clear that as the microstructure gets coarser, the relative values of Kt and D* increase and decrease, respectively. Although additional research will be needed for completely sustaining these correlations, these trends are highlighted here as an interesting finding, and they are speculated to arise from three different sources: 1) the prominence of crack deflection as another operative toughening mechanism intrinsic to medium- and coarse-grained carbides, 2) the higher transgranular carbide area fraction within crack paths, and 3) the expected higher effective ductility of the binder in the materials here studied.

Toughening mechanisms and crack growth resistance (R-curve) behavior

b

Subgroup

dwc (µm)

Kt (MPa*m1/2)

D*

Fine

<1

6.5

3.5

Medium

1-2

8.3

2.9

Coarse

>2

9.6

2.6

Fig. 5. Plot of experimentally measured critical strain energy release rate (in terms of KIc2) against an energy-related parameter given by the product of volume of deformed binder (per unit area of fracture surface) times the plastic work done (per unit volume of deformed binder): a) for the whole set of experimental data from current paper plus other published references (linear-regression fit — solid line; potential-like fit — dotted line), and b) for the same experimental dataset, but represented as a function of carbide grain size subgroups: fine, medium and coarse (polynomial fit — solid line; tangency condition for each subgroup — dotted lines).

There exists extensive information about the relationship between microstructure and resistance to the propagating crack for high-toughness ceramics and composites. From this perspective, cemented carbides have usually been considered as an academic example of toughening of ceramics (carbides) by reinforcements (ductile metallic ligaments). Under the consideration of potential failure origins as small processing flaws without any initial ductile ligament zone (e.g. pores, large carbides or binderless carbide agglomerates), it is clear that any stable crack growth (i.e. before failure) would imply variable resistance-curve (R-curve) characteristics. Indeed, it is this R-curve behavior which dictates the effective strength

Table 2 Best fitting Kt and D* parameters for the data plotted in Fig. 5a and b. Subgroup

dWC range (μm)

Kt (MPam1/2)

D*

Whole dataset (Fig. 5a) Fine-grained (Fig. 5b) Medium-grained (Fig. 5b) Coarse-grained (Fig. 5b)

0.4–5.5 b1.0 1.0–2.0 N2.0

6.9 6.5 8.3 9.6

3.2 3.5 2.9 2.6

Y. Torres et al. / Int. Journal of Refractory Metals and Hard Materials 45 (2014) 179–188

K deflection

  2 3 2 θ þ S D cos 6 7 2 7K appl ¼6 4 5 DþS

ð2Þ

where Kdeflection and Kappl are the effective and the nominal far-field stress intensity factors respectively, D is the length of the deflected span, S the distance over which mode I crack growth occurs, and θ is the tilt angle. Using as inputs λCo ~ S; dWC ~ D; and θ ~ 50–60°, the effective crack-tip stress intensity values resulting from Eq. (2) result to be between 80 and 90% of the applied nominal far-field stress intensity factor, i.e. the relative toughness contribution from crack deflection in the hardmetals investigated may be estimated as ranging from 2 to 4 MPa√m, in quite satisfactory agreement with the values assessed from the R-curve analysis above described.

S

D

for real hardmetal components, this topic being quite relevant from a practical viewpoint. Concerning R-curves, they are experimentally complicated to determine and time consuming. This is particularly true when dealing with brittle-like materials. However, a qualitative description of such a behavior may be attained from general ideas relating toughening mechanisms, toughness of “large” cracks (i.e. once crack-tip shielding is fully developed) and the intrinsic crack growth resistance of the matrix material. Following the analysis presented above, for the rock bit hardmetals studied such fundamental considerations may be sorted in a schematic R-curve (Fig. 6) on the basis of: 1) toughness enhancement (K R ) is mostly given by the energy expended in the constrained plastic stretching of the cobalt binder ligaments, over a region behind the crack tip of about 5dWC in length [5,6]; 2) toughness increases with crack extension up to a maximum steady-state level, which corresponds to the bridging length where the ligament zone is totally developed, i.e. K Ic for large cracks (as given in Table 1); and 3) the critical crack-tip stress intensity factor required for crack initiation (Kt ) includes not only the intrinsic toughness of the WC matrix (~ 6–7 MPa√m, as found from the best-fitting value for fine-grained carbides in Fig. 5b) but also additional toughening mechanisms, particularly crack deflection. It results in an effective Kt value of ~ 8–10 MPa√m (Table 2) that includes a toughening contribution from this crack–microstructure interaction of about 2–3 MPa√m. This statement may be supported by estimating crack deflection contribution to toughening associated with the microstructure length scale of the materials under consideration. This is done by recalling a simple micromechanical model (see idealized schemes shown in Fig. 7, for both crack–microstructure interaction observed and a crack with periodic tilts) proposed by Suresh [40], from which crack deflection effects on the nominal mode I stress intensity factor may be estimated according to:

183

θ

θ

S Fig. 7. Schematic view of crack deflection phenomena in the studied hardmetals and idealization of a small segment of a crack with periodic tilts.

Microstructural coarsening effects on the FCG behavior of cemented carbides FCG resistance assessment: kinetics description and controlling parameters FCG experiments were conducted using SENB specimens similar to the ones used for fracture toughness evaluation (see Fracture toughness assessment section). As indicated there, sharp cracks were generated by means of cyclic compression (under reverse cyclic bending) and further propagated under tensile fatigue, in order to get rid of any residual stress effects. FCG behavior was subsequently determined following a direct-measurement method using a high resolution telescope. In doing so, the sides of the fatigue specimens were previously polished to mirror-like finish in order to facilitate crack size measurement. Tests were carried out with a fourpoint bending using a fully articulating test jig with inner and outer spans of 20 and 40 mm respectively. They were conducted under load control with a sinusoidal waveform, at working frequencies ranging from 0.5 to 5.0 Hz, in a servohydraulic testing machine. FCG at two different load ratio (R) values (0.1 and 0.5) were investigated. Crack growth thresholds, defined at crack growth rates of 10−9 m/cycle, were attained following an incremental loading sequence corresponding to 0.1 MPa√m steps. At least two samples were evaluated for each material and testing variant. All tests were carried at room temperature and with humidity values of about 55%. The dependence of FCG rate on ΔK under R value of 0.1 is shown in Fig. 8 for all the hardmetal grades studied. As it is well-established for cemented carbides [12,13,16,18,24–26] there exists (1) stable crack extension already at ΔK values significantly lower than that at which Kmax begins to approach KIc, and (2) a large power-law dependence of da/dN on ΔK (n coefficients – slopes in Fig. 8 – within a relationship of type da/dN = C (ΔK) n, ranging from 11 to 15). As for the case of hardness and fracture toughness, only cemented carbide D shows a distinct (and improved) different FCG behavior from the other grades.

10

-6

KR,max = KIc

K (MPam1/2)

Ductile reinforcement toughening λ Co

α dWC

~ 8 -10

Crack deflection toughening

~6-7

da/dN (m/cycle)

17 - 20

A B C D

10-7

10

-8

Intrinsic toughness of WC 10

-9

5

Δ c (μ m) Fig. 6. Schematic R-curve behavior for coarse-grained cemented carbides.

10

15

ΔK (MPam ) 1/2

Fig. 8. FCG rates as a function of ΔK under a load ratio of 0.1.

18

184

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Attempting to evaluate load ratio effects (and consequently the controlling fatigue mechanics parameters) FCG experimental data are shown as a function of ΔK, separately for each hardmetal grade, in Fig. 9. The pronounced load ratio effects should be attributed, as is the case of other brittle-like materials such as structural ceramics or intermetallics [10,41–43], to a marked dependence of FCG kinetics on Kmax. The fact that it also applies to the cemented carbides studied here is confirmed by Fig. 10, where load ratio effects are barely visible when the FCG data is plotted against Kmax. However, it is interesting to highlight that the FCG experimental data gathering attained in Fig. 10 for all the grades investigated is more pronounced than expected from the literature results for lower-toughness hardmetals [18]. Such a finding, which is quantitatively accounted below, is a first indication of a possible departure of the fatigue behavior of these higher-toughness materials from fatigue sensitivity–microstructure trends previously postulated [18]. The relative dominance of Kmax and ΔK as the controlling fatigue mechanics parameters for the cemented carbides studied may be assessed by considering that K max ¼

ΔK 1−R

ð3Þ

and explicitly including both parameters in a modified Paris–Erdogan relationship of the form da m n ¼ CK max ΔK dN

10

ð4Þ

where C, m and n are constants. Hence, factoring out a constant (1 − R)n from each data set in Fig. 10, best fit n values may be determined for which the experimental data are collapsed onto nearly single curves, distinct for each hardmetal grade, as shown in Fig. 11. The obtained optimal n values are listed in Table 3 together with the corresponding C and m values, the latter readily calculated from the slope (= n + m) resulting from least-square regression of the data in Fig. 11. Consistent with the experimental results, the regression analysis indicates that m is much larger than n, i.e. Kmax influence on FCG is predominant over that of ΔK, the difference in the relative dominance of each parameter becoming less pronounced as fracture toughness rises. From this viewpoint, it should be noted that although such an inverse relationship between the ratio m/n and fracture toughness (the latter rather described in terms of effective ductility of the constrained metallic binder as directly proportional to binder mean free path, λCo) is expected [18], it is found to be displaced up and to the right with respect to the trend estimated from data attained on finer-grained hardmetals (Fig. 12). Considering that higher m/n ratios for a given λCo value may be physically interpreted as the material exhibiting a more ceramic-like fatigue behavior, the described trend for the cemented carbides here studied is speculated to come from a more pronounced interaction of the carbide phase with the propagating cracks (i.e. higher area fraction of transgranular fracture through coarse carbides) and/or the effective operation of toughening mechanisms that are not susceptible to degradation under cyclic loading. Such ideas will be further discussed below, once the threshold-related data and crack–microstructure interactions are presented.

-6

10

-6

b)

a)

10

10

-7

10

da/dN (m/cycle)

da/dN (m/cycle)

10

-8

10

10 3

5

10

15

-8

R = 0.1 R = 0.5

R = 0.1 R = 0.5

-9

-7

-9

3

20

5

1/2

10

-6

10

20

15

20

-6

d) da/dN (m/cycle)

da/dN (m/cycle)

10

15

Δ K (MPam )

c) 10

10 1/2

Δ K (MPam )

-7

-8

10

-7

R = 0.1 R = 0.5

10

-8

R = 0.1 R = 0.5 10

-9

10 3

5

10

15

20

-9

3

5

1/2

Δ K (MPam ) Fig. 9. Effect of R on the da/dN vs ΔΚ behavior: a) A; b) B; c) C; and d) D.

10 1/2

Δ K (MPam )

Y. Torres et al. / Int. Journal of Refractory Metals and Hard Materials 45 (2014) 179–188 -6

-6

10

10

b)

a) -7

10

da/dN (m/cycle)

da/dN (m/cycle)

-7

-8

10

10

-8

10

R = 0.1 R = 0.5

R = 0.1 R = 0.5

-9

-9

10

10 3

5

10

15

3

20

5

1/2

10

15

20

15

20

1/2

Kmax (MPam )

Kmax (MPam ) -6

-6

10

10

c)

d)

-7

da/dN (m/cycle)

da/dN (m/cycle)

185

10

-8

10

-7

10

-8

10

R = 0.1 R = 0.5

R = 0.1 R = 0.5 -9

-9

10

10 3

5

10

15

20

1/2

3

5

10 1/2

Kmax (MPam )

Kmax (MPam )

Fig. 10. Effect of R on the da/dN vs Kmax behavior: a) A; b) B; c) C; and d) D.

FCG threshold–toughness ratio and crack–microstructure interactions Besides the FCG kinetics analysis developed above, further relevant information on FCG behavior may be extracted from the experimental data shown in Figs. 8 to 11. This may be done by considering the measured FCG threshold (Kth) and its dependence on microstructure, as compared to the one observed when assessing fracture toughness.

da/dN/(1-R) n (m/cycle)

10

10

10

10

-6

A C B D

-7

-8

-9

7

10

15

18

K max (MPam1/2 ) Fig. 11. R-normalized FCG rate as a function of Kmax for the hardmetal grades studied.

FCG threshold values for each hardmetal grade and load ratio studied are listed in Table 3. For comparison purposes, KIc values and the corresponding Kth/KIc ratios, a parameter used here for describing fatigue sensitivity, are also included in Table 3. In general, it may be discerned that Kth tends to increase as KIc rises. This experimental finding is different from the one previously observed in lower-toughness hardmetals with different microstructures [18], but similar to that evidenced by Fry and Garret, particularly on the coarse-grained hardmetal side [13, 24]. This very interesting result is clearly shown in Fig. 13, where the fatigue sensitivity data for lower- and higher-toughness hardmetals are plotted as a function of microstructure. As may be seen, fatigue sensitivity for the rock bit cemented carbides here investigated (as given by a higher value of Kth /K Ic ratio) is not only lower than expected from extrapolation of the data determined for finer-grained hardmetals, but also seems to decrease – or at least to remain constant – with effective ductility of the binder. This is in direct agreement with the m/n–microstructure correlations discussed above and are again speculated to be a direct consequence of the distinct specific degradation or inhibition under cyclic loading of the prevalent toughening mechanisms observed for coarser-grained hardmetals, i.e. ductile ligament bridging on the one hand, and crack deflection and transgranular WC fracture on the other one. Regarding crack–microstructure interactions, the experimental observation by both optical and scanning electron microscopies are not much different from the ones already described for stably growing cracks under monotonic loading in Microstructural coarsening effects on the fracture toughness of cemented carbides section. As is clearly evidenced in Figs. 14 and 15, crack growth proceeds through a complex

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Table 3 Experimental FCG data for each hardmetal and testing variant investigated. Grade

A B C D

%wtCo

9.5 10.0 14.0 12.0

dWC (μm)

λCo (μm)

KIc (MPam1/2)

C

2.21 1.96 1.92 4.11

0.64 0.61 0.80 1.29

17.1 17.9 17.7 20.4

4 7 2 9

m

× × × ×

10−18 10−20 10−19 10−26

scenario: crack deflection (attempting to run through metallic binder paths) and/or transgranular carbide cracking, particularly for large carbides. As a result, crack paths are rather tortuous (non planar) mainly following metallic binder, but exhibiting pronounced and frequent deflections. Finally, the fracture surfaces of all the samples tested were inspected by SEM. As it has been observed before for other cemented carbides [18,44–46], a transition within binder-phase regions from brittle mode – cleavage-like with little plastic deformation and characterized by surface markings of crystallographic nature – to sharper resolved ductile dimple fracture develops as loading conditions go from cyclic to monotonic with increasing R (for instance, Fig. 16). On the other hand, no differences are seen, as a function of monotonic or cyclic loading, regarding the failure micromechanisms of the carbide phase.

8 10 9 13

1 1 2 3

R = 0.1

R = 0.5

Kth (MPam1/2)

Kth/KIc (%)

Kth,max (MPam1/2)

Kth/KIc (%)

8.4 8.2 8.2 10.7

49 46 46 53

8.5 8.4 9.0 11.6

50 47 51 57

mechanism discerned for hardmetals is rather extrinsic: ductile ligament reinforcement. On the other hand, as the microstructure gets coarser an additional toughening mechanism has also been identified in this investigation, i.e. crack deflection. Although in the literature

Experimental data (rock-bit hardmetals)

Fatigue sensitivity as the result of toughening effectiveness under cyclic loading Critical and subcritical extension of a crack can be considered to be a result of the mutual competition of two classes of mechanisms: 1) ones associated with intrinsic microstructural damage that promote crack growth ahead of the crack tip; and 2) others that shield the crack tip from the far-field driving forces and are referred to as extrinsic [8]. Intrinsic mechanisms are an inherent property of the material (e.g. processes that create microcracks or voids leading to classical failure micromechanisms), whereas extrinsic mechanisms are usually a main factor for the development of resistance-curve (R-curve) behavior by acting within the crack wake (e.g. transformation, brittle fiber/whisker or ductile particle toughening). For materials with little or no ductility in tension, as it is the case for cemented carbides, the role of extrinsic mechanisms is far more important in their toughening. Thus, it is not surprising that the main toughening

n

Experimental data (fine-grained)

Extrapolated trend (coarse-grained)

+ Effective ductility of constrained binder Fig. 13. Fatigue sensitivity–microstructure relationship for WC–Co cemented carbides.

before

Ceramic-like

5 μm

Experimental data (rock-bit hardmetals) Experimental data (fine-grained)

after Extrapolated trend (coarse-grained)

Metallic-like

Effective ductility of constrained binder Fig. 12. Influence of microstructure on the m/n ratio for WC–Co cemented carbides.

Fig. 14. D grade — optical micrograph showing typical crack–microstructure interaction during stable FCG. The previous microstructural aspect of the subsequently cracked region is shown for comparison purposes.

Y. Torres et al. / Int. Journal of Refractory Metals and Hard Materials 45 (2014) 179–188

187

5 μm 10 μm Fig. 15. D grade — SEM micrograph showing typical crack–microstructure interaction during stable FCG. Note the local competition between crack deflection and transgranular carbide cracking within the detailed image at the right.

crack deflection is usually listed as an extrinsic toughening mechanism too, such description should be carefully taken as: 1) it is not directly associated with any R-curve toughening promotion; and 2) its induced toughening effect comes from deviations of the crack tip from the mode I growth plane, rather than from direct shielding and reduction of the effective applied mode I stress intensity factor. The implications of such considerations from a fatigue viewpoint are discussed below. Under cyclic loading, mechanistically FCG in extrinsically-toughened brittle materials (e.g. advanced ceramics, intermetallics and hardmetals) is conceptually different from that in intrinsically-toughened ductile alloys [10]. Essentially, fatigue crack advance mechanisms in the former materials are given by the degree of degradation of crack-tip shielding behind the crack tip, as far as there is a toughening mechanism susceptible to be “degraded” or “inhibited” during cyclic loading. In the case of cemented carbides, given that ductile metallic binder is a relevant toughening agent through reinforcement behind the crack-tip, its susceptibility to be degraded or become inoperative under cyclic loads would imply a significant fatigue sensitivity for grades exhibiting such toughening mechanism exclusively. On the other hand, the presence and effectiveness of crack deflection as a toughening mechanism should not be affected by cyclic loading. Hence, a hardmetal exhibiting crack deflection as a complementary toughening mechanism, besides the ductile ligament reinforcement one, should be less sensitive to fatigue. Within this context, and considering the toughening contribution estimated from crack deflection above (Microstructural coarsening effects on the fracture toughness of cemented carbides section), it should be expected that coarser-grained hardmetals exhibit noticeable lower (from 20 to 30% down) fatigue sensitivity values (i.e. higher Kth/KIc ratios) than finer-grained ones. This indeed is in satisfactory agreement with the experimental findings presented in Fig. 13 and supports the conclusion that microstructural coarsening effectively reduces the fatigue susceptibility of WC–Co cemented carbides. Finally, it should be pointed out that the conclusions attained from the crack deflection

R = 0.1

analysis are also in complete concordance with previous findings by the authors where mixed-mode loading (a scenario intrinsic to any crack deflection mechanism) was found to lessen the assessed fatigue sensitivity of hardmetals [47]. Conclusions Based on the analysis of the experimental findings presented in this paper on fracture toughness assessment and FCG behavior evaluation for four different rock-bit hardmetal grades, the following conclusions may be drawn: (1) Rock bit cemented carbides exhibit high fracture toughness values (between 17 and 20 MPa√m) in direct association with their specific microstructural characteristics, i.e. medium/coarse carbide grain size and medium cobalt content. The corresponding influence of microstructural coarsening on the measured fracture toughness may be accounted by considering the effective operation of ductile ligament bridging and crack deflection, beyond the intrinsic crack growth resistance of WC, as the prominent toughening mechanisms. (2) Microstructural coarsening of WC–Co cemented carbides implies that relative increments in toughness are partly retained under cyclic loading (in terms of crack growth threshold); thus, fatigue sensitivity for coarser-grained grades is found to be lower than it could be extrapolated on the basis of trends determined for finergrained hardmetals. From FCG kinetics perspective, Kmax influence and relative m/n values (as related to monotonic and cyclic loading-related phenomena, respectively) are also higher than expected. (3) The FCG behavior of rock bit cemented carbides may be understood on the basis that the prevalent toughening mechanisms (ductile ligament bridging and crack deflection) exhibit distinct susceptibility to be degraded or inhibited under cyclic

R = 1.0

R = 0.5

2 μm Fig. 16. C grade — SEM micrographs showing fractographic aspects associated with stable cyclic (at R values of 0.1 and 0.5) and unstable monotonic crack growth (R = 1).

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