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Powder Metallurgy Progress, Vol.15 (2015), No 2 218

FATIGUE RESPONSE AND FRACTOGRAPHY OF SINTERED MATERIALS

H. Danninger, B. Weiss, A. Betzwar-Kotas, G. Khatibi, Ch. Sohar, M. Dlapka, Ch. Gierl-Mayer

Abstract Fatigue is one of the main failure mechanisms in materials and components, involving also powder metallurgy products. In the present study the typical failure modes and fracture surfaces are shown for sintered steels, PM tool steels and WC-Co hardmetals, with particular focus on crack initiation. It is shown that for sintered steels, mostly microstructure-controlled initiation occurs, i.e. numerous microcracks are formed at different locations, and through crack growth and coalescence final fracture occurs. With PM tool steels, in contrast, defect controlled failure is observed, cracks starting from the few remaining nonmetallic inclusions, while ingot metallurgy tool steels fail through cleavage of large carbides, i.e. microstructure-controlled. Hardmetals, finally, exhibit microstructure controlled failure, mainly through binder decohesion from larger WC grains. A fatigue “limit”, i.e. a horizontal branch of the S-N curve, is not observed in any of the materials but the graphs consistently drop up to N > 10E9. Keywords: Fatigue, fractography, sintered steels, tool steels, hardmetals, crack initiation

INTRODUCTION Although material data sheets contain primarily properties such as tensile and

yield strength, elongation to fracture etc., the main loads that result in failure of metallic components are corrosion, wear and in particular fatigue. This holds also for powder metallurgy materials which are frequently found in applications where fatigue is the main loading variant. Typically the loading is in the high N range, also for hardmetal tools N up to 10E6 being quite common, and automotive parts can easily encounter N > 10E8 cycles.

From this reason, there have been numerous studies on the fatigue behaviour of powder metallurgy materials (e.g. [1-16]. Here it must however be considered that there are quite different variants of “PM materials”. The main and most evident characteristic is the inherent porosity: while for classical, mostly ferrous, PM precision parts produced by the press-and-sinter route, porosity is a typical and more or less inevitable feature, ranging from 6 to 20% [17], other PM products such as hardmetals or PM tool steels as well as high performance light alloys, porosity has to be avoided at any rate, and these materials are thus usually fully dense. There are some intermediate groups in which slight porosity, typically present as isolated, more or less spherical, pores is tolerated, as many MIM products, but in general the presence of evenly distributed and generally open and interconnected porosity [18] is a distinguishing feature between these groups of PM products, which also has Herbert Danninger, Agnieszka Betzwar-Kotas, Golta Khatibi, Christian Sohar, Magdalena Dlapka, Christian Gierl-Mayer, Technische Universität Wien, Institut für Chemische Technologien und Analytik, Vienna, Austria Brigitte Weiss, Agnieszka Betzwar-Kotas, Golta Khatibi, Universität Wien, Fakultät für Physik, Vienna, Austria Dedicated to Dr. Andrej Šalak at the occasion of his 90th birthday

Powder Metallurgy Progress, Vol.15 (2015), No 2 219 profound implications on the fatigue behaviour. Most studies have been performed on sintered steels since these are employed for structural, load bearing parts [19] e.g. in automotive engines and transmissions, but also fully dense PM products have been studied thoroughly, mostly with respect to the effect of singular defects [20, 21]. In the present study, the fatigue fracture behaviour is described both for fully dense materials such as sintered hardmetals and HIPed (i.e. pressure sintered) tool steels as well as for classical sintered steels as used for PM precision parts.

Fatigue initiation In general, one main question regarding the fatigue behaviour of materials is the

failure initiation mode, if it occurs microstructure-controlled or defect-controlled. This is related to the question if there is a large amount of potentially crack-initiating features in the microstructure – such as numerous pores, inclusions etc. – or only a few of them [20]. E.g. for high purity steels, as employed for springs or bearings, it has been shown that failure is usually linked to a few slag inclusions in the loaded volume [22, 23]. The effect of defects can be assessed using different concepts (e.g. [24]); it can be graphically depicted in the Kitagawa-Takahashi diagrams [25] (Fig.1): there is a certain critical defect size for each material and loading condition [26] above which the largest defect in the loaded volume causes fatigue failure.

When applying this to sintered materials it can be stated that the homogeneously distributed porosity can be defined as “integral defects”, i.e. in fact the numerous sintering contacts are of relevance which of course follow a size (or area) distribution but a continuous one; in general it is the sum of the cross sections of all contacts relative to the nominal specimen cross section that determines the fatigue properties, which parameter is described as “load bearing cross section” Ac [27]; a related approach is to use the nominal cross section minus the contacts, i.e. 1-Ac, which is defined as “plane porosity” after Slesar et al. [28, 29].

Fig.1. Kitagawa-Takahashi diagram (schematic).

Also with fully dense PM materials such “integral defects” can be found, as e.g. large carbides etc., as will be shown below. Since however in the absence of porosity the horizontal line in the Kitagawa diagram is generally at higher levels while the angled branch – defined by the effective threshold stress intensity factor – is not too much affected, the intersection of both lines, which gives the critical defect size, is shifted to the left, i.e. to lower diameters, which implies that materials which higher inherent strength are more sensitive to singular defects.


In the following, some powder metallurgy materials are presented with regard to fatigue behaviour and fracture surfaces, sintered steels and hardmetals as materials produced through the classical press-and-sinter route; however, for comparison also pressure-sintered, i.e. HIPed, materials are presented. As for wrought metallic materials [30, 31], also for PM types fractographic investigations are the main tool for studying the crack path and in particular identifying the initiation site (see e.g. [32]).

There is however a further aspect to be considered: In particular for complex-shaped precision parts, under service conditions fatigue failure may be neither microstructure- nor defect-controlled but may originate from geometrical features such as notches, bores, blind holes etc., i.e. areas of stress concentration. Here, more or less the same design rules as for components from wrought steels apply. The effect of notches on fatigue of sintered steels has been extensively studied by Beiss et al. who also described the role of the effective loaded volume and its correlation to notch geometry. This topic, as well as that of mean stress effects, is not discussed here; the reader is referred to the respective publications (e.g. [33, 34]).

Sintered steels PM steels produced by pressing and sintering are characterized by residual

porosity, typically in the range of 6 .. 15%, which is a clear difference to the fully dense materials described above. Porosity is usually an integral property, i.e. the volume content of the pores is more or less the same in all areas of a PM part (with some variations in case of complex shaped parts). Since porosity lowers the effective load bearing cross section of the material – only the sintering contacts being able to transfer mechanical load – it is clear that the porosity is the main parameter for affecting the mechanical, thus also the fatigue, properties of PM materials (e.g. [1, 4, 6, 10, 35]).

Fig.2. Kitagawa-Takahashi diagram for sintered steels with varying porosity

When plotting the Kitagawa-Takahashi diagram for a structural steel, it is clear that the horizontal line drops with increasing porosity. On the other hand, fracture mechanical studies have shown [35] that the threshold stress intensity factor hardly varies with porosity, at least up to about 12%, i.e. the angled line in the Kitagawa-Takahashi diagram remains more or less in place. This means that with increasing porosity the intersection between the two lines is moved to the right, to larger defect sizes (see Fig.2), the material becomes less performing but also more tolerant to singular defects. Therefore it

Powder Metallurgy Progress, Vol.15 (2015), No 2 221 can be expected that PM steels tend to microstructure-controlled failure rather than to defect controlled one, except in case of relatively large defects. This of course depends on the total porosity: the high density-high performance steels today used for engine and transmission parts are more sensitive to singular defects than standard components with e.g. 6.8 g.cm-3 density.

This also agrees with experience. Furthermore, as also described for hardmetals (see below), in case of microstructure-controlled failure it is more difficult to identify the crack initiation sites than in the case of defect-controlled one in which latter case the defects mostly stand out as singularities. Usually the initiation sites are the more easily found the higher the loading cycle number was; therefore, ultrasonic fatigue testing is a useful tool also here, but also more conventional fatigue test methods, as e.g. rotating beam fatigue, show this effect. Also the differentiation between the areas of stable fatigue crack growth and final fracture is easier at higher N, in part also since the former area is larger. Fig.3 shows RBF specimens that have failed after different N; the larger proportion of the area of stable crack growth is evident in the right image, as is the more pronounced different between the former area and that of final fracture.

a) Nf = 1.3 . 10E5 b) Nf = 7.0 . 10E7

Fig.3. RBF surfaces of Fe-1.5%Mo-0.7%C, 7.4 g.cm-3, heat treated. 100 Hz, R = -1

a) overview b) Crack-initiating small defect

Fig.4. RBF surface of sintered steel Fe-Mo-C, heat treated. Nf = 7.0 . 10E7.100 Hz, R = -1

The difference between the initiation area, that of stable crack growth and of final fracture is usually visible also with sintered steels. The initiation area is characterized by a smooth surface (similar to the “mirror” area described in hardmetals); occasionally the

Powder Metallurgy Progress, Vol.15 (2015), No 2 222 initiating feature is discernible, as shown in Fig.4 where a small secondary pore has initiated fatigue cracking (as an example of defect-control also with sintered steels).

The difference between the various areas of the fatigue fracture surface is particularly well visible in thermochemically treated steels. In Figure 5 a plasmanitrided specimen is shown that failed at rather high load and moderate N. The smooth area of initiation is clearly evident (Fig.5a), with transgranular cleavage, and also the nitrided area stands out clearly as a dark rim. The area of stable crack growth (Fig.5b) is still quite smooth, with a fine, striation-like structure, but without orientation. In the transition zone between the area of stable crack growth and that of final fracture (Fig.5c), this morphology is still present, but also dimples can be seen. In the area of final fracture, the dimple morphology is absolutely dominant (Fig.5d), the fracture surface resembling that of impact tested specimens (and final fracture is in fact something like impact loading).

With regard to striations it can be stated that they may be present in sintered steels, but in many cases they are not observed; therefore the absence of striations cannot be regarded as proof that failure has not occurred through fatigue loading.

a) initiation area b) area of stable crack growth

c) transition area to final fracture d) area of final fracture Fig.5. RBF surface of Fe-Mo-C, plasmanitrided. Nf = 5.7 . 10E5. 100 Hz, R = -1


a) early stage of fatigue b) later stage of fatigue Fig.6. Surfaces of fatigue loaded specimens, slightly notched; different stages of fatigue

process. Loading direction: y axis

If a porous PM steel component is cyclically loaded, the stress is concentrated at the sintering contacts since there the cross section is minimum, and the first small cracks appear there. It has been shown [36] that in this first stage the orientation of the cracks is not so much determined by the applied stress than by the microstructure, and therefore also cracks which run almost parallel to the orientation of stress are observed (see e.g. Fig.6a). Only after extended loading the microcracks tend to join, and these larger cracks are progressively oriented perpendicular to the stress orientation (Fig.6b) and grow into that direction until one of them becomes supercritical, growing to final failure. This process can be followed e.g. by systematic studies on metallographic samples in the SEM; [36]; Figs.6 a, b are typical images of the early stages of fatigue cracking.

Also in sintered steels, initiation through singular defects may occur (e.g. [37]). Such defects may be slag inclusions, fully dense ones as shown in Fig.7a or partially reduced as in Fig.7b, large secondary pores (Fig.7c) generated e.g. from larger alloy element particles (Co, Mo, Cr,…), but also pore clusters (Fig.7d) or areas on increased porosity which may be caused by agglomeration of pressing lubricant. Clean processing conditions and optimized mixing, e.g. stepwise addition of the various components of the powder mix, are measures to eliminate such singular defects and thus enhance the reproducibility of the fatigue properties.

An intermediate variant of crack initiation, between microstructure- and defect-controlled initation, respectively, is the emergence of intergranular fracture zones. These effects have been observed with heat treated sintered steels, both sinter hardened and carburized ones; typical initiation sites are shown in Fig.8 [38]. As is clearly visible, the area of intergranular failure may vary in size, covering only a few grains, as visible in Fig.8a, or extending over a range of >100 µm (Fig.8b), which is definitely in the range of crack- initiating defects, at least for high strength sintered steels (see Fig.2). Such intergranular area are observed also in case hardened wrought steels [39]; in [38] it has been described that the effect occurs with Cr-Mo prealloyed steels with relatively high carbon content, and it is assumed that very thin, filmlike carbides are formed during cooling that are uncritical in case of monotonic or low-N caclic loading but tend to decohesion when loaded in the gigacycle range. Thus, carbon control (which in the case of sintered steels is always linked also to oxygen control) is an essential measure to prevent these unwelcome effects.


a) slag inclusion b) Partially reduced inclusion

c) Secondary pore d) cluster of secondary pores

Fig.7. Crack-initiating singular defects observed in sintered steels (Push-pull, 20 kHz, R = -1).

a) internal initiation b) initiation from the surface

Fig.8. Intergranular fatigue crack initiation areas in Cr-Mo prealloy steels, sinter hardened (Push-pull, 20 kHz, R = -1) [34]

In addition to crack initiation, also the crack path in a PM steel is of high interest, both in static [40] and in cyclic loading. The crack path is not only determined by the porosity but, in particular in the case of chemically heterogeneous sintered steels such as

Powder Metallurgy Progress, Vol.15 (2015), No 2 225 Ni-Cu-Mo diffusion bonded grades, also the microstructural constituents. The high relevance of the alloy element distribution in sintered steel has been known for a long time (see e.g. [41]); It is still a matter of discussion if chemically heterogeneous steels offer better mechanical, esp. fatigue, performance than homogeneous (prealloyed) steels of the same composition, due to crack blunting in austenitic regions, as proposed by Engdahl et al. in 1990 [42]. Studies are described e.g. in [43-45], also including the role of surface residual stresses [46].

PM tool steels Powder metallurgy tool steels form the “high end” of the tool steels commercially

available [47], both with regard to cost and to performance. PM enables introduction of much higher contents of carbides in the material, in particular of carbides that are difficult to process by ingot metallurgy such as e.g. VC. This gives increased wear resistance; furthermore, PM tool steels have an isotropic microstructure, thus their properties are identical in all orientations, which makes life for the designer easier since it is not necessary to think about how to machine a tool from the blank, as has to be done with IM tool steels.

PM tool steels have much finer microstructure than standard IM grades (see Figs.9a-c). Therefore also the crack-initiating microstructural components are different. With ingot metallurgy steels, both high speed steels and, still more pronounced, cold work tool steels, large carbides are present that are elongated in the direction of hot working. As stated above, this gives different static and also cyclic properties parallel and perpendicular to the carbide orientation [48], understandably the former being the preferred direction. Also from the fracture surfaces, this orientation effect is evident.

a) Ingot metallurgy, cross section

b) Ingot metallurgy, longitudinal section

c) PM tool steel

Fig.9. Microstructures of ingot metallurgy vs. powder metallurgy tool steels

When studying the fracture surfaces after testing in the parallel orientation, it can be observed that cracking invariably starts at large single carbides or larger carbide clusters [49], typically by transgranular failure of the carbides which are well bonded to the steel matrix [50]. As shown in Fig.10, this failure usually occurs near the surface (unless compressive residual stresses are present, in which case fisheye failure occurs, see Fig.10b [51, 52]). Around this initiation site, a hemispherical (in case of fisheye spherical) area – “granular (bright) area” - indicates a zone where the fine secondary carbides have been separated from the matrix by decohesion (Fig.10c, d). Failure is typically microstructure-controlled since there are always numerous coarse carbides than can act as crack initiation sites.

PM tool steels, in contrast, exhibit defect-controlled failure; here, the initiation sites are the few remaining nonmetallic inclusions in the material (and, in case of vacuum

Powder Metallurgy Progress, Vol.15 (2015), No 2 226 sintered tool steels, at pores) [53]. The reason is that the carbides in PM tool steels are much smaller than in IM grades, typically <5 µm as compared to up to 40 µm with IM steels. Therefore the carbides, even if they would fail, are too small to start a propagating crack in the material. The nonmetallic, typically oxidic, inclusions, in contrast, show decohesion from the matrix. In Fig.11, an example for such a crack-initiating inclusion is shown; from the composition of these inclusions it can be concluded that they originate from the lining of the melting furnace or the tundish. Since they are very fine, much finer than the carbides in IM steels, the fatigue endurance strength of PM steels is significantly higher; since however the size and location of the inclusion is a matter of statistics, the scatter of the data is somewhat wider than with the IM steels, as typical for defect-controlled vs. microstructure-controlled failure. This is evident from the respective S-N graphs if the experimental scatter is plotted (see Fig.12).

a) stress-free surface; cracking from the surface

b) compressive stresses at the surface; “fisheye” fracture

c) Fracture surface of ingot metalurgy tool

steel D2; SE image d) as Fig.10c, BSE image

Fig.10. Fatigue fracture surfaces of ingot metallurgy cold work tool steel AISI D2 (W.Nr.1.2379, Boehler K110). Push-pull, 20 kHz, R = -1


a) Overview; crack initiation from the

surface b) as Fig.10a; initiation site – nonmetallic

inclusion Fig.11. Fatigue fracture surfaces of PM cold work tool steel Boehler K390. Push-pull, 20

kHz, R = -1.

Fig.12. S-N curves for ingot metallurgy (K110) vs. powder metallurgy (K390) cold work

tool steels

Hardmetals The single PM product that is most important economically, causing the highest

turnover with about 10 GUSD worldwide annually, is hardmetal which is the most important tool material for metal cutting, rock drilling, wood cutting but also for rolling etc. Fatigue loading of hardmetals occurs in particular in interrupted cutting as e.g. in milling of metallic workpieces but also in percussion drilling of rock or concrete. In particular in interrupted metal cutting a typical feature is that not only cyclic mechanical loads but also thermal loads are applied; the latter effect is further aggravated in case of cooling fluids being applied since in this case the cutting edge is quenched immediately after each cutting run, thus being exposed to thermal shock. It should however be considered that also thermal cycling in fact means primarily mechanical cycling since thermal expansion and contraction

Powder Metallurgy Progress, Vol.15 (2015), No 2 228 are the main effects that lead to cracks. Typical indicators for fatigue failure of hardmetal tools are comb cracks growing into the tool perpendicular to the cutting edge [54].

Not surprisingly, the fatigue behavior of hardmetals has been studied thoroughly quite early (e.g. [55]), fracture mechanical data being described in [56, 57]; also the crack propagation behavior in case of fatigue loading has been investigated using e.g. EBSD imaging [58]. The damage during cutting itself has been studied by SEM, more recently also combined with FIB, 3-dimensional plots of the cracks being thus obtained [54, 59].

Laboratory testing of hardmetals is difficult due to its brittle character, standard push-pull testing in servohydraulic machines being hampered by even slight misalignments resulting in erroneous data. Therefore many groups have done bending tests [60, 61], the significance of which however is limited by the relatively small loaded volume which is problematic in particular in the case of defect-controlled failure.

Recently, ultrasonic resonance fatigue testing [62, 63] has been applied also to hardmetals [60]. This technique is particularly well suited for hardmetals since the specimen has to be fixed only on one end, the other one resonating freely; thus, slight misalignments do not affect the results. Ultrasonic testing also enables attaining very high loading cycle numbers in short times; 10E10 cycles can be attained in about 5 days compared to 8 years for standard fatigue testing. Such high N are useful when the initiation mode has to be assessed since the role of defects stands out the more clearly the higher N are attained during testing. Furthermore, also the existence of a “true fatigue limit”, i.e. a horizontal section of the S-N graph, can be checked fairly easily. For low-ductility materials such as hardmetals, also the frequency effect is insignificant as shown e.g. for high strength sintered steels in [64].

The main disadvantage is the high cost of specimen manufacturing; the hourglass test specimens that proved to be best suited have to be machined from a presintered bar that is subsequently liquid phase sintered, the thread for fixing it to the testbed being machined oversize to compensate for shrinkage during sintering. To assess the correct oversize to obtain e.g. M10x1 after sintering requires a lot of experience and very precise sintering.

A highly important factor is the stress state at the surfaces. If the surface are ground, as is common practice with hardmetals, compressive residual stresses are introduced which may easily exceed 2000 MPa. Those compressive RS “shield” the specimen surface from the loads, thus resulting in fatigue failure from the interior and formation of the typical “fisheye” fracture surfaces, similar to the “core crushing” known from case hardened steel components. This phenomenon has been described in detail with tool steels (see above).

In general it is recommended to check the crack initiation sites of fisheye failure very carefully. Such failure modes may be caused by defect initiation, e.g. from pores or inclusions; if however none such phenomena can be identified it is to be checked if compressive RS have been present at the surfaces; in this case not only the crack initiation sites are relocated but also the fatigue endurance strength is shifted to higher levels. This is a welcome effect in service – therefore shot peening to introduce compressive RS is done for steel parts but also for hardmetals – but it is undesirable if the fatigue properties of the material have to be determined. In this case it has to be granted that any compressive RS have been lowered to insignificant levels (typically <200 MPa), e.g. by annealing.

S-N curves of commercial hardmetal grades show surprisingly small experimental scatter for such a fairly brittle material [65], indicating reproducible quality, but it stands out clearly that there is no fatigue “limit” but the S-N graph drops consistently at least up to N = 10E10 cycles.


a) Nf = 3.20 . 10E6 b) Nf = 2.07 . 10E7 c) Nf = 5.71 . 10E8

Fig.13. Typical fracture surfaces of hardmetals tested with different stresses. Push-pull, 20 kHz, R = -1

Fatigue fracture surfaces of hardmetals are shown in Fig.13, the tests having been performed by ultrasonic resonance testing. The regular failure mode is from the surface; only in rare cases internal failure and “fisheye” appearance are recorded. In those rare cases a defect in the material has resulted in cracking from inside; generally however hardmetals are surprisingly reliable in failing from the surface, as can also be seen from the small “mirror”, i.e. the zone of slow crack growth. This “mirror” is more or less hemispherical while in case of internal cracking it would be a circle, the typical “fisheye”. Thus, hardmetals are an example for “microstructure-controlled failure”. In particular for ultrafine hardmetal grades which today are standard for metal cutting, it may be extremely difficult to identify the actual site of fatigue crack initiation since no single microstructural feature can be found that may be linked to failure.

For medium and coarse hardmetal grades, with WC grain size >1 µm, the initiation sites are more easily found; in most cases they are large WC facets that show decohesion from other WC grains or the binder phase, as shown in Fig.14a. In these areas there is virtually no indication of ductile failure of the matrix. When moving from the initiation site towards the area of final fracture, more and more dimples are found (see Fig.14b, c). This indicates the “metal-like” fatigue behavior according to Llanes [66]: the bridging effect of the binder areas is effective in case of static loading but much less or not in cyclic one.

a) Initiation zone at surface b) Transition area, 100 µm

from crack origin c) Final fracture

Fig.14. Fatigue fracture surfaces of coarse-grained hardmetal after ultrasonic fatigue testing. Push-pull, 20 kHz, R = -1

The main parameters for hardmetals are binder content and WC grain size. The former parameter seems to be not too significant while finer grained microstructure tends to improve the fatigue endurance strength [65, 67], which is not surprising when considering that crack initiation is caused at large WC grains; it can be assumed that even if the WC grain size distribution in hardmetals is fairly wide, in a fine grained grade even the largest

Powder Metallurgy Progress, Vol.15 (2015), No 2 230 WC grains are still markedly smaller than in a standard or coarse grained variant. An experimental parameters that should not be ignored is of course the surface roughness: grinding marks, as typically shown in Fig.15, may cause crack initiation prematurely; therefore grinding and polishing processes have to be carried out carefully to avoid surface damage.

a) Overview b) Detail

Fig.15. Fracture surfaces of fine grained material after ultrasonic fatigue; (a) crack initiation site; (b) deeper groove of the surface produced by sample machining. Push-pull, 20 kHz,

R = -1.

CONCLUSIONS Powder metallurgy products are frequently subjected to fatigue loading in service.

This holds for structural components from sintered steels as well as for PM tool materials. Since the loading is mostly in the high to ultra high cycle fatigue loading, the role of crack initiation is of high importance, and here both microstructure controlled and defect controlled failure may be encountered, the differentiation requiring thorough fractographic studies.

With sintered steels, the inherent porosity lowers the inherent fatigue endurance strength, thus lowering the sensitivity to defects and favouring microstructure-controlled failure. In this case, the initiation sites may be difficult to define, since the porosity is rather “integral”, but the surface is a probable area of initiation. Defect initiated failure may occur if impurities are introduced, e.g. during powder handling, or if coarse alloy element particles that form secondary pores. In practice, of course, also geometrical features such as notches or blind holes must be considered since they may be more relevant than the microstructure, but here it is the task of the designer rather than the materials scientist to find the optimum solution.

With PM tool steels, defect-control is observed, because the fine carbides are too small to act as initiation sites, and therefore the last remaining slag inclusions (which are quite small either) start cracks through decohesion. This is in pronounced contrast to wrought tool steels which contain numerous large carbides that act as initiation sites through cleavage fracture, thus giving a clear example of microstructure-controlled failure


In the case of hardmetals, microstructure-controlled initiation dominates, which is somewhat surprising with regard to the low toughness of these materials and indicates high cleanliness. Cracking starts from the surface and typically through binder decohesion from larger carbide grains. The initiation area is macroscopically visible as “mirror”, i.e. as a very smooth, usually semicircular, zone around the initiation site, and the size of this mirror is related to the loading cycle number to fracture.

With none of the materials investigated, a “fatigue limit”, i.e. a horizontal branch of the S-N (Wöhler) curve at high N has been found, i.e. the designer always has to work with a fatigue endurance strength that is related to the estimated loading cycle number to which the component is exposed in service.

Acknowledgement This work was in part supported by the Austrian FWF (project P17650-N02), in

part carried out within the project “Hoganas Chair” sponsored by Hoganas AB, Sweden. Furthermore, the authors want to thank the project partners in the EPMA UFTH group, in particular the industry partners for financing the studies on hardmetals: CERATIZIT, Durit, Element Six, Hartmetall AG, Hilti, Kennametal, Sandvik, Unimerco.

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