Powder Metallurgy Progress, Vol.8 (2008), No 3 230

MECHANISMS OF FRACTURE IN COMMERCIAL LAMP GRADE TUNGSTEN AT THE AMBIENT

I. Gaal, L. Bartha

Abstract Microcrack evolution in the semi-brittle region of lamp grade tungsten wires was investigated by means of torsion testing at ambient temperature followed by SEM observations. The fracture performance of this microalloy may be of general interest because it has high ambient tensile strength (2-3.5 GPa) in heavily drawn conditions, due to an ultrafine fibre structure (fibre widths are less than 300 nm), and high ambient tensile strength (1.8-2.1 GPa) in a specific recrystallized (at temperatures from 1300 to 1900°C) condition: submicron-sized fibre structure stabilised by long rows of nanosized potassium bubbles (d < 100 nm). The following features of the micro-crack evolution were revealed. (i) The microcrack initiation strain depends markedly on the processing conditions. (ii) The microcracks evolve at triple grain junctions and often coalesce into a helical macrocrack in a Lüders band. (iii) The triple junction hardening seems to play an essential role in microcrack generation. (iv) The low cycle fatigue performance is extremely sensitive to the processing conditions. Keywords: microcrack initiation strain, microcrack coalescence

INTRODUCTION Lamp grade tungsten is often referred to as non-sag (NS) tungsten, as fine size NS

wires (D < 390 μm) have excellent creep resistance even at very high temperatures (T > 2700°C), if the applied stress is less than 100 MPa. This performance is connected with a coarse, elongated, interlocking, grain structure that is often referred to as NS grain structure [e.g. 1]. During service life this structure is stabilised by an ensemble of nanosized potassium bubbles, provided that the oxygen and carbon activities of the atmosphere are sufficiently low. The actual potassium level depends on the intended applications. Table 1 gives some information on the composition of high quality NS ingots sintered from acid-leached doped tungsten powder, for which the temperature of the final sintering step is higher than for conventional sintering of technically pure tungsten [4,5].

Tab.1. Typical composition of high quality, sintered NS ingots (in μg/g).

Ca Fe Mo P S Al Si K O C <1 ≤4 <15 <1 <1 5-15 2-10 50-80 <50 <50

The actual O and C levels are usually lower (e.g. xO = 5 μg/g, xC = 30 μg/g), than

the upper limits of Table 1. The typical parameters of the dispersed potassium phase in high

István Gaal, László Bartha, Research Institute for Technical Physics and Materials Science, Budapest, Hungary

Powder Metallurgy Progress, Vol.8 (2008), No 3 231 quality wires are: VV : 0.2-0.3%, bubble diameter : <d> around 50 nm, and d < 100 nm [1-3].

This paper deals with the fracture and fatigue performance of fine size NS wires in torsion test at the ambient. The motivation of our material selection was as follows.

(i) "As-drawn", fine size NS wires have an extremely high tensile strength at the ambient: it ranges from 2 to 3.5 GPa. The high strength is provided by an ultrafine fibre structure, the fibre widths of which range from 100 to 300 nm. The high strength is combined with significant tensile ductility: the engineering strain at fracture ranges from 1 to 5% [6-9].

(ii) The submicron sized, recrystallized fibre structure (RF structure) evolves via continuous microstructural coarsening upon annealing at temperatures from 1300 to 1900°C [e.g. 7-11], and the RF structure retains the strong <110> texture of the as-drawn structure. The microstructural coarsening in the RF structure is very limited, even upon long annealing periods, if the heating rate to the temperature of the isothermal annealing is sufficiently high and the oxygen and carbon activities of the annealing atmosphere are sufficiently low. [e.g. 1,3,10,11]. The stability of the RF structure is provided by long rows of nanosized bubbles [e.g. 1,10,11] The ambient tensile strength of the RF structure is quite high (1.8-2.1 GPa) [4,6,11]. In addition, its bend ductility and torsional ductility are also exceptionally high at the ambient [e.g. 1, 12-16]. It is of significant importance that the RF structure allows us to follow the generation of microcracks in a sort of submicron-sized recrystallized tungsten. The thermal stability of the RF structure may be of general interest with respect to low cycle fatigue, because such kind of stability is of primary importance in this deformation-mode [17].

(iii) The processing route of the NS wires has many similar features with the recent experimental processing routes for high strengths and ultrafine microstructures in technically pure tungsten. Severe plastic deformation starts in both cases with full density, recrystallized stock-rods, that are fabricated from sintered ingots through hot rolling and/or swaging. In the NS processing, the stock-rods are subjected to severe plastic deformation with carefully adjusted rate of severe work hardening via successive steps of swaging and wire drawing at continuously decreasing working temperatures that range from 1300 to 600°C [e.g. 18, 19]. In the range of fine wire sizes, this process yields ultrafine fibre structures (fibre width 300 – 100 nm). In recent studies [20-24], coarse grained recrystallized stock-rods were subjected to severe plastic deformation by applying various combinations of swaging, equal-channel angular extrusion (ECAE), high-pressure torsion (HPT) and rolling at successively decreasing temperatures from 1300 to 600°C. The achieved microstructural length scales and the achieved compressive and tensile strengths were similar to those of NS wires.

The paper deals with the following topics. (i) Methodology for determination of the microcrack initiation strain in torsion test. (ii) Search for pieces of evidence clarifying the role of the triple junction hardening in the initiation of microcracks. (iii) Application potential of torsional low cycle fatigue testing in the qualification of NS wires.

EXPERIMENTAL

Material This work compares the performance of three types of commercial NS wires: (i)

shock resistant wires for vehicle halogen lamps (our code HV), (ii) wires for halogen lamps (our code HL, or W218), (iii) old type of NS wires for halogen lamps (code HGK). The HV and HL wires have higher purity with respect to Ca, P, and S, than the older type HGK

Powder Metallurgy Progress, Vol.8 (2008), No 3 232 material. HV and HL wires are produced from doped tungsten powder that is acid-leached before sintering, while the HGK technology used un-leached powders [4]. (The most important features of the processing and testing of HL wires and HV wires are given in [18] and [19].) The accumulated true plastic strain in our samples was high. The recrystallized stock-rod (diameter Dr) was subjected to cumulative axial elongation with a true strain (∈ = 2ln(Dr/Df) of 7 to attain HGK wires with a diameter (Df) of 174 μm. The stock-rod of HV samples was subjected to a cumulative true strain of 7.3 to attain wires with a diameter of 205 μm.

Mechanical characterisation The tensile strength of our samples with NS grain structure was 1±0.1 GPa, as it is

usual in high quality wires. The other mechanical parameters are listed in Table 2. Let us note that the strain at fracture in this Table is not given as the usually denoted engineering strain at fracture. The corresponding true strain, ∈f = ln[1+A/100], was given in Table 2, because the shear strain at fracture in torsion test is usually given as true strain, Γf. The value of Γf was measured far from the helical fracture surface, along which the sample separated into two parts. In the case of RF structures, the temperature of the V6 annealing characterises the sample. ∈f and Γf give the true strain at fracture in tensile and torsion test, respectively. The fibre widths LT and LTC were measured by means of TEM and SEM, respectively.

Tab.2. The tensile stress parameters of the investigated sorts of NS wires at the ambient.

Rm[GPa]

Rp0.2[GPa]

∈f Γf LT

[μm] LTF

[μm] as drawn 2.5 2.0 0.03 1.2 - - HGK3

D:320μm 1800°C 1.8 - 0.01 1.3 - 0.8 as drawn 2.8 2.5 0.03 1.0 0.13 - 1530°C 2.1 1.9 0.01 1.5 - -

HGK2 D:174μm

1800°C 1.9 1.85 0.01 1.7 0.75 0.7 HV

D:205μm as drawn 2.7 2.3 0.03 1.2 0.15 -

The difference between Rp1% and Rm was not significant, because the rate of work

hardening is very low in this strain region with respect to the tensile strength. In contrast, below 0.01 true strain the relative work-hardening rate is higher than 1, and the plastic elongation at Rm was always higher than 0.01 true plastic strain (Table 2). The fracture strain in tensile test decreased in the major part of the samples with increasing annealing temperature (Table 2). In contrast, the fracture strain in torsion test was very high, both in as drawn condition and in the RF-structures (Table 2).

Microstructural characterisation The as-drawn fibre structure consists of crystallites (i.e. fibres) that are elongated

in the drawing direction and are joined with three kinds of boundaries: (i) grain boundaries of the recrystallized stock-rod, which has been embedded into a dense dislocation network upon wire drawing, (ii) dense dislocation walls that are parallel to the drawing direction and (iii) incidental dislocation walls. (All the three kinds of these boundaries are usually referred to as fibre boundaries.). The fibre width, LT, is defined as the mean radial intercept

Powder Metallurgy Progress, Vol.8 (2008), No 3 233 among the fibre boundaries. The fibre widths of Table 2 are taken from previous TEM studies carried out on comparable samples [7,9,25].

The RF structure consists of elongated crystallites which are joined by high and low angle grain boundaries. The longitudinal generatrix of the elongated crystallites is in good approximation parallel to the central line of the wires (Fig.1c). The TEM contrast of the grain boundaries was typical for equilibrium grain boundaries, in which the density of extrinsic dislocation is low, and also the effects of other local stress sources (like oxygen or carbon clusters) are negligible [25]. It is a peculiarity of our RF structure, that its dislocation density is quite high (1010 cm/cm3) according to the X-Ray line profile analysis.

(a) (b)

(c)

Fig.1. Recrystallized fibre structure. (a) TEM micrograph taken on transversal thin film. (W218, V anneal at 1330°C for 30 s). (b) Longitudinal SEM fractograph (W218: V

anneal at 1330°C for 30 s). The black dots are traces of potassium bubbles. (c) Longitudinal SEM fractograph. (HL1: V6 anneal at 1800°C. Fibre width, LTF = 0.75 μm. The long traces of the grain boundary triple lines are nearly parallel to the central line of

the wire.

In order to reveal the transversal cross section of fibres, TEM micrographs were taken on thin foils being perpendicular to the central line of the samples [9] (Fig.1c). It turned out that the shape of the transversal fibre cross section changes upon annealing in W218 type “as-drawn” wires in the following way. The average ratio of the longest and shortest radial fibre-diameter, R, is about 4, while the upper limit of R amounts to 20. In contrast, upon a short anneal at 1800°C the mean value of R is 1.6, and the maximal value of R amounts to 4 [9].

Powder Metallurgy Progress, Vol.8 (2008), No 3 234

The fractographs of the RF structure have their own importance for two reasons. (i) It is difficult to distinguish the low and high angle grain boundaries in RF

microstructures by means of statistically relevant methods [26], and (ii) it is of primary importance to reveal in large areas those fibre boundaries,

that are prone to decohesion upon deformation at low temperatures. In order to get a suitable fracture path, double notched samples were bent in liquid

nitrogen bath, since in this case the fracture surface is nearly planar and lies parallel to the wire axis. The fracture pattern was depicted by means of high resolution scanning microscopy (LEO 1540 XB-Gemini FEG). The fractographs provide sharp traces of grain boundary triple lines and give useful hints for the three dimensional morphology of the RF structure (Fig.1). In order to determine the fibre widths, the grain boundary triple lines were counted along a radial test line of given length. The obtained apparent fibre width was corrected with the usual π/2 factor to get the fibre widths LTF listed in Table 2. (This correction assumes that the grain facets were randomly oriented with respect to the projection plane of the scanning imaging.)

The long traces of the grain boundary triple lines are nearly parallel to the central line of the wire. (In Fig.1c, the deviation from this direction is less than ±5°). In addition, the short grain boundary segments are far from being perpendicular to the central line of the wire (Fig.1c), as is assumed in various simple models. It is of importance that the triple joints connect short segments with long longitudinal grain boundary segments (Fig.1c), because this feature may lead to wing crack evolution via triple junctions hardening during plastic deformation, as was also observed in coarse grained recrystallized tungsten and molybdenum [5,27].

Sample preparation The route of sample preparation was designed to provide also information about

the ductility of the NS coils that are stress relieved on a molybdenum mandrel. It is well known that the applied stress relief anneal consists of two steps [28]. The first step intends to remove the W2C islands from the free surface by means of a heat treatment in hydrogen atmosphere with an adjusted dew point, while the a second heat treatment provides the required RF structure. (It is carried out in a sufficiently dry hydrogen.) After stress relief, the molybdenum mandrel is dissolved in a composite acid solution, in which the coils preserve the wire drawing grooves. In this condition, the coils have to show up a limited ductility upon stretching according to the industrial standards. Since stretching of a coil is composed of simultaneous bending and twisting, it of importance to have an insight into the parameters that govern the evolution of micro- and macrocracks in a torsion test [1,28]. (The ductility problems of tungsten coil have similar features to those of helical springs made of various sorts of premium grade high strength steels. The merits of torsion testing in the latter context have been summarised by Hallgarth [29].)

Since hydrogen occlusion might lead to embrittlement, too, we applied vacuum anneal. In order to attain high heating and cooling rates (∼500°C/s) the majority of the samples were annealed by self-resistance heating. (The samples were used also as their own resistance thermometers as described in [30]).

Before vacuum anneal, the lubricant layers were removed in hot aqueous solution of H2O2 (100 mg/g) [6]. A set of samples with code V were annealed by self-resistance heating in a Balzers high vacuum chamber having a residual vacuum pressure of 10-8 mbar. Code V5 and V6 refer to heat treatments for 15 min carried out at an adjusted P(O2) pressure of 5·10-5 and 10-6 mbar, respectively.

Powder Metallurgy Progress, Vol.8 (2008), No 3 235

The annealing code VM refers to samples that were annealed in two steps. The first treatment was performed at 1200°C for 10 min in the Balzers chamber at an adjusted P(O2) pressure of 10-4 mbar, while the second treatment was carried at 1800°C for 30 min in the vacuum chamber for Auger electron spectroscopy having a residual pressure of 10-10 mbar.

Samples with code F were electropolished in an aqueous solution of NaOH (50 mg/g) at room temperature in “as-drawn” condition to remove the lubricant layer [6]. They were furnace annealed for 15 min in dry H2 (having a dew point of -60°C) at 1500°C.

The profile of the drawing grooves was preserved during the heat treatments. (They are removed upon heat treatments at 2500°C for several hours [30].)

The samples selected for this study were split free in the "as-drawn" condition according to the conventional tests [19,28]. In order to reveal the surface quality in finer details, the wire surface was inspected after a V6 anneal at 1800°C. The sample surface was inspected by meansof SEM at a magnification of 1000 and 10 000, in order to reveal the following features: (i) long macrocrack along the grain boundaries with fine opening, (ii) microcracks at the grain boundary junctions, (iii) detection of the profiles of the drawing grooves (Figs.3 and 5a) and (iv) the loss of one or more grains from the surface, e.g. by corrosion processes upon cleaning the lubricant layer. The samples of this study did not have any of the said features before torsion test.

TORSION TEST An Amsler wire torsion testing machine was used to test the samples in free-end

torsion. The gauge length (L) over diameter (D) ratio was 500 in the unidirectional torsion tests. The applied axial stress amounted to 0.2 GPa. One full turn lasted 1 sec, i.e. the shear strain rate was 2.10-3/s at the free surface. The direction of rotation was manually changed in the low cycle torsion test at every full turn.

Fig.2. The drawing groove pattern reveals a diffuse Lüders band (LB) in the centre of the macrograph. The core of the LB goes over in both sides through transition bands into the

domains with axially homogeneous surface shear strain. The surface shear strain is axially homogeneous also in the LB core. The values of surface shear strain are different in the two sides of the LB. In this case, they amount to 0.10 and 0.12. The black marks are positions of

the orientation detector in the scanning microscope. (Sample HGK).

EVALUATION OF THE SURFACE SHEAR STRAIN If a cylindrical sample is subjected to axially homogeneous torsion, then the axial

straight marker lines turn to helices on the surface. The angle between the wire axis and the tangent of the helical marker line (i.e. the pitch angle) remains constant in every point of the helix, as long as the twist remains axially homogeneous. In this case, the shear strain in the vicinity of the free surface (the surface shear strain) is given by the relation

Powder Metallurgy Progress, Vol.8 (2008), No 3 236

Γ = =D ΩL

tg2

θ ( ) (1)

where D is the wire diameter, L is the gauge length at which the twist angle is θ. (θ is measured in radian units). Therefore, the pitch angle will be a proper measure of the local surface shear strain, as long as the pitch angle is constant in a well-defined interval of the cylinder. Of course, in order to test the homogeneity, one should have a great number of parallel maker lines at the start of the test. We used the drawing grooves as marker lines, since they are parallel to the drawing direction at the start of the torsion test.

Because an appropriate scanning image is an orthogonal projection of the cylindrical wire surface, the angle between the external contour of the wire and the tangent vector of the marker line will be equal to Ω at that point, in which the marker line cuts the central line of the projected wire surface. This angle can be measured in the scanning microscope by means of a position detector. To this end, it is adjusted to the external contour of the wire and to a drawing groove at that place, where it cuts the central line of the wire-macrograph. (The said positions of the detector are marked with black lines in Fig.2.)

RESULTS AND DISCUSSION

Microcracks evolving in torsion test The pattern of microcrack positions is quite simple in the RF structures, in which

the initiation strain for Lüders band formation is high. (The micrograph in Fig.3 was taken on a sample, which was free of Lüders bands at a surface shear stress of 0.46). In this case the microcracks are attached to previous triple grain boundary junctions and the damage is restricted to a quite low fraction of the triple junctions. In addition, the direction of the opened up grain boundary segments has no simple relation with the position of the principal directions of the external stress field. Of course, this is also true for the direction T and C, denoting in Fig.3 the compressive component of the stress field acting on the long boundary segments and the tensile component of the stress field acting on the short grain boundary segments being perpendicular to T. One may, therefore, conclude that microcrack formation should be sensitive both to the crystallographic parameters of the opened grain boundary segment and to the stress field of dislocations that are preferentially stored at the grain boundary junctions upon deformation in the semi-brittle temperature region. (Such a triple junction hardening has been extensively studied in coarse grained tungsten and molybdenum [27,32]).

We encountered the discussed microcrack pattern only in samples that were annealed at 1800°C. One may speculate that this is not accidental, because interstitial clusters on the grain boundaries should support grain boundary decohesion. One may, therefore, assume that interstitial clusters go into solid solution in our samples at 1800°C, and the rapid cooling after heat treatment may lead to more favourable cluster distribution, than those typical for the as drawn condition.

Powder Metallurgy Progress, Vol.8 (2008), No 3 237

Fig.3. Upon torsion microcracks evolve at the grain boundary triple junctions, when the

shear strain crosses a critical threshold. DR denotes the direction of the central line of the wire. T is the average direction of the elongated fibres. Since T is parallel to the drawing grooves, we made use of the grooves as a marker-lines for the strain evaluation. In this

micrograph the angle between T and DR (i.e. the pitch angle Ω) is 24.7°. (HL sample, V6 annealed at 1800°C.)

Lüders bands The surface strain in NS wires is often localised into diffuse Lüders bands (LB) in

a torsion test. The macrograph in Fig.2 shows the typical drawing groove pattern in and around an LB. It consists of three regions: core, transition bands, and matrix. The surface shear strain is axially homogeneous in the LB core, since the drawing grooves are parallel to each other at the centre line of the wire image. The surface shear strain in the core is high: in Fig.2 it amounts to 0.44. The LB core is joined via transition regions to the matrix, i.e. to those parts of the sample in which the surface shear strain is axially homogeneous and relatively low. (In Fig.2 the surface shear strain in the matrix, Γh, amounted to 0.1 and 0.12 in the two sides of the LB.) The length of a typical transition region is about one wire diameter. Therefore, the axial gradient of the shear strain is high in it, and Γ can not be measured in terms of eqn. (1).

Both the core and the transition band have their own crack patterns (Fig.4). The most typical feature of an LB core is a continuous helical split. Its trace covers the free surface usually along a single drawing groove (Fig.4a). The microcrack pattern of the transition band reveals the propagation mode of the LB, since in this region separate wing-type microcracks coalesce into a helical macrocrack that lies in a single drawing groove. In our case, the coalescence has a peculiar character: the long grain boundary segments of the triple grain junctions are continuously opened up, as the global twist on the sample increases (Fig.4b) (The oriented coalescence of microcracks is a well-documented feature also upon the macrocrack evolution in uniaxial compression of coarse grained recrystallized tungsten [5]). One may, therefore, conclude: the determination of the LB initiation strain, Γc, is a central point in the study of the crack evolution in NS tungsten in torsion.

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Tab.3. Crack initiation strain in torsion, Γc at the ambient in various NS fibre structures.

diameter [μm]

as drawn F anneal 1500°C

V5 anneal 1530°C

VM anneal 1800°C

HV1 205 0.05 - - 0.4 HL1 174 0.05 0.04 0.15-0.25 0.3 HL2 174 0.02 0.02 0.02-0.03 0.1 HGK 174 0.01-0.02 0.02 0.02-0.03 -

Fig.4a. The core of the LB. The surface roughness is higher in the core than in its own

transition bands. The core of the LB contains a well developed macrocrack. There are also few microcracks at the grain triple junctions outside the LB core.

Fig.4b. The transition band between the core of Lüders band and the domain with axially homogeneous surface strain. In the deep drawing groove the microcrack evolved at the

triple junction starts to propagate along the groove. In other domains, there are microcracks at three triple junctions. (Sample (HGK2).

The determination of the LB initiation strain is confronted with the following scenario. On the one hand, the torque-twist diagrams do not have inevitably a distinct maximum, when LBs evovle, while on the other hand, when such a maximum is observed,

Powder Metallurgy Progress, Vol.8 (2008), No 3 239 and the deformation is stopped, an LB with a helical macrocrack has been always observed. (In the latter case the surface shear strain outside the LB can be considered as the crack initiation strain, Γc).

When a distinctive maximum is absent in the torque-twist diagram, the initiation strain, Γc, has to be determined by successive approximation. To this end, we start with the observation of a short LB in the given batch. The corresponding matrix strain gives an upper limit for Γc, as the matrix strain can not be lower than initiation strain. Thereafter, another sample of this batch is twisted to attain a slightly lower surface strain as the already known upper limit of Γc. If the strain is axially homogeneous and the surface crack-free in this trial, a lower limit is obtained for Γc. In sample HL1, the observed upper and lower limits were 0.1 and 0.09. Therefore, one may, consider Γc = 0.1 as a fair value for the LB initiation strain.

(a) (b)

(c)

Fig.5. The giant difference in the fatigue performance of samples GK3 and H1. (a) There are no microcracks on the surface of sample H1 after 500 cycles with a strain amplitude of 0.01. (b) At the same cyclic deformation sample GK3 was divided by deep crevices into

fibre-bundles, (c) the surface of which was covered with a network of intercrystalline macrocracks.

The LB initiation strain was determined in various sorts of NS wires (Table 3) The results support the following conclusions. The crack initiation strain increased significantly, if the samples were annealed at high temperatures, although the fibre width increased also upon this treatment. In addition, the cooling rate has also to play a role, because furnace cooling resulted in lower crack initiation strains (Table 3).

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LOW CYCLE FATIGUE Low cycle fatigue test was carried out on two sorts of NS wires. To this end, (i) 5

HGK samples (diameter 340 μm, code GK3) were V6 annealed at 1530°C, and (ii) 3 samples of HV samples (diameter 205 μm, code H1) were VM annealed at 1800°C. The samples were cyclically twisted with one full turn in clockwise and one full turn in anti-clockwise direction. The gauge length was adjusted to obtain a surface strain amplitude (Γtot = Γel + Γpl) of 0.01, as long as Γ remains axially homogeneous. The cyclic torsion generated lattice defects in both cases, as evidenced by a significant increase of the residual resistance ratio R(4K)/R(300K). (In this context let us note that the resistance ratio is independent of the micro- and macrocrack generation [30].)

According to Fig.5a, there are no signs of microcracks on the surface of the H1 samples after 500 cycles. This performance supports the expectation that also the bulk of the H1 samples remained microcrack-free, since the cracks will be initiated in the vicinity of the free surface, as this domain is subjected to the largest shear stresses in torsion.

In contrast, deep radial fracture crevices appeared after about 300 cycles in the GK3 samples, as it is usual in low cycle torsion in various metals and alloys having a high tensile strength. The life cycle of the samples was 500 cycles. The vicinity of a fracture rosette is shown in Fig.5b. The radial fracture surfaces of the crevices are covered with a network of intercrystalline macrocracks (Fig.5c). The free surface is covered with a similar striation network, too.

Of course, the performance of the GK3 samples is in full accord with the expectation that tungsten is very prone to grain boundary decohesion. However, the message of Fig.5 highlights the possible merits of an appropriate technology, that may markedly improve the fatigue performance of commercial NS wires, as long as the fine scale of the microstructure is preserved by means of a proper nano-dispersed second phase.

It is difficult to appraise the performance of the H1 samples with respect to other published results, because we are merely aware of publications that reported on the high cycle fatigue performance of tungsten. However, let us mention one technical point. The number of life cycles lies between 300 and 3000 at a plastic strain amplitude of 0.01 in 8 sorts of commercial alloys (Fig.8 in [33]), and the H1 samples fit favourably into this scenario.

CONCLUSIONS The surface shear strain at fracture is much higher in the torsion test at the

ambient, than the surface shear strain at which microcracks are initiated at the free surface at the grain boundary triple junctions.

The evolution of the microcrack is random with respect to the geometrical position of the triple junction in the co-ordinate system of the principal axes of the external stress. This observation suggests that the evolution of the microcracks may be governed by the grain boundary character sensitive triple junction hardening also in submicron grained tungsten.

The performance in cyclic fatigue is markedly different in various types of NS tungsten. The material developed for shock resistant service seems to provide also a remarkably good cyclic fatigue performance in torsion.

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