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Materials Science and Engineering A 527 (2010) 6683–6689

Contents lists available at ScienceDirect

Materials Science and Engineering A

journa l homepage: www.e lsev ier .com/ locate /msea

he effect of prestrain on low and high temperature creep in Ti834

ark Whittakera,∗, Paul Jonesa, Cameron Pleydell-Pearcea, David Ruggb, Steve Williamsb

Swansea University, Singleton Park, Swansea SA2 8PP, United KingdomRolls-Royce plc, Moor Lane, Derby DE24 8BJ, United Kingdom

r t i c l e i n f o

rticle history:eceived 24 March 2010eceived in revised form 1 July 2010ccepted 2 July 2010

a b s t r a c t

The low and high temperature creep properties of the � + � titanium alloy Ti834 have been studied forvarious levels of applied plastic prestrain. Evidence is shown that the prestrain at room temperaturecauses the formation of quasi-cleavage facets in a manner described by the Evans–Bache model. These

eywords:restrainitaniumacets

facets are also produced when the material is creep tested at room temperature, although larger areasof faceting are found. The applied prestrain causes a drop in the strain at failure at room temperaturedue to a high dislocation density inhibiting further movement of dislocations. At high temperature thisincrease in applied prestrain results in an accelerated creep rate. This is attributed to a number of issues,including increased thermal energy allowing greater dislocation recovery and mobility. Swaged Ti834material is also considered and is shown to behave similar to highly prestrained material with evidence

educe

reep of strain hardening and r

. Introduction

The near � titanium alloy Ti834 was developed for high temper-ture applications and shows exceptional mechanical propertiesp to 630 ◦C. This fact, accompanied by the desirable strength toeight ratio for titanium alloys, has made Ti834 a popular choiceith designers of gas turbines for compressor disc and blade appli-

ations.To achieve an optimum compromise between high tempera-

ure creep and lower temperature fatigue, careful control of theicrostructure has been applied by alloy developers. However, pre-

ious work [1] has shown that Ti834 is one of a series of near � and+ � titanium alloys which show a sensitivity to dwell periods at

ow temperature, loosely termed ‘cold creep’. This type of failureas first appreciated in service during the early 1970s when twoncontained fan disc failures were shown to be as a result of dwellatigue in the near � alloy Ti685.

Subsequent research [2] has shown that this type of dwell sensi-ivity is related to the formation of ‘quasi-cleavage’ facets that oftenccur subsurface and lead to crack initiation, promoting prematureailure. The Evans–Bache model seeks to describe the formation ofnitial facets in terms of dislocation pile ups at interfaces between

ard and soft grains, or possibly even structural units. These facets,hich generally have a near basal plane orientation, develop due

o the gradual separation of intense slip bands under the action oftensile stress normal to the slip plane. The presence of strong and

∗ Corresponding author. Tel.: +44 1792 295573; fax: +44 1792 295693.E-mail address: m.t.whittaker@swansea.ac.uk (M. Whittaker).

921-5093/$ – see front matter © 2010 Elsevier B.V. All rights reserved.oi:10.1016/j.msea.2010.07.008

d cold creep rate.© 2010 Elsevier B.V. All rights reserved.

weak regions within the microstructure facilitates this mechanismthrough stress redistribution. Grains that have a basal plane orien-tation nearly perpendicular to the loading direction act as a ‘strong’grain, i.e. slip is difficult. However, if adjacent grains have easilyavailable slip systems, they can provide the mechanism for dislo-cation pile ups at the boundary. The pile up provides the requiredinduced shear (�f) and tensile (�) stress on the unfavourably ori-entated basal plane with �s the resolved shear stress on the ‘weak’grain, Fig. 1. More recent research has established that the angle atwhich these facets form is a function of loading condition [3].

Although the model aptly describes the process of facet gener-ation, which has more recently been extended [4,5] and supportedwith experimental evidence in the near alpha titanium alloy Ti685[6], the macroscopic loading conditions under which extensivefacet formation occurs have not been so well described. In particu-lar the formation of facets under application of a constant load (orstress) in a conventional creep test is an area of interest.

Furthermore, it is widely acknowledged that the application ofprestrain at room temperature can produce marked alterations inthe creep rate of titanium alloys [7]. However, the current workseeks to investigate how these creep rates and other mechanicalproperties are affected by the presence of quasi-cleavage facetswhich may be introduced through prestrain.

The paper seeks to address these issues through a series of dif-ferent loading conditions. Ti834 material was supplied in an as

received condition from a disc forging, although orientation wasnot recorded. These specimens were then prestrained to four dif-ferent levels, 2% total compressive prestrain (referred to as −2%),2% total tensile prestrain, 8% total tensile prestrain while a batch ofspecimens remained in the as received condition (referred to as 0%).

6684 M. Whittaker et al. / Materials Science and E

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ig. 1. Dislocation pile ups at grain boundaries induce a combination of shear andensile stresses which allow facet formation in titanium alloys, usually with a basalr near basal plane orientation.

he prestraining was used to simulate prior working operationsn the material. Specimens were subsequently tested under hotnd cold creep conditions along with stress relaxation and fatigueests, with SEM analysis used to explore fracture surface featuresncluding facet formation.

. Experimental method

The Ti834 was found to have a bimodal microstructure withrimary alpha grains with a semi-major axis ranging from 20o 200 �m in diameter, and grains encompassing transformedroduct with a semi-major axis ranging from 50 to 200 �m in

iameter. The crystallographic texture of the material was essen-ially random, which limited concerns about specimen orientationreviously mentioned, Fig. 2.

Prestraining was performed on a Mayes servo hydraulic testachine, under strain control, using an axial MTS extensometer

ig. 2. Microstructure and texture of TI834. A mixture of equiaxed and slightly elongarystallographic texture is of a transformation type with low intensity. Samples were alsignificant differences evident. (Texture maps produced over an area of 5 mm × 3 mm).

ig. 3. Microstructure and texture of Ti834 swaged material. The large amount of deformexture. (Texture maps produced over an area of 5 mm × 3 mm).

ngineering A 527 (2010) 6683–6689

with a gauge length of 12 mm. Specimens were loaded to totalstrains of −2%, 2% or 8% using a constant strain rate of 0.5%/s.They were then unloaded to zero load and removed from the testmachine prior to subsequent testing. The average modulus on load-ing into tension or compression was found to be 120 GPa, consistentwith published values [8].

Specimens from each prestrain batch were then tested underone of the following conditions:

(i) Ambient temperature (20 ◦C) creep testing at 950 MPa andunder constant load conditions. Loading rates were approx-imately 100 MPa/s undertaken on a servo hydraulic testmachine.

(ii) High temperature creep testing (600 ◦C) at 400 MPa and underconstant load conditions. Loading rates again were approx-imately 100 MPa/s undertaken on a servo hydraulic testmachine.

iii) Strain control fatigue under R = −1 conditions at a constantstrain rate of 0.5%/s until failure (classed as a 10% drop inload from the stabilised condition) using a trapezoidal 1-1-1-1waveform, according to BS7270 [9].

(iv) Stress relaxation for 3 h at a constant strain of 1%. These spec-imens were then cycled to failure under the same conditionsused in (iii).

To explore extreme prestraining a batch of the Ti834 also under-went swaging in which the bar was elongated by 108%. Fig. 3 showsthe microstructure and texture of the material. All tested specimenswere aligned with the direction of extrusion of the swaged bar.Where facetted areas have been quoted, results have been gainedthrough image analysis techniques of the fracture surface.

3. Results

Although the applied prestrains have been quoted as −2%, 2%and 8%, the amount of residual plastic prestrain is considerably less

ted primary alpha grains with a maximum diameter of 200 pm are evident. Theo taken for microstructure and texture analysis orthogonal to that shown, with no

ation is evident in the microstructure, and by the production of a well-defined fibre

M. Whittaker et al. / Materials Science and Engineering A 527 (2010) 6683–6689 6685

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Fig. 4. Loading curves for total prestrain of (a) 2% (b) −2

n each case. Assuming the material unloads elastically at the endf prestraining, the residual plastic strain in each case is approx-mately −1.25%, 1.25% and 7.25%.As such it is this residual plasticrestrain which is referred to throughout the paper. Although data

s available for the effect of prestrain on both the fatigue and creeproperties of titanium alloys [7,10], these are mainly based on sig-ificantly smaller levels of prestrain than have been considered inhis case.

Fig. 4 shows the stress–strain characteristics of the material dur-ng the application of the 1.25%, −1.25 and 7.25% prestrains. It islear that the yield stress is higher in compression than in tension.

Fig. 5 compares creep curves at 20 ◦C. The strains do not includetrain on loading, which differed significantly for the four levels ofrestrain. Fig. 6 shows these loading curves and it is clear that the.25% prestrain and particularly the 8% prestrain specimens strain

arden significantly. Because of prior loading in compression, the1.25% specimen displays a decrease in yield stress.

It is clear from Fig. 5 that prestraining causes a reduction in thetrain at failure, at all levels of prestrain. It is also noticeable that

ig. 5. Creep curves at 20 ◦C at an initial stress of 950 MPa. Increases in prestrainause a reduction in strain at failure, and it can be seen that the high dislocationensity in the 7.25% prestrain specimen results in a significant drop in creep strainate.

%. A higher yield stress can be observed in compression.

a significant reduction in the amount of primary creep occurs inthe prestrained specimens. The 7.25% prestrain specimen showsonly a small amount of primary creep and essentially no secondarycreep. This specimen was removed, unfailed, after 10 days withlittle evidence of creep. In all cases there is little or no evidence oftertiary creep. This is a consequence of the low temperature testing,with titanium alloys showing logarithmic type creep at ambienttemperature, and the fact that tests were conducted under constantload rather than constant stress, leading to rapid failure on necking.

Fig. 7 summarizes creep at 600 ◦C and under a constant loadcorresponding to an initial stress of 400 MPa. All specimens showa more substantial tertiary creep phase than was seen at 20 ◦C.

Significant differences are now clear between the data at 20and 600 ◦C. Perhaps the most easily recognisable is the fact thatthe 7.25% prestrain specimen, which showed essentially no creep

at 20 ◦C now shows the fastest rate, lowest strain at fracture andshortest time to fracture of all the specimens.

The 1.25% and −1.25% curves are quite similar although theonset of tertiary creep is earlier in the 1.25% specimen. A clear trend

Fig. 6. Loading curves for creep tests at 20 ◦C. A lower yield stress occurs in the−1.25% specimen due to earlier yield in compression on prestraining. Strain hard-ening is clearly evident in the 1.25% and 7.25% specimens.

6686 M. Whittaker et al. / Materials Science and Engineering A 527 (2010) 6683–6689

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interpreted by considering the generation of dislocations throughplastic deformation brought about by prestraining at room tem-perature. The presence of these dislocations can be inferred from

ig. 7. Creep curves at 600 ◦C. at an initial stress of 400 MPa. It is clear that increasesn prestrain level at this temperature cause a reduction in strain to failure andupture time and an increase in creep strain rate.

bserved is that at this temperature an increase in prestrain leadso a decrease in strain at fracture and time to fracture.

Fig. 8 shows a comparison of creep curves performed on thewaged material at 20 ◦C. The swaged material in each case haseen heat treated for 2 h at 450, 600 and 700 ◦C, respectively, beforeir cooling. It can be seen that although there is a larger amount ofrimary creep in the specimen heat treated at 700 ◦C, and that ithows a slight acceleration in creep rate, when compared with the% prestrain specimen, both the amount of creep strain and thereep rate are almost negligible. The situation is reversed howeverhen creep tests are performed on the swaged material at 600 ◦C,

ig. 9. In this case the swaged material shows a faster creep ratehan any of the prestrained material.

. Discussion

In order to evaluate the effect of different loading conditionsn facet formation, it is first necessary to isolate conditions whereacet formation is not promoted. In the current work, fatigue test-ng of as received (0% prestrain) specimens showed little evidencef facetting across the specimen. This is to be expected under fullyeversed strain control loading conditions, since the mean stress onhe specimen will be low and the reversible nature of the appliedtress means that dislocation build-ups are easily relieved. This pro-

ides a control condition, allowing for the isolation of the effectsf prestraining and stress relaxation on the fractographic featuresithin these specimens.

ig. 8. Creep curves for swaged material at 20 ◦C indicating that heat treatmentsp to 700 ◦C do not encourage significant recovery of dislocations, to allow lowemperature creep.

Fig. 9. Creep curves at 600 ◦C including swaged material. The swaged material fol-lows the same trend as the prestrained material, in that higher deformation levelsincrease creep rate at these temperatures.

The application of prestrain to specimens initially appeared to bea facet-forming process, with isolated facets being observed in both1.25% and −1.25% test specimens, which were subsequently cycledto failure under the strain control conditions described above.Image analysis techniques indicated that approximately 0.01% ofthe fracture surface was covered in facets in this case. Further-more, more extensive facetting (0.02% of the fracture surface) wasobserved in the 7.25% prestrain specimen, Fig. 10. Previous workhowever [11] has indicated that facet formation during monotonicloading is unlikely. Instead, when a period of stress relaxation isintroduced, the time dependent nature of the process means thatfacets form more readily. On closer inspection of Fig. 4(a) this hassignificant bearing on the current work. A small delay occurredbetween peak strain being achieved and unloading of the specimento zero load. Although this delay constituted only a few seconds, itis evident from the drop in stress at the end of the loading curvethat significant stress relaxation has occurred. This is more likely tobe the source of facet formation in these tests. At 7.25% prestrain,significantly more dislocations have been generated, and as suchmore facets form due to pile ups at grain boundaries.

The results from Fig. 5 for creep tests at room temperature can be

the strain hardening effect seen in the 1.25% and 7.25% prestrain

Fig. 10. Isolated quasi-cleavage facet in 7.25% prestrain specimen. The facets do notform during the prestrain loading, but during a brief period of stress relaxation priorto unloading. Specimen cycled to failure by R = −1 fatigue loading.

M. Whittaker et al. / Materials Science and Engineering A 527 (2010) 6683–6689 6687

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Baushinger effect (−1.25% prestrain) which cause the tests to reach

ig. 11. Creep strain rates in prestrained material. As would be expected during lowemperature (logarithmic) creep a continuously decaying primary phase ends withshort accelerating tertiary phase and failure.

pecimens on loading, Fig. 6. It is interesting first to note that aecaying primary phase is observed in each test, which is followed

n the 0%, 1.25% and −1.25% tests by a short tertiary phase beforenal failure. The 7.25% prestrain test was removed from test after0 days because of its extremely slow strain rate. Fig. 11 illustrateshat apart from a significant decrease in the 7.25% test, the pre-train had little effect on creep strain rate. It is clear, however, thathe prestrain had a significant effect on the strain at failure, withignificant reductions in the 1.25% and −1.25% tests. The disloca-ions generated by the prestrain process restrict further dislocation

ovement and final failure occurs at greatly reduced strains asecking occurs. The extremely high dislocation density generatedy 7.25% prestrain is such that almost all dislocation movement isifficult, and the strain rate falls dramatically.

It should also be acknowledged that some of the differencesn creep curve shape will be due to the change in yield stressf the prestrained material. Through considering the 0.05% prooftress (0.05�PS), the different behaviour of the primary stage ofhe creep tests can be identified, remembering that tests wereoaded to an initial stress of 950 MPa. In 0% (0.05�PS = 880 MPa)nd −1.25% (0.05�PS = 550 MPa) the yield stress is clearly exceedednd new dislocations are generated, leading to accelerated pri-ary rates. However in the 1.25% (0.05�PS = 1025 MPa) and 7.25%

0.05�PS = 1000 MPa) tests creep occurs by the movement of exist-ng dislocations and a reduced primary rate is evident.

The fracture surfaces of creep specimens showed extensivereas of quasi-cleavage facetting providing evidence that disloca-ion creep facilitates this facet formation, Fig. 12. Previous researchnto this alloy [12] has shown evidence of macrozone formationlarge areas of common crystallographic orientation) that may beelated to the cold dwell effects that are often evident under fatigueesting. EBSD analysis of the section studied however has not shownny evidence of such macrozones, Fig. 13. It can be seen that theres little ordered texture over lengths greater than approximately00 �m. It is acknowledged that the individual quasi-cleavageacets making up the larger faceted areas seen in the low tempera-ure creep tests correlates well with areas of common orientationhown in the crystallographic orientation map, and therefore it isikely that these faceted areas are of a common orientation. It is alsonteresting to note that these areas show approximately the same

imensions as the grains of transformed product or primary alpharains seen in the microstructure.

It is interesting to consider, however, the stress relaxation testserformed as part of the programme of work. Since the mechanisms

Fig. 12. Faceted areas in 0% prestrain cold creep specimen. The large amount ofplasticity in these tests allows for cleavage on a large scale.

involved in stress relaxation will be similar to the strain accumu-lation evident in the cold creep tests, it is expected that similarfractographic features would be evident. In fact, the cleavage facetsthat are produced show distinct differences to those produced inthe creep specimens. Rather than the large areas of facetting in thecreep specimens, facet areas are now small and isolated from eachother, with facets seeming to be related to primary alpha grainswhich have previously been reported to have basal or near basalorientations [3].

To illustrate the reason for the larger facetted areas in the creepspecimen, it is useful to consider the accumulation of plastic strainduring a stress relaxation test where elastic strain is exchanged forplastic strain

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and comparing these values to the levels of plasticity seen in thecreep tests, Fig. 5, it becomes clear that the significantly lower levelsof plastic strain accumulated in the stress relaxation tests limit thenumber of facets produced. Image analysis techniques have beenapplied to illustrate this point, with stress relaxation specimensshowing 0.5% of the fracture surface covered by facets, whereasthis value was as high as 3.8% for the creep specimens.

As described earlier, it is stress relaxation of this type, at theend of the prestrain period which generates facets, rather than theprestrain loading itself. Fig. 14 shows the stress relaxation charac-teristics of the four different prestrain levels. The vertical axis isplotted as normalized stress, i.e. stress divided by the peak stressreached during loading. This gives a clearer indication of the pro-cesses occurring during the period of relaxation, by removing theissues of strain hardening (1.25% and 7.25% prestrain) or reverse

various peak stresses [13]. Similar trends are seen to the room tem-perature creep tests, in that for the lower amounts of prestrain,little variation is seen to the as received material. However, for the7.25% prestrain specimen, the high dislocation density again pro-

6688 M. Whittaker et al. / Materials Science and Engineering A 527 (2010) 6683–6689

Fig. 13. Crystal orientation map of Ti834 specimen indicating th

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ig. 14. Stress relaxation of prestrained specimens, plotted as normalized stress, i.e.tress divided by peak stress achieved in that test. Significant differences are seenetween the 7.25% prestrain specimen and other prestrain levels.

ibits the movement of further dislocations, and as such little stress

elaxation is seen.

At 600 ◦C significantly different results are found in creep tests tohose performed at room temperature, Fig. 7. Clearly, increased pre-train now results in an increase in creep strain rate and a decrease

ig. 15. Ductile voids evident in Ti834 at 600 ◦C. This type of feature is typical ofigh temperature creep phenomena.

at no significant large areas of common orientation occur.

in time to rupture and strain at rupture. At these temperatures inTi834, due to the increase in thermal energy, dislocations are farmore mobile and processes such as climb and cross slip are pro-moted, resulting in a higher creep strain rate and hence reducedrupture time. It is clear that the mechanisms operating at low tem-perature are quite different to those at high temperature, a viewreinforced by the now prolonged tertiary period of deformation ineach specimen and extensive void formation at 600 ◦C.

The swaging process clearly has a marked effect on the mate-rial. The vertical axis of the micrograph in Fig. 3 is parallel to thedirection of elongation. It is clear that the primary alpha grains haveexperienced considerable deformation in this direction. The crys-tallographic texture of the material has also been severely affected.The as received material showed a very weak texture but that hasnow been replaced by a well-defined fibre texture. The (0 0 0 1)basal pole figure is taken from the end of the cylindrical bar andshows that basal planes lie in a fan shaped distribution, approxi-mately parallel to the length of the bar. This is consistent with otherwork [14]. Clearly such a dramatic change in structure impacts onthe mechanical properties and this was evident from two tensiletests. A dramatic drop in the modulus value was found with theswaged material having a value of approximately 100 GPa, com-pared with 120 GPa for the as received material. The change in themodulus is a consequence of the strong texture of the material. Pre-vious work [15] has shown that basal plane orientation can causedifferences in the modulus by as much as ±20 GPa and this is thecase here with loading near parallel to the majority of basal planes.

Also clear was a significant increase in the UTS of the mate-rial, from 1100 MPa in the as received material to approximately1350 MPa in the swaged material. This increase is a consequenceof the large dislocation density induced and associated with strainhardening.

The resultant high dislocation density led to an even slowercreep rate than in the 7.25% prestrain specimen. However, thematerial was heat treated to investigate any evidence of recovery.Fig. 8 shows the creep curves of the swaged material following heattreatments of 2 h at 450, 600 and 700 ◦C, respectively. The curve forthe 0% prestrain material is included for comparison purposes.

Following the 450 and 600 ◦C treatments the creep rate is stillessentially zero. At 700 ◦C there is a slight increase in the creeprate, particularly during primary creep. This suggests that at thistemperature a certain amount of recovery is occurring. It is likelythat had the time of the heat treatments been increased, further

increases in the creep rate would have been seen as the level ofrecovery increased.

Indeed when the swaged material was creep tested at 600 ◦C,Fig. 9, it shows a faster creep rate than any of the prestrained spec-imens, indicating that dislocation recovery at prolonged periods at

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M. Whittaker et al. / Materials Scienc

igh temperature is playing a major role. A number of other factorsay influence the creep rate at this temperature. Dynamic recrys-

allisation is a possibility under these conditions and may causeicrostructural variations that increase creep rate. It is clear that

xtensive ductile void formation, Fig. 15, occurs at this temperaturend is probably a consequence of enhanced dislocation mobility.he voids will certainly accentuate the onset of tertiary creep. Fur-hermore, facets formed at 20 ◦C during the prestrain process willtill exist in the material and will increase the rate of strain accumu-ation. These combined factors lead to a situation where an increasen prestrain clearly leads to a faster rate and reduced time to failure.

. Conclusions

Stress relaxation or creep at room temperature causes the for-mation of quasi-cleavage facets through dislocation piles ups atthe boundaries of suitably orientated grains. Creep leads to largerfaceted areas.Significant prestrain leads to a reduction in creep rate at roomtemperature as dislocation stress fields inhibit the movement offurther dislocations.At higher temperatures the increased availability of mechanismssuch as cross slip and climb mean that dislocation movement

occurs more easily and increased levels of prestrain lead to anincrease in creep rate but lower strain at failure.Swaged material behaves in a similar manner to heavily pre-strained material. However, material deformation through thisprocess can significantly alter material behaviour, and processes

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ngineering A 527 (2010) 6683–6689 6689

such as texture development and strain hardening require con-sideration for mechanical properties evaluation.

Acknowledgements

The authors would like to thank Rolls-Royce plc for their helpand input during the course of this work.

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[3] V. Sinha, M.J. Mills, J.C. Williams, Metall. Mater. Trans. A 37A (2006) 2015–2026.

[4] V. Hasija, S. Ghosh, M.J. Mills, D.S. Joseph, Acta Mater. 51 (2003) 4533–4549.

[5] F.P.E. Dunne, D. Rugg, A. Walker, Int. J. Plasticity 23 (2007) 1061–1083.[6] E.E. Sackett, L. Germain, M.R. Bache, Int. J. Fatigue 29 (2007) 2015–

2021.[7] B.C. Odegard, A.W. Thompson, Metall. Mater. Trans. B 5 (1973) 1207–

1213.[8] Timet website, www.timet.com.[9] BS7270, British Standards Institution, 1990.10] D.B. Lanning, T. Nicholas, G.K. Haritos, Mech. Mater. 34 (2002) 127–134.

12] L. Germain, M. Bache, Proceedings of 11th International Conference on Tita-nium, 2007, pp. 953–956.

13] M.T. Whittaker, W.J. Evans, Int. J. Fatigue 31 (2009) 1751–1757.14] G. Lutjering, Mater. Sci. Eng. A 243 (1998) 32–45.15] A.W. Bowen, Mater. Sci. Eng. 29 (1977) 19–28.