Low cycle fatigue behaviour of Ti alloy Timetal 834 at 873 K

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International Journal of Fatigue 29 (2007) 843–851

JournalofFatigue

Low cycle fatigue behaviour of Ti alloy Timetal 834 at 873 K

Nidhi Singh, Gouthama 1, Vakil Singh *

Department of Metallurgical Engineering, Centre of Advanced Study, Institute of Technology, Banaras Hindu University, Varanasi 221 005, India

Received 15 February 2006; received in revised form 29 August 2006; accepted 10 September 2006Available online 24 October 2006

Abstract

Low cycle fatigue behaviour of a near a titanium alloy Timetal 834 was studied, for a bimodal microstructure with �14 vol% of pri-mary a in the matrix of transformed b at different total strain amplitudes (Det/2) ranging from ±0.75% to ±1.25%, at 873 K. Cyclic soft-ening was exhibited at all the strain amplitudes, however, the rate as well as degree of softening was found to decrease with increase instrain amplitude. In general the rate of softening was high at all the strain amplitudes during the initial �10 cycles, however, it decreasedin subsequent cycles and shortly the cyclic stress became almost stable and remained so for the major fraction of life, until the onset ofmacrocrack propagation to cause rapid drop in the stress leading to fracture. Coffin–Manson (C–M) plot was found to exhibit a singleslope at 873 K, in contrast to dual slope at room temperature. Also fatigue life at 873 K was considerably higher than that at room tem-perature. The results are discussed in terms of increased homogeneity of deformation and higher stability of Ti3Al precipitates at theelevated temperature of 873 K than that at room temperature.� 2006 Elsevier Ltd. All rights reserved.

Keywords: Titanium alloy Timetal 834; Coffin–Manson relationship; Low cycle fatigue; Cyclic stress response; Planar slip

1. Introduction

Ti alloy Timetal 834 is a near- a alloy with high specificstrength, designed for high temperature application up to873 K as compressor disk and blade of gas turbines ofadvanced jet engines, to replace the heavy nickel base superalloys and increase the pay load. The b transus temperatureof this alloy is 1333 K [1]. It develops a wide variety ofmicrostructures, depending upon the solution treatmentand the subsequent rate of cooling. It has been establishedthat a good combination of creep strength, low cycle fati-gue properties and resistance against crack propagation isexhibited by a bimodal microstructure with �15 vol% ofprimary a in the matrix of fine grained transformed b withlamellar structure [1].

Recently the authors have studied low cycle fatigue(LCF) behaviour of this alloy in the bimodal microstruc-

0142-1123/$ - see front matter � 2006 Elsevier Ltd. All rights reserved.

doi:10.1016/j.ijfatigue.2006.09.006

* Corresponding author. Tel.: +91 542 2575403; fax: +91 542 2369478.E-mail address: [email protected] (V. Singh).

1 Presently with Department of Materials and Metallurgical Engineer-ing, Indian Institute of Technology, Kanpur 208016, India.

ture with �14 vol% of primary a in the matrix of trans-formed b, at different total strain amplitudes (Det/2) from±0.75% to ±1.7%, at room temperature [2]. It wasobserved that there was continuous cyclic softening, start-ing from the first cycle to the last cycle of failure, at all thestrain amplitudes referred to above, and the tendency forsoftening increased with strain amplitude. Coffin–Mansonrelationship (Dep/2 vs 2 Nf), however, was found to displaydual slope behaviour like that observed in several other fer-rous and non-ferrous alloys [3]. Several investigations havebeen carried out on one or other aspect of LCF behaviourof the alloy 834 at elevated temperatures; like the influenceof prior creep deformation at 873 K on cyclic yield stressand fatigue life [4], influence of oxidation and creep oncrack initiation and fatigue life at 773 K and 903 K [5],crack propagation at 873 K under creep-fatigue condition[6], cyclic stress response at a constant plastic strain ampli-tude (Dep/2) of ±0.2% over a wide range of temperaturefrom 623 to 923 K [7] and the effect of aging treatmenton cyclic stress response at a total strain range (Det) of0.97% at 773 K [8]. Maier et al. [9] have recently developeda macro crack propagation model to predict thermo

mailto:[email protected]

Fig. 1. Optical micrograph showing dual phase microstructure in the(a + b) ST-AC-A condition. The bright phase is primary a.

844 N. Singh et al. / International Journal of Fatigue 29 (2007) 843–851

mechanical fatigue (TMF) life from the isothermal data.However, no systematic investigation has been reportedon isothermal LCF behaviour of this alloy at its servicetemperature of 873 K over a wide range of strainamplitudes.

The present investigation was therefore undertaken tostudy cyclic stress response and LCF behaviour of the alloy834 at its service temperature of 873 K with the same bimo-dal micro-structure condition of �14 vol% primary a in thematrix of transformed b as that studied at room tempera-ture [2], over a wide range of total strain amplitudes (Det/2) from ±0.75% to ±1.25% and to compare with thatobserved earlier at the room temperature [2].

2. Experimental

Alloy 834 was procured from M/s. Timet, UK, in theform of rods of 20 mm diameter. It contained by weightpercent 5.8Al–4.0Sn–3.7Zr–0.7Nb–0.5Mo–0.34Si–0.05C–0.105O–0.004H and balance Ti. The rods had been solutionannealed in the a + b phase field at 1293 K for 2 h cooledin air and subsequently aged at 973 K for 2 h followed bycooling in air. This heat treatment is designated as(a + b) ST-AC-A. Microstructure of the material wasexamined under optical microscope. The specimens weremechanically polished and etched with a solution of 10%HF (specific gravity 1.13, purity 39–43% redistilled), 5%HNO3 (specific gravity 1.42, purity 69.0–70.5%) and 85%H2O (vol%). Volume fraction of the primary a phase wasdetermined by VIDS-III semiautomatic image analyzer.

Fine microstructural details and the deformation behav-iour of the material under LCF was studied by transmis-sion electron microscope, at 160 kV. TEM foils wereprepared by electrolytic thinning. Transverse slices of�100 lm thickness were sectioned from the as heat treatedrod and the gauge section of the LCF tested specimensusing a slow speed isomet cutter of thin disc shape. Theslices were thinned down to �50 lm by mechanical polish-ing and discs of 3 mm diameter were punched from them.TEM foils were prepared by electrolytic thinning of thesediscs in the electrolyte containing 59% methanol (specificgravity 10, purity 32%), 35% n-butanol (specific gravity:0.809–0.811) and 6% perchloric acid (specific gravity 1.67,purity 71–73%). Cooled to 223 K in a twin jet electropo-lisher at 12.5 V. TEM foils were examined in JEOL 200CX transmission electron microscope at 160 kV.

Cylindrical tensile specimens with gauge length anddiameter of 16 mm and 4.5 mm respectively, weremachined from quarter sections resulting from double lon-gitudinal sectioning of 100 mm long blanks of 20 mmdiameter. Tensile tests were performed using an Instronof 50 kN capacity, at a nominal strain rate of5.6 · 10�4 s�1. LCF tests were performed using cylindricalspecimen of gauge length and diameter of 15 mm and5.5 mm respectively, shoulder radii of 25 mm, and threadedends of 35 mm length and 12 mm diameter. LCF test spec-imens were machined from blanks of 100 mm length and

20 mm diameter. Gauge section of machined LCF speci-mens was mechanically polished with emery papers of 2/0to 4/0 grades and finally by sylvet cloth, soaked with sus-pension of 0.3 lm alumina powder in water, to reducethe roughness and minimize the surface effect, if any, onfatigue life. LCF tests were conducted on servohydraulicMTS of 50 kN capacity in total strain control mode bymounting high temperature MTS axial extensometer of12 mm gauge length (Model 632), at 873 K, in air. Testtemperature of 873 ± 2 K was achieved by a three zoneelectric resistance heating furnace. LCF tests were con-ducted under completely reversed loading (R = �1) withtriangular waveform at a constant cyclic frequency of0.2 Hz. Once the test temperature of 873 K was attainedthe samples were exposed at this temperature for 15 minto homogenize the temperature through out gauge section,prior to onset of the test. Fracture surfaces of the testedsamples were examined by JEOL 810 scanning electronmicroscope.

3. Results

3.1. Microstructure

Optical microstructure of the alloy Timetal 834 in the(a + b) ST-AC-A condition is shown in Fig. 1. It is obviousthat it is a duplex type of microstructure consisting of theprimary a phase and transformed b. The mean interceptlength of primary a and transformed b grains was deter-mined as 46 and 140 lm respectively and volume fractionof the primary a phase was estimated to be �14%. Finerdetails of the microstructure like platelets of secondary aand b retained between a platelets along with precipitatesof silicides are shown by TEM micrograph in Fig. 2. Sili-cides in the alloy 834 in the a + b solution treated and agedcondition have been characterized as (TiZr)6Si3 [10]. Lat-tice parameters of the titanium–zirconium silicides have

Fig. 2. Transmission electron micrographs showing microstructure ofalloy 834 in (a + b) ST-AC-A condition: (a) bright field TEM micrographshowing precipitates of silicide at inter platelet boundary of transformed aand (b) selected area diffraction pattern showing superlattic spots. Zoneaxis: ½01�11�.

Fig. 3. Optical micrographs showing longitudinal sections of the tensilespecimens tested at: (a) room temperature and (b) 873 K.

N. Singh et al. / International Journal of Fatigue 29 (2007) 843–851 845

earlier been estimated as a = 0.703 ± 0.004 nm andc = 0.360 ± 0.015 nm [11].

3.2. Tensile behaviour

Tensile properties of the alloy Timetal 834 in the (a + b)ST-AC-A condition both at room temperature and at873 K are presented in Table 1. It may be seen that thereis marked influence of the test temperature on tensile prop-erties in particular, the ductility of the alloy. Ductility isconsiderably higher at 873 K than that at room tempera-ture. While there is small increase in the uniform elonga-tion (eu), there is an appreciable increase in the totalelongation (ef) at 873 K. Also there is a marked increase

Table 1Tensile properties of the alloy Timetal 834 in the (a + b) ST-AC-A condition,

Temperature ry (MPa) rUTS (MPa) rUTS/ry

RT 945 1012 1.07873 K 528.5 628 1.19

in other ductility parameters, viz. reduction in area (RA)and true strain to fracture (ef) at 873 K. Generally thedegree of work hardening (rUTS/ry) decreases with increasein temperature, however, it may be seen that work harden-ing is relatively higher at 873 K than that at roomtemperature.

Microstructural features from longitudinal sections ofthe tensile specimens tested at room temperature and873 K are shown by optical micrographs in Fig. 3a and brespectively. It may be seen that there are transverse cracksmostly confined in the primary a phase, at room tempera-ture, whereas there are no cracks in the primary a phase at873 K, rather there are voids mostly associated with inter-

at room temperature and 873 K

eu (%) ef (%) RA (%) ef

7.9 14.5 14.1 0.15211.5 26.3 45 0.263


face of the primary a and transformed b at 873 K. Voidsmay be seen also within some transformed b grains. Itmay further be noted that volume fraction of interfacialvoids/cracks in the specimen tested at 873 K is consider-ably higher than the transgranular cracks in the primarya phase at room temperature, however, in spite of that duc-tility of the alloy is considerably higher at 873 K than thatat room temperature. The cracks and voids were not con-fined only to necked region but were spread to �2 mmon either side of the neck.

Fig. 4. Cyclic stress response of the alloy 834 in the (a + b) ST-AC-Acondition tested at: (a) room temperature and (b) 873 K.

Table 2LCF data for the (a + b) ST-AC-A condition at 873 K

±Det/2 (%) ±Dep/2 (%) ±Dec/2 (%)

N1 Nf/2 Nf N1 Nf/2

0.75 0.189 0.262 0.33 0.547 0.4850.80 0.252 0.309 0.359 0.545 0.4900.90 0.310 0.359 0.449 0.581 0.5411.00 0.409 0.447 0.531 0.583 0.5511.10 0.493 0.579 0.576 0.600 0.5231.25 0.619 0.621 0.721 0.613 0.624

3.3. LCF behaviour

3.3.1. Stress response

The variation of average cyclic stress amplitude (Dr/2)with number of cycles at room temperature and 873 K isshown in Fig. 4a and b respectively. It may be seen thatthere is cyclic softening at both the temperatures, however,there is a marked difference in the behaviour of softening atthe two temperatures. While at room temperature there is acontinuous softening starting from the first cycle to the lastcycle of failure particularly at higher strain amplitudes(Det/2 P ±1%) at 873 K there is rapid softening duringthe 10 initial cycles and the rate of softening becomes rela-tively slower in the subsequent cycles. The cyclic stressremains stabilized for the major fraction of fatigue life atall the strain amplitudes. Further, both the rate as well asdegree of softening is higher at room temperature than at873 K at Det/2 P ±1% and both of these increase withincrease in strain amplitude at room temperature. On theother hand, there is a opposite trend at 873 K and thedegree of softening decreases with increase in strain ampli-tude. LCF data at 873 K in terms of elastic/plastic strainamplitudes and the stress amplitude at first cycle (N1), halflife (Nf/2) and the last cycle to failure (Nf) for all the totalstrain amplitudes investigated are presented in Table 2. Thecyclic and tensile stress strain curves are shown in Fig. 5. Itmay be seen that level of the cyclic stress–strain curve islower than that of the monotonic one and thus there is cyc-lic softening at 873 K. The value of the cyclic work harden-ing exponent (n 0) determined from the plot of true cyclicstress and true cyclic plastic strain at half life was foundto be 0.254.

3.3.2. Fatigue life

The dependence of fatigue life on strain amplitude wasanalyzed using the Coffin–Manson relationship betweenthe plastic strain amplitude (Dep/2) and number of reversalsto failure (2Nf).

Dep=2 ¼ e0fð2N fÞc ð1Þwhere e0f and c are fatigue ductility coefficient and exponentrespectively. The C–M plots both for room temperatureand 873 K are shown in Fig. 6. It may be seen that whilethere is a dual slope in the C–M plot for the tests conductedat room temperature, there is only a single slope for the

±Dr/2 (MPa) Fatigue life (Nf) (Cycles)

Nf N1 Nf/2 Nf

0.413 519.5 461.8 392.5 16820.436 518.4 465.6 414.8 12100.455 539.9 502.2 423.1 9090.469 547.0 518.1 440.8 6860.524 557.7 538.2 486.9 5350.529 564.6 574.3 486.6 270

100 1000

10-3

10-2

RT 873K

Pla

stic

Str

ain

Am

plitu

de (

Δεp/

2)

Reversals to Failure (2Nf)

Fig. 6. Coffin–Manson plots for LCF test at room temperature and873 K.

Fig. 7. TEM micrograph showing widely spaced slip traces and intersec-tion of slip traces, in the LCF specimen tested at 873 K and Det/2: ±0.80%.

Fig. 5. Monotonic and cyclic stress strain curves at 873 K.


tests conducted at 873 K. The variation of fatigue life withelastic strain amplitude (Dee/2) was analyzed using the Bas-quin equation:

Dee=2 ¼ r0f=Eð2N fÞb ð2Þwhere r0f and b are fatigue strength coefficient and expo-nent, respectively and E is the modulus of elasticity. The

Table 3Cyclic work hardening and LCF parameters for the (a + b) ST-AC-A conditi

Temperature n0 bexp bcal cexp

RT 0.05 �0.036 �0.04 �0.36873 K 0.254 �0.1362 �0.112 �0.75

values of the important fatigue parameters are presentedin Table 3. The value of the fatigue ductility coefficient(e0f ) at 873 K is markedly higher than that at room temper-ature. The transition fatigue life (Nt), was determined fromthe intersection of Dee/2 vs 2Nf and Dep/2 vs. 2Nf lines. Itmay be seen that transition fatigue life at 873 K is signifi-cantly higher than that at room temperature.

3.3.3. Deformation behaviour

Deformation behaviour of the alloy 834 in LCF at873 K at different total strain amplitudes (Det/2) was stud-ied by transmission electron microscopy of the fatigue-tested samples. TEM micrograph in Fig. 7 shows slip tracesin the specimen tested at low total strain at amplitude (Det/2) of ±0.8%. It may be seen that there is intersection of sliptraces. Electron micrographs in Fig. 8 show deformationbehaviour of the specimen tested at a relatively higherstrain amplitude (Det/2: ±1%). There is a marked increasein slip activity (Fig. 8a). Further, there is easy transfer ofslip across the interfacial boundary of transformed a plate-lets and it is obvious that the small volumes of b, retainedbetween the a platelets, do not offer any resistance againstthe glide of dislocations from one platelet of a to the otherone (Fig. 8b). In fact, regions of retained b were notablydisintegrated due to glide of dislocations on multiple slipplanes within a platelets. Disintegration of retained b atinter platelet boundary of a under the action of cyclic load-ing may clearly be seen in Fig. 8c. The features with dark

on at room temperature and 873 K

ccal e0f r0f (MPa) Nt (Cycles) r0ys (MPa)

�.8 0.049 1200 60 850�0.440 1.0 1316 700 455

Fig. 8. TEM micrographs of the fatigue sample tested at 873 K, at Det/2:±1.0% showing: (a) high activity of planar slip, (b) ease of transfer ofplanar slip across the inter platelet boundary of a and (c) disintegration ofretained b, at the inter platelet boundary of a.


contrast in Fig. 8c are fragmented pieces of retained b. Itmay be seen that there is no pile up of dislocations at theinterface. Deformation behaviour of the specimen tested

in LCF at 873 K at the highest strain amplitude, Det/2:±1.25%, is shown in Fig. 9. It is obvious that in the regionin which orientation of a grains is unfavourable, there isactivation of only one set of slip system even at such a highstrain amplitude, and this is essentially planar in nature(Fig. 9a). SAD patterns in Fig. 9b clearly show Ti3Alreflections with underlined indices. The orientationbetween primary a and Ti3Al was established as:½01�11�ap==½01�12�Ti3Al. However, at favourable orientationof a platelet there is activation of more than one slip sys-tems (Fig. 9c). The slip traces may be seen to run fromtop to bottom across the a platelet boundaries withoutany difficulty. TEM micrograph in Fig. 9d shows completedisintegration of retained b into very small pieces of a fewnanometers. Also the interplatelet boundary of a isdisintegrated.

3.3.4. Fracture behaviour

Deformation of the material under LCF at room temper-ature and 873 K was analyzed from microstrctures of longi-tudinal sections of the tested samples (Fig. 10). As shown byarrows in the optical micrograph in Fig. 10a deformation atroom temperature has been quite heterogeneous and con-fined in narrow slip bands in the primary a phase. On theother hand the deformation at 873 K is much more homo-geneous and no such features are seen. Rather, there areinterfacial voids/cracks distributed throughout the longitu-dinal section (Fig. 10b). Surface cracks at elevated temper-ature are due to environmental effects. Fracture behaviourof the specimens tested in LCF at room temperature and873 K is shown in Fig. 11. While the facture surface of thespecimen tested at room temperature is quite rough(Fig. 11a), distinct fatigue striations may be seen in caseof the specimen tested at 873 K (Fig. 11b). The striationsin Fig. 11b are quite close to the crack initiation site.

4. Discussion

The cyclic stress response of the alloy 834 at 873 K, exhib-iting softening during the initial �10% of fatigue life, fol-lowed by stability of stress for the major fraction offatigue life, in the present investigation, is similar to thatreported earlier by Propotzky et al. [7] at the same test tem-perature of 873 K, at a plastic strain amplitude (Dep/2) of±0.2%. They have attributed the observed stress responseto three factors: (1) planar mode of dislocation slip, (2)coarsening of silicides, and (3) disintegration of lamellarboundaries of secondary a. Planarity of slip leads to shear-ing of the coherent ordered Ti3Al precipitates. The coarsen-ing of silicides, however, does not cause any strengthening asthe silicides are already totally incoherent to matrix; rather itleads to weakening of the material because of depletion ofsilicon from the matrix. It is important to mention here thatcoarsening of silicides does not cause depletion of only Sibut also Zr, which is an important solid solution strength-ener. The third factor, the disintegration of interplateletboundaries of secondary a and retained b also reduces the

Fig. 9. TEM micrographs of the sample tested in LCF, at 873 K, at Det/2: ±1.25% showing: (a) activity of only one set of planar slip due to unfavourableorientation of the a grain, (b) selected area diffraction pattern wowing Ti3Al reflections (indices underlining). Zone axis ½01�11�, (c) operation of more thanone slip system and (d) complete disintegration of inter platelet boundary and also retained b due to cyclic loading at high strain amplitudes.


resistance of the material against the transfer of planar slipfrom one platelet of a to the other adjacent one.

Thus, it is obvious that all these three factors, mentionedabove, lead to only softening of the material. Softening,however, was found to be limited only to the initial stageof cycling, up to less than 10% of the fatigue life. Thuswhile the initial softening of the material could be under-stood in the light of the above three factors, the constancyof the cyclic stress response observed over the major frac-tion of fatigue life (>90%), in the present investigation, can-not be understood in terms of the only above three factors,suggested by Propotzky et al. [7]. The observed stability ofthe cyclic stress response could either be due to a change inthe mode of interaction of the dislocations and the orderedTi3Al precipitates, from shearing to by passing, or fromactivation of additional hardening processes to counteractthe effect of softening resulting from the above three pro-cesses. It is evident from the TEM micrographs of theLCF tested samples that the mode of deformation at873 K remains essentially planar dislocation slip (Figs. 7–9) thus the observed stability of the stress response cannotbe attributed to change in the mode of interaction between

the dislocations and Ti3Al precipitates, from shearing tolooping. Therefore, there is much likelihood of activationof additional hardening processes apart from the usualwork hardening, to counteract the effect of softening.

It has been established in our earlier investigation thatcyclic softening in the alloy 834, in the (a + b) ST-AC-Acondition, at room temperature, results essentially fromshearing of the ordered Ti3Al precipitates [2]. Further,the tendency for softening, particularly at Det/2 P ±1.0%, may be seen to be much higher at room tem-perature than that at 873 K (Fig. 1). Thus it is obvious thatsoftening at 873 K resulting from the combined effect of theabove three processes is less effective than that resultingfrom the single process of shearing of the ordered Ti3Alprecipitates, at room temperature. The relatively less effec-tiveness of shearing of the Ti3Al precipitates on softeningat 873 K may be attributed to partial recovery of order inthe disordered region resulting from the leading disloca-tions, like that observed in the age hardening Ti–Al alloy[11]. Since the dislocations in the planar slip bands result-ing from LCF at 873 K are mostly unpaired, it may beinferred that there would have been partial recovery of

Fig. 10. Optical micrographs showing longitudinal section and fracturebehaviour of the LCF specimen, tested at Det/2: ±1.25%: (a) roomtemperature and (b) 873 K.

Fig. 11. SEM fractographs showing fracture behaviour in LCF at: (a)room temperature and (b) 873 K.


order in the sheared Ti3Al precipitates at 873 K and thiscould have led to hardening [12].

Further, since the LCF tests were carried out in labo-ratory air, absorption of oxygen and nitrogen from theatmosphere would have caused solid solution strengthen-ing, up to the depth of their absorption. The enrichmentof oxygen and nitrogen in this alloy above its Al equiva-lence would have also led to precipitation of Ti3Al in thematrix [13]. LCF testing at 873 K would also lead togrowth of the preexisting precipitates of Ti3Al throughvacancy generated during fatigue testing. Both, precipita-tion of fresh Ti3Al precipitates and the growth of theexisting Ti3Al precipitates would cause strengthening ofthe material as Ti3Al precipitates remain coherent evenup to 120 nm [14]. However, depletion of Al and Zr fromthe matrix would result in softening of the material asthey are strong solid solution strengtheners. The stabilityof cyclic stress at 873 K may thus be attributed to the bal-ance between the various softening and hardening effectsresulting from the different processes referred to above.

It is evident from TEM micrographs of the LCF sampletested at 873 K that even at the low strain amplitude(Det/2) of ±0.8%, there is intersection of planar slip tracesand the slip traces are closely spaced in some regions(Fig. 7). Thus, the deformation is much more homoge-neous at 873 K even at the low strain amplitudes. Itmay be noted that in sharp contrast deformation washighly heterogeneous at low strain amplitudes at roomtemperature and it remained localized only in a fewwidely spaced planar slip bands and enhanced the processof crack initiation [2]. The spacing between the slip bandsdecreases with increase in strain amplitudes (Figs. 8 and9). The increase in the tendency of intersection of slipbands with increase in strain amplitude is quite obviousfrom Fig. 8b. Thus it may be seen that at 873 K, theapplied cyclic strain is distributed in relatively more num-ber of slip bands than that at room temperature andhence the cyclic strain in the individual slip bands at873 K is relatively less than that at room temperature.Further, since part of the applied strain at 873 K is


accommodated also by the interfacial processes of voidformation and cracking at the interface of primary aand transformed b, only the remaining part of the appliedstrain needs to be accommodated by the usual process ofdislocation slip. Thus at 873 K, the effective strain in theindividual slip band is further reduced and the process offatigue crack initiation, resulting from planar slip at sur-face, is further delayed and fatigue life is increased.

The elimination of bilinearity from the C–M plot at873 K may also be understood in terms of increased homo-geneity of deformation, combined with partial recovery ofordering in the sheared Ti3Al precipitates and solid solu-tion strengthening due to oxygen and nitrogen absorption.Thus the process of fatigue crack initiation, even at lowstrain amplitudes, is delayed and fatigue life is increased.This results in elimination of dual slope from the C–M rela-tionship at 873 K.

5. Conclusions

Following conclusions are drawn from the presentinvestigation:

1. Alloy Timetal 834, in the (a + b) ST-AC-A heat treatedcondition, exhibits rapid cyclic softening in the initialstage (<10% of fatigue), and saturation of stress forthe major fraction of fatigue life (>90%), at 873 K andthe total strain amplitude (Det/2) from ±0.70% to±1.25%. The tendency for softening, however, decreaseswith the increase in strain amplitude. This behaviour isconsiderably different from that observed at room tem-perature, where continuous softening occurs withoutany saturation and the tendency for softening increaseswith increase in the strain amplitude.

2. At 873 K, initial softening occurs due to planarity ofslip, shearing of the ordered Ti3Al precipitates andcoarsening of silicides. The saturation in cyclic stressoccurs from the balance between the softening and hard-ening effects. Strengthening occurs due to partial recov-ery of ordering in the disordered Ti3Al precipitates,

sheared by leading dislocations; solid solution hardeningof the matrix by oxygen and nitrogen enrichment andconsequent precipitation of Ti3Al precipitates.

3. In contrast to the dual slope in C–M plot at room tem-perature, only a single slope is exhibited at 873 K. Alsofatigue life is considerably higher at 873 K than that atroom temperature.

4. The above effects are attributed to an increase in homoge-neity of deformation at 873 K, less strain in individual slipbands and consequent delay in the process of fatiguecrack initiation, even at lower strain amplitudes. Thusfatigue life at 873 K is enhanced and is not reduced evenat low strain amplitudes unlike that at room temperature.

References

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[2] Singh Nidhi, Gauthama, Singh V. Mater Sci Eng A 2002;325:324.[3] Radhakrishnan VM. Int J Fatigue 1992;14:305.[4] Hardt S, Maier HJ, Christ HJ. Int J Fatigue 1999;21:779.[5] Hardy MC. In: Rie KT, Portella PD, editors. Low cycle fatigue and

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[8] Kestler H, Mugharabi H, Renner H. In: Blenkinshop PA, Evans WJ,Flower HM, editors. Titanium 95, science and technology. London:The Institute of Materials; 1996. p. 1171.

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Low cycle fatigue behaviour of Ti alloy Timetal 834 at 873 K