Materials Science & Engineering A 576 (2013) 217–221

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Materials Science & Engineering A

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Determination of critical strain for rapid crack growth during tensiledeformation in aluminide coated near-α titanium alloy usinginfrared thermography

Sony Punnose n, Amretendu Mukhopadhyay, Rajdeep Sarkar, Zafir Alam, Dipak Das, Vikas KumarDefence Metallurgical Research Laboratory, Kanchanbagh, Hyderabad 500058, India

a r t i c l e i n f o

Article history:Received 27 November 2012Received in revised form12 March 2013Accepted 26 March 2013Available online 10 April 2013

Keywords:Tensile deformationFractureProtective coatingTitanium alloyThermal analysis

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esponding author. Tel.: +91 40 24586786; faxail addresses: snbyrec@yahoo.co.in, sony@dmr

a b s t r a c t

Determination of strain for initiation of rapid crack growth is vital for designing coated components foraerospace applications. Knowledge of such strain is useful for prevention of catastrophic failure of coatedcomponents. In the present study this critical strain has been determined during tensile deformation ofaluminide coated near-α titanium alloy using infrared thermography. A single step on-line method ofdetermination of stored energy change as a function of true plastic strain, that incorporates conductionheat loss correction in a simple way, has been used to determine such strain level. It is shown thatbeyond a strain level the crack ensemble that form becomes unstable and further straining leads to rapidcrack growth in the coating that penetrates into the substrate material. This manifests in anunprecedented trend in the stored energy change that has been identified with the strain for inceptionof rapid crack growth.

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1. Introduction

Near α-Ti alloys such as IMI-834 and Timetal 1100 are widelyused for the fabrication of compressor parts of advanced gasturbine engines. However, the use of these alloys is limited to atemperature of about 600 1C because of their poor oxidationresistance [1]. Oxidation causes the formation of a brittle α casingwhich adversely affects the mechanical properties of these alloys[1,2]. Increasing the high temperature use of near α-Ti alloys isbeing attempted by application of oxidation resistance coatingssuch as TiAlN, Al3Ti type aluminides, silicides and MCrAlY typecoatings [3–5]. Al3Ti-based diffusion aluminide coatings have beenreported to greatly enhance the oxidation resistance of IMI-834alloy up to 800 1C [3]. Despite their good oxidation resistance,Al3Ti coatings are inherently brittle and develop through-thickness cracks during coating formation as well as during cyclicoxidation exposure [5–7]. Furthermore, due to the differences inmechanical properties of the coating and base material, strainmismatch arises between the coating and the substrate duringservice that causes development of cracks in the coating. Thecoating cracks can potentially extend into the substrate and causepremature/catastrophic failure of the coated components. In thiscontext, it is important to understand the mechanism of crackformation in the coating and also to determine the critical strain, if

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any, beyond which cracks in the coating can easily grow into thebase material with minimal deformation. Determination of suchcritical strain levels is often difficult because of the absence of anyreliable technique to do so. In the present study, using an onlineinfrared thermography (IRT) technique the above purpose hasbeen examined.

The process of energy storage and heat release in metals duringdeformation has been widely studied since the work of Farren andTaylor [8]. Several experimental studies [9,10] have shown that theprocess of energy accumulation in metals depends on variousfactors. The energy storage in a material during deformation canbe determined by studying the heat dissipation process with thehelp of IRT technique. Oliferuk et al. [10–12] have studied theprocess of energy storage during tensile deformation in case ofaustenitic steel and armco iron using IRT technique. In the presentstudy, the same technique has been used to identify the level ofstrain at which an ensemble of cracks forms during tensiledeformation of a near-α Ti alloy coated with Al3Ti. The ensembleof cracks reaches a critical density at this strain beyond whichtheir rapid growth takes place. The results from the coated speci-mens have been compared with those from the uncoatedspecimens.

2. Experimental

A near-α Ti alloy of nominal composition (wt%) Ti–5.5Al–4Sn–4Zr–0.3Mo–1Nb–0.5Si–0.06C was solution annealed at 1020 1C for

S. Punnose et al. / Materials Science & Engineering A 576 (2013) 217–221218

2 h followed by oil quenching. The alloy was subsequently aged at700 1C for 2 h followed by air cooling. The heat treatment givesrise to a bi-modal microstructure consisting of about 15 vol %primary α phase in a transformed β matrix. Two sets of flat tensilespecimens having 18 mm gauge length were prepared from theheat treated rod of 20 mm diameter. The first set of specimens wastensile tested in uncoated condition. The second set of specimenswas coated with 50 μm thick Al3Ti coating prior to tensile testing.The coating was applied using pack aluminising method. Thepowder pack used consisted of (wt%) 15Al as the aluminiumsource, 2NH4Cl as activator and balance Al2O3 fillers. The coatingprocess was carried out at 800 1C for 4 h in an argon atmosphere.Tensile testing of the uncoated and coated specimens was carriedout at room temperature using a computer controlled INSTRON5500R testing machine at a crosshead speed of 1.0 mm/min.Interrupted tensile tests were carried out for four strain levels(true plastic), namely 0.18%, 0.5%, 0.8% and 1.2%. A few specimenswere also tested up to fracture.

An infrared camera was used to map temporal and spatialtemperature evolution along the gauge length of the specimensduring tensile deformation. The spectral range of the camera was3–5 μm. Image acquisition was done at a frequency of 200 Hz.Thermal resolution of the camera was 20 mK at 298 K. The gaugelength of the specimens was painted black to obtain uniformemissivity. From the thermal evolution pattern, the area of highestthermal activity along the gauge length was marked in all speci-mens. Recorded temperature was averaged over the area of thegauge where highest thermal activity was noted.

Using the principle of energy conservation, the energy balanceduring tensile deformation can be written as in Eq. (1) givenbelow [13]:

ρc∂T∂t

−k∂2T∂x2

¼ βs∂εp

∂t−αET

∂εe

∂tð1Þ

The term x in the above equation represents the distancemeasured along gauge length and T the absolute temperature.The term s, εp and εe are the components of stress, plastic strainand elastic strain respectively. Further, ρ, c, k, and E are density,specific heat, thermal conductivity and Young's modulus of thematerial respectively. During an irreversible plastic deformationprocess, certain fraction of the expended energy dissipates as heatwhile the balance is retained in the material as stored energy. Theterm β in the above equation is the ratio of the rate at whichenergy is being expended during deformation to the rate ofgeneration of heat generated due to plastic deformation.

Fig. 1. Average temperature change as a function of true plastic

To estimate the heat loss due to thermal conduction, ð∂2T=∂x2Þas function of time was obtained from a sequence of thermograms.The total heat loss was then calculated by numerical integration.After correcting for conduction heat loss and thermo-elastic cool-ing, the dissipated energy during tensile deformation was calcu-lated assuming a constant heat capacity value for the base material(0.525 J/g). Plastic energy expended was calculated from the truestress–true strain plot that was derived from the load-elongationdata. The stored energy as a function of true plastic strain was thencalculated as the difference between expended plastic energy andthe corresponding dissipated energy.

Specimens for scanning electron microscopy (SEM) were pre-pared by conventional metallographic techniques and etched withKroll's reagent. Specimens for transmission electron microscopy(TEM) were prepared by cutting slices from the location where thehighest thermal activity had been noted. Specimens were electro-polished by twin jet electro-polisher in a 5% H2SO4+balancemethanol solution (by volume) at −50 1C. Surface crack densityas a function of strain for the coated specimens was obtained bydrawing equally spaced horizontal and vertical lines on the SEMimages taken from the gauge surfaces of the interrupted speci-mens. The number of intersections between the crack and thelines was then counted. The crack density was then calculatedusing the following expression [14]:

D¼ 1A

∑nv

∑Lvþ∑nh

∑Lh

� �ð2Þ

The term A represents the area of the scanning, nv and nhrepresent the number of vertical and horizontal intersectionsrespectively and Lv and Lh represent the total length of the verticaland horizontal lines respectively. The crack density was furthernormalised with respect to the minimum crack density valuenoted for undeformed specimen.

3. Results and discussion

The cross-sectional microstructure of the coating in as-coatedcondition revealed a single-layer coating consisting of Al3Ti phase.The thickness of the coating was approximately 50 μm and itcontained several through-thickness cracks. These cracks weregenerated during cooling of the coated specimens from coatingformation temperature, due to mismatch of the coefficient ofthermal expansion between the coating and the substrate [5,6].Further cracking occurred during the subsequent tensile testing ofthe coated specimens. Fig. 1 shows the average temperature

strain for (a) uncoated specimen, and (b) coated specimen.

Fig. 2. Variation in stored energy vis-a-vis strain hardening rate as a function of true plastic strain for (a) uncoated specimen, and (b) coated specimen.

Table 1Tensile properties of coated and uncoated specimens.

Specimens 0.2%YS (MPa) UTS (MPa) % Elongation

Uncoated 1080710 1160710 15.572Coated 1050720 1140710 371.5

S. Punnose et al. / Materials Science & Engineering A 576 (2013) 217–221 219

change (averaged over the area of the gauge where highestthermal activity was noted) as a function of true plastic strainfor one typical uncoated (Fig. 1a)) and coated specimen (Fig. 1(b)).It can be noted that total temperature change for uncoated speci-menwas higher (∼4 1C) than the coated specimen. Fig. 2 shows thevariation in stored energy as a function of true plastic strain for thesame uncoated and coated specimens for which variation intemperature plots have been shown in Fig. 1. With increasingstrain stored energy of the uncoated specimens increased mono-tonically (Fig. 2(a)), while for coated specimens the stored energyincreased up to a strain level beyond which the rate of increasedropped, the stored energy value became saturated and remainedalmost constant up to fracture (Fig. 2(b)). The monotonic increasein the stored energy in the case of uncoated alloy is typical of anyductile material such as the present near-α Ti alloy. A similar trendwas also observed during the initial deformation in case of thecoated alloy. However, beyond a certain strain (∼0.7%), almost allthe energy expended for deformation of the specimen wasdissipated as heat because of the formation of cracks in thecoating and their subsequent growth into the substrate. As aresult, beyond the critical strain, the stored energy virtuallyremained constant. Apart from this, formation and growth ofcracks became localised that led to localised energy dissipationand corresponding high rate of increase in the local temperature.Hence, the local temperature rise recorded by the IR camerawas more.

Beyond the critical strain, cracks in the coating grow rapidlyinto the base material because of the stress concentration gener-ated at the crack tips. Eventually, such crack growth led topremature failure of the coated specimen. Table 1 shows thetensile properties of both uncoated and coated specimens. It isevident that the coated specimens showed lower ductility com-pared to the uncoated specimens. This decrease in ductility can beascribed to the propagation of coating cracks into the substrate[15].

To validate the above aspect, microstructural analysis of bothcoated and uncoated specimens were carried out. From the long-itudinal sections of failed coated and uncoated specimens (Fig. 3),it can be clearly seen that micro-voids nucleated during deforma-tion at the interface of primary α and transformed β phases. The

volume fraction of micro-voids close to fracture end was muchlower in the case of failed coated specimen (Fig. 3(a)) as comparedto the uncoated specimen (Fig. 3(b)). The uncoated specimen failedby conventional mechanism of voids nucleation and their subse-quent coalescence. In case of the coated specimen micro-voidsnucleated in the substrate during early stages of deformation.However, the formation of a critical density of cracks in the coatingas deformation progressed, led to instability in deformation in thelater stages. Up to the critical strain the coated system deformedby deformation of the substrate and by the formation of cracks inthe coating simultaneously. Beyond this strain cracks in the coat-ing grew rapidly and penetrated into the substrate. Thus, theformation and growth of cracks in the coating altered the failuremechanism of the substrate alloy. The fractograph shown in Fig. 3(c)clearly shows the crack initiation point in the base material fromthe coating that leads to the change in fracture mode of the basematerial in the coated specimen. The fracture surface of the coatedspecimen, as shown in Fig. 3(c), demonstrates cleavage featureswhile in the uncoated specimen quasi-cleavage failure had occurred(Fig. 3(d)). Dislocation substructure was examined in the 0.8%deformed both coated and uncoated specimens. Fig. 4 presentsthe sub-structures formed in the two cases. It can be qualitativelystated from Fig. 4(a) and (b) that up to a certain strain level thedislocation density of the base material in case of coated specimenis marginally lower compared to the uncoated specimen. Thusbefore the inception of the unstable crack growth in the coatedspecimen the base material of the coated system deforms almost inthe similar manner as the uncoated specimen.

Fig. 5 shows the normalised surface crack density change as afunction of true plastic strain in case of coated specimens. Asevident, the rate of rise in crack density with strain decreasessomewhat beyond a strain of about 0.2%. One would expect thisstrain to be close to the strain value at which the drop in the storedenergy was observed i.e. at about 0.7% (Fig. 2(b)). However, thiswas not the case and the difference between the critical strainvalues arrived from stored energy and crack density can beexplained as follows. Ideally, the crack density values should havebeen determined from the single coated specimen at various levelsof strain during the tensile test. However, because of the limitationin experimentation, the crack density values were actually deter-mined from the observation of surface of multiple specimens, eachdeformed to a different degree of deformations. In such a case,since each coated specimen that was tested can be expected to beinherently different from the other specimens despite the samesubstrate alloy and coating characteristics, the crack density vs.strain results obtained using multiple specimens would not besame as that obtained using a single specimen. It is worthmentioning here that due to this anisotropic nature of the coating

Fig. 3. Scanning electron micrograph of the fractured specimen showing (a) the formation of micro-voids at the primary α and transformed β interface for coated specimen,(b) for uncoated specimen, (c) fractograph of the coated specimen showing the crack initiation point and cleavage fracture (enlarged image in the inset), and (d) Fractographof the uncoated specimen showing the quasi-cleavage fracture (enlarged image in the inset).

Fig. 4. Transmission electron micrograph showing dislocation substructure inside transformed β lath for specimen interrupted at 0.8% true plastic strain (a) uncoated and,(b) coated.

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the trend of the stored energy curve and also the critical strainvalue observed was different in case of different coated specimens.In the present study the average value obtained from threespecimens was 0.63% with a range of (0.40–0.80). In case wherethe localised heat evolution was not high the change in the slope

of stored energy curve beyond the critical strain was less andwhere heat evolution was more the curve showed a distinctvariation as compared to the uncoated specimens.

It can be noted that, while the stored energy showed a distinctchange in trend beyond the critical strain in the coated specimen

Fig. 5. Variation of normalised crack density as a function of true plastic strain.

S. Punnose et al. / Materials Science & Engineering A 576 (2013) 217–221 221

such an abrupt change was not observed in the strain hardeningrate vs. true plastic strain plot (Fig. 2(b)). This is because thecoating dimension (50 μm) was much small than that of substrate(7 mm). As a result, in the strain hardening curve the crackformation and growth beyond the critical strain is not captured.However, since the local temperature rise caused by such crackgrowth could be easily captured by the infrared camera, thischange gets prominently reflected in the stored energy curves.

The tensile deformation behaviour and the correspondingstored energy change for the coated specimen would dependupon various factors such as type of coating (chemical composi-tion, microstructure), coating thickness, substrate material, andgeometry of the specimen. These factors would dictate the strainat which instability would occur and the density of crack ensemblethat would form. The present coating had near-zero ductilitybecause of the limited number of slip systems available in theAl3Ti phase constituting the coating. However, it has been reportedthat addition of elements such as Cr, Mn and Li to bulk Al3Ti canenhance its ductility by changing its crystal structure from DO22 toL12 [16–19]. If such ductilization can also be achieved in the Al3Ticoating, the tensile deformation behaviour and the correspondingstored energy change for the coated specimen is expected to bedifferent from that observed in the present study.

4. Conclusion

The study shows a method to determine the critical strainbeyond which crack growth become rapid and localised during

tensile deformation of aluminide coated near α titanium alloysystem. Determination of this strain is important to prevent anycatastrophic failure of the coated system. It is shown that suchlocalised crack growth process can be easily captured by theinfrared camera while the same information cannot be obtainedfrom the global stress–strain values. The findings of the study thus,can enhance the understanding of the behaviour of coated systemconsiderably and can also be applied to find the location of localcrack growth for further crack growth study.

The work highlighted a simple way of calculating stored energyof deformation as a function of strain. The sequence of evolution ofthermal patterns obtained during tensile test enables a single stepin-situ method of determination of the stored energy that incor-porates conduction heat loss correction in the calculation of heatdissipation. The method can be used for studying the deformationbehaviour vis-a-vis the stored energy change of any material muchmore accurately.

Acknowledgements

The authors wish to acknowledge the financial support fromDRDO. The support rendered by the members of TAG, MEG, HTCGand EMG, of DMRL are gratefully acknowledged.

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