In Situ Synchrotron Studies of Reversible and IrreversibleNon-elastic Strain in a Two-Phase TiAl Alloy

FRANCISCO ALFREDO GARCIA-PASTOR, HUI JIANG, DAVID HU, XINHUA WU,PHILIP J. WITHERS, and MICHAEL PREUSS

This paper contrasts the cyclic tensile loading of two-phase titanium aluminide Ti-44Al-8Nb-1Bmicrostructures, namely fully lamellar (FL) and duplex (DP). The former, in contrast to thelatter, shows premature yielding and hysteresis loops on cycling. These phenomena were studiedby in situ cyclic loading of these specimens using high energy synchrotron X-ray diffraction. Theresults show significant differences in the micromechanics of deformation. Load transferbetween the c and a2 phases has been identified at stresses below the macroscopic yield point inthe FL microstructure, while the DP microstructure showed no signs of such load transferbefore its well-defined macroscopic yield point. A partially reversible pseudo-plastic deforma-tion mechanism seems to be operating at relatively low stresses in the FL specimen. Thismechanism is believed to be twinning/partial reversible twinning of the c phase. The presence oftwins at relatively low applied stresses has been confirmed for the FL microstructure by electronchanneling contrast imaging. Further support of twinning/partial reversible twinning in FL butnot DP microstructures below the macroscopic yield point was obtained by following theevolution of the integrated intensity of particular diffraction peaks measured during the in situsynchrotron X-ray experiments.

DOI: 10.1007/s11661-013-2024-0� The Minerals, Metals & Materials Society and ASM International 2013

I. INTRODUCTION

TWO-PHASE (a2+ c) titanium aluminides arepromising intermetallic alloys for aerospace applicationsbecause of their excellent mechanical properties at hightemperature.[1,2] However, they exhibit extremely lowductility at room temperature, which has been a factorin limiting their use.[3] Current engineering alloys candisplay a wide range of microstructures depending onthe thermo-mechanical processing route followed. Asexpected, the mechanical properties are very differentdepending on the microstructure, with duplex (DP)microstructures showing an increased ductility butreduced strength at high temperatures,[4,5] while fullylamellar (FL) and nearly fully lamellar (NFL) micro-structures are relatively brittle at room temperaturewhile retaining their strength at high temperature.[2]

DP microstructures (generated by slow cooling fromthe a transus) comprise an even distribution of lamellarcolonies and equiaxed grains. In contrast, lamellarmicrostructures (generated at higher cooling rates thanthose for DP microstructures) contain at least 80 pct ofthese colonies in the NFL form, and approach 100 pct oflamellar colonies in the FL microstructure.[6] Lamellarcolonies are platelets of twinned tetragonal c phase andof the hexagonal a2 phase as described in.[7] This semi-coherent interface between the c and a2 follows thecrystallographic relationship described by Blackburn[8]

ð0001Þa2 111f gc and ½11�20���

��½1�10�c ½1�

giving six crystallographic variants of the c phase from asingle a2 grain. These variants can be distinguished bytheir stacking order into three matrix-variants and threetwin-variants. The presence of these variants gives riseto three different c/c interfaces.[7]

It is well known that the mechanical response oflamellar colonies is highly orientation-dependant, withcolonies oriented with the (0001) planes at 45 deg to thestress axis considered to be soft while colonies oriented at0 and 90 deg are considered to be hard.[9–11] Thisorientation dependence originates from the a2 phasebecause slip of dislocationswith non-basal Burgers vectorrequires a far higher stress than slip of dislocations with bparallel with h11�20i. The c phase deforms mainly by glideof ordinary dislocations on the {111} planes, with Burgersvector b = 1/2h110i and superdislocations with Burgersvectors b = h101i and b = 1/2h11�2i. Plastic deformationof this phase can also be realized by mechanical twinningwith b = 1/6 h11�2i on the {111} planes.[9,11]

FRANCISCO ALFREDO GARCIA-PASTOR, formerly Ph.D.Student with the School of Materials, The University of Manchester,Grosvenor Street, Manchester, M1 7HS, U.K., is now Lecturer with theCinvestav Unidad Saltillo, Av. Industria Metalurgica No. 1062, ParqueIndustrial, 25900 Ramos Arizpe, COAH, Mexico. Contact e-mail:francisco.garcia@cinvestav.edu.mxHUI JIANG, formerly Post-Doctor-al Associate with the IRC in Materials, The University of Birmingham,Edgbaston, B15 2TT,U.K., is now Scientist with the Oxford Instrumentsplc,HighWycombe,Bucks,HP123SE,U.K.DAVIDHU,Post-DoctoralAssociate, is with the IRC in Materials, The University of Birmingham.XINHUA WU, formerly Professor with the IRC in Materials, TheUniversity of Birmingham, is now Professor with the Department ofMaterials Engineering, Monash University, Clayton, VIC 3800, Aus-tralia. PHILIP J. WITHERS and MICHAEL PREUSS, Professors, arewith the School of Materials, The University of Manchester.

Manuscript submitted June 1, 2012.Article published online October 2, 2013

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Although twinning is unidirectional, it has beenproposed that an anti-twinning vector does exist in theordered structure with b = �1/3 h11�2i.[10,12] It has beensuggested that this anti-twinning shear could onlyoperate when an already formed twin is effectivelyannihilated.[10] To date, this vector has not beenconfirmed by transmission electron microscopy(TEM), but its proposed existence opens the possibilityof twinning of the c phase being a semi-reversibledeformation mechanism.

Previous tensile experiments on Ti-44Al-8Nb-1B(Ti4481) have shown that near and FL microstructuresexhibit an early deviation from strain linearity, atstresses well below the expected yield point.[6] Incontrast, Ti4481 with a DP microstructure shows analmost perfectly elastic response until the yield point isreached. The existence of early plasticity has beenassociated with pre-yield cracking,[13] although exper-iments carried out on a chromium-containing alloy(Ti-46.5 at. pct Al-4 at. pct (Cr, Nb, Ta, B)) suggestthat it could be related to massive twinning of the cphase.[14] In situ acoustic measurements during tensileloading carried out by Hu et al.[6] revealed that theearly deviation from strain linearity in the NFLmicrostructure is associated with several acousticevents, whereas in the DP microstructure such eventswere only detectable at much higher stresses when thematerial started yielding.

In the present study, the occurrence of early plasticityduring loading was studied by cyclic tensile loadingusing increasing stress levels with each cycle. Thisexperimental methodology has been used previously inmagnesium alloys to investigate a similar early deviationof linear strain response during mechanical load-ing.[15,16]

Previous efforts to analyze the complex deformationmechanisms of the Ti4481 alloy using neutron radi-ation faced several problems, since titanium is a poorneutron scatterer and the presence of B in thealloy.[17] Because of these limitations, the cyclic tensileloading response was analyzed at the micro-scale byundertaking in situ loading experiments using highenergy synchrotron X-ray diffraction, a techniquewhich has been previously used for a range ofmicromechanical and phase transformation stud-ies.[18–26] Shuleshova[25] has previously studied thesolidification routes of TiAl alloys using synchrotronradiation, while other researchers such as Novoselovaet al.[21] and Clemens et al.[23] have published resultson high temperature transformation kinetics of thistype of alloys using a similar technique. Regardingin situ studies of deformation mechanisms usingsynchrotron X-ray radiation, Liss et al. have doneseveral studies using a similar technique to the onereported in this paper. They have used the spottinessof the Debye–Scherrer diffraction rings to analyze notonly the developing texture,[22] but also to studyintragranular strains, grain rotations, and mosaicspread in these alloys.[24] In contrast, this paperaverages diffraction rings in order to analyze themicromechanics of intergranular and interphasestrains, as described by Dye et al.[18]

II. EXPERIMENTAL

A. Sample Preparation and Cyclic Tensile Experiments

Specimens of Ti4481 were prepared by plasma meltingmaster alloys and elemental feedstock to produce 25 to50 kg ingots, followed by isothermal forging at1423.15 K (1150 �C) with an initial strain rate of5 9 10�3 s�1 to a reduction in height of 70 pct. Thedetails of the thermomechanical processing route for thematerial with FL and DP microstructure are given inTable I.Typical micrographs of the FL and DP microstruc-

tures are presented in Figure 1. The FL microstructure(Figure 1(a)) is formed by relatively fine (50 to 70 lm)lamellar colonies with less than 2 vol pct equiaxed cgrains (around 10 lm). In contrast, the DP microstruc-ture (Figure 1(b)) is a much finer microstructure. Thelamellar colonies are in the order of 10 to 20 lm in size,which puts them in the same size range as the surround-ing equiaxed c grains. The fraction of lamellar coloniesand c grains is approximately equal in this microstruc-ture.Cyclic tensile loading using increasing stress levels

with each cycle was carried out on a MTS Alliance RT/100 universal testing machine fitted with a 50 kN loadcell. The specimens had a cylindrical gage section of4 mm diameter and a gage length of 25 mm. A contactextensometer (MTS 634.31F.24) with a 10 mm gagelength and an accuracy of 6 9 10�4 strain was attachedin the middle of the gage section. The tensile test pieceswere first pre-loaded to 10 MPa, which was also kept asthe minimum applied stress during cyclic loading. Thestress level was increased by 50 MPa with each cycle andthe strain rate was set at 3.4 9 10�4s�1 during bothloading and unloading.

B. Synchrotron X-ray Diffraction Experiments

In situ loading experiments using high energy syn-chrotron X-ray diffraction were carried out at beam lineID15b of the European Synchrotron Radiation Facility(ESRF) in Grenoble. Flat dog-bone specimens having across-sectional area of 1 9 1 mm2 of Ti4481 wereprepared by electro-discharge machining. The loadingexperiments were carried out using an Instron electro-thermal/mechanical tester (ETMT) fitted with a 3 kNload cell. The samples were first preloaded to 30 MPa tominimize any rigid-body rotations followed by cyclicloading in the same fashion as described in the previousparagraph. The strain rate was 5 9 10�3 s�1 duringboth loading and unloading.The experimental setup was typical of the

in situ diffraction experiments carried out at the ESRF

Table I. Processing Routes for the Analyzed Specimens of

Ti4481 Alloy

Specimen Processing Route

FL forge+1583.15 K (1310 �C)/1h/furnacecooled to 1423.15 K (1150 �C)/air cooled

DP forge+1513.15 K (1240 �C)/4h/furnace cooled

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beamlines.[26] The incoming X-ray beam had a wave-length of 0.143 A (~87 keV) and was collimated to aspot size of 250 9 250 lm2. Because of the high beamenergy and flux, measurements can be carried out intransmission mode, which allows recording of completeDebye–Scherrer diffraction rings from a defined volumewhile the specimen is mechanically tested. The diffractedDebye–Scherrer rings were captured using a two-dimen-sional Pixium CCD detector with an array of1910 9 2480 pixels and pixel size of 154 9 154 lm2.[27]

In order to increase the number of grains contributingto the diffraction rings, the specimen together with theETMT was continuously oscillated by ±1.5 mm. Theoscillation period was 4 seconds during which time 80(0.2 second exposure) images were added together andsaved as a single image. The stress level associated withthe averaged image was the one recorded for the 80thimage.

An azimuthal integration was carried out on theaveraged diffraction rings using the software Fit2D.[28]

This integration was made on a ±5 deg arc segmentaligned to the loading axis. Consequently, the integrateddiffraction patterns comprised reflections of all the grainfamilies, which had their plane normal parallel to theloading direction and satisfied Bragg’s law as describedin Reference 29. It is important to keep in mind that thegrain families analyzed in this way do not represent theentire volume of the studied material. Each diffractionpeak was fitted to a Pseudo-Voigt function using aMatLab script and in this way it was possible to followvariations in peak intensities, full-width half maximumand peak positions during the entire cycling experi-ments. The lattice strain variation during mechanicalloading was evaluated using the standard engineeringstrain equation, where d0 represents the initial d-spacingin the unloaded condition[19,29]

e ¼ d� d0ð Þ=d0: ½2�

C. Microstructural Observations

Microstructural analysis of tensile samples tested todifferent stress levels were carried out by first grindingand polishing the gage section followed by electropol-ishing. The specimens were analyzed in a field emissiongun scanning electron microscope (FEG/SEM) in back-scattered electron (BSE) and electron channeling con-trast modes (ECCI). ECCI mode has been previouslyused to assess the occurrence of twinning in the cphase.[30,31] In order to maximize the channeling con-trast effect, the working distance was 3 mm and theaccelerating voltage 8 kV.

III. RESULTS

A. Macroscopic Strain Response During Cyclic TensileLoading

Figure 2 shows the stress–strain curves obtained bycyclic loading of DP and FL Ti4481. Incremental cyclicloading resulted in no deviation from strain linearitybefore the macroscopic yield point for the specimen withDP microstructure. In contrast, the Ti4481-FL specimenshowed an early deviation from the expected elasticstrain behavior at relatively low applied stresses. Thisearly deviation seems to be at least partially reversible,as evidenced by the hysteresis loops.In addition to this early reversible inelastic strain, the

FL specimen exhibited irreversible yielding at stresseswhich were close to the expected macroscopic yieldpoint. These two phenomena are defined in the insertpresented in Figure 2. Notice that the elastic moduliduring loading and unloading are the same (~157 GPa).The irreversible strain detected after unloading isidentified as the plastic strain (ep), whereas the reversibledeformation shall be called here the pseudo-plasticstrain (epsp). Prangnell et al.

[32,33] assessed a similar early

Fig. 1—SEM/BSE images of Ti4481 with (a) FL and (b) DP microstructure. The a2 appears light, whereas the c phase appears in a darker shadeof gray. The bright white particles are titanium borides, which precipitate during processing.

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pseudo-plastic deviation by plotting the first derivativeof stress with respect to the strain, dr/de. This derivativeis shown in Figure 3 for the last three loading cycles(550, 600 and 650 MPa*) of the FL microstructure. It

can be seen that during loading dr/de shows a smalldrop from the value of Young’s modulus at a stress ofabout 275 ± 25 MPa,. Furthermore, the 1st derivativeof the stress–strain unloading curve exhibits a similardrop from the expected Young’s modulus, confirmingthe reversible nature of the pseudo-plastic strain. Similarbehavior has been reported by in particulate metal-matrix composites (MMCs),[32,33] in Mg alloys[16] inshape memory alloys.[34] In the case of MMCs, thestrain hysteresis has been attributed to the reversiblemovement of dislocations generated by internal stressesin the particles; for Mg alloys, it has been explained byreversible twinning; and in shape memory alloys, it hasbeen associated with stress-induced phase transforma-tions.

By studying 1st derivative curves such as those inFigure 3, it is possible to quantify the extent of theelastic response as a function of applied stress as inFigure 4(a). It can be seen that upon loading at500 MPa pseudo-plasticity starts at a stress of~325 MPa and restarts on unloading at an appliedstress of 500 to 310 = 175 MPa. With increasing plasticstrain (and increasing applied stress), the pseudo-plas-ticity starts at lower stress upon unloading and lowerstresses on reloading. In last full cycle, the elastic rangeis just 150 MPa during unloading, indicating that thepseudo-plastic deformation mechanism starts soon afterunloading begins. As the extent of the elastic responserange reduces, the width of the hysteresis loop increases.This is presented in Figure 4(b), in which the pseudo-plastic strain is plotted against the maximum appliedstress. This curve suggests that pseudo-plasticity would

not occur for Ti4481 with a FL microstructure forstresses below 450 MPa.

B. Lattice Strain Response During Cyclic LoadingMeasured by Synchrotron X-ray Diffraction

Previous efforts to assess the micromechanics ofdeformation in TiAl alloys using synchrotron X-raydiffraction were hindered by the relatively coarse FLmicrostructure (in the order of 50 to 75 lm). Suchmicrostructures generally yield spotty diffraction rings,which results in large uncertainties of peak shape andposition after caking the diffraction image using an arcrange that represents a certain loading direction.[24,35]

Oscillating the FL sample and averaging 80 imagesgreatly improved the statistics of the diffraction rings,thus reducing the aforementioned uncertainties. Duringmechanical loading, the diffraction peaks, measured forinstance in the loading direction, will shift and in thisway allow recording the elastic strain experienced by thegrain families associated with each diffraction peak(based on crystallographic orientation). This elasticstrain is generally referred to as intergranular latticestrain. While during pure elastic deformation the inter-granular strain response is linear with the applied stress,it might deviate during plasticity. Figure 5(a) displaysthe maximum longitudinal intergranular lattice straindetermined from the c reflections measured duringthe synchrotron X-ray experiment when loaded tothe maximum applied stress for each cycle. Typically,the largest intergranular lattice strains were observed forthe (220)c and the (200)c reflections. As expected, withincreasing applied stress, the maximum intergranular

Fig. 2—Stress–strain curves obtained from cyclic loading of speci-mens of Ti4481 with FL and DP microstructures. ep is the plasticstrain whereas epsp is the pseudo-plastic strain. Notice the early devi-ation from linearity during both loading and unloading which givesrise to the hysteresis loops in the response.

Fig. 3—First derivative of stress with respect to strain (dr/de) plot-ted against applied stress for a specimen of Ti4481 FL tested at (a)550 MPa, (b) 600 MPa, and (c) 650 MPa. Notice the onset of earlyplasticity when the derivative has a slope different than zero, indicat-ing that the stress–strain function is no longer linear. The dotted lineindicates the Young’s modulus. Data from the final stages of loadinghave been removed from the figure to make it clearer.

*Data from the final stages of loading have been removed from thefigure to make it clearer.

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lattice strain within the c phase increases in bothmicrostructures. A small residual intergranular latticestrain is seen in the FL (but not the DP) microstructureafter unloading from stresses in excess of 500 MPa.Apart from that, the maximum intergranular latticestrains of the c phase seem to be very similar whencomparing the two microstructures. By averaging thelattice strain response of each phase and subtracting theaveraged c lattice strain (including diffraction peaksfrom the (200), (002), (220), (202), and (111) grainfamilies) from the averaged a2 lattice strain (includingdiffraction peaks from the ð20�20Þ and ð20�21Þ grainfamilies), it was also possible to estimate the interphaselattice strain in the loading direction providing infor-mation on load transfer from one phase to the other(Figure 5(b)). It is important to mention that due torelatively small a2 phase volume fraction (~20 pct), therewere only two peaks which could be properly fitted andanalyzed.

A clear difference can be observed between the twoexamined specimens. In contrast to the DP microstruc-ture, the FL microstructure developed very significantinterphase lattice strains once the applied stress exceeds500 MPa. These strains did not disappear upon unload-ing. A more detailed presentation of the lattice strainresponse for the FL microstructure along the longitu-dinal direction of the (220)c and the ð20�20Þa2 peaks is

shown in Figure 6(a). These two reflections exhibited thelargest difference in terms of elastic strain duringloading. This is expected since the (220)c grain familiesbelong to soft oriented colonies, with the c/a2 interfaceshowing an angle of ~35 deg with respect to the loadingaxis, whereas the ð20�20Þa2 reflection belong to hardoriented colonies, with the interface at angle of 90 degwith respect to the loading axis. The (220)c grain familydisplayed an essentially linear elastic response below anapplied stress of 450 MPa reaching 3000 9 10�6 strain.Beyond this applied stress no significant increase in theelastic strain carried by the (220)c grain family isobserved. The ð20�20Þa2 peak on the other hand showsa complementary response: showing an acceleratedaccumulation of elastic strain beyond an applied stressof 450 MPa. Being a hard grain family, the ð20�20Þa2, isexpected to take more elastic strain to compensate forthe early yielding of the soft oriented c lamellae. Thisload transfer only partially reverses during unloading,leaving the (220)c grain family in residual compressionand the ð20�20Þa2 grain family in tension. The occurrenceof early load transfer between the a2 and c phases duringdeformation of the FL microstructure suggests that amicro-plastic deformation mechanism is operating atrelatively low stresses.The behavior displayed by the analyzed reflections in

the FL specimen contrasts strongly with the response ofthe same reflections recorded in the DP microstructure

Fig. 4—Following the analysis of dr/de, two features of the hystere-sis loops were assessed, (a) the range of stress over which an elasticresponse occurs and (b) the width in strain of the hysteresis loops.

Fig. 5—(a) intergranular strain for the c phase during loading andunloading and (b) interphase strain between the a2 and c phases forspecimens for both FL and DP microstructures.

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when exposed to the same cycling history as the FLmicrostructure, see Figure 6(b). In this microstructure,there is no indication of load transfer in the two grainfamilies (phases) at any stress level. This shows that, assuggested by the macroscopic behavior presented inFigure 2, there are no plastic deformation mechanismsoperating at stresses below 600 MPa in the DP micro-structure.

C. Peak Area Evolution During Cyclic LoadingMeasured by Synchrotron X-ray Diffraction

Since twinning reorients regions of the grains, itwould be expected to have an effect on the integrateddiffraction peak intensity for grain families affected bytwinning. The peak integrated intensity for each reflec-tion was estimated by multiplying the fitted FWHMwith the peak intensity. Three reflections were analyzedfor each specimen: (202)c, (220)c, and (111)c, in order toassess the possibility of early plasticity occurring in the cphase. Figure 7 shows a composite plot of the evolutionof the integrated intensity of these peaks, plotted againstapplied stress. The integrated intensity was normalizedwith respect to the pre-load condition, to allow an easycomparison between the different peaks. The fittingerror was extremely low, in the order of 0.0025normalized integrated intensity units. However, themeasuring uncertainty, estimated by averaging several

measurements in the pre-loaded condition was found tobe of ~±0.04 normalized integrated intensity units.The behavior of the integrated intensity of the (220)c

peak (Figure 7(a)) is particularly interesting as it showsthe most significant changes in integrated peak intensitywith increasing number of cycles/maximum appliedstress. For each loading stage, the integrated intensityof the (220)c reflection first increases but starts to fallbefore the maximum stress level is reached. Uponunloading, the integrated intensity drops lower than itwas at the starting point of the cycle. This trendcontinues until the integrated intensity of the (220)creflection has lost ~10 pct of its original intensity at theend of the last loading cycle. It is very important to notethat during the unloading part of the cycles, theintegrated intensity of the (220)c reflection appears firstto fall during unloading before regaining somewhat justbefore reaching zero load although this does not occurat 500 MPa. The lowest points of (220)c integrated peakintensity are highlighted by the arrows in Figure 7(a). Itis interesting to note that the stress level at which theintegrated intensity starts increasing during unloadingcoincides with the stress level at which the macroscopicelastic strain response during unloading becomes non-linear (Figure 4(a)).In contrast to the (220)c reflection, the integrated

intensities of the (202)c and (111)c reflections show verylittle or no signs of irreversible change in integratedintensity during cycling (Figures 7(b) and (c)). In bothcases, the integrated intensity rises slightly duringloading and falls during unloading. However, for thelast two cycles with maximum applied stresses of 600and 650 MPa, hardly any variation of integrated inten-sity is observed in these two peaks. The peak behaviordescribed here will be analyzed in detail in theSection IV.The integrated intensities of the diffraction peaks

obtained for the DP microstructure are presented inFigure 8. The intensities of these reflections vary onlyby 5 pct during cyclic loading with a tendency ofthe (202)c and the (111)c reach a maximum at peakload.

D. Electron Channeling Contrast Imaging (ECCI)

In previous experiments, ECCI has successfully beenused to assess the occurrence of twinning in titaniumaluminides.[28,30] In the present paper, it was used toconfirm the appearance of deformation twins in the cphase of the FL specimen. Figure 9 shows typical ECCIimages obtained for the FL microstructure tested aboveand below the proportional limit (200 and 500 MPa), aswell as for the DP microstructure tested to 500 MPa(Figure 9(c). The crystallographic variants of the cphase are clearly visible by contrast variations insidethe c lamella of the specimen tested at 200 MPa(Figure 9(a)) but no indications of slip lines or twinsare seen at this stress level. In the sample tested to500 MPa (Figure 9(b)), very thin lines that are obliqueto the c/a2 interface are clearly visible across the clamellae. In addition to these oblique lines, there aresome lines, which appear to be parallel to the interface,

Fig. 6—Lattice strain response during the last three loading for the(220)c and ð20�20Þa2 peaks for (a) the FL microstructure and (b) theDP microstructure.

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particularly visible in the bottom of Figure 9(b). Incontrast, there are no indications of thin lines in thelamellae or equiaxed c grains of the DP microstructuretested to 500 MPa, as shown in Figure 9(c).

TEM investigations have demonstrated that ordinarydislocations in the c phase appear in a loop configura-tion, whereas twins appear as short, thin lines eitheroblique or parallel to the interface.[36] Hence, theoblique lines detected by electron channeling contrastimaging are expected to represent twins. Due to theextremely thin width of these twins it was impossible to

confirm them using electron backscatter diffraction(EBSD).

IV. DISCUSSION

The main observations meriting discussion are sum-marized below.

1. In contrast to the DP microstructure, cyclic tensileloading of the FL microstructure revealed the

Fig. 7—Normalized integrated intensity plotted against applied stress during cyclic loading of the FL specimen for the (a) (220)c peak, (b) (202)cpeak, and (c) (111)c peak. The scatter detected in the (111)c peak is caused by the effect of the underlying (0002) a2 peak, which effectivelymakes it a composite peak.

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occurrence of hysteresis loops, initiating below themacroscopic yield stress.

2. The first derivative of the loading and unloadingcurves confirmed the early deviation from an elasticresponse for the FL microstructure. Additionally, itwas found that the elastic range decreased duringsuccessive (increased) loading cycles.

3. While the FL microstructure showed load transferbetween c and a2, starting at a stress level of~500 MPa, the DP microstructure showed no evi-dence of load transfer.

4. The integrated intensity of the (220)c peak variedsignificantly during cycling for the FL microstruc-ture but not for the DP microstructure. In the caseof the FL microstructure, the peak area of the(220)c reflection varied out-of-phase with the load-ing cycle (but in accordance with the onset of pseu-do-plasticity in Figure 4) and displayed an overalldecrease of integrated peak intensity at the end ofeach cycle.

5. The occurrence of twinning in the FL microstruc-ture when loaded in excess of 500 MPa has been

Fig. 8—Normalized integrated intensity plotted against applied stress during cyclic loading of the DP specimen for the (a) (220)c peak, (b) (202)cpeak, and (c) (111)c peak.

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confirmed by electron channeling contrast imaging.By contrast, no twins were evident in the DP micro-structure when loaded to the same stress level.

As the incidence of stress-induced phase transforma-tion in titanium aluminides can be safely discounted, theoccurrence of strain hysteresis loops during cyclicloading of the FL microstructure can only be relatedto either reversible movement of dislocations as it is thecase in metal-matrix composites studied by Prangnellet al.[32,33] or reversible twinning as reported in Mgalloys.[15,16]

As discussed in the results section, the stress at whichthe drop from the expected Young’s modulus occurs isprogressively lower with each additional loading cyclewhile the width of the increases. This behavior could berelated to the presence of twinning at relatively lowstresses. On the previously cited works from Cacereset al.[15] and Mann et al.[16] on Mg alloys, it is suggestedthat the plastic pre-strain acts as nucleating step fortwinning. Successive loading cycles would only increasetwinning width with few new twins being nucleated. Inorder to determine if this is also the case with Ti4481alloy, additional loading cycles with no further loadingincrements would be needed. Although the possibility ofpre-yield cracking has been previously discussed in thisalloy,[6] it has only been reported on NFL microstruc-tures, with an equiaxed gamma phase volume fraction ofat least 15 pct. Furthermore, no early crack indicationswere found in the FL specimen after cyclic loading.

Regarding the lattice strain measurements, the mostnoticeable observation is the significant load transfer ofthe (220)c reflection measured in the loading direction(Figure 6(a)). Once this particular grain family hasreached about 3000 9 10�6 lattice strain (at ~450 MPa),lattice straining appears to exhaust even though thematerial is further loaded to 650 MPa. A simplifiedcalculation of the Schmid factors for slip of this grainfamily by ordinary dislocations of the [1�10](111) typeis presented in Table II. It is evident that any slip alongthe expected [220] direction gives an Schmid factor ofzero making this slip mode impossible for this particularc lamellae soft orientation. Table II also shows thatordinary slip is possible along the [202] direction. It is

important to highlight, however, that because of thec-phase tetragonality, each of these normal slip direc-tions belong to a different grain family as described byKishida et al.[37]

Twinning, on the other hand, is a mechanism thatcould be operating in the (220) grain family as Table IIIshows high Schmid factors for two variants of theh112i{111} twinning system along the [220] normaldirection.A closer look at Figure 6(a) reveals that the loading

and unloading slopes of the lattice strain of the (220)cpeak are different. This change in slope and the residualstrain show that there is a partially reversible pseudo-plastic deformation mechanism operating in the (220)cgrain family. It is also important to highlight that theresidual lattice strain after the last loading cycle is~1 9 10�3 which is in good agreement with the macro-scopic residual stress displayed in Figure 2. This residualstrain could be the result of twins which are notcompletely reverted after unloading.The integrated intensity evolution for the (220)c

reflection described in the results section, could also be

Fig. 9—Electron channeling contrast images of a Ti4481 sample with FL microstructure (a) 200 MPa and (b) tested at 500 MPa. (c) Shows asample of Ti4481 with DP microstructure tested to 500 MPa. The slip lines detected by channeling contrast in the FL specimen are consistentwith the morphology of twins in the c phase. Notice the absence of such lines in the DP specimen. The loading axis is horizontal in all cases.

Table II. Schmid Factors for Ordinary Slip Alongthe [220] and [202] Directions

Slip System Type m[220]c m[202]c

½1�10�ð111Þ ordinary dislocation 0 0.41½1�10�ð�111Þ ordinary dislocation 0 0½110�ð1�11Þ ordinary dislocation 0 0.41½110�ð�111Þ ordinary dislocation 0 0

Table III. Schmid Factors for Twinning

Along the [220] Direction

Slip System Type m[220]c

½11�2�ð111Þ twinning 0.47½�11�2ð�111Þ twinning 0½1�1�2�ð1�11Þ twinning 0½�1�1�2�ð�1�11Þ twinning 0.47

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the result of a twinning/partially reversible twinningmechanism. This is supported by the evolution of the(220)c FWHM and peak intensity, shown inFigures 10(a) and (b). It can be seen that even thoughthe variation of the (220)c FWHM for each cycle startsrelatively in phase with the applied stress, the peakintensity decreases dramatically halfway through thecycle that takes the material to 550 MPa. Duringunloading the peak intensity only starts recovering oncea threshold stress has been passed. Interestingly, thisthreshold stress coincides with the stress at which themacroscopic strain response becomes non-linear duringunloading (Figure 4). This behavior continues duringthe following loading cycles. These data suggest that thestrong variation of the (220)c integrated peak intensity(Figure 7(a)), which is out-of-phase with the loading

cycles, is largely affected by very significant variations ofthe peak intensity, which is best explained by partialreversible reorientation of grains, i.e., a twinning/det-winning mechanism. As discussed above, the lowSchmid factor for ordinary slip for the [220] direction,coupled with the high Schmid factor for twinning forthis direction, make the effects of twinning moreapparent in this grain family. By comparison, thevariation in the (220)c FWHM for the DP microstruc-ture is shown in Figure 10(c). It is notable that for theDP microstructure the FWHM of the (220)c reflectionvaries significantly less during each cycle compared tothe FL microstructure. This indicates that the (220)cgrain family is exposed to larger variations in intergran-ular strain for the FL microstructure compared to theDP microstructure, which could be the root cause for

Fig. 10—Evolution of (a) FWHM and (b) intensity during cyclic loading of a Ti4481FL specimen for the (220)c peak of the FL specimen, andchanges in FWHM for the (220)c peak of the DP specimen. It can be clearly seen that both the intensity and FWHM of the (220)c peak are notin phase with the applied load. Additionally, the peak seem to be sharpening irreversibly after unloading with a change of 10 pct of its originalwidth, suggesting a mechanism that relieves strain heterogeneity. In contrast, the FWHM of the (220)c reflection appears to be impervious toapplied stress. This indicates that the (220)c grain family is exposed to a larger variation of intergranular strain in the FL microstructure com-pared to the DP microstructure, which could be the root cause for twinning in the FL specimen.

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twinning in the FL microstructure. Skrotski[38] demon-strated that as twinning results in very significant shearstraining locally (5 to 10 pct), it requires substantialstrain accommodation by dislocations around the twin.Since TiAl is an intermetallic with limited ductility, it isdifficult to see how the neighborhood accommodatessuch significant shear strain.

In spite of the strong evidence supporting twinning/partially reversible twinning as an important mechanismto explain the macroscopic hysteresis loops, it isimportant to highlight the limitations of the in situloading experiment using synchrotron X-ray diffraction.The information obtained by this technique is limited tofavorably (in terms of Bragg’s law) oriented colonies.Additionally, it is possible that other soft colonies(which XRD reflections could not be detected) aresuitably oriented for slip of ordinary dislocations andsuperdislocations and contribute to the early deviationfrom linear strain response detected macroscopically.However, ordinary dislocations do not affect the peakintegrated intensity and thus cannot be detected by thisanalysis, as shown in Figure 7(b).

It is likely that a combination of ordinary disloca-tions and twinning are responsible for the early yieldingon the specimen with FL microstructure. However,ordinary dislocations cannot explain the hysteresisloops that the FL microstructure exhibits during cyclicloading. This behavior is consistent with reversibletwinning such as the one described by Caceres et al.[15]

and Mann et al.[16] on Mg alloys. Finally, the case fortwinning is supported by the fact that the onset of loadtransfer from c to a2 (Figure 5) and significant varia-tions of the (220)c integrated peak intensity(Figure 7(a)) are both detected at stresses similar tothose at which the FL microstructure starts to displaystrain hysteresis loops.

V. CONCLUSIONS

Cyclic tensile loading with increasing the maximumload for each cycle was carried out on Ti4481 with twodifferent microstructures, namely FL and DP micro-structures, which revealed hysteresis stress strain loopsin the case of the FL but not for the DP microstructure.

The analysis of lattice strain and peak area carried outby analyzing data from synchrotron X-ray diffractionexperiments with an averaging technique support thetheory that the hysteresis loops found in the FLmicrostructure are caused by a twinning/partiallyreversible twinning mechanism. It is important tomention that both interphase load transfer and peakintensity loss for the (220)c reflection occur at approx-imately the same stress levels as the onset of hysteresisduring macroscopic tensile testing. XRD data did notprovide any indications of twinning for the DP micro-structure.

The residual macroscopic and lattice strains are alsoof approximately the same value after cyclic loading thespecimen with FL microstructure. This residual strain isthought to be caused by twins which did not completelyrevert to their original orientations.

ACKNOWLEDGMENTS

The authors would like to thank Dr. Matthew Peelfor his help during the synchrotron experiments at theESRF. This work has been supported by EPSRC(through Grant No. GR/T19704/01). Financial sup-port for the Ph.D. study of F.G.-P. at the Universityof Manchester through the Mexican science and tech-nology council (Conacyt) is also acknowledged.

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