Al-Ti Particulate Composite: Processing and Studieson Particle Twinning, Microstructure, and ThermalStability

DEVINDER YADAV, RANJIT BAURI, ALEXANDER KAUFFMANN,and JENS FREUDENBERGER

The present investigation shows that alternate to the ceramic particles, hard metallic particlescan be used as reinforcement in an aluminum matrix to achieve a good strength–ductilitycombination in a composite. Titanium particles were incorporated into aluminum by frictionstir processing (FSP) to process an Al-Ti particulate composite. FSP led to uniform distributionof the particles in the stir zone without any particle–matrix reaction, thereby retaining theparticles in their elemental state. Fracture and twinning of the Ti particles with differentfrequency of occurrence on the advancing and retreating sides of the stir zone was observed.Twinning of the particles was studied by focused ion beam-assisted transmission electronmicroscopy. The processed Al-Ti composite exhibited a significant improvement in strength andalso retained appreciable amount of ductility. The thermal stability of the fine-grained structureagainst abnormal grain growth (AGG) was improved by the Ti particles. The AGG in the Al-Ticomposite occurred at 713 K (440 �C) compared to 673 K (400 �C) in the unreinforcedaluminum processed under the same conditions. On the other hand, the particle–matrix reactionoccurred only at 823 K (550 �C), and hence the Ti particles were thermally more stable com-pared to the matrix grain structure.

DOI: 10.1007/s11661-016-3597-1� The Minerals, Metals & Materials Society and ASM International 2016

I. INTRODUCTION

ALUMINUM and its alloys are one of the mostimportant classes of lightweight materials which findtheir applications in structural, automotive, and aero-space industry owing to their high specific strength, hightoughness, and good corrosion-resistant properties. Theever-increasing demands of higher strength-to-weightratio led to the development of new techniques andtechnologies to improve the mechanical properties ofaluminum by various means. Some of the methodswidely adopted to achieve this include grain refinement,strain hardening, precipitation hardening, solid solu-tion strengthening, and thermo-mechanical treatments.Metal matrix composite (MMC) technology is anotherwidely used method for strength improvement of Alalloys.[1] However, although the strength is increased,

retaining the ductility is a major challenge in most of theabove processes.In the MMC technology, high-strength and

high-modulus ceramic particles such as Al2O3, SiC,B4C, AlN, TiB2, and TiC are added to aluminum toimprove its mechanical properties. These particles areinherently brittle and hence the improvement of strengthis accompanied by a large drop in the ductility.[2–5] Poorwetting, particle clustering, and interfacial debondingare the other main reasons attributed to the drop inductility in MMCs.[6–9]

Harder metallic particles can be used as reinforce-ments in an aluminum matrix as an alternative approachto retain the ductility of the composite. The issues ofresidual stress and wettability can be minimized, too.However, retaining metallic particles as reinforcement isdifficult. Metals with low solid solubility such as Ni, Fe,and Ti readily react with Al by means of the formationof brittle aluminides (NiAl, Al3Ni, Al3Fe, Al3Ti, etc.)with rapid kinetics.[10–15] On the other hand, high-solidsolubility metals such as Cu, Zn, Mg, and Li easilydissolve into Al to form solid solutions. Hence, the mainchallenge here is to incorporate metallic particles andretain them in elemental state as reinforcements. Inother words, the reaction or dissolution of the metallicparticles during processing is to be prevented which is adifficult task in conventional composite processingroutes.Ti is an attractive choice for reinforcement because of

its high specific strength, high modulus, and goodfatigue properties. There have been attempts to process

DEVINDER YADAV, Research Associate, and RANJIT BAURI,Associate Professor, are with the Department of Metallurgical andMaterials Engineering, Indian Institute of Technology Madras,Chennai 600036, India. Contact e-mail: deven3042@gmail.comALEXANDER KAUFFMANN, Research Associate, is with theInstitute for Applied Materials, Karlsruhe Institute of Technology,76131 Karlsruhe, Germany. JENS FREUDENBERGER, Professor, iswith the Institute for Metallic Materials, IFW Dresden, PO Box270117, 01171 Dresden, Germany, and also with the Institute ofMaterials Science, Technische Universitat Bergakademie Freiberg,Gustav-Zeuner-Str. 5, 09599 Freiberg, Germany.

Manuscript submitted September 10, 2015.

METALLURGICAL AND MATERIALS TRANSACTIONS A

Ti-reinforced Al matrix composites by other processessuch as disintegrating melt deposition (DMD),[16] accu-mulative roll bonding (ARB),[17] or accumulative swag-ing and bundling.[18] Although the reaction timebetween Al and Ti is minimized in the DMD process,the formation of intermetallic could not be preventedcompletely leading to a reduction in the ductility of thecomposite.[16] The Al/Ti composite processed by ARBalso exhibited a large drop in the ductility, andaccording to Liu et al., aluminum-based MMCs withhigh strength and good ductility are yet to be producedby ARB.[19]

Friction stir processing (FSP) is fast emerging as amicrostructure modification tool.[20,21] The process hasbeen widely used for grain refinement, microstructurehomogenization, and in situ composite fabrication.[22–30]

The process uses a specially designed tool having a pinand a shoulder. The frictional heat generated at thetool–plate interface takes the temperature of the mate-rial to a range where it can be easily deformed andprocessed. Due to the extensive material flow, theprocess also offers the possibility of incorporatingreinforcement particles in order to create a composite.Recently, it has been shown that incorporatinghigh-density metallic particles such as Ni and Cu inaluminum by FSP is feasible to process metal parti-cle-reinforced composites.[31,32]

The aim of the present investigation is to explore theutility of FSP to incorporate lighter metallic particlessuch as Ti as reinforcement, alternate to the brittleceramic particles, to process a composite. The otherobjective is to understand the microstructure developedand evaluate the thermal stability of the compositeagainst abnormal grain growth and particle–matrixreaction. An interesting feature that came out in thisinvestigation was the twinning of Ti particles and thisaspect was studied by focused ion beam (FIB)-assistedtransmission electron microscopy. The mechanical prop-erties were also evaluated.

II. EXPERIMENTAL DETAILS

Commercially pure aluminum (99.5 pct) was chosenas the matrix material. Grooves, 1 mm wide, 2 mm deep,and 50 mm in length were made on the plate and werefilled in open air with Ti powder (99.5 pct, 325 Mesh,Alfa Aesar). Figures 1(a) and (b) show the SEM imageof Ti powder and the particle size distribution, respec-tively. Several images were used to calculate the sizedistribution. The particle size distribution was wide withthe average particle size being 20 lm. FSP was carriedout on the groove with a tool made of M2 steel with ashoulder diameter, pin diameter, and a pin length of 15,5, and 3.5 mm, respectively. The pin was flat (with nothreads) and tapered at an angle of 5 to 7 deg. Theshoulder of the tool was made concave to have a betterconfinement of the plasticized material. A tool rotationspeed of 1000 rpm and a traverse speed of 30 mm/minwere used for FSP. It should be noted that this was theminimum ratio required to obtain a defect-free stir zone.When the ratio was lowered (by decreasing the rotation

speed or increasing the traverse speed), it either led todefects in the stir zone or insufficient particle incorpo-ration and distribution. FSP was also carried out on thepure aluminum plate without Ti particles (sampledesignated as FSPed Al) with the same parameters toevaluate the effect of Ti particles on microstructureevolution and thermal stability of the microstructure.The composite was characterized by X-ray diffraction

(XRD), scanning electron microscope (SEM), electronbackscattered diffraction (EBSD), and transmissionelectron microscope (TEM). Phase analysis was carriedout in a PANalytical Diffractometer using Cu Karadiation. The metallographically polished samples wereobserved in an FEI Quanta FEG SEM equipped withTEAM (texture and elemental analysis microscopy)EDS detector. For EBSD, the metallographically pol-ished samples were electropolished in a mixture ofperchloric acid and methanol at 263 K (�10 �C) and 10V and observed in an FEI Quanta FEG SEM equippedwith TSL-OIM software with a step size of 250 nm. Thesamples were subjected to ion milling in Gatan PECS(precision etching coating system) using a beam energyof 90 to 100 lA for 5 to 10 minutes before EBSD. Grainorientation spread (GOS) map was constructed from theEBSD scans. GOS is defined as the average of themisorientation of individual pixel from the averageorientation of the grain. For TEM analysis, 3-mm diskswere taken from the surface of the sample, metallo-graphically polished to a thickness of 90 lm, andsubjected to twin-jet polishing using a mixture ofperchloric acid and methanol at 253 K (�20 �C) and10 V. In order to analyze the twins in the Ti particles, thetarget preparation by means of lift-out was performedusing an FEI Helios 600i dual-beam SEM/FIB. The Titarget particle was imaged by backscattered electron(BSE) imaging with orientation contrast at 10 kV inimmersion mode in order to evaluate a suitable beamdirection for TEM investigations within the twinlamella. The lift-out was performed at 30 kV andvarious Ga ion currents. TEM observations were madein a Philips CM 12 microscope operating at 120 kV.Thermal stability of the composite microstructure and

particles was evaluated by heating the samples in atubular furnace in flowing argon atmosphere at differenttemperatures for different periods of time. The sampleswere quenched to freeze the microstructure and wereobserved in SEM after each heat treatment operation.Tensile tests were carried out on ASTM standard

samples (1 mm in thickness, 2.5 mm in width, and 10mm in gage length). Samples were sliced from the stirzone, parallel to the surface and the tool traversedirection by electrical discharge machining (EDM), andtests were carried out on an Instron (Model 3367)machine at a constant strain rate of 10�3 s�1.

III. RESULTS AND DISCUSSION

A. XRD Phase Analysis

The maximum solid solubility of Ti in Al is only 1.3wt pct (at 983 K) and hence Al3Ti intermetallics readily

METALLURGICAL AND MATERIALS TRANSACTIONS A

form when Ti is added to Al.[33] The XRD pattern of theprocessed composite in Figure 2 shows peaks primarilybelonging to Al and Ti. Weak peaks of Al3Ti wereobserved at 2h ~ 42.5 and 58.5 deg. No peaks corre-sponding to any other additional phase or reactionproducts were observed confirming that majority of theTi was in the elemental state. The presence of the weakAl3Ti peaks is discussed in the next section.

B. Composite Microstructure

Figures 3(a) through (c) show the SEM (BSE) imageof the surface of the processed composite on theadvancing side, center, and the retreating side of thestir zone, respectively. A collage of SEM (BSE) imagestaken in the cross section is shown in Figure 3(d). Auniform distribution of Ti particles in the Al matrix canbe observed. The area fraction of the Ti particles wasfound to be 7 pct using ‘Image J’ image analyzer.

Several images were used to calculate the area fraction.With images taken at 1009, the minimum particle sizethat was detectable with the analyzer was 4 lm and forimages taken at 509 it was 2 lm. The interface of the stirzone with the base metal on the advancing side is sharpand that on the retreating side is somewhat diffuse. It canalso be seen from the cross-section image that theparticles have reached up to the full depth (2 mm) ofthe groove. The particle–matrix interface was carefullylooked for several particles in SEM. The high-magnifi-cation SEM (BSE) image in Figure 4(a) shows a cleanparticle–matrix interface with no sign of any reactionproduct forming at the interface. There were very fewsmall Ti particles which had a diffuse interface with thematrix on one side as shown in Figure 4(b). It appearsthat there is some diffusion of Ti into the Al matrix;however, the formation of equilibrium reaction productdoes not appear as the contrast change is gradual. Thus,XRD and SEM results indicate that the particle–matrixreaction is prevented and the Ti particles are retained intheir elemental state in the Al matrix. The weak Al3Tipeaks in the XRD pattern (Figure 2) could come fromtwo sources. First, the initial Ti powder had a wideparticle size distribution with some very fine particles ofsize less than 1 lm which could react with Al to formAl3Ti intermetallic during FSP. Secondly, there could bea very thin layer of Al3Ti formed around the particles.This layer, however, is too thin to be captured by SEMand EDS. Moreover, if present, such a layer will onlyimprove the bonding between particle and matrix with-out affecting the bulk region. The Al3Ti peaks couldcome from one of these two sources or from both.Previous studies on FSP of Al-Ti system reported auniform dispersion of fine Al3Ti particles in the Almatrix. It is worth mentioning here that the Al3Tiintermetallic formed during sintering of the powdercompact before FSP was carried out on the compact.[26]

It can be seen from the Al-Ti phase diagram thatAl3Ti is the first stable phase that forms after thesolubility limit of Ti in Al is surpassed.[34] A heat inputor mechanical energy is needed to overcome the

Fig. 1—(a) SEM image showing the particle morphology of Ti pow-der and (b) particle size distribution.

Fig. 2—XRD pattern of the processed Al-Ti composite.

METALLURGICAL AND MATERIALS TRANSACTIONS A

activation barrier for the reaction between Al and Ti(Activation energy = 225 kJ/mol at 823 K[35]). DuringFSP, the ratio of tool rotation speed to the traversespeed determines the amount of heat input into thematerial with higher ratio leading to higher heat input.In the present case, a minimum ratio (1000/30), whichwas just sufficient to obtain a defect-free stir zone, wasused to minimize the heat input into the material. Asonly weak peaks of Al3Ti were detected, it appears thatthe corresponding heat input is not high enough to causesignificant Al-Ti reaction. There are experimental diffi-culties in placing thermocouples in the stir zone tomeasure the exact temperature the material experiencesduring FSP. The temperature can be theoreticallycalculated by the equation given by Arbegast andHartley:[36]

T

Tm¼ k

x2

#� 104

� �a

;

where x is the tool rotation speed and 0 is the tooltraverse speed. The exponent a is in the range from 0.04to 0.06, the constant K is between 0.65 and 0.75, and Tm

(�C) is the melting point of the alloy. Temperature in therange of 725 K to 807 K (452 �C to 534 �C) is expectedwith the parameters used (1000 rpm and 30 mm/min) inthe present case. This temperature corresponds to 0.77to 0.86 Tm for Al and 0.37 to 0.42 Tm for Ti. Hence,none of the phases go to liquid state during the process.It should be noted that Ti has very low diffusivity inAl.[37] Also it has been reported that the materialexperiences the peak temperature for very short dura-tion of time during FSP[21] which may be too small tohave any significant diffusion of atoms to form somereaction product. It can be argued that Al3Ti has anegative heat of formation (�36.9 kJ/ mol) and the largeheat release is sufficient to cause local melting at the

interface and accelerate the reaction.[34] Since weobserved weak peaks of Al3Ti, it appears that theparticle–matrix reaction is not significant. The heatinput into the material can be minimized by controllingthe process parameters of FSP and the reaction canbe controlled and even can possibly be preventedcompletely.

C. Ti Particle Analysis

The incorporated Ti particles were also characterizedin detail. An interesting feature was the twinning of Tiparticles along with particle fracture as shown inFigures 4(c) and (d), respectively. Several twin variantswith lamellar width in the range of 0.15 to 1 lm can beseen. It is interesting to note that the frequency oftwinned particles was more on the advancing side,whereas more fractured particles were observed on theretreating side. Twinned and fractured particles weremore pronounced on the surface than in the crosssection, too. Ti has HCP crystal structure at roomtemperature and a limited number of slip systems.Twinning plays a major role in deformation of Ti andthe extent of twinning increases with an increase instrain rate and a decrease in temperature.[38] However,the effect of temperature is less pronounced on twinningthan the strain rate.When carefully looking at the fractured particles,

most of them were found to be twinned. The role oftwins in the fracture of the particles, if any, needs somediscussion at this point. First, it can be seen from thetwinned and fractured particles that the twins spreadacross the particles from one end to the other(Figure 4(d)) and a careful observation at high magni-fication revealed micro-cracks in some of the twins asshown in Figure 4(e). It is well known that twins areregions of high stress concentration and can act as crack

Fig. 3—SEM (BSE) images of Al-Ti composite on the (a) advancing side, (b) center, (c) retreating side, and (d) the cross section of the stir zone.

METALLURGICAL AND MATERIALS TRANSACTIONS A

nucleation sites especially at twin–twin interactionsites.[38,39] Twinned regions are harder than the non-twinned regions due to Basinski mechanism whichinvolves the transformation of glissile dislocations tosessile dislocation inside them.[40] Twin boundaries alsoact as a barrier to dislocation motion. Overall, twinningstrain hardens the material and cracks can generateinside them or at the twin–twin intersections. Therefore,in the present case it appears that Ti particles undergotwinning first to accommodate the imposed deformation

during FSP and the cracks nucleate inside them duringfurther deformation causing the particles to fracture.The Ti twins could not be indexed by EBSD as the

particles were severely deformed. The FIB lift-outtechnique, as described before, was carried out forlocation-specific TEM in order to analyze the twins.Therefore, the lamella was cut perpendicular to the twintrace visible by BSE imaging. Figure 5(a) shows a TEMimage of the Ti twins in which twin boundaries asstraight lines can be seen (indicated by arrows).

Fig. 4—SEM (BSE) images showing the (a) particle–matrix interface, (b) small Ti particle with diffuse interface on one side, (c) twinned Ti parti-cle, (d) fractured particle with twins spreading across it, and (e) micro-cracks in the twins.

METALLURGICAL AND MATERIALS TRANSACTIONS A

Figure 5(b) shows the diffraction pattern of the centralpart of the image revealing a ½1 0 �1 �2� zone axis. Thus,the according pole of the twin interface from thebright-field image can be assigned to a ð1 0 �1 1Þ plane.For a f1 0 �1 1g h1 0 �1 �2i compression twin system forHCP metals with near-ideal c/a ratio, the observed zoneaxis should not vary by changing from adjacent region

into the twin as it can be seen in the stereographicprojection (Figure 5(c)) which includes the matrix aswell as the twin orientation with respect to the deter-mined zone axis. In contrast, the experimental diffrac-tion patterns of the adjacent regions are different fromthose of the center region. The diffraction patternsreveal similarities to a ½1 1 �2 �3� zone axis which is slightlyoff (~15 deg) from the initial zone axis. This might becaused by the severe plastic deformation of the Tiparticles during the FSP process and the high disloca-tion density visible in the adjacent region of the twinboundary. Paton and Backofen[41] reported that theobserved twin system is active in Ti in the temperaturerange of 673 K to 1073 K (400 �C to 800 �C). Theparticles experience the same temperature as that of thematrix but less strain during the process. The Tiparticles flow with the plasticized Al and get shearedwhen they come into contact with the shoulder and pinof the tool. However, the effect of shoulder appears tobe more pronounced as there were more twinned andfractured particles on the surface than in the crosssection of the stir zone.

D. Microstructure Development

The microstructure developed in the composite wasstudied in detail by EBSD. The analysis was carried outon the surface of the stir zone at the advancing side,center, retreating side, and also in the cross section ofthe stir zone. Figures 6(a) through (d) show the EBSD(IPF + grain boundary) map of the composite atdifferent locations. The black particles in the images arethe Ti particles which could not be indexed as they wereheavily strained. A confidence index cut-off of 0.1 wasused to remove the poorly indexed points. Fine grains inthe range of 4 to 8 lm are observed in the stir zonecompared to the grain size of 80 lm of the base plate.Thus, apart from incorporating Ti particles into the Almatrix, FSP also caused grain refinement of the matrix.The microstructure is characterized by fine and equiaxedgrains with narrow grain size distribution at all thelocations, and this homogeneity in the microstructureis required for homogeneous mechanical propertiesthroughout the volume of the stir zone. The microstruc-tural details such as grain size, fraction of high-anglegrain boundaries (HAGBs) (>15 deg misorientation),and grain orientation spread (GOS) are summarized inTable I.The first thing to be noticed from Table I is that there

is some variation in the microstructural features atdifferent locations within the stir zone. This variationindicates the presence of a possible strain gradientwithin the stir zone. This can be attributed to differenttemperatures and the degree of deformation theselocations experience due to variation in material flowpattern. During FSP, the material flows from theretreating side to the advancing side, and as the toolmoves ahead in the advancing side, the deformedmaterial is deposited behind the tool in the retreatingside. Also, the material is pushed downward on theadvancing side and moves upward on the retreating side

Fig. 5—(a) TEM bright-field image of a twinned Ti particle (twinboundaries are highlighted by arrows), (b) diffraction pattern fromthe region within the circle of (a) including an indexation accordingto a ½1 0 �1 �2�; and (c) stereographic projection presenting the orienta-tion change by f1 0 �1 1g h1 0 �1 �2i compression twinning of the deter-mined orientation of the region within the circle of (a).

METALLURGICAL AND MATERIALS TRANSACTIONS A

within the pin diameter.[42] It is generally believed thatthe plasticized material flows several times around thepin before being consolidated. Therefore, the materialflow is not symmetric around the tool and as a result the

material experiences different combinations of strainand temperature within the stir zone. In addition, thematerial experiences peak temperature and strain on thesurface. This gives rise to non-uniform plastic

Fig. 6—EBSD map of the composite at the (a) advancing side, (b) center, (c) retreating side, and (d) the cross section of the stir zone.

Table I. Summary of Microstructure Developed in Al-Ti Composite and FSPed Al, Data Taken from EBSD Analysis

Al-7Ti composite Advancing Side Center Retreating Side Cross Section

Avg. Grain size (lm) 6 8 4 8Fraction of HAGBs in pct 54 42 61 75GOS (deg) 2.1 3.1 1.8 1.6

FSPed Al Advancing Side Center Retreating Side Cross Section

Avg. Grain size (lm) 9 9 7 10Fraction of HAGBs in pct 63 64 72 77Avg. GOS (deg) 1.9 2.1 1.4 0.8

METALLURGICAL AND MATERIALS TRANSACTIONS A

deformation or strain gradient in the stir zone during theprocess.[43,44] This in turn results in a variation of themicrostructure (grain size, fraction of HAGBs, andGOS) and the degree of fracture and twinning of theparticles at these four locations of the stir zone. Thesmaller grain size and the presence of more fractured Tiparticles on the retreating side are an indication that thisside of the stir zone experiences more strain during FSP.

The fine grains during FSP are formed by dynamicrecrystallization (DRX) process. Various forms of DRXsuch as continuous, discontinuous, and geometric havebeen reported to be occurring during FSP of various Alalloys.[21,44–47] It has also been shown that the incorpo-ration of micron-sized (>10 lm) metallic particles doesnot affect the mechanism of dynamic recrystallizationduring FSP.[30,48] Fine particles (<1 micron), on theother hand, can affect the recrystallization kinetics bothat low- and high-temperature deformation by parti-cle-stimulated nucleation (PSN). However, at hightemperatures, PSN may become less viable as the storedenergy is reduced and the dislocations may be able tobypass the particles without forming deformation zonearound them.[49]

It can be also seen from Table I that the incorpora-tion of Ti has some effect on the nature of grainboundaries and the remnant deformation (GOS value)inside the grains when compared with FSPed Al. As the

deformation behavior of Al and Ti is different, theincorporation of Ti leads to the creation of additionalsources of dislocation at the particle–matrix interfaces.Also, the difference in the coefficient of thermal expan-sion of Ti (8.6 9 10�6 K�1) and Al (23.1 9 10�6 K�1)[50]

causes dislocation generation as the material experiencesa thermal cycle during FSP. Hence, the grains in theAl-Ti composite have higher dislocation density and,consequently, a higher GOS value compared to FSPedAl. As the dislocations rearrange by dynamic recovery(due to high stacking fault energy of Al) and formsubgrain boundaries, a higher fraction of such bound-aries (hence, lower pct HAGB) is expected in thecomposite compared to FSPed Al (Table I).

E. Mechanical Properties

Tensile tests were carried out on Al-Ti composite andFSPed Al to evaluate the mechanical properties.Figure 7 shows the engineering stress–strain curvesand the properties are summarized in Table II. TheAl-Ti composite exhibited an yield strength (0.2 pctproof stress) of 118 MPa which is almost 3.4 times(240 pct) higher than that of the base Al. The ultimatetensile strength (UTS) also improved 2 times. Theimportant point to note is that even after this significantimprovement in strength, the composite retained anappreciable amount of ductility. The mechanical prop-erties of some of the ceramic particle-reinforcedAl-based composites[2,3,5] and FSPed composites[51,52]

are shown in Table II. It should be noted that the matrixmaterial in References 51 and 52 is the same as that usedin the present study (commercially pure Al). It can beseen that Ti particles are more effective in enhancing thestrength of the composite than the ceramic particles andcause a drop in the ductility. The improvement instrength arises from a combined effect of grain refine-ment and Ti particle incorporation. Ti particles are wellbonded with the matrix and the load is effectivelytransferred to the Ti particles through the interface. Thegrains have high dislocation density due to differentialdeformation and thermal mismatch between Al and Ti(shown in Section III–F–1). Hence, high dislocationdensity and subgrain boundaries observed inside thegrains also contribute to the strength of the composite asthey act as a hindrance to the dislocation motion. Theductility of the composite was obviously lower com-pared to pure Al, which is due to the presence of the Ti

Fig. 7—Engineering stress–strain curves of base Al, FSPed Al, andAl-Ti composite.

Table II. Tensile Test Properties of Base Al, FSPed Al, and Al-Ti Composite (Standard Deviations are Shown in Parenthesis)

0.2 Pct Proof Stress (MPa) UTS (MPa) Pct Elongation

Base Al 35 (0.6) 72 (1) 39 (0.6)FSPed Al 82 (2) 89 (3) 35 (2)FSPed Al-7 pct Ti composite 118 (2) 149 (0.7) 16 (0.9)Al-3.5Cu/Al2O3 (powder metallurgy route)[2] 238 (134) 374 (261) 2.2 (14)Al356-5.5wt pct TiB2 (casting route)[3] 240 (228) 303 (287) 7 (13)Al2024-5vol pct SiC (casting route)[5] 214 (75) 320 (185) 3.2 (21)FSPed Al-7Ni composite[51] 104 (35) 127 (72) 25 (39)FSPed Al-5TiB2 composite[52] 124 (35) 154 (70) 20 (40)

The values of unreinforced composite are given in brackets.

METALLURGICAL AND MATERIALS TRANSACTIONS A

particles. Ti being an HCP metal has a limited numberof slip systems and displays lower ductility than Al.

Microstructure of the Ti particles also plays a role inthe mechanical properties of the composite. Astrain-hardened and twinned particle is less ductile andwill lower the overall fracture strain of the composite.However, it would improve the strength of the compos-ite as the particle strength itself increases due to strainhardening. The twinning of particles also played a rolein their fracture generating smaller particles of irregularmorphology. The irregular particles will generate morestress concentration centers in the composite. Theseconditions may lead to early fracture of the particlesduring tensile loading leading to fracture of the com-posite. Equiaxed and strain-free particles are desirablefor better ductility in the composite. If the particles werenot twinned and fractured during FSP, the compositewould have possibly exhibited higher ductility.

The fracture surface of the tensile samples wasobserved under SEM. Figures 8(a) and (b) show theSE and BSE images, respectively, of the fracture surface

of the composite. All the samples showed deep dimplesindicating the ductile mode of failure. Ti particleslocated inside the dimples can be seen from the BSEimage. The high-magnification image shows that there isno de-cohesion between the Ti particle and the Almatrix. This indicates superior interfacial bondingbetween the particle and the matrix.

F. Thermal Stability of the Composite

1. Thermal stability of matrix microstructureA fine-grained microstructure such as that obtained in

FSP can undergo abnormal grain growth (AGG) whenexposed to higher temperatures.[53,54] Therefore, a priorknowledge of the time and temperature up to which themicrostructure in the processed materials is stable is ofparamount importance. This would also help find thetemperature range in which the processed composite canbe safely used or further processed for shaping ormanufacturing. The composite was subjected to thermalexposure at different temperatures to evaluate thethermal stability of the microstructure. The microstruc-ture underwent abnormal grain growth after a thermalexposure at 713 K (440 �C) for 10 minutes, as shown inFigure 9(a). Another 10 minutes of exposure led tofurther growth of the abnormally growing grain and thefine grains were completely consumed as shown inFigure 9(b). There was no sign of any reaction of Tiparticle with Al at this temperature and the particle–ma-trix interface was clean. The processed pure Al (FSPedAl) was also given similar thermal treatment andcompared with the composite. When compared to themicrostructure before thermal exposure (Figure 9(c)),FSPed Al underwent AGG at 673 K (400 �C) after10 minutes of thermal exposure (Figure 9(d)). AGGoriginates from the preferential growth of a few grainswhich have some special growth advantage over theirneighbors. Grains oriented favorably for AGG consumethe other grains leaving behind many island grains in themicrostructure (Figure 9(d)).AGG occurs suddenly and it is difficult to capture the

initial events. Also, AGG is harmful as it leads to asudden drop of the mechanical properties of thematerial. The main factors governing AGG are sec-ond-phase particles, texture and surface effects.[49]

According to the very recent works carried out byTaleff and Pedrazas[55] and Omori et al.,[56] the mech-anism that causes AGG are still not well understood.The GOS map of the composite in Figure 10 shows

different degrees of spread in the orientation inside theneighboring grains. The spread in the orientation insidesmaller grains was lower compared to the bigger grains.White color grains have a spread of more than 5 deg.The spread in the orientation comes from the presenceof dislocations and dislocation boundaries inside thegrains. TEM analysis of the composite microstructurealso showed that the smaller grains were mostly free ofdislocations and the bigger ones had various dislocationsubstructures inside them (Figure 11(a)). As shown inFigure 11(b), free dislocations and dislocation tangleswere observed in some of the grains and other adjacentgrains were virtually free of dislocations. All these

Fig. 8—(a) SEM image of the fracture surface of the composite and(b) BSE image showing superior particle–matrix bonding.

METALLURGICAL AND MATERIALS TRANSACTIONS A

features indicate a varying degree of stored energy inneighboring grains that resulted in different GOS valuesacross them. Srolovitz et al. showed theoretically thatthe difference in stored energy between grains may leadto excessive grain growth.[57] Thus, the varying storedenergy inside the neighboring grains could be the drivingforce for AGG during the thermal exposure in thepresent case.

It should be noted that the microstructure of thecomposite was stable up to 120 minutes of thermalexposure at 623 K (350 �C) and no grain growth wasobserved. This confirms the stability of the microstruc-ture. It can also be noticed that the presence of Tiparticles raised the temperature of AGG by 40 Kcompared to pure Al. This can be attributed to thepinning effect exerted by the Ti particles, especially thefiner ones (size £ the grain size), on the grain boundaries.FSP also led to fracture of some of the particles creatingfiner ones. Similar improvement of thermal stability of

FSPed AZ31 alloy from 573 K to 673 K (300 �C to400 �C) after incorporating SiC particle was reported byMorisada et al.[58]

2. Thermal stability of Ti particlesThe Ti particles in the composite are in a non-equi-

librium state and may form stable reaction productsunder favorable thermal conditions. Various reactionproducts of Al and Ti are reported to form during theirhigh-temperature annealing.[59–61] Hence, the stability ofthe particles was also evaluated by thermal exposurestudies. A thermal exposure at 823 K (550 �C) for20 minutes led to the development of a thin layer(500 nm to 1 lm) of a reaction product around the Tiparticles as shown in Figure 12(a). The same particlebefore heat treatment is shown in the inset. Thecomposition of the layer was found to be close to thatof Al3Ti intermetallic from the EDS analysis. Thefully grown continuous intermetallic layer formed a

Fig. 9—EBSD (IPF + grain boundary) map of Al-Ti composite after heat treatment at 713 K (440 �C) for (a) 10 min, (b) after 20 min showingabnormal grain growth, (c) FSPed Al before heat treatment, and (d) after heat treatment at 673 K (400 �C) for 10 min.

METALLURGICAL AND MATERIALS TRANSACTIONS A

core–shell type structure around the particles as shownin Figure 12(b). The TEAM-EDS phase mapping inFigure 12(c) revealed the presence of three phases. The

phase in yellow color that formed the shell is rich in bothAl and Ti. It can also be seen that some of the particlesgot completely converted to intermetallic (Al3Ti).Increasing the exposure time to 120 minutes led to

the growth of the intermetallic and the thickness of theshell layer increased. Large pores were observed in Almatrix and very fine pores were observed in the Tiparticles after the heat treatment as shown inFigure 12(d). As Al and Ti atoms migrate across theinterface, they leave behind vacancies. The atoms reactto form intermetallic phase at the interface and thevacancies left behind condense to form pores byKirkendall effect.[62] Al-Ti intermetallics have highlyordered structure and different molar volume capacitythan Al and Ti. Three atoms are needed per Ti atom toform Al3Ti kind of intermetallics and the atomic radiusof Al is also smaller (1.432 A) compared to Ti(1.47 A). Hence, more Al flux toward the interfaceand consequently more vacancy flux toward Al areexpected to occur giving rise to larger pores in Al. It isworth mentioning here that the kinetics of Al-Tireaction will be faster for a strain-hardened Ti particlethan a strain-free particle as high dislocation contentwill enhance the diffusion rate.It should be noted that the particle–matrix reaction

occurred at a much higher temperature [823 K (550 �C)]than the AGG temperature [713 K (440 �C)], and henceTi particles are thermally more stable compared to thegrain structure. Therefore, the presence of the Tiparticles, though in non-equilibrium condition, doesnot pose any hindrance to the high-temperature appli-cation or secondary processing of the composite.Further studies, however, are needed in the future

scope of this investigation to look at the particle–matrixinterface atomistically in high-resolution TEM to getmore insight into the wettability and interface reactionand ascertain whether there is a diffusion layer or anyother reaction product formed at the interface.

IV. CONCLUSIONS

In the present study, it is shown that FSP led touniform distribution of Ti particles in the aluminummatrix, retaining majority of the particles in theirelemental state. The study can be extended to Al alloysto achieve good strength–ductility combination. It willbe interesting to see if finer and stable particles canfurther raise the temperature of AGG. The followingspecific conclusions can be drawn from the presentstudy:

1. Alternate to brittle ceramic particles, hard metallicparticles, such as Ti, can be incorporated into an Almatrix and can be retained in their elemental stateby a solid-state process such as FSP to process moreductile composites.

2. Strain during the process caused fracture andtwinning of the Ti particle. The twins were com-pressive type, and being the regions of high stressconcentration, they act as a source of crack nucle-ation leading to the fracture of the particles.

Fig. 10—GOS map of the composite showing variation in remnantdeformation inside the grains.

Fig. 11—TEM image of the composite showing (a) dislocation sub-structure inside the grains and (b) dislocation tangles inside thegrains.

METALLURGICAL AND MATERIALS TRANSACTIONS A

3. The combination of high strain and temperaturecaused dynamic recrystallization during FSP andrefined the grain size of the Al matrix. Themicrostructure was characterized by equiaxed finegrains in the range of 4 to 8 lm with narrow grainsize distribution.

4. The composite exhibited 3.4 times higher yieldstrength (118 MPa) compared to that of base Al (35MPa). The improvement in strength is attributed tothe combined effect of grain refinement and particleaddition. The composite also retained 16 pct duc-tility which is higher than those of many conven-tional ceramic particle-reinforced composites.

5. The composite microstructure exhibited better sta-bility against abnormal grain growth (AGG) as itoccurred at 713 K (440 �C) in the compositecompared to 673 K (400 �C) in the FSPed pure Al.

6. The Ti particles were thermally more stable than themicrostructure of the matrix. Particle–matrix reac-tion occurred only after a thermal exposure of morethan 20 minutes at 823 K (550 �C). Ti particlesform a diffusion couple with the matrix and theoccurrence of Kirkendall effect lead to the forma-

tion of pores in the Al matrix and intermetallic layeraround the particles giving rise to a core–shell typestructure.

ACKNOWLEDGMENTS

The authors would like to thank the Faculty at theMaterials Joining Laboratory, IIT Madras, for provid-ing access to the NRB-supported FSP facility. Theauthors thank T. Sturm for experimental assistancewith the FIB lift-out.

REFERENCES1. T.W. Clyne and P.J. Withers: An Introduction to Metal Matrix

Composites, 1st ed., Cambridge University Press, Cambridge,1993.

2. B. Ralph, H.C. Yuen, and W.B. Lee: J. Mater. Process. Technol.,1997, vol. 63, pp. 339–53.

3. J.V. Wood, P. Davies, and J.L.F. Kellie: Mater. Sci. Technol.,1993, vol. 9, pp. 833–40.

Fig. 12—SEM (BSE) image of (a) Ti particle after heat treatment at 823 K (550 �C) for 10 min (inset shows the same particle before heat treat-ment), (b) core–shell type structure around the Ti particle, (c) EDS phase mapping showing the formation of Al-Ti-rich phase around particles(yellow color), and (d) pores in Al after 120 min of thermal exposure.

METALLURGICAL AND MATERIALS TRANSACTIONS A

4. B.F. Luan, N. Hansen, A. Godfrey, G.H. Wu, and Q. Liu: Mater.Des., 2011, vol. 32, pp. 3810–17.

5. S.J. Hong, H.M. Kim, D. Huh, C. Suryanarayana, and B.S. Chun:Mater. Sci. Eng. A, 2003, vol. 347, pp. 198–04.

6. R.J. Arsenault and N. Shi: Mater. Sci. Eng., 1986, vol. 81,pp. 175–87.

7. K.B. Lee, H.S. Sim, S.Y. Cho, and H. Kwon: Mater. Sci. Eng. A,2001, vol. 302, pp. 227–34.

8. J.J. Lewandowski and C. Liu: Mater. Sci. Eng. A, 1989, vol. 107,pp. 241–55.

9. Z. Wang, M. Song, C. Sun, D. Xiao, and Y. He: Mater. Sci. Eng.A, 2010, vol. 527, pp. 6537–42.

10. S.C. Deevi and V.K. Sikka: Intermetallics, 1996, vol. 4, pp. 357–75.11. W.Y. Yang and G.C. Weatherly: J. Mater. Sci., 1996, vol. 31,

pp. 3707–13.12. J.C. Rawers and H.E. Maupin: J. Mater. Sci. Lett., 1993, vol. 12,

pp. 637–39.13. D. Ozyurek, S. Tekeli, T. Tuncay, and R. Yilmaz: Powder Metall.

Met. Ceram., 2012, vol. 51, pp. 491–95.14. M. Sujata, S. Bhargava, and S. Sangal: Iron Steel Inst. Jpn., 1996,

vol. 36, pp. 255–62.15. S. Mousheng, Z. Mengxian, H. Bin, Z. Shuguang, and L. Jianguo:

Rare Met. Mater. Eng., 2008, vol. 37, pp. 1570–74.16. S.K. Thakur and M. Gupta: Compos. A, 2007, vol. 38, pp. 1010–18.17. D. Yang, P. Cizek, P. Hodgsona, and C. Wen: Scripta Mater.,

2010, vol. 62, pp. 321–24.18. T. Marr, J. Freudenberger, D. Seifert, H. Klaub, J. Romberg, I.

Okulov, J. Scharnweber, A. Eschke, C.G. Oertel, W. Skrotzki, U.Kuhn, J. Eckert, and L. Schultz: Metals, 2011, vol. 1, pp. 79–97.

19. C.Y. Liu, B. Zhang, P.F. Yu, R. Jing, M.Z. Ma, and R.P. Liu:Mater. Sci. Eng. A, 2013, vol. 580, pp. 36–40.

20. R.S. Mishra, M.W. Mahoney, S.X. McFadden, N.A. Mara, andA.K. Mukherjee: Scripta Mater., 2000, vol. 42, pp. 163–68.

21. R.S.Mishra andZ.Y.Ma:Mater. Sci. Eng. R, 2005, vol. 50, pp. 1–78.22. C.I. Chang, X.H. Du, and J.C. Huang: Scripta Mater., 2007,

vol. 57, pp. 209–12.23. J.Q. Su, T.W. Nelson, and C.J. Sterling: Scripta Mater., 2005,

vol. 52, pp. 135–40.24. J.Q. Su, T.W. Nelson, and C.J. Sterling: J. Mater. Res., 2003,

vol. 18, pp. 1757–60.25. Z.Y. Ma: Metall. Mater. Trans. A, 2008, vol. 39A, pp. 642–58.26. C.J. Hsu, C.Y. Chang, P.W. Kao, N.J. Ho, and C.P. Chang: Acta

Mater., 2006, vol. 54, pp. 5241–49.27. R.S. Mishra, Z.Y. Ma, and I. Charit: Mater. Sci. Eng. A, 2003,

vol. 341, pp. 307–10.28. C.F. Chen, P.W. Kao, L.W. Chang, and N.J. Ho: Metall. Mater.

Trans. A, 2009, vol. 41A, pp. 513–22.29. M.S. Khorrami, M. Kazeminezhad, and A.H. Kokabi: Metall.

Mater. Trans. A, 2015, vol. 46A, pp. 2021–34.30. H.S. Arora, H. Singh, and B.K. Dhindaw: Int. J. Adv. Manuf.

Technol, 2012, vol. 61, pp. 1043–55.31. D. Yadav and R. Bauri: Mater. Lett., 2010, vol. 64, pp. 664–67.32. D. Yadav and R. Bauri: Mater. Sci. Technol., 2015, vol. 31,

pp. 494–500.33. T.B. Massalski and H. Okamoto: Binary Alloy Phase Diagrams,

ASM International, Materials Park (OH), 1993.34. J.C. Schuster and M. Palm: J. Phase Equilib. Diffus., 2006, vol. 27,

pp. 255–77.

35. Q. Zhang, B.L. Xiao, and Z.Y. Ma: Mater. Chem. Phys., 2013,vol. 139, pp. 596–602.

36. W.J. Arbegast, and P.J. Hartley: Proceedings of the FifthInternational Conference on Trends in Welding Research, PineMountain, USA (1998) 541.

37. S.K. Das and L.A. Davis:Mater. Sci. Eng., 1988, vol. 98, pp. 1–12.38. J.W. Christian and S. Mahajan: Prog. Mater. Sci., 1995, vol. 39,

pp. 1–157.39. M.H. Yoo: Metall. Mater. Trans. A, 1981, vol. 12A, pp. 409–418.40. S.R. Kalidindi, A.A. Salem, and R.D. Doherty: Adv. Eng. Mater.,

2003, vol. 5, pp. 229–32.41. N.E. Paton and W.A. Backofen: Metall. Mater. Trans. B, 1970,

vol. 1B, pp. 2839–47.42. T.U. Seidel and A.P. Reynolds: Metall. Mater. Trans. A, 2001,

vol. 32A, pp. 2879–84.43. S. Agarwal, C.L. Briant, L.G. Hector, Jr, and Y.L. Chen: J.

Mater. Eng. Perform., 2007, vol. 16, pp. 391–403.44. J.Q. Su, T.W. Nelson, and C.J. Sterling: Mater. Sci. Eng. A, 2005,

vol. 405, pp. 277–86.45. P.B. Prangnell and C.P. Heason: Acta Mater., 2005, vol. 53,

pp. 3179–92.46. C.G. Rhodes, M.W. Mahoney, W.H. Bingel, and C. Calabrese:

Scripta Mater., 2003, vol. 48, pp. 1451–55.47. J.Q. Su, T.W. Nelson, and C.J. Sterling: Phil. Mag., 2006, vol. 86,

pp. 1–24.48. R. Bauri, D. Yadav, C.N.S. Kumar, and B. Balaji: Mater. Sci.

Eng. A, 2015, vol. 620, pp. 67–75.49. F.J. Humphreys and M. Hatherly: Recrystallization and Related

Annealing Phenomena, 2nd eds., Elsevier, Oxford, 2004, pp. 448,368.

50. E.A. Brandes and G.B. Brook (Eds.): Smithell’s Metals ReferenceBook, 7th Ed., Butterworth and Heinemann, Oxford, 1992, pp. 14.

51. D. Yadav and R. Bauri: Mater. Sci. Eng. A, 2011, vol. 528,pp. 1326–33.

52. D. Yadav and R. Bauri: J. Mater. Eng. Perform., 2015, vol. 24,pp. 1116–24.

53. I. Charit and R.S. Mishra: Scripta Mater., 2008, vol. 58,pp. 367–71.

54. S. Mironov, K. Masaki, Y.S. Sato, and H. Kokawa: Metall.Mater. Trans. A, 2013, vol. 44A, pp. 1153–57.

55. E.M. Taleff and N.A. Pedrazas: Science, 2013, vol. 341,pp. 1461–62.

56. T. Omori, T. Kusama, S. Kawata, I. Ohnuma, Y. Sutou, Y. Araki,K. Ishida, and R. Kainuma: Science, 2013, vol. 341, pp. 1500–02.

57. D.J. Srolovitz, G.S. Grest, and M.P. Anderson: Acta Metall.,1985, vol. 33, pp. 2233–47.

58. Y. Morisada, H. Fujii, T. Nagaoka, and M. Fukusumi:Mater. Sci.Eng. A, 2006, vol. 433, pp. 50–54.

59. T. Marr, J. Freudenberger, A. Kauffmann, J. Romberg, I. Okulov,R. Petters, J. Scharnweber, A. Eschke, C.G. Oertel, U. Kuhn, J.Eckert, W. Skrotzki, and L. Schultz: Metals, 2013, vol. 3,pp. 188–201.

60. L. Xu, Y.Y. Cui, Y.L. Hao, and R. Yang: Mater. Sci. Eng. A,2006, vols. 435–436, pp. 638–47.

61. D. Yang, P. Hodgson, and C. Wen: Intermetallics, 2009, vol. 17,pp. 727–32.

62. A.D. Smigelskas and E.O. Kirkendall: Tran. Am. Inst. Min.Metall. Pet. Eng., 1947, vol. 171, pp. 130–42.

METALLURGICAL AND MATERIALS TRANSACTIONS A