H O S T E D B Y

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Materials International

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Progress in Natural Science: Materials International 26 (2016) 1–12

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Strategies for improving ductility of ordered intermetallics

Z.B. Jiao, J.H. Luan, C.T. Liun

Center for Advanced Structural Materials, Department of Mechanical and Biomedical Engineering, College of Science and Engineering, City University of HongKong, Hong Kong, China

Received 25 October 2015; accepted 4 January 2016Available online 11 February 2016

Abstract

Ordered intermetallics possess attractive high-temperature properties; however, low ductility and brittle fracture limit their use as engineeringmaterials in many cases. This paper provides a comprehensive review on the recent progress in the development of ductile ordered intermetallicsand summarizes the strategies used to improve the tensile ductility of ordered intermetallics, including control of ordered crystal structures,engineering grain-boundary structure and chemistry, eliminating environmental embrittlement, microstructure optimization, control of phasestability, and promoting transformation-/twining-induced plasticity. The basic ideas and related mechanisms underlying these ductilizingstrategies are discussed. In addition, a brief mention of the current use of intermetallic alloys for structural and corrosion applications is made.& 2016 The Authors. Production and hosting by Elsevier B.V. on behalf of Chinese Materials Research Society. This is an open access articleunder the CC BY-NC-ND license (http://creativecommons.org/licenses/by-nc-nd/4.0/).

Keywords: Ordered intermetallics; Ductility; Ductilizing strategy; Microstructure; Alloy design

1. Introduction

Ordered intermetallics constitute a unique class of metallicmaterials that form long-range-ordered crystal structures belowa critical ordering temperature. This class of materials are nowbeing widely investigated as potential high-temperature struc-tural materials due mainly to their attractive properties,including good high-temperature strength, resistance to oxida-tion and corrosion, relatively low material density and highmelting point [1–9]. Unlike conventional or disordered alloyswhose yield strengths decrease with increasing test tempera-ture, many ordered intermetallics exhibit an increase, ratherthan a decrease, in yield strength with increasing test tempera-ture. In addition, their long-range order produces strongerbonding and closer packing between atoms, which restricts theatom mobility and generally leads to slower diffusion pro-cesses and better creep resistance in ordered lattices [10].These unique properties of ordered intermetallics making them

10.1016/j.pnsc.2016.01.01416 The Authors. Production and hosting by Elsevier B.V. on behalflicense (http://creativecommons.org/licenses/by-nc-nd/4.0/).

g author. Tel.: þ852 34427213; fax: þ852 34420172.ss: chainliu@cityu.edu.hk (C.T. Liu).nder responsibility of Chinese Materials Research Society.

very attractive for high-temperature structural and corrosion

applications [11–13].Most intermetallic alloys, however, exhibit brittle fracture

and low tensile ductility at ambient temperatures, whichseverely restricts their practical applications as structuralmaterials [14]. In addition, many intermetallics based onaluminides are sensitive to moisture in the environment atlower temperatures [15,16]. For the past several decades, agreat deal of effort has been devoted to understanding thebrittle fracture and embrittlement mechanism of the interme-tallic alloys. As a result, both intrinsic and extrinsic factors thatembrittle the intermetallic alloys have been identified. Inparallel, a substantial body of work has attempted to improvethe ductility of ordered intermetallics through application ofphysical metallurgical principles, which leads to the develop-ment of a variety of advanced ordered intermetallics with agood combination of high strength and high ductility for manyindustrial applications.The purpose of this paper is to review the recent progress

in the development of ductile ordered intermetallics andsummarizes the scientific strategies used to improve thetensile ductility of ordered intermetallics, including control

of Chinese Materials Research Society. This is an open access article under the

Fig. 1. Effect of electron concentration (e/a) on the stability of ordered crystalstructures in the Ni3V–Co3V–Fe3V alloys [21].

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of ordered crystal structures, engineering grain-boundarystructure and chemistry, eliminating environmental embrit-tlement, microstructure optimization, control of phasestability, and promoting transformation-/twining-inducedplasticity. These strategies can serve as useful guidelinesfor the alloy design of new high-strength and high-ductilityintermetallic alloys in the future. Finally, a brief mentionof the current use of intermetallic alloys for structural andcorrosion applications in industries is given at the end ofthis review paper.

2. Control of ordered crystal structures

One of the principal obstacles to the room-temperaturetensile ductility of many ordered intermetallics is the low-symmetry crystal structures that have an insufficient number ofindependent slip systems. It has been found that the ductility ofthese intermetallics can be substantially improved by control ofthe ordered crystal structures, that is, changing the crystalstructure from low symmetry (such as ordered hexagonalstructure) to high symmetry (such as ordered cubic structures)through composition modifications.

One classic example is control of the ordered crystalstructures of the quasiternary system of Ni3V–Co3V–Fe3Vthrough adjusting the electron concentrations. It is known thatthe ordered structures of A3B intermetallics are built from theregular stacking of the close packed ordered layers, eitherbeing of triangular (T) type or rectangular (R) type. Stacking ofthe T layers gives ordered structures with a cubic or hexagonalsymmetry depending on the stacking sequence, while stackingof R layers generally form ordered structures with a tetragonalsymmetry. Beck and Dwight [17,18] noted a general correla-tion between the stacking character of the T layers with theelectron concentration ratio (e/a, the number of valenceelectrons per atom) in the A3B intermetallic alloys. With anincrease in e/a, the ordered structure changes from predomi-nantly cubic to predominantly hexagonal stacking. Furtherincrease in e/a leads to a change in the basic layer structurefrom T to R type [17–19]. Liu and Inouye [20,21] found thatthe ordered crystal structure of the Ni3V–Co3V–Fe3V alloyscan be altered systematically by controlling the electronconcentration ratio, where the e/a is taken as 9, 8, 7, and5 for Fe, Co, Ni, and V, respectively. As shown in Fig. 1, thebinary Co3V alloy forms a six-layered hexagonal orderedstructure with a hexagonality of 33.3% [21]. The electronconcentration ratio in Co3V (e/a¼8) can be increased bypartial replacement of Co (e/a¼9) with Ni (e/a¼10), formingternary (Co,Ni)3V alloys. With increase of e/a from 8 to 8.54,the hexagonality of the ordered structure increases step by stepfrom 33.3% to 66.7% and finally to 100%. Further increase ine/a above 8.54 produces a change in the basic layer structurefrom T to R type, which gives rise to a tetragonal orderedstructure similar to Ni3V (DO22) [20,21]. On the other hand,the e/a in Co3V can be reduced by partial substitution of Fe (e/a¼8) for Co (e/a¼8), forming ternary (Co,Fe)3V alloys.When the e/a is below 7.89, the L12 ordered cubic structurehaving the stacking sequence ABC is stabilized in the (Co,

Fe)3V alloys. Thus, control of the electron concentration ratioprovides an elegant way for alloy design. As Co is anexpensive elements, the cost of the (Co,Fe)3V alloys can besignificantly reduced by replacing Co with less expensiveelements (e.g., Fe and Ni). Considering that the e/a of Fe, Co,and Ni are 7, 8, and 9, respectively, Co in the (Co,Fe)3V alloycan be replaced by an equal amount of an equiatomic mixtureof Ni and Fe, which alters the alloy composition but not e/a.As long as the electron concentration falls roughly in the samerange as (Co,Fe)3V, the ordered cubic structure remains stablein the (Ni,Co,Fe)3V alloys. Eventually, all Co atoms can bereplaced by Ni and Fe atoms, leading to the formation of L12

ordered cubic alloys of the composition (Ni,Fe)3V without anyCo elements [20,21].In addition, control of the ordered crystal structures offers an

effective way for tuning the room-temperature ductility ofseveral ordered hexagonal and cubic alloys. Room-temperaturetensile elongation of the Ni3V–Co3V–Fe3V alloys with differ-ent crystal structures are presented in Fig. 2 [21]. Thehexagonally ordered alloys Co3V and (Ni,Co)3V are verybrittle with less than 2% elongation at room temperature[20,21]. Their low ductility can be attributed inherently totheir hexagonal crystal structure, which has limited numbers ofslip systems. By contrast, the L12 ordered cubic alloys with thecompositions of (Fe,Co)3V, (Fe,Co,Ni)3V, and (Fe,Ni)3V areall very ductile, showing a tensile elongation of above 40%.The deformation behavior of these L12 ordered cubic alloys issimilar to that of face-centered cubic (fcc) materials havingtwelve {111} slip systems. Therefore, the crystal structure ofthe ordered intermetallics alloys can be altered from hexagonalCo3V and (Ni,Co)3V to cubic (Fe,Co)3V, (Fe,Co,Ni)3V, and(Fe,Ni)3V via a macro-alloying approach, which significantlyenhances the ductility of the ordered intermetallics. The controlof ordered crystal structures and ductility by macro-alloyingrepresents a major advance in alloy design of ductile orderedintermetallics. At the present time, the study of high-entropyalloys (HEAs) with close to equiatomic compositions has

Fig. 2. Comparison of room-temperature tensile elongations of cubic andhexagonal alloys [21].

Fig. 3. Effect of boron additions on the tensile elongation and fracturebehavior of Ni3Al (24 at% Al) tested at room temperature [29].

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become a hot research topic [22–25]. Actually, the series ofalloys based on (Fe,Co,Ni)3V can be regarded as HEAs.

3. Engineering grain-boundary structure and chemistry

Many ordered intermetallics exhibit severe brittleness thatoriginates at grain boundaries. A classic example is Ni3Al withan ordered cubic structure (L12), which is highly ductile insingle-crystal forms but very brittle in polycrystalline formseven though there are 12 independent slip systems [26,27]. Inmost metals and alloy, the grain-boundary fracture is related tothe segregation of impurities, such as S and P, at grainboundaries, which causes intergranular brittleness [28]. Studiesof fracture in high-purity polycrystalline Ni3Al using Augerelectron spectroscopy (AES), however, revealed brittle inter-granular fracture even without any detectable segregation ofimpurities at grain boundaries [29–31]. The grain boundary is,therefore, considered to be intrinsica1ly weak in the poly-crystalline Ni3Al alloys. It has been found that the brittleintergranular fracture of polycrystalline Ni3Al is due mainly tothe poor grain-boundary cohesion resulting from the highordering energy and large difference in electronegativity andvalence electrons between Ni and Al [26].

Micro-alloying processes were used to enhance the grain-boundary cohesion and overcome the grain-boundary brittle-ness problems. Among various doping agents, boron has beenfound to be the most effective one in ductilizing the Ni3Alalloys. Aoki and Izumi [26] first discovered that a smallamount of boron additions can suppress intergranular fractureand enhance tensile ductility of Ni3A1 at room temperature.Subsequently, considerable effort has been devoted to thedevelopment of ductile Ni3A1 alloys by boron micro-alloying[29,32–34]. The variation of the room-temperature tensileductility as a function of boron addition in Ni3Al (24 at%Al) is shown in Fig. 3 [29]. The boron-free Ni3Al alloyexhibits a negligible ductility at room temperature and thefracture surface is characterized by completely intergranularfracture. Micro-alloying with a minor amount of boron sharplyincreases the ductility and changes the fracture mode fromfully intergranular fracture to predominantly transgranularfracture. In particular, a tensile ductility of as high as 50%

was achieved in Ni3Al by micro-alloying with 0.1 wt% B. Thestriking effect of boron on ductility attracted considerableattention to understanding the ductilization mechanism ofboron. Both Auger microprobe [29] and atom probe [35]studies showed that boron segregates strongly to the grainboundaries in Ni3Al with a depth of a few atomic layers,although the extent of segregation appears to vary fromboundary to boundary [29].More interestingly, alloy stoichiometry was found to have a

large effect on the ductility and fracture behavior of boron-doped Ni3Al [29]. Liu et al. found that room-temperatureductility and fracture behavior of boron-doped Ni3Al isdependent critically on the deviation from alloy stoichiometry.The fracture surfaces of the boron-doped Ni3Al alloys isillustrated in Fig. 4 as a function of bulk Al concentration,which clearly shows the effect of alloy stoichiometry on thefracture behavior of Ni3Al at room temperature. Boron is veryeffective in improving the ductility and suppressing intergra-nular fracture in the Ni3Al alloys containing less than 24 at%Al. As the Al concentration is increased from 24 to 24.8 at%,the ductility decreases sharply from �49.4% to �6.0%, andthe failure mode changes from transgranular to mixed modeand then to mainly intergranular fracture (Fig. 4) [29]. Thetensile properties of the aluminides with 425 at% Al couldnot be measured because the ingots cracked during fabrication[29]. Auger studies of freshly fractured surfaces of boron-doped Ni3Al samples indicate that the level of boron segrega-tion to the grain boundaries decreases significantly and thegrain-boundary aluminum concentration increases moderatelywith increasing bulk Al concentration [29]. The enrichment ofgrain-boundary aluminum may reduce the electronic interac-tion between nickel and boron atoms at the grain boundaries,

Fig. 4. SEM fractographs of Ni3Al doped with 0.05 wt% B showing the effect of alloy stoichiometry on the fracture behavior at room temperature: (a) 24 at% Al,tensile fractured, (b) 24.5 at%, Al, tensile fractured, (c) 24.8 at% Al, tensile fractured, and (d) 26 at%, Al, fractured by bending [28].

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thereby lowing the ductilization effect of boron. As a result,the boron micro-alloying becomes less effective in ductilizingNi3Al when there is less than a critical amount of boron(�10 at% B) presented at the grain boundaries. These resultssimply suggest that deviations from alloy stoichiometryinfluence the grain-boundary chemistry, which, in turn, affectsthe grain-boundary cohesion and the overall ductility of theNi3Al alloys. In addition, the beneficial effect of boron onductility has also been observed in other ordered intermetallics,such as L12-ordered Ni3Si [36] and Ni3Fe [37], B2-orderedFeAl [38] and L10-ordered TiA1 alloys [39], although theeffect of boron is most pronounced in the L12 intermetallics.

4. Eliminating environmental embrittlement

One of the most important findings in the study ofdeformation and fracture of ordered intermetallics is thediscovery of the moisture-induced hydrogen embrittlement in

nickel and iron aluminides. Iron aluminides based on Fe3Aland FeAl exhibit very attractive properties, including lowdensity, low cost, and excellent oxidation and corrosionresistance, which make them promising for elevated-temperature structural applications in hostile environments. Amajor handicap in pursuing these developments is their poorductility and fracture resistance at ambient temperatures.Although aluminides had been known to be brittle for a longtime, the major cause of this phenomenon was not identifieduntil 1989 [40]. Liu et al. [40,41] first demonstrated that FeAland Fe3Al containing less than 40 at% Al are intrinsically quiteductile and that the poor ductility commonly observed in airtests at ambient temperatures is caused mainly by an extrinsiceffect – moisture induced environmental embrittlement. Thetensile stress–strain behavior and fracture surfaces of Fe-36.5Al (at%) in various testing environments at room tem-perature are illustrated in Fig. 5 [40], which vividly demon-strates the environmental embrittlement. The aluminide shows

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a ductility of �2% when tested in air, which increases to�6% in vacuum and further to �17.6% in dry oxygen. Theimprovement in ductility is accompanied by a change in thefracture mode from transgranular cleavage in air/water tomainly intergranular in dry oxygen, suggesting that cleavageplanes in FeAl are more susceptible to embrittlement than thegrain boundaries [40].

Mechanistically, the environmental embrittlement in ironaluminides, as well as in other aluminides, can be explained bya chemical reaction involving the interaction of aluminumatoms with moisture in air and the generation of hydrogen, asdescribed by the simple chemical reaction [40,41]:

2Alþ3H2O-Al2O3þ6H (1)

The high-fugacity atomic hydrogen, resulted from the abovereaction between aluminum atoms and H2O in moist air,penetrates into crack tips and causes brittle cleavage crackpropagation and premature failure. This embrittlement issimilar to hydrogen embrittlement, except that hydrogen hereis generated by decomposition of moisture in air. The ductilityin the vacuum test is higher than that in the air test but lowerthan that in the oxygen test, indicating that even the residualmoisture in the vacuum test (say, 10�5 Torr) is enough tocause significant embrittlement as compared with the oxygentest. The observation that the best ductility was obtained in dryoxygen can be understood by considering the direct reaction ofaluminum atoms with oxygen [40,41]:

4Alþ3O2-2Al2O3 (2)

The reaction of aluminum atoms with oxygen competes withthe moisture reaction, thereby reducing the generation ofatomic hydrogen from Eq. (1). Similarly, Fe3Al alloys alsoshowed severe environmental embrittlement at room tempera-ture [41–43]. A Fe–25Al (at%) alloy shows a high ductility of�13% in dry oxygen and a reduced ductility of �4% in moistair at room temperature [42]. In comparison, the ductilityexhibits a significant drop to as low as 1% when hydrogen ischarged simultaneously with tensile testing. In addition, bothFeAl and Fe3Al alloys showed an increase in ductility with

Fig. 5. Stress–strain behavior and fracture surfaces of Fe-36.5 at% Al at roomtemperature in various environments [40].

increasing strain rate, in support of the moisture-inducedhydrogen embrittlement mechanism [42].Note that many other aluminides, such as Ni3Al, also suffer

severely from the environmental embrittlement [32,44,45].Polycrystalline samples of a boron-free Ni3Al alloy exhibitedroom temperature tensile ductilities of 6–9% in water, 11–13%in air, and 48–51% in oxygen, demonstrating that the grainboundaries in this alloy are not intrinsically brittle; theybecome embrittled only in the presence of moisture. It shouldbe pointed out that moisture embrittlement paths in the nickeland iron aluminides are different. In the case of Ni3Al, thedecomposed hydrogen goes into the materials from grainboundaries, while for the Fe3Al and FeAl, hydrogen goes intothe materials both from grain boundaries and bulk materials,resulting in embrittlement. The difference is due to the fastdiffusion in both grain boundaries and bulk in iron aluminideswith open body-centered cubic (bcc) structures.In parallel to the mechanistic understanding of the embrit-

tlement mechanism in aluminides, significant alloy designefforts have been devoted to alleviating or reducing theembrittlement through application of physical metallurgicalprinciples. The schemes used to improve the ductility of thealuminides include [16,46]: (1) surface oil coating; (2) forma-tion of protective oxide scales on surfaces; (3) refinement ofgrain structure by thermo-mechanical treatments or second-phase particles; (4) enhancement of grain-boundary cohesionby micro-alloying with boron; and (5) reducing hydrogensolubility and diffusivity by alloying additions (possiblyboron). In particular, a significant improvement in ductilitycan be achieved when multiple schemes are used simulta-neously. For example, the ductility of the binary Fe3Al alloycan be significantly improved by a combination of thermo-mechanical treatments and alloying additions. The tensilecurves of the binary Fe3Al and multi-component Fe3Al-based(28 at% Al) alloys are compared in Fig. 6 [47]. The ductileFe3Al-based alloys shows a ductility of as high as �16%when tested in air at room temperature, much high than thatobtained in the binary Fe3Al alloy (�4%) tested in the samecondition [47].

5. Microstructure optimization

In general terms, it is well established that the microstructureoptimization, such as the grain-size refinement, is a potent wayto improve ductility of many materials [48,49]. Gammatitanium aluminides have attracted a great deal of attentionfrom the aerospace community due mainly to their attractiveproperties for high-temperature structural applications, includ-ing low density, good oxidation resistance, and high-temperature strength [3]. Generally, the room-temperaturetensile elongation of gamma alloys ranges from 0.3%–4%,and yield and tensile strengths vary from 250–600 MPa and300–700 MPa, respectively, depending on the alloy composi-tion, processing method, microstructure, and grain size. It hasbeen found that the mechanical properties of the gammatitanium aluminides are very sensitive to the microstructures,which in turn, depends strongly on the control of processing

Fig. 6. Improvement in tensile properties of Fe3Al in air by changing thecomposition and processing [47].

Fig. 7. (a) Room-temperature tensile elongation as a function of grain size and(b) room temperature yield strength as a function of interlamellar spacing inTiAl alloys [50].

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and appropriate alloy modifications. In order to identify keymicrostructural parameters governing the mechanical proper-ties, the effects of grain size and lamellar spacing on themechanical properties of a Ti–47Al–2Cr–2Nb (at%) alloy wereinvestigated. The yield strength and tensile elongation at roomtemperature are illustrated in Fig. 7 as a function of grain size(d) and interlamellar spacing, respectively [50]. In Fig. 7a, thetensile elongation at room temperature is strongly dependenton the grain size, showing an increase in ductility withdecreasing grain size [50]. The yield strength of TiAl alloysat room temperature was analyzed by the Hall–Petch equationbased on both grain size and interlamellar spacing. No linearrelationships existed for the grain size plot, but an excellentrelationship was found for the interlamellar spacing plot, asshown in Fig. 7b by Liu et al for the yield strength at roomtemperature [50]. The plot clearly indicates that the interla-mellar spacing is the key parameter controlling the yieldstrength of these two-phase TiAl alloys [50]. The inverserelationship between the ductility and grain size in Fig. 7a canbe explained by assuming the presence of a grain-size micro-crack and subsequently estimating the stress intensity of thecrack and the resultant ductility due to crack propagation [51].It is known that the critical stress concentration to create amicro-crack is closely related with the grain size of the alloys.Micro-cracks are expected to form due to the strain incompat-ibility at or across the lamellar grain boundaries and theanisotropic tensile properties [51]. In this process, ductility isprimarily controlled by the amount of general yielding, whichwould be greater for small grain sizes [51].

In addition, the effect of the lamellar structure on themechanical properties of single-orientation TiAl crystals (Ti-49.3 at% Al), referred to as polysynthetically twinned (PST)crystals, has been studied [11,52]. The yield stress andelongation-to-fracture of these PST specimens are verysensitive to their lamellar orientation with respect to thetensile axis. The room-temperature yield strength of the PSTspecimens is shown in Fig. 8a as a function of the angle Φbetween the lamellar boundaries and tensile axis [52]. The

yield strength is strongly dependent upon the angle Φ,although the orientation dependence is not exactly symme-trical with respect to Φ¼451. The highest yield strength ofover 400 MPa was obtained for the Φ¼901, while the lowestyield strength of as low as 100 MPa was obtained for theangle Φ in the range of 30–601. The variation of the yieldstrength with Φ can be explained by the two main deforma-tion modes (i.e., the true twinning of the {111}o112]-typeand slip along o110]), both of which are operational for thehard and easy modes of deformation. Furthermore, the tensileelongation-to-fracture is also strongly dependent on Φ, asillustrated in Fig. 8b [52]. The tensile elongation is also notsymmetrical with respect to Φ¼451, and the largest tensileelongation of as high as 20% was observed for a specimen

Fig. 8. (a) Yield stress and (b) tensile elongation of PST crystals of TiAl as afunction of the angle Φ [52].

Fig. 9. Tensile stress–strain curves of polycrystalline YAg and YCu tested inair at 22 1C [55]. The stress–strain plot for commercial 3105 aluminum alloy isplotted for comparison.

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with the Φ¼311, which is far larger than any other reportedvalues of room-temperature ductility of TiA1 intermetallics.A low ductility of �1% was found for the N orientation with

Φ¼901. It has been suggested that reducing the N orientation(see Fig. 11b) in polycrystalline aggregates should improveductility [52].

6. Control of phase stability

B2-ordered intermetallics have attracted a significant interest duemainly to their strength, sometimes even at high temperatures.However, their room-temperature polycrystalline ductility is oftenlacking due either to extrinsic effects, such as the moisture-inducehydrogen embrittlement, or to intrinsic behavior, such as thelacking of sufficient slip modes [1]. It is known that the Von Misescriterion for polycrystalline deformation requires five independentplastic deformation modes [53]. For B2-ordered intermetallics, thecommon o1004 slip mode is insufficient to satisfy the VonMises criterion and the o1114 slip mode is highly needed.However, the o1114 slip mode is rarely active in B2-orderedintermetallics with a high anti-phase boundary (APB) energy [54].In recent years, significant tensile ductility has been

achieved in highly ordered B2-type RM compounds (whereR¼a rare earth and M¼a main group or transition metal),such as YAg, YCu, and DyCu [55]. These RM phases are linecompounds that do not deviate from ideal stoichiometry andcrystallize in the CsCl (B2)-type structure. The stress–straincurves for YAg and YCu are shown in Fig. 9, and the curve fora common commercial Al alloy is also included for compar-ison [55]. The YAg alloy exhibits an elongation of as high as20% before fracture, which is extraordinary for an intermetalliccompound. The room-temperature plane-strain fracture-toughness (KIC) of DyCu and YAg are determined to be ashigh as 25.5 and 19.1 MPa m1/2, respectively, the valuescomparable to those of commercial aircraft aluminum alloys[56]. For comparison, KIC is 5–6 MPa m1/2 for the B2-orderedNiAl intermetallic [57].Transmission electron microscope (TEM) studies shows that

dislocations in both polycrystalline and single-crystal YAg arepredominantly the o1114 type, with a smaller number ofo1104 -type dislocations [55], which is really surprising.First, no slip lines corresponding to the o1114 slip wereobserved in the YAg single-crystal tensile-test specimens eventhough a relatively high resolved force to initiate o1114 slipon either the {110} or {112} planes was applied in allorientations. Second, room-temperature slip in the o1114direction of B2-ordered intermetallics is normally seen only incompounds with low ordering energies [55], because thepassage of a single ½o1114 dislocation introduces anAPB in the lattice, which raises the crystal energy. The first-principles calculation indicates that the unstable stacking-faultenergies, defined as the maximum energy for bicrystalsslipping homogeneously [58], associated with slip along theo1004 direction in YCu and YAg (300–600 mJ m�2) arecomparatively lower than the APB energies of these materials(700–1100 mJ m�2) [59,60]. In contrast, for NiAl, theunstable stacking fault energies (41300 mJ m�2) are greaterthan the APB energies (800–1000 mJ m�2) [59,60], suggest-ing a qualitatively different behavior of the RM materials.

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7. Promoting transformation-/twining-induced plasticity

In designing crystalline ceramics and metals, the concepts oftransformation-induced plasticity (TRIP) and twinning-inducedplasticity (TWIP) have often been adopted to increase theductility and work-hardening properties [61–64]. Similarly,attempts have been devoted to introducing the TRIP and TWIPeffects into ordered intermetallics.

Among B2-ordered intermetallics, near-equiatomic Ti–Nialloys exhibit superior shape memory and superplastic

Fig. 10. Stress–strain curves of the Zr50Co50, Zr50Co39Ni11 and Zr50Co36Ni14alloys at room temperature [68].

Fig. 11. Bright field images and corresponding electron diffraction patterns of the

properties due to a martensitic transformation from a cubicprimitive B2 to a monoclinic B190 phase [65]. It is thereforeexpected that the tensile ductility of other intermetalliccompounds at room temperature can be enhanced by thedeformation-induced martensitic transformation. The polycrys-talline B2-ordered ZrCo exhibits ductile behavior with a tensileelongation of �7% at room temperature, while the B33-ordered ZrNi is quite brittle [66]. It was found that themartensitic transformation from B2 to B33 structures takesplace by substitution of 14–50 at% Ni for Co [67]. Therefore,it is expected that the tensile ductility at room temperature ofternary Zr–Co–Ni alloys can be further enhanced by themartensitic transformation. The tensile stress–strain curvesfor the Zr–Co–Ni alloys are as shown in Fig. 10 [68]. TheZr50Co50 alloy exhibits a yield strength of 240 MPa and a totalelongation of 6%, while the total elongation of allZr50Co40Ni10, Zr50Co39Ni11, Zr50Co38Ni12 and Zr50Co37Ni13alloys is over 10%, much larger than that observed for theZr50Co50 alloy. Especially, the Zr50Co39Ni11 alloy exhibited anexcellent ductility with an elongation of as high as 23%. Fig.11 shows the TEM micrographs of the Zr–Co and Zr–Co–Nialloys before and after deformation [68,69]. Before deforma-tion (Fig. 10 a and b), the Zr50Co39Ni11 alloy consists of asingle phase with B2 structure, the average grain size beingapproximately 5 mm. After deformation at room temperature(Fig. 10 c and d), lenticular martensite was observed in the B2

Zr50Co39Ni11 alloy (a and b) before and (c and d) after deformation [68,69].

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parent phase, in addition to dislocations with the o1004B2-type Burgers vector. The martensite variants grow along the{100}B2 planes in the parent grain. The orientation relationshipbetween the B2 parent phase and the B33 martensite wasdetermined to be [001]B2//[100]B33, (010)B2//(021)B33, and(110)B2//(010)B33 [68,69].

Mechanical twinning in TiAl alloys has been observed tooccur under tensile and compressive mechanical testing con-ditions [3]. Interestingly, Mn was found to be beneficial forenhancing the bending ductility of ordered TiAl alloys bypromoting the generation of deformation twins. As shown inFig. 12a, the Ti–48.4Al–1Mn (at%) alloy is able to bendsignificantly at room temperature [70], and the elongation atthe tensile surface of the specimen was estimated to be �5%.The microstructure of the polished tension-side surface of thespecimen after bending is shown in Fig. 12b [70]. A plenty ofstraight lines were observed, which are the deformation twins.Note that the deformation twins was not be observed abso-lutely in Al-rich TiAl alloys and found only in a small amountfor the Ti-rich TiAl alloys.

TEM microstructures of the Mn-free and Mn-added TiAlalloys after deformation are shown in Fig. 13a and b,respectively [71]. From the TEM observations, it is seen thatthe principal mechanism of room-temperature deformation inMn-added TiAl is twinning (Fig. 13b), which is completelydifferent from the deformation mechanism of the Mn-free TiAl

Fig. 12. (a) Room-temperature bend work of the Ti–48.4Al–1Mn (at%) alloy,and (b) microstructure of the polished tension-side surface of the specimen[70].

specimen, whose mechanism is considered to be an interactionof super-dislocations (Fig. 13a) [71]. The origins for theenhancement of twinning by the Mn additions in the room-temperature deformation can be twofold [71]. First, Mn playsan important role in stabilizing the twin structures by pinningthe twin dislocations at the twin boundaries, which makes thenucleation process of twin dislocations easier and is critical forthe twin deformation process. Second, the Mn additionsdecrease the stacking fault energy of the TiAl alloys, whichmakes the twin deformation energetically easier than plasticdeformation by super-dislocations [71].

8. Conclusion remarks and future work

Ordered intermetallics have attractive properties for struc-tural applications at elevated temperatures in hostile environ-ments; however, low ductility and brittle failure generally limittheir use as engineering materials. This paper reviews therecent progress in designing ductile ordered intermetallics andsummarizes the mechanistical strategies used to improve thetensile ductility of ordered intermetallics, including control ofordered crystal structures, engineering grain-boundary struc-ture and chemistry, eliminating environmental embrittlement,microstructural optimization, control of phase stability, andpromoting transformation-/twining-induced plasticity. Thesestrategies enable the ordered intermetallics to be tailored tooffer enhanced tensile ductility, which is critical in manyprocessing or load-bearing applications. It is predicted thatsome of these approaches are also applicable to designing newhigh-strength and high-ductility ordered intermetallics.Recently, ordered intermetallic alloys with improved proper-

ties have drawn a great deal of attentions from industries, andmany intermetallics alloys have been successfully applied forstructural and corrosion applications. One prominent exampleis the use of the Ti–(44–46)Al–(2–6)Nb–(0.2–4)(Cr,W,B,Y)(wt%) alloy for Boeing 787 applications [72,73]. Since 2009,this low-density and high-strength TiAl alloy has been used asstages 6 and 7 turbine blades in Boeing 787-Genx, which hasled to enormous benefits, including reducing the weight by 400lb, cutting the fuel consumption by 15%, and cutting thegreenhouse gas emission by 50% [72,73]. Another example isthe use of Ni, Fe, and Ti-based aluminides as porous materialsfor filtering gas and liquid solutions used for productions ofacids and gases [12,13]. An industrial filtration application ofthe porous TiAl alloys has been carried out for the removal ofsuspended substances from liquid titanium tetrachloride. Byapplying the TiAl filter, a continuous filtering service for overthree months with the solid-to-liquid ratio dropping from alevel of 5% down to less than 0.1% has been achieved [12,13].In order to further speed up the development of new ductile

ordered intermetallics and eventually achieve the geneticdesign of advanced ordered intermetallics in different alloysystems, future efforts should be focused on the combinationof the ductilizing strategies with the state-of-the-art computa-tional calculations, such as thermodynamic and first-principlescalculations as well as molecular dynamic simulations, and thecutting-edge material characterizations using Cs-corrected

Fig. 13. TEM images of the (a) Mn-free and (b) Mn-added TiAl aftercompression [71].

Z.B. Jiao et al. / Progress in Natural Science: Materials International 26 (2016) 1–1210

transmission electron microscope and 3D atom probe tomo-graphy. In addition, from the application point of view, it iscritical to improve the ductility of the intermetallics withoutscarifying other useful properties, including the room- andelevated-temperature strength, creep resistance, welding prop-erties, castability, machinability, corrosion resistance, oxida-tion resistance and fatigue properties. Future studies maycontribute more to an in-depth understanding of the composi-tion–microstructure–property relationship and provide possibleapproaches to improve the overall properties of the orderedintermetallics for engineering applications.

Acknowledgments

This research was supported by the internal funding fromCity University of Hong Kong, Hong Kong (account No.CityU 9380060) and the Research Grant Council (RGC) of theHong Kong government, through the General Research Fund(GRF) with the account numbers of CityU 11209314 and11205515.

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Dr. C. T. Liu (Chain T. Liu), an academician ofNational Academy of Engineering (NAE), USA, anda foreign member of Chinese Academy of Engineer-

ing (CAE). He was employed at the Oak RidgeNational Laboratory (ORNL) from 1967 to 2005,served as a Professor and Distinguished ResearchProfessor in the Department of Materials Science

and Engineering at the University of Tennessee from2005-2008, and is now a University Distinguished

professor at City University of Hong Kong.

Z.B. Jiao et al. / Progress in Natural Science: Materials International 26 (2016) 1–1212

Dr. Liu is a world authority in the field of structural intermetallics. His workon the microalloying effect of boron and ductility improvement of polycrystal-line Ni3Al and other intermetallics greatly stimulated a worldwide interest inordered intermetallics in the early 1980s. In recent years, Dr. Liu has devotedhis major effort on the glass formation and alloy design of bulk metallic glasses(BMGs) with superior glass forming ability.He was awarded with the 1988 E. O. Lawrence Award in the category of

Materials Research by the Department of Energy (DOE), USA.

Dr. Liu is an author and a co-author of more than 500 technical publications,and his papers have been cited extensively by the materials community. Interms of the impact of scientific papers, Dr. Liu was identified by ISI in 1995as one of the top five individuals in the world, whose papers were highly citedin journals of materials sciences and engineering.