Materials Science and Engineering A 371 (2004) 113–118

Microstructural evolution of superplasticity in large-grainedNi–48Al intermetallics

Jing Hua,b, Dongliang Lina,∗a School of Materials Science and Engineering, Open Laboratory of Education Ministry of China for High Temperature Materials and Tests,

Shanghai Jiao Tong Univerisity, Shanghai 200030, PR Chinab Department of Mechanical Engineering, Jiangsu Polytechnic University, Changzhou 213016, Jiangsu PR China

Received 17 June 2003; received in revised form 1 October 2003

Abstract

Superplastic behavior has been found in large-grained Ni–48Al intermetallics with grain size of 200�m. The large-grained Ni–48Al alloyexhibit deformation characteristics of most fine-grained superplastic alloys without the usual pre-requisites of a fine grain size and grain bound-ary sliding (GBS). Metallographic examination (OM) has shown that the average grain size of large-grained intermetallics decreases duringsuperplastic deformation and a much finer grain size could be obtained after superplastic deformation. Transmission electron microscopy(TEM) and electron back-scattered diffraction (EBSD) observations have shown that there are great numbers of subgrain boundaries which forma network and among which the proportion of low and high angle boundaries increases with the increase of strain. The observed superplasticphenomenon is explained by continuously dynamic recovery and recrystallization (CDRR). During superplastic deformation, an unstable sub-grain network forms and these subboundaries absorb gliding dislocations and transform into low and high angle grain boundaries. A dislocationgliding and climbing process accommodated by subgrain boundary sliding, migration and rotation, enables the superplastic flow to proceed.© 2003 Elsevier B.V. All rights reserved.

Keywords: NiAl intermetallics; Large grains; Superplasticity; Continuously dynamic recovery and recrystallization; High temperature deformation

1. Introduction

The intermetallic compound NiAl which crystallizes inthe B2 (CsCl) structure is regarded as a promising candidatefor high temperature applications because of its low density,high melting point, and good oxidation resistance. In thepast decades, significant improvement in high temperaturestrength, processing, and design methodology for NiAl hasbeen achieved. However, limited room temperature ductilityand fracture toughness as well as poor elevated temperaturestrength still seriously hinder its commercial use. Mean-while, the poor ductility at low temperature makes NiAlvery difficult and tedious to fabricate, which also impedesits commercial use.

Superplastic deformation has become an important wayto fabricate the metals with poor ductility, which is tra-ditionally found in fine-grained metals and alloys, wheregrain boundary sliding (GBS) is considered to be respon-

∗ Corresponding author. Tel.:+86-21-62932544;fax: +86-21-62820892.

E-mail address: dllin@sjtu.edu.cn (D. Lin).

sible for the superplasticity. In recent years, superplastic-ity has been found in large-grained Fe3Al [1,3,5,6], FeAl[2,4–6], Ni3Al [7], �-TiAl [8], Ni–40Al [9], and Ni–45Al[10] alloys in our laboratory. Meanwhile superplasticity wasalso found in stoichiometric NiAl alloys[11] and borondoped Fe3Si [12] with large grains. It has been shown[1–12]that large-grained intermetallics exhibit most deformationcharacteristics of conventional fine-grained superplastic ma-terials, without the prerequisites of a fine grain size andgrain boundary sliding. The observed superplastic behav-ior was explained by continuously dynamic recovery andrecrystallization (CDRR). Recently, we found superplastic-ity in large-grained Ni–48Al intermetallics at initial strainrates ranging from 1.25× 10−4 to 2× 10−3 s−1, and testingtemperatures ranging from 1025 to 1100◦C. The maximumelongation of 188.2% to fracture was obtained under an ini-tial strain rate of 1.125× 10−3 s−1 at 1100◦C [13]. It hasbeen shown that large-grained Ni–48Al alloy exhibits thesimilar deformation characteristics as large-grained Ni–40Aland Ni–45Al intermetallic alloys. However, the maximumelongation to fracture increases with the increase of Ni con-tent. We are interested to know if the difference results

0921-5093/$ – see front matter © 2003 Elsevier B.V. All rights reserved.doi:10.1016/j.msea.2003.10.245

114 J. Hu, D. Lin / Materials Science and Engineering A 371 (2004) 113–118

from the different Ni content in Ni-rich NiAl alloys, which,in turn, leads to different content of anti-site point defects[14,15]. In our laboratory, transmission electron microscopy(TEM), electron back-scattered diffraction (EBSD) observa-tions, and point defect analysis have been used to investigatethe microstructural evolution during superplastic deforma-tion and the mechanisms of superplasticity in Ni-rich NiAlintermetallic alloys. A part of the results presented in thispaper to show the microstructural evolution and the mecha-nism of superplastic deformation in large-grained Ni–48Alintermetallic alloy.

2. Experimental

The NiAl alloy bars with a size of∅ 20 mm× 200 mmwere prepared by induction melting under vacuum con-dition using high purity starting materials (99.98%Ni and99.99%Al). After a controlled thermo-mechanical process-ing (CTMP), the samples were cooled slowly inside thefurnace. Finally, to obtain the initially large grain size, thesamples were heat-treated at 1000◦C for 1.5 h and cooledslowly inside the furnace. The flat tensile specimens with agauge section of 8 mm× 4 mm× 1.5 mm were cut from thesamples with electro-discharge machine.

Tensile tests were performed on a Shimadzu AG-100KNAtesting machine equipped with a three-zone resistance fur-

Fig. 1. Optical micrographs of the homogeneous deformation zone of Ni–48Al sample deformed at 1075◦C andε = 1.5×10−3 s−1: (a) before deformation,(b) ε = 40% ,(c)ε = 96%, and (d)ε = 158%.

nace at a constant rate of cross-head displacement, initialstrain rates ranging from 1.25× 10−4 to 2× 10−3 s−1, andtesting temperatures ranging from 1025 to 1100◦C. Thetensile tests were interrupted when the elongation reached40, 96, and 158%, and specimens were fast cooled by wa-ter quenching to maintain the deformed structure at hightemperature. The metallographic specimens were etchedwith a solution of 8 CH3COOH+ 4 HNO3 + HCl (volumeportion). Automatic generation and indexing of electronback-scattered diffraction patterns were carried out on anorientation imaging microscope produced by Oxford In-struments Microanalysis Group, England, and equippedwith a back-scattered electron detector and Opal analysissoftware. Beam scan mode was adapted with a step spac-ing of 1�m over an area of 115�m × 115�m–161�m ×161�m. The image of microstructure was reconstructed bycreating grain boundary maps from the EBSD pattern mea-surements. Designation of grain boundaries was based ona grain boundary criteria,ω, given by the researcher. Mis-orientation angleθ is calculated between grid points in thescan field and compared withω. In the paper, four criteria,3◦ < ω < 5◦, 6◦ < ω < 10◦, 10◦ ≤ ω < 15◦, andω ≥ 15◦were considered. By employing these criteria during thegeneration of the grain boundary maps, different imagesof the microstructure were constructed. TEM foils wereprepared by the conventional twin jet polishing techniqueusing an electrolyte of 10% perchloric acid in methanol

J. Hu, D. Lin / Materials Science and Engineering A 371 (2004) 113–118 115

at about−30◦C. The foils were examined in JET-200CXelectron microscopes operating at 200 kV.

3. Microstructure evolution

Metallographic examination (OM) shows that the re-sulting average grain size in tensile specimens of inves-tigated intermetallic alloy after tensile testing graduallydecreases with the increase of superplastic deformation.Fig. 1 shows the optical micrograph of Ni–48Al before andduring superplastic deformation under an initial strain rateof 1.5×10−3 s−1 at 1075◦C, Fig. 1(a)shows that the initialgrain size is about 200�m and the shape of each grain isequiaxial before deformation. The grain boundaries appearsplintery and the average grain size increases slightly afterdeformation of 40% shown inFig. 1(b). There are two kindsof grain boundary morphologies shown inFig. 1(c), one isdeeply etched, and the other is etched slightly, where thepossible reason is that different misorientation exists amongdifferent grains. The average grain size is refined from about200�m before deformation to about 20�m after deformedto 158% as shown inFig. 1(d). No cavities were found in themicrostructures after deformation, even in the near fracturearea.

Fig. 2. Distribution of misorientation in Ni–48Al during superplastic deformation at 1075◦C: (a) original, (b) 40%, (c) 96%, and (d) 158%.

The most recently developed scanning electron mi-croscope (SEM)-based electron back-scattered diffractiontechnique has been used to investigate the microstructureevolution and the type of grain boundaries and the misori-entation between grains in Ni–48Al intermetallics duringsuperplastic deformation. From the Kikuchi patterns, thetype of grain boundaries can be identified and the misorien-tation between grain boundaries can be measured[16]. Themisorientation between grains as small as 3◦ is detectable[17]. It is found that the density of low angle grain bound-aries increases with the elongation in Ni–48Al, indicatingthat the dislocation glide and climb play an important rolein the deformation and this is supported by transmissionelectron microscopy observations.

Fig. 2 shows the distribution of misorientation in thesample of Ni–48Al alloy before and after deformation.In all these diagrams, misorientation angle is divided intofour parts according to gray scales, which are 3–5◦, 5–10◦,10–15◦, and >15◦.

Fig. 2(a) indicated that in the initial microstructure themajority of grain boundaries are HAGBs with misorientationangles above 15◦. When deformed to 40%, the number ofgrain boundaries withθ < 10◦, especiallyθ < 5◦ increasedprominently (shown inFig. 2(b)). Fig. 2(c) and (d)indicatedthat when deformed to 96 and 158%, the number of LAGBs

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Table 1The total number of grain boundaries with various misorientations

Misorientation Total number

Original ε = 40% ε = 96% ε = 158%

≤5◦ 4422 8813 8363 89856–10◦ 549 1342 2000 351911–15◦ 404 297 812 1674>15◦ 10411 6802 8850 11667

with both θ > 5–10◦ and θ=10–15◦ increased, but thosewith θ < 5◦ changed very little, or it could be considered toremain constant (Table 1).

This phenomenon resulted from the evolution ofsubboundaries. During deformation, the newest grainboundaries produced are subboundaries or LAGBs withmisorientation angleθ < 5◦, which transformed continu-ously to those withθ > 5◦, the LAGBs withθ > 5◦ trans-formed continuously to those with higher misorientationangle, and eventually transformed to HAGBs. The grainboundary number in each misorientation angle scale dependson the relationship between the speed of grain boundariesnewly produced to that transformed to higher misorientation.When the deformation is small, the former speed is higherthan the latter one, but with the increase of deformation,both speeds tend to be equal with the strain. Therefore, atthe larger deformation of 96 and 158%, the number of grainboundaries withθ < 5◦ keeps dynamically stable, and thosewith θ > 5◦ increased. Meanwhile, at larger strain, HAGBsincreased, which resulted from the transformation of un-stable LAGBs withθ < 15◦. Such a change correspondedto the formation of new grains, and led to refine the grainsize.

In order to have a clear comparison of the proportionof low and high angle grain boundary in Ni–48Al duringsuperplastic deformation, the number of misorientation inFig. 2 is transformed into frequency as shown inFig. 3.

Fig. 3. Frequency of misorientation in Ni–48Al during superplastic defor-mation at 1075◦C and initial strain rate of 1.5 × 10−3 s−1: (a) original,(b) 40%, (c) 96%, and (d) 158%.

The fact that the number of low angle grain boundaryis much higher in the sample after deformation than thatin the sample before deformation clearly indicates thatdislocation glide and climb play an important role duringdeformation, and subgrain boundaries and low angle grainboundaries divided the initial large grains into subgrainsduring deformation.

Subgrain boundaries with a misorientation lower than 2◦can be detected by TEM analysis. The results of TEM obser-vation on the dislocation configurations in superplasticallydeformed Ni–48Al intermetallics showed that there exist agreat number of subgrain boundaries, which divided the ini-tial large grains into fine subgrains during deformation. Thecommon feature of these boundaries is that they are com-posed of dislocation walls and dislocation networks.

TEM images of substructure are shown inFig. 4. Be-fore deformation, the grainboundaries are mostly straightand dislocation free (Fig. 4(a)) and few dislocations arefound within the grains (Fig. 4(b)). When the samples weredeformed to 96%, dislocation arrays and long segmentsconstituted the main deformation substructure, and somesubboundaries were observed, as shown inFig. 4(c) and (d).When the elongation was up to 158%, more subboundariesappeared, and well-developed subgrains could be found.In the microstructure of the specimen deformed to fracturewith elongation of 186.3%, it was found that the initialgrains are divided into subgrains with an average size ofabout 3–8�m. In our TEM observations, dislocation arraysand subboundaries can be found easily in all deformedspecimens with the elongation from 96 to 186.3%, and theordinary recrystallizing nucleus with high angle boundariesare hardly observed in all deformed specimens.

4. Discussion

The flow behavior in superplastic deformation of thelarge-grained Ni–48Al alloy is similar to that in conven-tionally superplastic alloys. However, there exist somesignificant differences in microstructural evolution andmechanism between the large-grained Ni–48Al alloy andconventionally fine-grained superplastic materials. Metal-lographic examination has shown that the average grainsize of the large-grained Ni–48Al alloy decreased dur-ing superplastic deformation and a much finer grain sizecould be obtained after superplastic deformation, while infine-grained super plastic alloys the grain size keeps nearlyconstant or has a slight growth.

EBSD and TEM studies indicate that the superplastic de-formation in Ni–48Al alloy is controlled by dislocation glideand climb. The gliding dislocations interacted and arrangedthemselves as subgrain boundaries through their climbingprocesses. Superplastic deformation needs not only disloca-tion glide but also the movement between subgrains, such assubgrain boundary sliding and migration, which should becontrolled by dislocation climb in the subgrain boundaries.

J. Hu, D. Lin / Materials Science and Engineering A 371 (2004) 113–118 117

Fig. 4. Typical TEM images of Ni–48Al during deformation, (a) and (b) before deformation, (c) and (d) 96% deformation, (e) and (f) 158% deformation,(g) and (h) 188% deformation.

The movement between subgrains results in the migrationof the initial high angle grain boundaries of the alloys. Dur-ing deformation, gliding dislocations can be trapped andabsorbed by subboundaries, and then subboundaries evolvecontinuously towards LAGBs, in further, to HAGBs.

Similar results in Fe3Al, FeAl, and TiAl alloys have beenobserved by OM, EBSD, and TEM[1–8], which suggest that

the same microstructural evolution in large-grained Ni–48Alintermetallics during superplastic deformation exists as inother large-grained intermetallic alloys.

In our previous work[1–10], we had proposed a contin-uously dynamic recovery and recrystallization model to ac-count for the superplastic behavior of initially large-grainedFe3Al and FeAl alloys. Microstructural evolution observed

118 J. Hu, D. Lin / Materials Science and Engineering A 371 (2004) 113–118

in this study clearly indicates that CDRR is the process oc-curring during superplastic deformation in the large-grainedNi–48Al alloy. At the early stage of deformation, disloca-tion gliding is dominant. With further deformation, climbprocess occurred and the dislocations arrange themselves assubgrain boundaries. The subgrain boundaries divide the ini-tial large grains into fine subgrains. The subgrain boundariesabsorb the dislocations coming from the subgrains and thusthe misorientation between subgrains increases. The con-clusion that the subgrain boundaries can absorb the disloca-tions was based on not only the analysis on the dislocationsin the grain boundaries but also the fact that the density ofdislocations in the grains or subgrains is very low.

In the present study, the microstructure evolution dur-ing superplastic deformation underwent the same process inboth Fe3Al and FeAl alloys[1–6]. The initial large grainsin large-grained Ni–48Al alloys were refined during defor-mation. Most of the refined grains are bordered by subgrainboundaries or low angle grain boundaries. The number ofsubgrain boundaries and LAGBs, as well as HAGBs in-creased with the strain.

5. Conclusions

The flow behavior during superplastic deformation ofthe large-grained Ni–48Al alloy is similar to that in con-ventionally superplastic alloys. However, there exist somesignificant differences in microstructural evolution andmechanism between the large-grained Ni–48Al alloy andconventionally fine-grained superplastic materials. Metallo-graphic examination has shown that the average grain sizeof the large-grained Ni–48Al alloy decreased during super-plastic deformation and a much finer grain size could be

obtained after superplastic deformation. EBSD observationshas shown the density of LAGBs and HAGBs increasedduring superplastic deformation, and TEM observationshas shown the formation of subboundaries and also theevidence of dislocation glide in the interior of grains. Theexperimental results support our previous model of contin-uous recovery and recrystallization for the mechanism ofsuperplastic deformation in large-grained intermetallics.

Acknowledgements

This work is supported by the National Natural ScienceFoundation of China.

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