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Materials Science and Engineering A 441 (2006) 142–148

Superplasticity of single phase Ni–42Al intermetallics with large grains

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

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

Received 24 March 2006; received in revised form 26 July 2006; accepted 3 August 2006

bstract

Superplasticity was found in single phase Ni–42Al alloy with initial grain size of 200 �m under an initial strain rate of 1.25 × 10−4 s−1 to× 10−3 s−1 in temperatures ranging from 1000 ◦C to 1100 ◦C. A maximum elongation of 306% was obtained under an initial strain rate of× 10−3 s−1 at 1075 ◦C. Optical metallography (OM) observation showed that the average grain size was refined during superplastic deformation

rom initial 200 �m to less than 20 �m. Transmission electron microscopy (TEM) observation showed that an unstable subgrain boundary networkas formed during superplastic deformation. The subgrain boundaries were transformed into low and high angle grain boundaries by absorbingliding dislocations. The large-grained superplastic phenomenon could be explained by continuous dynamic recovery and recrystallization (CDRR).

2006 Elsevier B.V. All rights reserved.

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eywords: Binary NiAl intermetallics; Large-grained; Superplasticity; Continu

. Introduction

The intermetallic compound NiAl which crystallizes in the2 (CsCl) structure is regarded as a promising candidate forigh temperature applications because of its low density, highelting point and good oxidation resistance. In the past decades,

ignificant improvement in high temperature strength, process-ng and design methodology for NiAl alloy has been achieved.owever, limited room temperature ductility and fracture tough-ess as well as poor elevated temperature strength still seriouslyinder its commercial use. Meanwhile, the poor ductility at lowemperature makes NiAl very difficult and tedious to fabricate,hich also impedes its commercial use.Superplastic forming has been proven to play an important

ole in the fabrication of metals with poor ductility in indus-ry. The superplastic deformation of fine-grained metals usuallyequires fine grain size of less than 6 �m because grain bound-ry sliding (GBS) is considered to contribute to the superplasticeformation. Recently, in our laboratory superplasticity was

ound in large-grained Fe3Al and FeAl alloys [1–5], Ni3Al-ased intermetallic alloy [6–8], �-TiAl alloy [9] with the initialrain sizes being 100–300 �m, 10–30 �m and 95 �m, respec-

∗ Corresponding author. Fax: +86 21 62803241.E-mail address: dllin@sjtu.edu.cn (D. Lin).

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921-5093/$ – see front matter © 2006 Elsevier B.V. All rights reserved.oi:10.1016/j.msea.2006.08.011

ynamic recovery and recrystallization; High temperature deformation

ively. It was shown that large-grained aluminides [3–5] and-TiAl alloy [9] exhibit all deformation characteristics of con-entional fine-grained superplastic materials, without the pre-equisites of fine grain size and grain boundary sliding. Thebserved superplastic behavior can be explained by continuousynamic recovery and recrystallization (CDRR) [3–5].

Succeeding in research on large-grained superplasticity ineAl, Fe3Al and TiAl, the same phenomenon was found ini–40Al [10], Ni–45Al [11] and Ni–48Al [12–15] by our lab-ratory, and in stoichiometric NiAl by other laboratory [16]. Inrder to clarify the effect of Ni content and anti-site defects on thelongation of Ni-rich NiAl alloys during superplastic deforma-ion, another Ni-rich NiAl alloy, Ni–42Al is selected for study.n this paper, the superplastic deformation behavior in Ni–42Alith grain sizes of 200 �m is reported. The tensile behaviorf Ni–42Al under an initial strain rate of 1.25 × 10−4 s−1 to× 10−3 s−1 at a temperature range of 1000–1100 ◦C was exam-

ned. The microstructure evolution during and after deformations revealed. Combined with previous work in FeAl, Fe3Al andiAl the mechanism of superplastic deformation is discussed.

. Experimental

The NiAl alloy bars with a size of Ø 20 mm × 200 mmere prepared by induction melting under vacuum condition

nd Engineering A 441 (2006) 142–148 143

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The strain rate sensitivity index, m value, was measured usingthe cross-head speed change test. On the equation of

σ = Kε̇m (1)

J. Hu, D. Lin / Materials Science a

sing high purity starting materials (99.98% Ni and 99.99%l). After homogenizing for 5 h at 1100 ◦C, the bars were put

nto stainless steel pipes and hot pressed. A total strain of30% was calculated from the ratio of the initial and final

eights of the samples. After compression, the samples wereooled to 500 ◦C, held for 5 h, and then cooled inside the fur-ace. The samples were heat treated at 1000 ◦C for 1 h andooled in vacuum. Tensile specimens with a gauge section ofmm × 4 mm × 1.5 mm were cut from the samples with elec-

rodischarge machine. The tensile tests were performed undermbient conditions using a Shimazu AG-100KNA test machinequipped with a three-zone resistance furnace at a constantate of cross-head displacement. Initial strain rates range from× 10−4 s−1 to 2.5 × 10−2 s−1, and the testing temperaturesere from 1000 ◦C to 1100 ◦C. Specimens were held at test

emperature for 15 min before deformation and water quenchedfter deformation or fracture. The metallographic specimensere etched with a solution of 8CH3COOH + 4HNO3 + 1HCl

volume portion). Automatic generation and indexing of elec-ron back-scattered diffraction (EBSD) patterns were carriedut on an orientation imaging microscope produced by Oxfordnstruments Microanalysis Group, England, and equipped withback-scattered electron detector and Opal analysis software.eam scan mode was adapted with a step spacing of 1 �m overn area of 115 �m × 115 �m to 161 �m × 161 �m. The imagef microstructure was reconstructed by creating grain bound-ry maps from the EBSD pattern measurements. Designationf grain boundaries was based on a grain boundary criteria, ω,iven by the researcher. Misorientation angle θ is calculatedetween grid points in the scan field and compared with ω. Inhe paper, four criteria, 3◦ < ω < 5◦, 6◦ < ω < 10◦, 10◦ ≤ ω < 15◦,nd ω ≥ 15◦ were considered. By employing these criteria dur-ng the generation of the grain boundary maps, different imagesf the microstructure were constructed. TEM foils were pre-ared by twin jet polishing in the electrolyte of a 5% perchloriccid solution in methanol at −30 ◦C. TEM observations wereerformed on a JEM-200CX transmission microscope.

. Results

.1. Tensile behavior

Table 1 shows the effect of the testing temperature and initialtrain rate on the elongation. It is shown that the large-grained

able 1ffect of temperature and initial strain rate on elongation of Ni–42Al

nitial strain rate (s−1) Temperature (◦C)

1000 1025 1050 1075 1100

.25 × 10−4 234 200

.875 × 10−4 229

.5 × 10−4 183 185 276 237 219× 10−4 250 256.5 × 10−4 175 208 280 258× 10−3 220 270 300 306 220.5 × 10−3 218× 10−3 200

Fig. 1. Macrograph of specimens tested at 1075 ◦C.

i–42Al alloy possesses good ability of superplastic deforma-ion. In addition, it is also shown that the ability of superplasticeformation of Ni–42Al is lower than that of Ni–40Al [10], butigher than that of Ni–45Al [11] and Ni–48Al [12]. The max-mum elongation of 306% was obtained under an initial strainate of 1 × 10−3 s−1 at 1075 ◦C.

Fig. 1 shows a macrograph of the specimens tested in ten-ion at 1075 ◦C. The gauge section of specimen is still uniformt an elongation above 200%. Obvious necking could not bebserved until the specimen was near to fracture. This is alson evidence for the potential of superplastic deformation for thearge-grained Ni–42Al intermetallic.

Fig. 2. Stress–strain rate relationship for Ni–42Al.

144 J. Hu, D. Lin / Materials Science and Engineering A 441 (2006) 142–148

. true

m

m

Fr

Fig. 3. True stress vs

value can be calculated as

= ∂ ln σ

∂ ln ε̇(2)

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ig. 4. Optical micrograph of large-grained Ni–42Al before and after deformationespectively.

strain for Ni–42Al.

or every two strain rates m can be calculated by

= ln(σ1/σ2)

ln(ε̇1/ε̇2)(3)

: (a) before deformation; (b)–(e) deformed to 40%, 100%, 200% and 306%,

J. Hu, D. Lin / Materials Science and Engineering A 441 (2006) 142–148 145

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Fig. 5. The distribution of misorientation for Ni–42Al: (a

Fig. 2 shows the ln σ versus ln ε̇ over a wide strain rate rangef 2.5 × 10−5 s−1 to 2.5 × 10−2 s−1. The m value can be deducedrom the slope of the curves.

It can be seen from Fig. 2 that the m value varies from 0.25 to.33 under different testing conditions. Generally, it decreasesith increasing strain rate at a given temperature. However, theaximum elongation of 306% was found at a strain rate of× 10−3 s−1 with the m value of 0.27 at 1075 ◦C instead oft the lower strain rate of 1.25 × 10−4 s−1 corresponding to theaximum m value 0.33.To examine the flow behavior of Ni–42Al alloy, the curves

f true stress versus true strain were plotted, which were basedn the assumption of uniform deformation. Fig. 3 illustrateshe typical true stress versus true strain curve of the Ni–42Alnder different testing condition. It clearly shows that the maxi-um flow stress decreases as the temperature increases under a

iven initial strain rate, and the maximum flow stress increasess the initial strain rate increases at a given testing temperature.nd it also shows that the flow stress keeps relatively steady

n a narrow range of strain at some testing conditions, whicheans that a balance between dynamic hardening and softening

xists.

.2. Microstructure

Optical metallography (OM) shows that the resulting averagerain size in tensile specimens of investigated Ni–42Al alloyfter tensile testing gradually decreases with the increase of

(dai

re deformation; (b)–(d) deformed to 40%, 100%, 250%.

uperplastic deformation. Fig. 4 shows the optical micrograph ofi–42Al alloy during superplastic deformation under an initial

train rate of 2.5 × 10−4 s−1 at 1050 ◦C. The initial grain sizes about 200 �m and the shape of each grain is equiaxial beforeeformation (Fig. 4(a)). The grain boundaries appear splinterynd grain size increases slightly deformed to 40% (Fig. 4(b)).here are two kinds of grain boundary morphology shown inig. 4(c)–(e), one is deeply etched (marked A), and the other istched slightly (marked B), whose possible reason is that dif-erent misorientation exists among different grains. The averagerain size changes from about 200 �m before deformation tobout 20 �m after deformation to 306% (Fig. 4(e)).

No cavities are found in the microstructures after deforma-ion, even in the near fracture area.

Fig. 5 shows the distribution of misorientation in the samplesf Ni–42Al alloy before and after deformation. In all these dia-rams, misorientation angle is divided into four parts accordingo gray grades, which are 3–5◦, 5–10◦, 10–15◦ and >15◦.

Fig. 5(a) indicated that in the initial microstructure beforeeformation the majority of grain boundaries are HAGBs (highngle grain boundaries) with misorientation angles above 15◦.

hen deformed to 40%, the number of grain boundaries with< 10◦, especially θ < 5◦ increased prominently and that ofrain boundaries with misorientation angles above 10◦ decrease

shown in Fig. 5(b)). Fig. 5(c) and (d) showed that wheneformed to 100% and 250%, the number of LAGBs (lowngle grain boundaries) with both θ > 5◦–10◦ and θ = 10◦–15◦ncreased, but those with θ < 5◦ changed very little, or it could

146 J. Hu, D. Lin / Materials Science and Engineering A 441 (2006) 142–148

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Fig. 6. Typical TEM images of Ni–42Al during deformation: (a) and (b) bef

e considered to remain constant. On the contrast, the number ofAGBs increases gradually with the deformation and the mis-rientation distribute much more uniformly between θ = 10◦ and0◦when deformed to 250% than that before deformation.

The results of EBSD show that the density of low angle grainoundaries increases at the early stage of superplastic defor-ation and keeps constant during the deformation. The high

ngle grain boundaries increase gradually during the deforma-

ion, which indicate that subgrain boundaries were transformednto low and high angle grain boundaries by absorbing glid-ng dislocations during superplastic deformation. The resultsf EBSD also show that the dislocation glide and climb play

Ts

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eformation; (c) and (d) ε = 40%; (e) and (f) ε = 100%; (g) and (h) ε = 306%.

n important role during deformation and this is supported byransmission electron microscopy (TEM) observations.

The results of TEM observation on the dislocation configura-ions in superplastically deformed Ni–42Al intermetallics showhat there exists a great number of subgrain boundaries, whichivided the initial large grains into fine subgrains during defor-ation. The common feature of these boundaries is that they are

omposed of dislocation walls and dislocation networks. Typical

EM images of substructure of Ni–42Al intermetallics duringuperplastic deformation are shown in Fig. 6.

Fig. 6(a) and (b) show that the grain boundary is straightnd very few dislocations exist in the grains in the specimen

J. Hu, D. Lin / Materials Science and En

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Acknowledgment

ig. 7. Subgrain boundary rotation during superplastic deformation in Ni–42Al.

efore deformation. When the sample was deformed to 40%,islocation arrays and long segments as well as a few subbound-ries were observed, as shown in Fig. 6(c) and (d). Fig. 6(e)nd (f) show that when the strain was increased to 100%, moreubboundaries appeared, and well-developed subgrains coulde found. In the microstructure of the specimen deformed toracture with elongation of 306%, it was found that the ini-ial grains are divided into subgrains with an average size ofbout 3–8 �m, as shown in Fig. 6(g) and (h). Meanwhile, newubboundaries are still formed continuously due to dislocationslimbing and reaction. In our TEM observations, dislocationrrays, subboundaries can be found in all deformed specimensith the strain from 40% to 306%, and the classic recrystalliza-

ion nucleus with high angle boundary could not be observed inll deformed specimens.

The results of TEM observation also show that subgrainoundary gliding, migration and rotation took place duringuperplastic deformation in Ni–42Al intermetallics, as shown inig. 7, which can accommodate plastic deformation and enable

he superplastic flow to proceed.

. Discussion

As mentioned before, the flow behavior in the superplas-ic deformation of the Ni–42Al is similar to those in conven-ionally fine-grained superplastic material. But there are some

arked differences between two kinds of superplastic deforma-ion. Firstly, conventionally fine-grained superplasticity resultsrom the GBS, and this requires that the grain size of the alloys no more than 6 �m. But in the Ni–42Al, the mean grain sizes 200 �m. This suggests that the grain size is not the dominantactor for the observed superplasticity phenomenon. Moreover,he superplastic deformation of alloys with fine grains, the grainize keeps nearly constant or increases slightly. On the contrary,n the Ni–42Al, the grain size is refined during deformation.

herefore, we can deduce that a new mechanism different fromBS is operating during the superplastic deformation in large-rained Ni–42Al alloys. F

gineering A 441 (2006) 142–148 147

The facts that there are subboundaries in the deformed spec-mens and the grain size decreases during deformation indicatehat subboundaries are changing during superplastic deforma-ion. In addition, the recrystallization nucleus cannot be foundn all deformed specimens. Therefore, we can deduce that aontinuous recovery and recrystallization takes place duringuperplastic deformation in the Ni–42Al.

EBSD and TEM examinations indicate that dislocation playsn important role during the superplastic deformation. At thearly stage of deformation, dislocation sliding is dominant. Withurther deformation, dislocation climbing occurs and the dislo-ations arrange themselves as subgrains. The subgrain bound-ries absorb the gliding dislocations coming from the subgrainsnd thus the misorientation between the subgrains increases, ando they are transformed into low and eventually high angle grainoundaries. This process absorbs deforming energy and enablesuperplastic deformation available. The conclusion that the sub-rain boundaries can absorb the dislocations was based on notnly the analysis on the dislocation arrangement in the grainoundaries but also the fact that the density of dislocations inhe grains or subgrains is very low.

Similar phenomena have also been observed in FeAl [2],e3Al [3] and TiAl [5]. They show the same characteristics ofubgrain evolution during deformation as Ni–42Al in this work.o it is reasonable to deduce that the superplastic mechanism is

he same. We can deduce that a continuous dynamic recovery andecrystallization took place dynamically during the deformationn the large-grained Ni–42Al as happens in iron aluminides [3–5]nd TiAl alloys [9] and this is probably a dominant mechanismor large-grained superplastic deformation.

The maximum tensile elongation of Ni–42Al alloy is lowerhan that of Ni–40Al but higher than that of Ni–45Al andi–48Al alloys possibly because of the different content of anti-

ite point defects in different Ni-rich NiAl alloys [17].

. Conclusions

1) The Ni–42Al intermetallics with a mean grain size of200 �m exhibits superplasticity under a strain rate of1.25 × 10−4 s−1 to 2 × 10−3 s−1 in a temperature rangeof 1000–1100 ◦C. The maximum elongation to fracture of306% with m value of 0.27 was obtained at 1075 ◦C underan initial strain rate of 1 × 10−3 s−1.

2) The grain size was refined during deformation from 200 �mto less than 20 �m. The mechanism of superplasticity inthe large-grained Ni–42Al was proposed to be continuousdynamic recovery and recrystallization.

3) The maximum tensile elongation of Ni–42Al alloy is lowerthan that of Ni–40Al but higher than that of Ni–45Al andNi–48Al alloys possibly because of the different content ofanti-site point defects in different Ni-rich NiAl alloys.

This work was supported by the National Natural Scienceoundation of China.

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48 J. Hu, D. Lin / Materials Science a

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