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High tensile elongation of a directionally solidified NiAl multiphase alloy at high temperatures
Materials Science and Engineering A 396 (2005) 194–201
High tensile elongation of a directionally solidified NiAl multiphase alloy at high temperatures C.Y. Cui ∗ , J.T. Guo, Y.H. Qi, H.Q. Ye Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, China Received 27 October 2004; received in revised form 6 January 2005; accepted 28 January 2005
Abstract The ductility of a directionally solidified Ni–30Al–5Mo–0.5Hf (at.%) alloy was examined in various experimental conditions of temperatures and strain rates. Tensile tests revealed that the present alloy exhibited tensile elongation from room temperature to high temperature, in particular, the large elongations, 140–160%, were observed at high temperatures, depending on the temperature and strain rate. The value of strain rate sensitivity index, m, of 0.27 and the apparent activation energy of 413 kJ/mol were measured. Based on the observed mechanical behavior and microstructures, the mechanism responsible for this large elongation operative at high temperatures was suggested. © 2005 Elsevier B.V. All rights reserved. Keywords: Directional solidification; NiAl alloy; Ductility; Dynamic recovery and recrystallization
1. Introduction NiAl is a potential high temperature structural material for applications in gas turbine engines and has been extensively studied during the last decade [1]. However, the low ductility and toughness at room temperature and poor elevated temperature strength limit its application as an engineering material. Recent advances in single crystal alloy development and eutectic composite technologies make them promising for improving the creep strength of NiAl [2–4]. Creep strength comparable to superalloy Rene’ 80 has been achieved in a NiAl–Cr(Mo) composite with Hf addition [4]. Creep strength can be further improved by controlling the microstructure [5]. Unfortunately, this promising NiAl–Cr(Mo) alloy is unsuitable for a structural material due to its poor ductility at room temperature. Thus, the issue of obtaining ductility and toughness at room temperature remains a major block in the development of a structural NiAl alloy. To overcome this ductility deficiency, two possibilities exist. One approach is alloying NiAl to lower the ordering energy of the material so that it is easier to initiate alternate slip ∗ Corresponding author. Present address: National Institute for Materials Science, Tsukuba, Ibaraki 305-0047, Japan. Tel.: +81 29 859 2544. E-mail address: [email protected] (C.Y. Cui).
0921-5093/$ – see front matter © 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2005.01.043
systems during deformation. Previous results [6] indicate that ductility improvement can be achieved by the addition of Fe, Ga, and Mo in the NiAl single crystal. Another approach is introducing a ductile second phase to increase the ductility and toughness of the composite. The mechanism of ductile phase toughening can be generally divided into the intrinsic mechanism, such as slip transfer, and the extrinsic mechanism, such as crack blunting, deflection and bridging [7,8]. The latter approach is the focus in this study. According to the Ni–Al phase diagram [7], the Ni3 Al phase seems to be a good candidate for ductile phase toughening of NiAl, because NiAl–Ni3 Al alloy can be readily formed in situ from the melt or by heat treatment. In single crystal form, Ni3 Al exhibits over 98% elongations at room temperature for [0 0 1] oriented crystals [9], while in polycrystal form, Ni3 Al exhibits about 10% tensile elongation at room temperature in air [10]. The potential of the Ni3 Al phase in enhancing the fracture resistance of NiAl has been demonstrated by several investigators using different processing methods [8,11,12]. Misra et al. [8] showed that room temperature tensile ductility up to 9% is achieved in directionally solidified (DS) NiAl–Ni3 Al two-phase alloy with the continuous ␥ , whereas the ductility of DS single-phase [0 0 1] NiAl is negligible. The enhancement in ductility is attributed to a combination of slip transfer from the ductile ␥ to the brittle  and extrinsic mechanism
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such as crack blunting, deflection and bridging. However, the creep strength of these alloys with discontinuous morphology of ␥ phase is marginally higher than that of single-phase Ni–40Al and significantly lower than that of typical superalloys at the 1000–1200 K temperature regime. Consequently, a high temperature reinforcing phase will have to be incorporated into this ductile phase toughened system. For example, instead of ␥ , Mo might be used as a reinforcing element. Mo not only has a high melting point (2890 K), but certain Mo alloys are very creep resistant. Further, Mo is thermodynamically compatible with NiAl, a factor that has to be taken into account when producing artificial composite with NiAl alloy. The main purpose of this paper is to examine the ductility of Ni–30Al–5Mo–0.5Hf alloy at room temperature and high temperature. Also, in order to identify the operative mechanism responsible for the large elongation at high temperature, the emphasis was placed on the evaluation of constitutive equation involving the strain rate sensitivity index and the activation energy and also on the observation of the deformation microstructure.
2. Experimental A vacuum induction melted and drop cast ingot of nominal composition (at.%) Ni–30Al–5Mo–0.5Hf alloy was directionally solidified (DS) in the Al2 O3 –SiO2 ceramic mold under an Ar atmosphere by modified Bridgman method. The solidification rate was about 180 mm/h. The thermal gradient through the solid–liquid interface was estimated to be 7–9 K/mm. The analyzed chemical composition is wt.%: Al 15.6, Hf 1.39, Mo 9.49, Ni bal. Tensile specimens having a gauge section of 2.5 mm × 2 mm × 16 mm were cut from the DS alloy along the growth direction by electron discharge machine (EDM). Tensile tests were performed at initial strain rates ranging from 1.0 × 10−4 to 6.2 × 10−3 s−1 on a Shimadzu AG-250KNE testing machine. The tensile tests were performed at temperatures ranging from room temperature to 1373 K in air. Elongation was calibrated from the change between the two markers on the gauge surface after testing. The load data were collected at a frequency of 5 Hz using a computer. Three-point bending technique was used to determine the fracture toughness of the present alloy. Single U-notched specimens with dimensions of 5 mm × 10 mm × 50 mm were cut by EDM with the crack propagation direction perpendicular to the growth direction. A fatigue pre-crack was not introduced at the notch tip prior to testing. The specimens then were loaded in flexure over a 40 mm span on a MTS880 machine. The crosshead speed was 0.3 mm/min. Scanning electron microscopy equipped with energy dispersive X-ray spectroscopy (EDX) was used to characterize the as-grown microstructure, and to investigate the fracture behavior of the alloys. The specimens for trans-
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mission electron microscopy (TEM) observation were prepared by conventional ion-thinning procedure. The foils were made from the deformed tensile specimens by sectioning normal to the growth direction. The TEM observations were performed on a JEM2010 high resolution electron microscopy.
3. Result 3.1. Microstructure Fig. 1 shows the typical microstructure along both the longitudinal and transverse sections of the DS Ni–30Al–5Mo–0.5Hf alloy, showing that the alloy exhibits a dendritic structure. EDX analysis indicates that Ni and Al are enriched in the dendritic arm with dark contrast (Table 1), showing that the dendritic arm is the NiAl phase. SEM observation at high magnification, as shown in Fig. 1(b), showed that there also exist the Ni3 Al precipitates with strip or round shape and Mo particles. The interdendritic region with bright contrast consists of Ni3 Al and Mo phase, as listed in Table 1. NiAl precipitates are also found in the Ni3 Al phase, as shown in Fig. 1(c). The observation on the transverse direction also confirmed that the present alloy consists of three phases, that is, NiAl, Ni3 Al and Mo phases. Here, 0.5 at.% Hf is considered to be completely soluble in some of the constituent phases, because no Hf-concerned phases are observed. X-ray analysis showed that the growth direction of the NiAl phase is [0 0 1], and that of the Ni3 Al phase is mainly [0 0 1] and [0 1 1]. Selected area electron diffraction analysis revealed that the NiAl and Ni3 Al phases have the following orientation relationship: [1¯ 1 1] //[0 1 1]γ ,
(1 1 0) //(1 0 0)γ
This microstructure characteristic of the present alloy is consistent with that of Misra’s work [8]. Slow directional solidification of the alloy at 5 mm/h results in an aligned, continuous rod like ␥ microstructure in the NiAl matrix. The DS of the alloy at faster rates, 45–150 mm/h, produces cellular growth, resulting in a discontinuous morphology of the ␥ rods. At much faster rate 550 mm/h, the ␥ phase exist as a continuous interdendritic phase along the entire length of the crystal with two-phase ␥ +  dendrites. Thus, the addition of Hf and Mo has little effect on the microstructure of Ni–30Al alloy. Table 1 Chemical composition of the NiAl, Ni3 Al and Mo phases in the NiAl–Mo(Hf) alloy (at.%) Location
Ni
Al
Mo
Hf
Comments
Dendritic core Interdendritic
59.3 73.9 2.7
40.0 21.2 2.1
0.7 2.9 95.0
– 2.0 0.2
NiAl Ni3 Al Mo
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Fig. 1. The microstructure along the longitudinal (a)–(c) and transverse (d) sections of the DS NiAl–Mo(Hf) alloy. Table 2 Room temperature tensile properties of the DS NiAl–Mo(Hf) alloy Alloy
Strain rate (s−1 )
Yield stress (MPa)
Fracture stress (MPa)
Elongation (%)
NiAl–Mo(Hf) Ni–30Al [8]
1.0 × 10−4 1.4 × 10−4
NiAl–Cr(Mo, Hf) [13]
1.0 × 10−4
440 420a 500b –
509 600 840 255
1.3 4.4 9.0 –
a b
DS at rate of 5 mm/h. DS at rate of 45 mm/h.
3.2. Ductility at room temperature Table 2 lists the mechanical properties of the present alloy tested at room temperature. The room temperature tensile ductility is evaluated to be 1.3%, which is slightly lower than that of DS Ni–30Al alloy with a discontinuous ␥ phase [12], indicating that the addition of Mo and Hf has little effect on the room temperature ductility of the Ni–30Al alloy. It should be noted that this ductility is still a considerable improvement over the DS NiAl–Cr(Mo)–Hf alloy [13], which was found to have zero tensile ductility up to 1173 K. The
fracture toughness of the present alloy is calculated by: a S f (1) K Q = pQ 3/2 w BW where pQ is the fracture load, S the spanning length, B the specimen thickness, W the specimen depth and f(a/w) the dimensional parameter. The fracture toughness is evaluated √ to be 16.8 MPa m. Though the fatigue crack was not introduced in this study, the notch radius was smaller than 0.12 mm, thus KQ is slightly higher than KIC , due to the ductility of the alloy at room temperature. This value is
Fig. 2. Fracture surface observation of (a) the tensile specimen, arrows represent the void and (b) the fracture toughness specimen. Arrows show the crack propagation; (a) shows the dendritic core; (b) shows the interdendritic area.
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about three√to four times higher than that of monolithic NiAl (4–6 MPa m) [14], two times higher than that of whisker or particulate reinforced NiAl composites [14,15]. These results showed that the present alloy possesses not only good tensile ductility, but also the higher fracture toughness, which are promising for cold working at room temperature. Fig. 2 shows the fracture surface of the (a) tensile and (b) fracture toughness specimens. A lot of voids, together with cleavage fracture, are clearly observed in the fracture surface (Fig. 2(a)), indicating that fracture occurs by void coalescence and cleavage fracture of NiAl. As shown in Fig. 2(b), the fracture of dendritic cores is quasi-cleavage-like while the interdendritic ␥ phases exhibit great degree of plastic stretching. It is suggested that the ductility improvement of the present alloy with discontinuous ␥ phase may be attributed to the intrinsic ductility of the  phase afforded by dislocation generation from the interface [8]. However, since the ductility of the present alloy is much lower than that of DS Ni–30Al alloy with continuous ␥ phase, the extrinsic ductile phase toughening process, such as crack blunting and bridging, cannot be expected in the present alloy. 3.3. Ductility at high temperatures 3.3.1. Flow behavior The stress–strain curves were investigated as a function of initial strain rate particularly at high temperatures where the large elongations were observed. Fig. 3(a) shows the variation of the true stress–strain curves with the initial strain rate, which were deformed at 1323 K. At a high strain rate where low value of elongation is shown, the stress–strain curve exhibits a rapid work hardening, and after reaching the peak stress, the stress decreases monotonously. With decreasing the strain rate, the flow and peak stress decrease, showing that this large elongation is rate-dependent. Fig. 3(b) shows the temperature effect on the stress–strain curves of the alloy tested at an initial strain rate of 5.2 × 10−4 s−1 . The temperature effect is similar to that of strain rate, that is, an increase in temperature can obtain the similar stress–strain curve to that by a decrease in strain rate. 3.3.2. Constitutive equation Generally, it is well accepted that high temperature deformation of materials in relation to the large elongation is characterized by high value of strain rate sensitivity index, m, in the empirical equation: σ = K˙εm
(2)
where σ is the flow stress, ε˙ the strain rate and K the material constant. Peak stress is plotted against strain rate in Fig. 4 as a function of temperature. The linear relation holds between flow stress and strain rate, where the slope, m, the strain rate sensitivity index, is 0.27, independent of temperature. Also, the high temperature deformation of material can be repre-
Fig. 3. True stress–true strain curves of the DS NiAl–Mo(Hf) alloy deformed at (a) 1323 K and (b) 5.2 × 10−4 s−1 .
sented by power law equations: −Q n ε˙ = Aσ exp RT
(3)
where ε˙ is the strain rate, A the material constant, σ the flow stress, n the stress exponent (n = 1/m), R the gas constant, T the absolute temperature and Q the apparent activation energy. From analysis of the experimental data, the Q value is estimated to be 413 kJ/mol. The activation energy for deformation or creep of metallic materials generally corresponds to that for self-diffusion at temperatures above 0.5 Tm , where Tm is the absolute melting temperature. This value of the activation energy is much higher than that measured for lattice diffusion of Ni (220–300 kJ/mol) and creep (250–320 kJ/mol) in a NiAl [16], but it is very close to the activation energy of
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Fig. 4. Relationship between peak stress and strain rate at various temperatures and strain rates for the DS NiAl–Mo(Hf) alloy.
Fig. 5. Change of the elongation as a function of strain rate for the DS NiAl–Mo(Hf) alloy.
DS NiAl–Fe(Nb) alloy [17] and DS NiAl–25Cr alloy [18], which will be discussed later.
3.3.4. Deformed microstructure Specimens for microstructure observations were waterquenched immediately after deformed to a strain of ε = 0.1, 0.5 and fracture at 1323 K and initial strain rate of 5.2 × 10−4 s−1 . The observed microstructures are shown in Fig. 6. It is clearly shown that the original dendritic morphology becomes elongated along the tensile axis. Specifically, at a strain of 0.1 (Fig. 6(a)), most of Mo and block Ni3 Al phases coexist in this stage. And small amounts of elongated Mo phases are also found among the coexisting Mo and block Ni3 Al. Compared with Fig. 1(b), Mo also deforms during the tensile test and thus Mo may be very beneficial
3.3.3. Elongation Fig. 5 shows the tensile elongation obtained at various temperatures and strain rates. It is found that the elongation at each deformation temperature increases with decreasing the strain rate. This strain rate dependence of elongation in the present alloy is similar to that of monolithic NiAl alloy [19]. The maximum elongation, 160%, is obtained at 1323 K and initial strain rate of 5.2 × 10−4 s−1 under the deformation conditions studied in the present study.
Fig. 6. SEM observation on the specimen surfaces deformed to (a) ε = 0.1, (b) ε = 0.5 and (c) fracture at 1323 K and the strain rate of 5.2 × 10−4 s−1 . DRX grains are shown by arrows in (b).
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Fig. 7. TEM micrographs on the specimens deformed to (a) ε = 0.1 and (b) ε = 0.5 at 1323 K and the strain rate of 5.2 × 10−4 s−1 .
to the deformation. At a high strain of 0.5 (Fig. 6(b)), the NiAl and Ni3 Al phases are obviously elongated along the tensile axis, and only small amount of cavities are found on the interface, showing that these two phases can be deformed accommodately. In addition, small isolated grains existing in the Ni3 Al phase, as shown by arrows in Fig. 6(b) are also observed. By comparing the precipitation size and morphology in microstructure observation conducted before and after the tensile test, the isolated grains are considered to be introduced by the dynamic recrystallization (DRX). It should be noted that DRX mainly takes place in the Ni3 Al phase. The DRX is responsible for the strain softening observed in the stress–strain curves. For the fracture specimens (Fig. 6(c)), many cavities existing on the NiAl and Ni3 Al interface are observed, showing that the crack initiates from the interface and propagates along the interface, leading to fracture. Fig. 7 shows TEM micrographs of the specimens deformed to (a) ε = 0.1 and (b) ε = 0.5 at 1323 K and the strain rate of 5.2 × 10−4 s−1 . At the strain of 0.1, which corresponds to the strain just after showing the peak stress, high densities of dislocations are formed in the Ni3 Al phase, resulting in the work-hardening, as shown in Fig. 3. At the strain of 0.5, the microstructure in Fig. 7(b) consists of fine grain with a few microns and large grains with high dislocation density. The fine grains are thought to be DRX grains, while the large grains are the prior grain. Furthermore, Mo particles are also found near the DRX grains. Therefore, Mo may impede the dislocation movement and thus stabilize the microstructure during deformation at high temperatures.
4. Discussion It turned out that the DS NiAl–Mo(Hf) alloy possesses tensile elongation from room temperature to high temperature. That is, at room temperature, the alloy possesses 1.3%
tensile ductility and the ductility improvement is ascribed to the intrinsic mechanism, i.e. the generation of mobile a 1 0 0 dislocation, which has been interpreted in detail by Noebe et al. [7] and Misra et al. [8]. At high temperature, ductility of NiAl alloy is not a problem, since the alloy undergoes a brittle-to-ductile transition (BDT). For example, the NiAl–Cr(Mo)–Hf alloy exhibits elongation higher than 5% above 1173 K [13]. However, in this study, a large tensile elongation over 160% is observed in the present alloy at high temperature. Next, we will discuss the large elongation in terms of superplastic deformation. It is generally believed that superplastic deformation in the polycrystals occurs when the strain rate sensitivity index, m, is closer or more than 0.3. The value of m measured in this study, 0.27, is very close to 0.3 at the tested temperature, showing that the large elongation may be explained in term of m. However, previous studies showed that there are insufficient for explaining superplasticity mechanism of the DS alloy only by m. For example, Kim and Hanada [20] found that the m value in Fe3 Si single crystal depends on the orientation and thus it cannot fully explain the superplastic mechanism. Takasugi et al. [21] also found that the anomalously large elongation of single crystal (SC) NiAl alloy at intermediate temperatures depends on the crystal orientation and strain rate, the value of m is only 0.12. At higher temperature with large m value, the expected large elongation cannot be obtained. On the other hand, this result cannot also be explained by the grain-sliding based mechanism applied to many metals, alloys, ceramics and intermetallic alloys, exhibiting finegrained superplasticity, because, for the present DS alloy, the observed grain size is too large to cause grain boundary sliding. In fact, TEM observation showed that dislocation movements play an important role on this large elongation. At the early stage of deformation, gliding and multiplication of dislocation are dominant, which is clearly shown in Fig. 7(a). In
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Table 3 Deformation characteristics of some DS or SC alloys with large elongation Alloy
Method
Elongation (%)
Temperature (K)
m
Activation energy (kJ/mol)
Mechanism
NiAl–Mo(Hf) NiAl–Cr
DS DS [18] HEa [24] DS [17] HE [23] SC [21] SC [20]
160 160 480 260 470 170 150
1323 1323 1123 1323 1253 773 1173
0.27 0.27 0.50 0.29 0.52 0.12 0.30
413 397 463 390 224 120 –
DRX DRX GBSb DRX GBS + DRX Glide + climb [1 1 1] or [1 0 0] glide
NAi–Fe(Nb) NiAl Fe3 Si a b
Represents the hot extrusion. Represents the grain boundary sliding.
fact, Lautenschlager et al. [22] predicted that {1 0 0} 1 0 0, {1 1 0} 1 0 0 and {1 1 0} 1 0 0 dislocations are operative in NiAl deformed at high temperature by the limited slip trace investigation. The increase of dislocation density is accompanied by work hardening, which is clearly shown in the stress–strain curves in Fig. 3. However, with further deformation, dislocation climb occurs and dislocations rearrange themselves as subgrain boundary, and then forms subgrain, as shown in Figs. 6 and 7(b), resulting in the strain softening. Therefore, the comparable work hardening (by the glide motion of the dislocations) with the softening (by dynamic recovery and recrystallization) takes place at high temperature and offers steadily decrease flow stress, but not so significant softening, resulting in the large elongation. At last, we briefly discuss some DS NiAl alloys with the large elongation. Table 3 compares the superplastic characteristics such as strain rate sensitivity index, m, elongation, the activation energy and mechanism between the present alloy and other DS or SC alloys [17,18,20–24]. The three DS NiAl-based alloys with the large elongation have the similar  + ␥ microstructure with different precipitates, that is, the present alloy consists of , ␥ and Mo phases; the DS NiAl–Fe(Nb) alloy consists of , ␥ and ␥ [17]; while the DS NiAl–25Cr alloy consists of , ␥ and ␣-Cr [18], thus these DS NiAl alloys exhibit almost the same deformation characteristic, such as m, elongation, the activation energy and the deformation mechanism, as listed in Table 3. In addition, the higher activation energy of these DS alloys is ascribed to the solid solution element, Mo, Hf, Nb, because these elements are believed to affect the onset of dislocation climb by forming the solution atmosphere around the dislocation and effectively pinning them [13], resulting in the higher activation energy. On the other hand, the NiAl–Fe (Nb) and NiAl–Cr alloys, whether fabricated by DS or hot extrusion (HE), showed superplastic deformation. As listed in Table 3, the elongation of the alloy fabricated by HE shows much higher elongation than that of the alloy fabricated by DS. This implies that the mechanism of the superplastic deformation in the NiAl–Fe(Nb) and NiAl–Cr polycrystals is not only attained by the glide and climb of the dislocations. One possible explanation of this difference is that the additional grain boundary sliding between the grains occurs during the deformation process, as suggested previously [23,24]. DRX may occur more easily in the polycrystal than the DS alloy.
5. Conclusion The tensile properties were examined in the DS NiAl–Mo(Hf) alloy at room temperature and high temperature and the following results were obtained: 1. The DS NiAl–Mo(Hf) alloy was composed of NiAl, Ni3 Al and Mo phases. 2. The DS NiAl–Mo(Hf) alloy exhibited about 1.3% tensile elongation at room temperature and this ductility improvement was ascribed to the intrinsic mechanism of ductile phase toughening. 3. The DS NiAl–Mo(Hf) alloy showed large elongation up to 160% at 1323 K and initial strain rate of 5.2 × 10−4 s−1 . The comparable work hardening (by the glide motion of the dislocations) with the softening (by dynamic recovery and recrystallization) took place at high temperature and offered steadily to decrease the flow stress, resulting in the large elongation.
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High tensile elongation of a directionally solidified NiAl multiphase alloy at high temperatures C.Y. Cui ∗ , J.T. Guo, Y.H. Qi, H.Q. Ye Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, China Received 27 October 2004; received in revised form 6 January 2005; accepted 28 January 2005
Abstract The ductility of a directionally solidified Ni–30Al–5Mo–0.5Hf (at.%) alloy was examined in various experimental conditions of temperatures and strain rates. Tensile tests revealed that the present alloy exhibited tensile elongation from room temperature to high temperature, in particular, the large elongations, 140–160%, were observed at high temperatures, depending on the temperature and strain rate. The value of strain rate sensitivity index, m, of 0.27 and the apparent activation energy of 413 kJ/mol were measured. Based on the observed mechanical behavior and microstructures, the mechanism responsible for this large elongation operative at high temperatures was suggested. © 2005 Elsevier B.V. All rights reserved. Keywords: Directional solidification; NiAl alloy; Ductility; Dynamic recovery and recrystallization
1. Introduction NiAl is a potential high temperature structural material for applications in gas turbine engines and has been extensively studied during the last decade [1]. However, the low ductility and toughness at room temperature and poor elevated temperature strength limit its application as an engineering material. Recent advances in single crystal alloy development and eutectic composite technologies make them promising for improving the creep strength of NiAl [2–4]. Creep strength comparable to superalloy Rene’ 80 has been achieved in a NiAl–Cr(Mo) composite with Hf addition [4]. Creep strength can be further improved by controlling the microstructure [5]. Unfortunately, this promising NiAl–Cr(Mo) alloy is unsuitable for a structural material due to its poor ductility at room temperature. Thus, the issue of obtaining ductility and toughness at room temperature remains a major block in the development of a structural NiAl alloy. To overcome this ductility deficiency, two possibilities exist. One approach is alloying NiAl to lower the ordering energy of the material so that it is easier to initiate alternate slip ∗ Corresponding author. Present address: National Institute for Materials Science, Tsukuba, Ibaraki 305-0047, Japan. Tel.: +81 29 859 2544. E-mail address: [email protected] (C.Y. Cui).
0921-5093/$ – see front matter © 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2005.01.043
systems during deformation. Previous results [6] indicate that ductility improvement can be achieved by the addition of Fe, Ga, and Mo in the NiAl single crystal. Another approach is introducing a ductile second phase to increase the ductility and toughness of the composite. The mechanism of ductile phase toughening can be generally divided into the intrinsic mechanism, such as slip transfer, and the extrinsic mechanism, such as crack blunting, deflection and bridging [7,8]. The latter approach is the focus in this study. According to the Ni–Al phase diagram [7], the Ni3 Al phase seems to be a good candidate for ductile phase toughening of NiAl, because NiAl–Ni3 Al alloy can be readily formed in situ from the melt or by heat treatment. In single crystal form, Ni3 Al exhibits over 98% elongations at room temperature for [0 0 1] oriented crystals [9], while in polycrystal form, Ni3 Al exhibits about 10% tensile elongation at room temperature in air [10]. The potential of the Ni3 Al phase in enhancing the fracture resistance of NiAl has been demonstrated by several investigators using different processing methods [8,11,12]. Misra et al. [8] showed that room temperature tensile ductility up to 9% is achieved in directionally solidified (DS) NiAl–Ni3 Al two-phase alloy with the continuous ␥ , whereas the ductility of DS single-phase [0 0 1] NiAl is negligible. The enhancement in ductility is attributed to a combination of slip transfer from the ductile ␥ to the brittle  and extrinsic mechanism
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such as crack blunting, deflection and bridging. However, the creep strength of these alloys with discontinuous morphology of ␥ phase is marginally higher than that of single-phase Ni–40Al and significantly lower than that of typical superalloys at the 1000–1200 K temperature regime. Consequently, a high temperature reinforcing phase will have to be incorporated into this ductile phase toughened system. For example, instead of ␥ , Mo might be used as a reinforcing element. Mo not only has a high melting point (2890 K), but certain Mo alloys are very creep resistant. Further, Mo is thermodynamically compatible with NiAl, a factor that has to be taken into account when producing artificial composite with NiAl alloy. The main purpose of this paper is to examine the ductility of Ni–30Al–5Mo–0.5Hf alloy at room temperature and high temperature. Also, in order to identify the operative mechanism responsible for the large elongation at high temperature, the emphasis was placed on the evaluation of constitutive equation involving the strain rate sensitivity index and the activation energy and also on the observation of the deformation microstructure.
2. Experimental A vacuum induction melted and drop cast ingot of nominal composition (at.%) Ni–30Al–5Mo–0.5Hf alloy was directionally solidified (DS) in the Al2 O3 –SiO2 ceramic mold under an Ar atmosphere by modified Bridgman method. The solidification rate was about 180 mm/h. The thermal gradient through the solid–liquid interface was estimated to be 7–9 K/mm. The analyzed chemical composition is wt.%: Al 15.6, Hf 1.39, Mo 9.49, Ni bal. Tensile specimens having a gauge section of 2.5 mm × 2 mm × 16 mm were cut from the DS alloy along the growth direction by electron discharge machine (EDM). Tensile tests were performed at initial strain rates ranging from 1.0 × 10−4 to 6.2 × 10−3 s−1 on a Shimadzu AG-250KNE testing machine. The tensile tests were performed at temperatures ranging from room temperature to 1373 K in air. Elongation was calibrated from the change between the two markers on the gauge surface after testing. The load data were collected at a frequency of 5 Hz using a computer. Three-point bending technique was used to determine the fracture toughness of the present alloy. Single U-notched specimens with dimensions of 5 mm × 10 mm × 50 mm were cut by EDM with the crack propagation direction perpendicular to the growth direction. A fatigue pre-crack was not introduced at the notch tip prior to testing. The specimens then were loaded in flexure over a 40 mm span on a MTS880 machine. The crosshead speed was 0.3 mm/min. Scanning electron microscopy equipped with energy dispersive X-ray spectroscopy (EDX) was used to characterize the as-grown microstructure, and to investigate the fracture behavior of the alloys. The specimens for trans-
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mission electron microscopy (TEM) observation were prepared by conventional ion-thinning procedure. The foils were made from the deformed tensile specimens by sectioning normal to the growth direction. The TEM observations were performed on a JEM2010 high resolution electron microscopy.
3. Result 3.1. Microstructure Fig. 1 shows the typical microstructure along both the longitudinal and transverse sections of the DS Ni–30Al–5Mo–0.5Hf alloy, showing that the alloy exhibits a dendritic structure. EDX analysis indicates that Ni and Al are enriched in the dendritic arm with dark contrast (Table 1), showing that the dendritic arm is the NiAl phase. SEM observation at high magnification, as shown in Fig. 1(b), showed that there also exist the Ni3 Al precipitates with strip or round shape and Mo particles. The interdendritic region with bright contrast consists of Ni3 Al and Mo phase, as listed in Table 1. NiAl precipitates are also found in the Ni3 Al phase, as shown in Fig. 1(c). The observation on the transverse direction also confirmed that the present alloy consists of three phases, that is, NiAl, Ni3 Al and Mo phases. Here, 0.5 at.% Hf is considered to be completely soluble in some of the constituent phases, because no Hf-concerned phases are observed. X-ray analysis showed that the growth direction of the NiAl phase is [0 0 1], and that of the Ni3 Al phase is mainly [0 0 1] and [0 1 1]. Selected area electron diffraction analysis revealed that the NiAl and Ni3 Al phases have the following orientation relationship: [1¯ 1 1] //[0 1 1]γ ,
(1 1 0) //(1 0 0)γ
This microstructure characteristic of the present alloy is consistent with that of Misra’s work [8]. Slow directional solidification of the alloy at 5 mm/h results in an aligned, continuous rod like ␥ microstructure in the NiAl matrix. The DS of the alloy at faster rates, 45–150 mm/h, produces cellular growth, resulting in a discontinuous morphology of the ␥ rods. At much faster rate 550 mm/h, the ␥ phase exist as a continuous interdendritic phase along the entire length of the crystal with two-phase ␥ +  dendrites. Thus, the addition of Hf and Mo has little effect on the microstructure of Ni–30Al alloy. Table 1 Chemical composition of the NiAl, Ni3 Al and Mo phases in the NiAl–Mo(Hf) alloy (at.%) Location
Ni
Al
Mo
Hf
Comments
Dendritic core Interdendritic
59.3 73.9 2.7
40.0 21.2 2.1
0.7 2.9 95.0
– 2.0 0.2
NiAl Ni3 Al Mo
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Fig. 1. The microstructure along the longitudinal (a)–(c) and transverse (d) sections of the DS NiAl–Mo(Hf) alloy. Table 2 Room temperature tensile properties of the DS NiAl–Mo(Hf) alloy Alloy
Strain rate (s−1 )
Yield stress (MPa)
Fracture stress (MPa)
Elongation (%)
NiAl–Mo(Hf) Ni–30Al [8]
1.0 × 10−4 1.4 × 10−4
NiAl–Cr(Mo, Hf) [13]
1.0 × 10−4
440 420a 500b –
509 600 840 255
1.3 4.4 9.0 –
a b
DS at rate of 5 mm/h. DS at rate of 45 mm/h.
3.2. Ductility at room temperature Table 2 lists the mechanical properties of the present alloy tested at room temperature. The room temperature tensile ductility is evaluated to be 1.3%, which is slightly lower than that of DS Ni–30Al alloy with a discontinuous ␥ phase [12], indicating that the addition of Mo and Hf has little effect on the room temperature ductility of the Ni–30Al alloy. It should be noted that this ductility is still a considerable improvement over the DS NiAl–Cr(Mo)–Hf alloy [13], which was found to have zero tensile ductility up to 1173 K. The
fracture toughness of the present alloy is calculated by: a S f (1) K Q = pQ 3/2 w BW where pQ is the fracture load, S the spanning length, B the specimen thickness, W the specimen depth and f(a/w) the dimensional parameter. The fracture toughness is evaluated √ to be 16.8 MPa m. Though the fatigue crack was not introduced in this study, the notch radius was smaller than 0.12 mm, thus KQ is slightly higher than KIC , due to the ductility of the alloy at room temperature. This value is
Fig. 2. Fracture surface observation of (a) the tensile specimen, arrows represent the void and (b) the fracture toughness specimen. Arrows show the crack propagation; (a) shows the dendritic core; (b) shows the interdendritic area.
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about three√to four times higher than that of monolithic NiAl (4–6 MPa m) [14], two times higher than that of whisker or particulate reinforced NiAl composites [14,15]. These results showed that the present alloy possesses not only good tensile ductility, but also the higher fracture toughness, which are promising for cold working at room temperature. Fig. 2 shows the fracture surface of the (a) tensile and (b) fracture toughness specimens. A lot of voids, together with cleavage fracture, are clearly observed in the fracture surface (Fig. 2(a)), indicating that fracture occurs by void coalescence and cleavage fracture of NiAl. As shown in Fig. 2(b), the fracture of dendritic cores is quasi-cleavage-like while the interdendritic ␥ phases exhibit great degree of plastic stretching. It is suggested that the ductility improvement of the present alloy with discontinuous ␥ phase may be attributed to the intrinsic ductility of the  phase afforded by dislocation generation from the interface [8]. However, since the ductility of the present alloy is much lower than that of DS Ni–30Al alloy with continuous ␥ phase, the extrinsic ductile phase toughening process, such as crack blunting and bridging, cannot be expected in the present alloy. 3.3. Ductility at high temperatures 3.3.1. Flow behavior The stress–strain curves were investigated as a function of initial strain rate particularly at high temperatures where the large elongations were observed. Fig. 3(a) shows the variation of the true stress–strain curves with the initial strain rate, which were deformed at 1323 K. At a high strain rate where low value of elongation is shown, the stress–strain curve exhibits a rapid work hardening, and after reaching the peak stress, the stress decreases monotonously. With decreasing the strain rate, the flow and peak stress decrease, showing that this large elongation is rate-dependent. Fig. 3(b) shows the temperature effect on the stress–strain curves of the alloy tested at an initial strain rate of 5.2 × 10−4 s−1 . The temperature effect is similar to that of strain rate, that is, an increase in temperature can obtain the similar stress–strain curve to that by a decrease in strain rate. 3.3.2. Constitutive equation Generally, it is well accepted that high temperature deformation of materials in relation to the large elongation is characterized by high value of strain rate sensitivity index, m, in the empirical equation: σ = K˙εm
(2)
where σ is the flow stress, ε˙ the strain rate and K the material constant. Peak stress is plotted against strain rate in Fig. 4 as a function of temperature. The linear relation holds between flow stress and strain rate, where the slope, m, the strain rate sensitivity index, is 0.27, independent of temperature. Also, the high temperature deformation of material can be repre-
Fig. 3. True stress–true strain curves of the DS NiAl–Mo(Hf) alloy deformed at (a) 1323 K and (b) 5.2 × 10−4 s−1 .
sented by power law equations: −Q n ε˙ = Aσ exp RT
(3)
where ε˙ is the strain rate, A the material constant, σ the flow stress, n the stress exponent (n = 1/m), R the gas constant, T the absolute temperature and Q the apparent activation energy. From analysis of the experimental data, the Q value is estimated to be 413 kJ/mol. The activation energy for deformation or creep of metallic materials generally corresponds to that for self-diffusion at temperatures above 0.5 Tm , where Tm is the absolute melting temperature. This value of the activation energy is much higher than that measured for lattice diffusion of Ni (220–300 kJ/mol) and creep (250–320 kJ/mol) in a NiAl [16], but it is very close to the activation energy of
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Fig. 4. Relationship between peak stress and strain rate at various temperatures and strain rates for the DS NiAl–Mo(Hf) alloy.
Fig. 5. Change of the elongation as a function of strain rate for the DS NiAl–Mo(Hf) alloy.
DS NiAl–Fe(Nb) alloy [17] and DS NiAl–25Cr alloy [18], which will be discussed later.
3.3.4. Deformed microstructure Specimens for microstructure observations were waterquenched immediately after deformed to a strain of ε = 0.1, 0.5 and fracture at 1323 K and initial strain rate of 5.2 × 10−4 s−1 . The observed microstructures are shown in Fig. 6. It is clearly shown that the original dendritic morphology becomes elongated along the tensile axis. Specifically, at a strain of 0.1 (Fig. 6(a)), most of Mo and block Ni3 Al phases coexist in this stage. And small amounts of elongated Mo phases are also found among the coexisting Mo and block Ni3 Al. Compared with Fig. 1(b), Mo also deforms during the tensile test and thus Mo may be very beneficial
3.3.3. Elongation Fig. 5 shows the tensile elongation obtained at various temperatures and strain rates. It is found that the elongation at each deformation temperature increases with decreasing the strain rate. This strain rate dependence of elongation in the present alloy is similar to that of monolithic NiAl alloy [19]. The maximum elongation, 160%, is obtained at 1323 K and initial strain rate of 5.2 × 10−4 s−1 under the deformation conditions studied in the present study.
Fig. 6. SEM observation on the specimen surfaces deformed to (a) ε = 0.1, (b) ε = 0.5 and (c) fracture at 1323 K and the strain rate of 5.2 × 10−4 s−1 . DRX grains are shown by arrows in (b).
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Fig. 7. TEM micrographs on the specimens deformed to (a) ε = 0.1 and (b) ε = 0.5 at 1323 K and the strain rate of 5.2 × 10−4 s−1 .
to the deformation. At a high strain of 0.5 (Fig. 6(b)), the NiAl and Ni3 Al phases are obviously elongated along the tensile axis, and only small amount of cavities are found on the interface, showing that these two phases can be deformed accommodately. In addition, small isolated grains existing in the Ni3 Al phase, as shown by arrows in Fig. 6(b) are also observed. By comparing the precipitation size and morphology in microstructure observation conducted before and after the tensile test, the isolated grains are considered to be introduced by the dynamic recrystallization (DRX). It should be noted that DRX mainly takes place in the Ni3 Al phase. The DRX is responsible for the strain softening observed in the stress–strain curves. For the fracture specimens (Fig. 6(c)), many cavities existing on the NiAl and Ni3 Al interface are observed, showing that the crack initiates from the interface and propagates along the interface, leading to fracture. Fig. 7 shows TEM micrographs of the specimens deformed to (a) ε = 0.1 and (b) ε = 0.5 at 1323 K and the strain rate of 5.2 × 10−4 s−1 . At the strain of 0.1, which corresponds to the strain just after showing the peak stress, high densities of dislocations are formed in the Ni3 Al phase, resulting in the work-hardening, as shown in Fig. 3. At the strain of 0.5, the microstructure in Fig. 7(b) consists of fine grain with a few microns and large grains with high dislocation density. The fine grains are thought to be DRX grains, while the large grains are the prior grain. Furthermore, Mo particles are also found near the DRX grains. Therefore, Mo may impede the dislocation movement and thus stabilize the microstructure during deformation at high temperatures.
4. Discussion It turned out that the DS NiAl–Mo(Hf) alloy possesses tensile elongation from room temperature to high temperature. That is, at room temperature, the alloy possesses 1.3%
tensile ductility and the ductility improvement is ascribed to the intrinsic mechanism, i.e. the generation of mobile a 1 0 0 dislocation, which has been interpreted in detail by Noebe et al. [7] and Misra et al. [8]. At high temperature, ductility of NiAl alloy is not a problem, since the alloy undergoes a brittle-to-ductile transition (BDT). For example, the NiAl–Cr(Mo)–Hf alloy exhibits elongation higher than 5% above 1173 K [13]. However, in this study, a large tensile elongation over 160% is observed in the present alloy at high temperature. Next, we will discuss the large elongation in terms of superplastic deformation. It is generally believed that superplastic deformation in the polycrystals occurs when the strain rate sensitivity index, m, is closer or more than 0.3. The value of m measured in this study, 0.27, is very close to 0.3 at the tested temperature, showing that the large elongation may be explained in term of m. However, previous studies showed that there are insufficient for explaining superplasticity mechanism of the DS alloy only by m. For example, Kim and Hanada [20] found that the m value in Fe3 Si single crystal depends on the orientation and thus it cannot fully explain the superplastic mechanism. Takasugi et al. [21] also found that the anomalously large elongation of single crystal (SC) NiAl alloy at intermediate temperatures depends on the crystal orientation and strain rate, the value of m is only 0.12. At higher temperature with large m value, the expected large elongation cannot be obtained. On the other hand, this result cannot also be explained by the grain-sliding based mechanism applied to many metals, alloys, ceramics and intermetallic alloys, exhibiting finegrained superplasticity, because, for the present DS alloy, the observed grain size is too large to cause grain boundary sliding. In fact, TEM observation showed that dislocation movements play an important role on this large elongation. At the early stage of deformation, gliding and multiplication of dislocation are dominant, which is clearly shown in Fig. 7(a). In
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Table 3 Deformation characteristics of some DS or SC alloys with large elongation Alloy
Method
Elongation (%)
Temperature (K)
m
Activation energy (kJ/mol)
Mechanism
NiAl–Mo(Hf) NiAl–Cr
DS DS [18] HEa [24] DS [17] HE [23] SC [21] SC [20]
160 160 480 260 470 170 150
1323 1323 1123 1323 1253 773 1173
0.27 0.27 0.50 0.29 0.52 0.12 0.30
413 397 463 390 224 120 –
DRX DRX GBSb DRX GBS + DRX Glide + climb [1 1 1] or [1 0 0] glide
NAi–Fe(Nb) NiAl Fe3 Si a b
Represents the hot extrusion. Represents the grain boundary sliding.
fact, Lautenschlager et al. [22] predicted that {1 0 0} 1 0 0, {1 1 0} 1 0 0 and {1 1 0} 1 0 0 dislocations are operative in NiAl deformed at high temperature by the limited slip trace investigation. The increase of dislocation density is accompanied by work hardening, which is clearly shown in the stress–strain curves in Fig. 3. However, with further deformation, dislocation climb occurs and dislocations rearrange themselves as subgrain boundary, and then forms subgrain, as shown in Figs. 6 and 7(b), resulting in the strain softening. Therefore, the comparable work hardening (by the glide motion of the dislocations) with the softening (by dynamic recovery and recrystallization) takes place at high temperature and offers steadily decrease flow stress, but not so significant softening, resulting in the large elongation. At last, we briefly discuss some DS NiAl alloys with the large elongation. Table 3 compares the superplastic characteristics such as strain rate sensitivity index, m, elongation, the activation energy and mechanism between the present alloy and other DS or SC alloys [17,18,20–24]. The three DS NiAl-based alloys with the large elongation have the similar  + ␥ microstructure with different precipitates, that is, the present alloy consists of , ␥ and Mo phases; the DS NiAl–Fe(Nb) alloy consists of , ␥ and ␥ [17]; while the DS NiAl–25Cr alloy consists of , ␥ and ␣-Cr [18], thus these DS NiAl alloys exhibit almost the same deformation characteristic, such as m, elongation, the activation energy and the deformation mechanism, as listed in Table 3. In addition, the higher activation energy of these DS alloys is ascribed to the solid solution element, Mo, Hf, Nb, because these elements are believed to affect the onset of dislocation climb by forming the solution atmosphere around the dislocation and effectively pinning them [13], resulting in the higher activation energy. On the other hand, the NiAl–Fe (Nb) and NiAl–Cr alloys, whether fabricated by DS or hot extrusion (HE), showed superplastic deformation. As listed in Table 3, the elongation of the alloy fabricated by HE shows much higher elongation than that of the alloy fabricated by DS. This implies that the mechanism of the superplastic deformation in the NiAl–Fe(Nb) and NiAl–Cr polycrystals is not only attained by the glide and climb of the dislocations. One possible explanation of this difference is that the additional grain boundary sliding between the grains occurs during the deformation process, as suggested previously [23,24]. DRX may occur more easily in the polycrystal than the DS alloy.
5. Conclusion The tensile properties were examined in the DS NiAl–Mo(Hf) alloy at room temperature and high temperature and the following results were obtained: 1. The DS NiAl–Mo(Hf) alloy was composed of NiAl, Ni3 Al and Mo phases. 2. The DS NiAl–Mo(Hf) alloy exhibited about 1.3% tensile elongation at room temperature and this ductility improvement was ascribed to the intrinsic mechanism of ductile phase toughening. 3. The DS NiAl–Mo(Hf) alloy showed large elongation up to 160% at 1323 K and initial strain rate of 5.2 × 10−4 s−1 . The comparable work hardening (by the glide motion of the dislocations) with the softening (by dynamic recovery and recrystallization) took place at high temperature and offered steadily to decrease the flow stress, resulting in the large elongation.
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