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Microstructure investigation of NiAl–Cr(Mo) interface in a directionally solidified NiAl–Cr(Mo) eutectic alloyed with refractory metal
Materials Science and Engineering A 373 (2004) 279–285
Microstructure investigation of NiAl–Cr(Mo) interface in a directionally solidified NiAl–Cr(Mo) eutectic alloyed with refractory metal Y.X. Chen∗ , C.Y. Cui, J.T. Guo, D.X. Li Solids Atomic Imaging Division, Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, PR China Received 30 October 2003; received in revised form 23 January 2004
Abstract The microstructure of a directionally solidified NiAl–Cr(Mo) eutectic alloyed with refractory metal in as-processed and heat-treated states has been studied by means of scanning electron microscopy and high resolution electron microscopy (HREM). The microstructure of the NiAl–Cr(Mo) eutectic was characterized by lamellar Cr(Mo) phases embedded within NiAl matrix with common growth direction of 1 1 1. The interface between NiAl and lamellar Cr(Mo) did not have any transition layers. Misfit dislocations were observed at the NiAl–Cr(Mo) interface. In addition to lamellar Cr(Mo) phases, coherent Cr(Mo, Ni, Al) precipitates and NiAl precipitates were also observed in the NiAl matrix and lamellar Cr(Mo) phases, respectively. After hot isostatic pressing and heat treatment, the NiAl–Cr(Mo) interfaces became smooth and straight. Square array of misfit dislocations was directly observed at the (0 0 1) interface between NiAl and Cr(Mo, Ni, Al) precipitate. The configuration of misfit dislocation network showed a generally good agreement with prediction based on the geometric O-lattice model. © 2004 Elsevier B.V. All rights reserved. Keywords: NiAl; TEM; Microstructure; O-lattice model; Misfit dislocation
1. Introduction Nickel aluminide (NiAl) has a number of attractive properties including excellent oxidation resistance, high melting temperature, high thermal conductivity and relatively low density. As a result, alloys based on the intermetallic compound NiAl are being considered as candidate materials to replace nickel-based superalloys in some high temperature structural applications [1]. However, to make NiAl a viable structural material, it is necessary to overcome some of its inherent problems. These include low ductility and fracture toughness at ambient temperature and inadequate strength and creep resistance at elevated temperature. Accordingly, significant efforts have concentrated on enhancing the mechanical properties of NiAl through grain refinement, micro-and macro-alloying as well as incorporating second phase reinforcement [2–6].
∗ Corresponding author. Present address: Advanced Materials Research Institute, University of New Orleans, Science Building 2023, New Orleans, LA 70148, USA. Tel.: +1-504-280-5570; fax: +1-504-280-3185. E-mail address: [email protected] (Y.X. Chen).
0921-5093/$ – see front matter © 2004 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2004.01.026
The idea of reinforcing NiAl with in-situ refractory metals by directional solidification (DS) of eutectic alloys was actively pursued in the early 1970s [7–9]. Recently, there has been renewed interest in developing NiAl based DS eutectics for high temperature structural applications. A number of eutectic alloys consisting of NiAl and refractory metal phases, such as NiAl–Cr [10] and NiAl–Cr(Mo) [11], showed both improved toughness and creep strength compared to single-phase NiAl. DS of these eutectic alloys resulted in in-situ composites where one or more phases are aligned to the growth direction. For instance, NiAl–34Cr eutectic was characterized by a fibrous microstructure consisting of Cr rods embedded within NiAl matrix. However, the addition of Mo to the NiAl–Cr eutectic changed the microstructure from a fibrous reinforcing morphology to a lamellar morphology. In terms of fracture resistance, a lamellar morphology is more desirable than a fibrous structure since the crack front does not have continuous access to the brittle matrix. Additional advantage of producing NiAl-refractory metal eutectic alloy is that the phases are thermodynamically stable even up to the melting point. Until now, very few papers concerning atomic structure of NiAl–Cr(Mo) interfaces were found. In this study, we report
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interfacial atomic structure and misfit dislocation configuration at the NiAl–Cr(Mo) interfaces in as-processed and heat-treated states by means of field emission gun scanning electron microscopy (FEG SEM) and FEG high resolution electron microscopy (HREM) equipped with energy dispersive X-ray spectroscopy (EDXS).
2. Experimental procedure A vacuum induction melted and drop cast ingot of NiAl–28Cr–5.5Mo–0.5Hf (at.%) was directionally solidified under Ar atmosphere in the Al2 O3 –SiO2 ceramic mold by the standard Bridgeman method. For studying the effect of heat treatment on the interface structure between NiAl and Cr(Mo), some samples were hot isostatically processed (HIP) at 1300 ◦ C/100 MPa for 2 h and then heat-treated at 1050 ◦ C for 40 h. Transmission electron microscopy (TEM) samples were prepared by conventional procedure, including cutting discs perpendicular to the growth direction with a thickness of 0.2 mm, mechanical polishing them to 50 m thickness, dimpling to 20 m and finally ion milling. TEM and HREM observations were performed with a JEM 2010 high resolution electron microscope. SEM analysis was performed with a JSM 6301F FEG SEM. EDXS analysis was performed in a Hitachi HF2000 FEG TEM using a focused beam about 1nm in diameter. The quantification of the EDXS spectra was carried out by means of the standard Cliff–Lorimer quantitative procedure.
3. Results and discussion 3.1. SEM observation of as-processed alloy A typical microstructure of the as-processed NiAl– Cr(Mo)–Hf alloy at transverse section is shown in Fig. 1. The as-processed alloy is composed of three phases, which are lamellar Cr(Mo) phase, NiAl matrix and discontinuously distributed white phase. The white phase was identified
as Heusler phase Ni2 AlHf and was discussed elsewhere [12]. The directionally solidified NiAl–Cr(Mo)–Hf alloy consists of multiple grains. The Cr(Mo) lamellae in each grain are arranged in a radial pattern normal to the grain boundary with the center of the grain often composed of a few faceted rods. At the periphery of the grain, the lamellae are thicker than at the center indicating a lower growth rate normal to the grain boundary in the cusp region compared with the growth rate in the center of the grain. In addition to the above microstructure characteristics, several growth defects were formed within each grain. These include the following: (1) termination of a single Cr(Mo) lamella followed by the formation of a new Cr(Mo) lamella (A in Fig. 1) or a necking point formed at certain region of Cr(Mo) lamella (B and C in Fig. 1); (2) near the grain boundary, an extra Cr(Mo) lamella is embedded in the relatively regular lamellar structure (D in Fig. 1). These growth defects were identified to be the main mechanism for adjusting the phase spacing to accommodate mismatch strain between NiAl and Cr(Mo) for their different growth conditions (growth rate, temperature, etc.). However, these growth defects should be the major damage initiation sites for the material under mechanical loading [7,13]. According to early reports about NiAl–Cr and NiAl– Cr(Mo) eutectic alloys [8,9], the addition of Mo to NiAl–Cr eutectic system was to control the morphology of the Cr from a fibrous to a lamellar shape. With increasing Mo, the lattice mismatch between NiAl and Cr(Mo) at room temperature decreased until, at 0.6 at.% Mo, the mismatch became zero, and this was the composition at which the transition from non-faceted Cr(Mo) rods to faceted rods occurred. Increasing the Mo content to 5 at.% resulted in a lamellar structure with almost no faceted plane present. The lattice parameter of NiAl did not change with increasing Mo content. However, the lattice parameter of Cr(Mo) increased continuously with increasing Mo. When the content of Mo was 5.5 at.%, the lattice parameter of Cr(Mo) was measured to be 0.2942 nm [7], which was larger by 1.99% than that of Cr. The lattice mismatch parameter δ, of NiAl and Cr(Mo) can be calculated using the equation: δ=
2(α2 − α1 ) α1 + α 2
where α1 and α2 represent lattice parameters of NiAl and Cr(Mo), respectively. This gave a lattice mismatch parameter of 1.88%. This small lattice mismatch favors stability of the microstructure at high temperature [8,9,14]. 3.2. TEM and HREM observations of NiAl and Cr(Mo) phases in as-processed state Fig. 1. SEM image showing typical microstructure of the as-processed NiAl–Cr(Mo)–Hf alloy at transverse section. Growth defects of lamellar Cr(Mo) are denoted by A–D.
A selected area diffraction pattern from NiAl and Cr(Mo), as shown in Fig. 2(a), establishes a cube-on-cube orientation
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Fig. 2. (a) Selected area diffraction pattern taken from NiAl and lamellar Cr(Mo). (b) HREM image viewed along the common [1 1 1] zone axes of NiAl and lamellar Cr(Mo). Misfit dislocations are denoted by black arrows. (c) HREM image showing core structure of the misfit dislocation. The black arrows indicate the interface.
relationship (OR) between NiAl and Cr(Mo) phases, i.e. ¯ NiAl ||(1 0 1) ¯ Cr(Mo) . The [1 1 1]NiAl ||[1 1 1]Cr(Mo) and (1 0 1) cube-on-cube OR is consistent with previous reports [7–9]. Diffraction pattern obtained parallel to the growth direction of the crystal also indicates a 1 1 1 growth direction for both phases. Fig. 2(b) is a HREM image viewed along the common [1 1 1] zone axes of NiAl and Cr(Mo). The image revealed that no transition layer was formed in the interface. The interface is roughly parallel to the (1¯ 1 0) lattice plane of NiAl and Cr(Mo). Misfit dislocations, which array periodically about one of every 60 atomic spacing, are clearly visible with lattice distortion visible in the vicinity of the dislocation core (denoted by black arrows). Extra planes of the misfit dislocation are present at elastically soft NiAl side of the interface. The NiAl–Cr(Mo) interface shown in Fig. 2(b) is not smooth and straight. The regions where misfit dislocations located protrude into NiAl side of the interface, indicating a relatively high growth rate of these
cusps during solidification process of the alloy. The curved interface observed here implies that the NiAl–Cr(Mo) interfaces are not in their equilibrium state. There should be a long rang stress field along the interfaces. In addition, stress concentration should be generated in the vicinity of the misfit dislocations. It is expected that a curved interface will enhance segregation of impurities to the interface and drive crack tip through the interface. Fig. 2(c) is an enlarged HREM image showing core structure of the misfit dislocation. It can be seen from the image that the atomic matching between NiAl and Cr(Mo) has been achieved locally by a slight distortion of the NiAl lattice in the vicinity of the dislocation core. An attempt was made to determine Burgers vector of the misfit dislocation. A Burgers circuit was drawn across the interface encircling the dislocation core. From the Burgers circuit, it can be seen ¯ that the Burgers vector of the misfit dislocation is 1/3[1 1 2]. ¯ It should be noted that 1/3[1 1 2] is a Burgers vector of par-
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Fig. 3. (a) BF TEM image showing fine NiAl precipitates within Cr(Mo) lamella. Coherency is indicated by the strain contrast, g = [1 1 0]. (b) Corresponding HREM image viewed along the common [1 1 1] zone axes of Cr(Mo) and fine NiAl precipitate.
Fig. 4. (a) BF TEM image showing fine Cr(Mo, Ni, Al) precipitates within NiAl matrix. (b) Corresponding HREM image viewed along the common [1 1 1] zone axes of NiAl and Cr(Mo, Ni, Al) precipitate.
tial dislocation in B2 -ordered NiAl. The smallest possible complete dislocation Burgers vector in B2 -ordered NiAl is 0 0 1, since it has, according to the calculation of Ball and Smallman [15], the lowest energy. Therefore, the actual ¯ Burgers vector of the misfit dislocations should be [0 0 1]. The line direction of the misfit dislocation, i.e. 1 1 1, is not ¯ indicating that the misfit dislocaperpendicular to 0 0 1, ¯ tions at the (110) interface are of mixed type. The projection ¯ to 1 1 2 ¯ is 1/31 1 2, ¯ which accommodates the of 0 0 1 lattice mismatch between NiAl and Cr(Mo). Fig. 3(a) is a bright field (BF) TEM image, showing fine precipitates within the lamellar Cr(Mo) phases. The average size of these fine precipitates was measured to be about 25 nm. The EDXS results of these fine precipitates were as following (at.%): 41.29 Al, 56.35 Ni, 1.93 Cr, 0.43 Mo, indicating that these fine precipitates were NiAl but rich in Ni. The zero contrast lines at middle of each NiAl precipitate implies that the fine NiAl precipitates should be coherent with the matrix. Further, a corresponding HREM image viewed along the common [1 1 1] zone axes of the lamel-
lar Cr(Mo) and NiAl precipitate confirms that the fine NiAl precipitates are indeed coherent with the lamellar Cr(Mo), as shown in Fig. 3(b). The small lattice mismatch between the two phases was therefore fully accommodated by elastic strain in both sides of the interface. The dark contrast along the NiAl–Cr(Mo) interface implies that lattice strains are indeed present at the NiAl–Cr(Mo) interface regions. The Cr(Mo) phase containing fine NiAl precipitates indicates some solubility of Ni and Al in Cr(Mo) at high temperature. In addition to lamellar Cr(Mo) phases, fine precipitates with a average value of 30 nm in size were also observed in the NiAl matrix (Fig. 4(a)). EDXS results of these fine precipitates were as followings (at.%): 87.16 Cr, 4.89 Mo, 5.15 Al, 2.80 Ni, indicating that these fine precipitates are Cr with small amount of Mo, Ni and Al. Fig. 4(b) is a corresponding HREM image viewed along the common [1 1 1] zone axes of NiAl and fine Cr(Mo, Ni, Al) precipitate. The fine Cr(Mo, Ni, Al) precipitate is coherent with NiAl matrix. The modulation waves observed in NiAl matrix are due to the B2 -L10 martensitic phase transformation of NiAl [2,16,17].
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state are eliminated, implying that the NiAl–Cr(Mo) interfaces are at their equilibrium state after HIP and aging treatment. 3.4. Misfit dislocation network at the NiAl–Cr(Mo) interface
Fig. 5. BF TEM image showing coarsened NiAl precipitates entangling with mobile dislocations in lamellar Cr(Mo) phase after HIP and aging treatment. Misfit dislocations are visible at the NiAl–Cr(Mo) interface.
3.3. TEM and HREM observation of NiAl and Cr(Mo) phases after HIP and aging treatment The morphology of NiAl and Cr(Mo) phases were almost the same as that of as-processed alloy after HIP and aging treatment except a few notable microstructural changes. The coarsening of NiAl precipitates was observed in Cr(Mo) lamellae, as shown in Fig. 5. Semi-coherent interfaces between the NiAl precipitates and the Cr(Mo) lamellae were therefore formed. The Misfit dislocations are visible at there interfaces. Mobile dislocations which are rarely seen in lamellar Cr(Mo) in as-processed state are entangled with NiAl precipitates, indicating a small plastic deformation of lamellar Cr(Mo) during HIP. Another notable change, as shown in Fig. 6, is that the interface between NiAl and Cr(Mo) become smooth and straight. The cusps which array periodically at the NiAl–Cr(Mo) interfaces in as-processed
Fig. 6. HREM image viewed along the common [1 1 1] zone axes of NiAl and lamellar Cr(Mo) after HIP and aging treatment.
Interfacial coherency is of great interest since the presence of misfit dislocations might be indicative of the relative bonding strength at the interfaces. The geometrical model proposed by Bollmann [18] is a useful method to predict misfit dislocation structure at low-indexed interface between two phases. The O-theory postulates that the structure of a crystal interface is determined by two sets of parameters. The first set is related to the structure and relative orientation of the two crystals neighboring the interface, which determined the O-lattice that is present. The second set of parameters is given by the path of the boundary, i.e. the boundary plane. The dislocation structure of the interface is then given by the intersection of the boundary plane with the O-lattice cell walls, i.e. the Wigner–Seitz cell, with the translation vector attributed to the respective cell wall being the Burgers vector of the dislocation. It was calculated that the O-lattice points formed a bcc lattice with lattice parameter of 15.89 nm by following calculation steps introduced in ref. [12]. The misfit dislocation network at the (0 0 1) interface was obtained by the intersection of the boundary plane with the Wigner–Seitz cell walls of the bcc O-lattice. The resulting dislocation network at the (0 0 1) interface is shown in Fig. 7. The misfit dislocation network at the (0 0 1) interface is a square array of edge dislocations aligned along 1 0 0 direction with Burgers vector of the type 0 1 0. The spacing between misfit dislocations was calculated to be 15.89 nm. Fig. 8 is a BF TEM image taken with a small tilt away from the [0 0 1] zone axes of NiAl and Cr(Mo, Ni, Al) precipitate. From the image, a square array of the misfit dislocations at the (0 0 1) interface is clearly visible. The
Fig. 7. Schematic drawing of prediction of misfit dislocation network at the NiAl–Cr(Mo) (0 0 1) interfaces based on the geometric O-lattice model.
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Fig. 8. BF TEM image showing a square array of misfit dislocation at the NiAl–Cr(Mo, Ni, Al) (0 0 1) interface. The image was taken with a small tilt away from the [0 0 1] zone axes of NiAl and Cr(Mo, Ni, Al).
line direction of the misfit dislocation is 1 0 0. The spacing between misfit dislocations is measured to be 21 nm. This value is larger than the geometrical spacing, i.e. 15.89 nm. According to previous report [7], the lattice constant of Cr(Mo) increased linearly with increasing of Mo content. As described in Section 3.2, the content of Mo in Cr(Mo, Ni, Al) particles is lower than that in lamellar Cr(Mo). It is thus expected that the lattice constant of Cr(Mo, Ni, Al) particles should be shorter than that of lamellar Cr(Mo), resulting in a smaller lattice mismatch between NiAl and Cr(Mo, Ni, Al) precipitates than that between NiAl and lamellar Cr(Mo). Since the calculated dislocation spacing has an inverse relation to the lattice mismatch parameter, it is expected that the calculated misfit dislocation spacing at the NiAl–Cr(Mo, Ni, Al) interface should be very close to the measured value, i.e. 21 nm. The observed configuration, line direction and spacing of misfit dislocations at the (0 0 1) interfaces between NiAl and Cr(Mo, Ni, Al) are in generally good agreement with the theoretical prediction based on the geometric O-lattice model. The O-lattice model has some relevance with misfit dislocation structure formed in highly symmetric NiAl–Cr(Mo, Ni, Al) interface. It was reported that the misfit dislocation network should provide a strengthening for the alloy. The strengthening has an inverse relation to the spacing between misfit dislocations. In the present case, it is expected that the strengthening from the misfit dislocation network should be very small since the lattice mismatch is very small [12]. Additional strengthening should result from the formation of imperfect dislocations as mobile dislocations in NiAl crystal cross misfit dislocation network into lamellar Cr(Mo) phases, as shown in Fig. 9. As more and more dislocations cross from NiAl to Cr(Mo), each should leaves behind an imperfect dislocation. These dislocations soon make it difficult for new dislocations to cross [19].
Fig. 9. BF TEM image showing mobile dislocations in NiAl matrix cross NiAl–Cr(Mo) interface into Cr(Mo) phase after HIP and heating treatment.
The slow cooling of NiAl–Cr(Mo)–Hf alloy from processing temperature where the mismatch is smaller to room temperature where the mismatch is larger provide an increase in the number of misfit dislocations at the interface to accommodate the increasing mismatch. Upon cooling, the interface stress resulting from changing lattice mismatch caused the matrix to deform. Since the matrix mobile dislocations had the same Burgers vector as misfit dislocation, the mobile dislocations generated in the matrix moved into the interface to form misfit dislocation network as observed here, resulting in a reduction in strain energy of the interface.
4. Conclusions The microstructure of a directionally solidified NiAl– Cr(Mo) eutectic alloyed with refractory metal was characterized by lamellar Cr(Mo) phases embedded within NiAl matrix with growth defects near the grain boundary regions. The common growth direction of NiAl and lamellar Cr(Mo) was 1 1 1. The interfaces between NiAl and lamellar Cr(Mo) did not have any transition layers. Mixed type misfit dislo¯ were observed at the cation with Burgers vector of 0 0 1 ¯ NiAl–Cr(Mo) (1 1 0) interface. Coherent Cr(Mo, Ni, Al) precipitates and NiAl precipitates were observed in NiAl matrix and lamellar Cr(Mo) phases, respectively. After HIP and aging treatment, the interfaces between NiAl and lamellar Cr(Mo) became smooth and atomically flat. The coherent NiAl precipitates in the as-processed state lost coherency with the lamellar Cr(Mo) phases. A square array of misfit dislocation at the NiAl–Cr(Mo, Ni, Al) (0 0 1) interface was directly observed. The configuration of misfit dislocation network at the NiAl–Cr(Mo, Ni, Al) (0 0 1) interface showed a generally good agreement with the prediction based on the geometric O-lattice model.
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Acknowledgements This work was supported by National Natural Science Foundation of China on grant nos. 59831020, 59871055, 59895156, and 59895152.
References [1] R. Darolia, JOM 43 (3) (1991) 44. [2] J.Y. Dai, Z.P. Xing, Y.G. Wang, D.X. Li, J.T. Guo, L.L. He, H.Q. Ye, Mater. Lett. 20 (1994) 23. [3] R.R. Bowan, A.K. Misra, S.M. Arnold, Met. Trans. 26 (1995) 615. [4] H.P. Chiu, J.M. Yang, R.A. Amato, Mater. Sci. Eng. A 203 (1995) 81. [5] C.Y. Liu, S.M. Jeng, J.M. Yang, R.A. Amato, Mater. Sci. Eng. A 191 (1995) 49. [6] L.G. Yu, J.Y. Dai, Z.P. Xing, D.X. Li, J.T. Guo, H.Q. Ye, J. Mater. Res. 12 (1997) 1790. [7] H.E. Cline, J.L. Walter, Met. Trans. 1 (1970) 2907.
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[8] H.E. Cline, J.L. Walter, E.F. Koch, L.M. Osike, Acta Metall. Mater. 19 (1971) 405. [9] H.E. Cline, J.L. Walter, E. Lifshin, R.R. Russell, Met. Trans. 2 (1971) 189. [10] D.R. Johnson, X.F. Chen, B.F. Oliver, R.D. Noebe, J.D. Whittenberger, Intermetallics 3 (1995) 99. [11] X.F. Chen, D.R. Johnson, R.D. Noebe, B.F. Oliver, J. Mater. Res. 10 (1995) 1159. [12] Y.X. Chen, C.Y. Cui, Z.Q. Liu, L.L. He, J.T. Guo, D.X. Li, J. Mater. Res. 15 (2000) 1261. [13] J.M. Yang, S.M. Jeng, K. Bain, R.A. Amato, Acta Metall. Mater. 45 (1997) 295. [14] J.L. Walter, H.E. Cline, E.F. Koch, Trans. TMS-AIME 245 (1969) 2073. [15] A. Ball, R.E. Smallman, Acta Metall. 14 (1966) 1517. [16] A. Murthy, E. Goo, Acta Metall. Mater. 41 (1993) 3435. [17] S. Muto, S.D. Scheryvers, N. Merk, L.E. Tanner, Acta Metall. Mater. 41 (1993) 2377. [18] W. Bollmann, Crystal Defects and Crystalline Interfaces, Spring, New York, 1970. [19] J. Friedel, Dislocations, Addison Publishing Co., 1964, p. 185.
Microstructure investigation of NiAl–Cr(Mo) interface in a directionally solidified NiAl–Cr(Mo) eutectic alloyed with refractory metal Y.X. Chen∗ , C.Y. Cui, J.T. Guo, D.X. Li Solids Atomic Imaging Division, Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, PR China Received 30 October 2003; received in revised form 23 January 2004
Abstract The microstructure of a directionally solidified NiAl–Cr(Mo) eutectic alloyed with refractory metal in as-processed and heat-treated states has been studied by means of scanning electron microscopy and high resolution electron microscopy (HREM). The microstructure of the NiAl–Cr(Mo) eutectic was characterized by lamellar Cr(Mo) phases embedded within NiAl matrix with common growth direction of 1 1 1. The interface between NiAl and lamellar Cr(Mo) did not have any transition layers. Misfit dislocations were observed at the NiAl–Cr(Mo) interface. In addition to lamellar Cr(Mo) phases, coherent Cr(Mo, Ni, Al) precipitates and NiAl precipitates were also observed in the NiAl matrix and lamellar Cr(Mo) phases, respectively. After hot isostatic pressing and heat treatment, the NiAl–Cr(Mo) interfaces became smooth and straight. Square array of misfit dislocations was directly observed at the (0 0 1) interface between NiAl and Cr(Mo, Ni, Al) precipitate. The configuration of misfit dislocation network showed a generally good agreement with prediction based on the geometric O-lattice model. © 2004 Elsevier B.V. All rights reserved. Keywords: NiAl; TEM; Microstructure; O-lattice model; Misfit dislocation
1. Introduction Nickel aluminide (NiAl) has a number of attractive properties including excellent oxidation resistance, high melting temperature, high thermal conductivity and relatively low density. As a result, alloys based on the intermetallic compound NiAl are being considered as candidate materials to replace nickel-based superalloys in some high temperature structural applications [1]. However, to make NiAl a viable structural material, it is necessary to overcome some of its inherent problems. These include low ductility and fracture toughness at ambient temperature and inadequate strength and creep resistance at elevated temperature. Accordingly, significant efforts have concentrated on enhancing the mechanical properties of NiAl through grain refinement, micro-and macro-alloying as well as incorporating second phase reinforcement [2–6].
∗ Corresponding author. Present address: Advanced Materials Research Institute, University of New Orleans, Science Building 2023, New Orleans, LA 70148, USA. Tel.: +1-504-280-5570; fax: +1-504-280-3185. E-mail address: [email protected] (Y.X. Chen).
0921-5093/$ – see front matter © 2004 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2004.01.026
The idea of reinforcing NiAl with in-situ refractory metals by directional solidification (DS) of eutectic alloys was actively pursued in the early 1970s [7–9]. Recently, there has been renewed interest in developing NiAl based DS eutectics for high temperature structural applications. A number of eutectic alloys consisting of NiAl and refractory metal phases, such as NiAl–Cr [10] and NiAl–Cr(Mo) [11], showed both improved toughness and creep strength compared to single-phase NiAl. DS of these eutectic alloys resulted in in-situ composites where one or more phases are aligned to the growth direction. For instance, NiAl–34Cr eutectic was characterized by a fibrous microstructure consisting of Cr rods embedded within NiAl matrix. However, the addition of Mo to the NiAl–Cr eutectic changed the microstructure from a fibrous reinforcing morphology to a lamellar morphology. In terms of fracture resistance, a lamellar morphology is more desirable than a fibrous structure since the crack front does not have continuous access to the brittle matrix. Additional advantage of producing NiAl-refractory metal eutectic alloy is that the phases are thermodynamically stable even up to the melting point. Until now, very few papers concerning atomic structure of NiAl–Cr(Mo) interfaces were found. In this study, we report
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interfacial atomic structure and misfit dislocation configuration at the NiAl–Cr(Mo) interfaces in as-processed and heat-treated states by means of field emission gun scanning electron microscopy (FEG SEM) and FEG high resolution electron microscopy (HREM) equipped with energy dispersive X-ray spectroscopy (EDXS).
2. Experimental procedure A vacuum induction melted and drop cast ingot of NiAl–28Cr–5.5Mo–0.5Hf (at.%) was directionally solidified under Ar atmosphere in the Al2 O3 –SiO2 ceramic mold by the standard Bridgeman method. For studying the effect of heat treatment on the interface structure between NiAl and Cr(Mo), some samples were hot isostatically processed (HIP) at 1300 ◦ C/100 MPa for 2 h and then heat-treated at 1050 ◦ C for 40 h. Transmission electron microscopy (TEM) samples were prepared by conventional procedure, including cutting discs perpendicular to the growth direction with a thickness of 0.2 mm, mechanical polishing them to 50 m thickness, dimpling to 20 m and finally ion milling. TEM and HREM observations were performed with a JEM 2010 high resolution electron microscope. SEM analysis was performed with a JSM 6301F FEG SEM. EDXS analysis was performed in a Hitachi HF2000 FEG TEM using a focused beam about 1nm in diameter. The quantification of the EDXS spectra was carried out by means of the standard Cliff–Lorimer quantitative procedure.
3. Results and discussion 3.1. SEM observation of as-processed alloy A typical microstructure of the as-processed NiAl– Cr(Mo)–Hf alloy at transverse section is shown in Fig. 1. The as-processed alloy is composed of three phases, which are lamellar Cr(Mo) phase, NiAl matrix and discontinuously distributed white phase. The white phase was identified
as Heusler phase Ni2 AlHf and was discussed elsewhere [12]. The directionally solidified NiAl–Cr(Mo)–Hf alloy consists of multiple grains. The Cr(Mo) lamellae in each grain are arranged in a radial pattern normal to the grain boundary with the center of the grain often composed of a few faceted rods. At the periphery of the grain, the lamellae are thicker than at the center indicating a lower growth rate normal to the grain boundary in the cusp region compared with the growth rate in the center of the grain. In addition to the above microstructure characteristics, several growth defects were formed within each grain. These include the following: (1) termination of a single Cr(Mo) lamella followed by the formation of a new Cr(Mo) lamella (A in Fig. 1) or a necking point formed at certain region of Cr(Mo) lamella (B and C in Fig. 1); (2) near the grain boundary, an extra Cr(Mo) lamella is embedded in the relatively regular lamellar structure (D in Fig. 1). These growth defects were identified to be the main mechanism for adjusting the phase spacing to accommodate mismatch strain between NiAl and Cr(Mo) for their different growth conditions (growth rate, temperature, etc.). However, these growth defects should be the major damage initiation sites for the material under mechanical loading [7,13]. According to early reports about NiAl–Cr and NiAl– Cr(Mo) eutectic alloys [8,9], the addition of Mo to NiAl–Cr eutectic system was to control the morphology of the Cr from a fibrous to a lamellar shape. With increasing Mo, the lattice mismatch between NiAl and Cr(Mo) at room temperature decreased until, at 0.6 at.% Mo, the mismatch became zero, and this was the composition at which the transition from non-faceted Cr(Mo) rods to faceted rods occurred. Increasing the Mo content to 5 at.% resulted in a lamellar structure with almost no faceted plane present. The lattice parameter of NiAl did not change with increasing Mo content. However, the lattice parameter of Cr(Mo) increased continuously with increasing Mo. When the content of Mo was 5.5 at.%, the lattice parameter of Cr(Mo) was measured to be 0.2942 nm [7], which was larger by 1.99% than that of Cr. The lattice mismatch parameter δ, of NiAl and Cr(Mo) can be calculated using the equation: δ=
2(α2 − α1 ) α1 + α 2
where α1 and α2 represent lattice parameters of NiAl and Cr(Mo), respectively. This gave a lattice mismatch parameter of 1.88%. This small lattice mismatch favors stability of the microstructure at high temperature [8,9,14]. 3.2. TEM and HREM observations of NiAl and Cr(Mo) phases in as-processed state Fig. 1. SEM image showing typical microstructure of the as-processed NiAl–Cr(Mo)–Hf alloy at transverse section. Growth defects of lamellar Cr(Mo) are denoted by A–D.
A selected area diffraction pattern from NiAl and Cr(Mo), as shown in Fig. 2(a), establishes a cube-on-cube orientation
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Fig. 2. (a) Selected area diffraction pattern taken from NiAl and lamellar Cr(Mo). (b) HREM image viewed along the common [1 1 1] zone axes of NiAl and lamellar Cr(Mo). Misfit dislocations are denoted by black arrows. (c) HREM image showing core structure of the misfit dislocation. The black arrows indicate the interface.
relationship (OR) between NiAl and Cr(Mo) phases, i.e. ¯ NiAl ||(1 0 1) ¯ Cr(Mo) . The [1 1 1]NiAl ||[1 1 1]Cr(Mo) and (1 0 1) cube-on-cube OR is consistent with previous reports [7–9]. Diffraction pattern obtained parallel to the growth direction of the crystal also indicates a 1 1 1 growth direction for both phases. Fig. 2(b) is a HREM image viewed along the common [1 1 1] zone axes of NiAl and Cr(Mo). The image revealed that no transition layer was formed in the interface. The interface is roughly parallel to the (1¯ 1 0) lattice plane of NiAl and Cr(Mo). Misfit dislocations, which array periodically about one of every 60 atomic spacing, are clearly visible with lattice distortion visible in the vicinity of the dislocation core (denoted by black arrows). Extra planes of the misfit dislocation are present at elastically soft NiAl side of the interface. The NiAl–Cr(Mo) interface shown in Fig. 2(b) is not smooth and straight. The regions where misfit dislocations located protrude into NiAl side of the interface, indicating a relatively high growth rate of these
cusps during solidification process of the alloy. The curved interface observed here implies that the NiAl–Cr(Mo) interfaces are not in their equilibrium state. There should be a long rang stress field along the interfaces. In addition, stress concentration should be generated in the vicinity of the misfit dislocations. It is expected that a curved interface will enhance segregation of impurities to the interface and drive crack tip through the interface. Fig. 2(c) is an enlarged HREM image showing core structure of the misfit dislocation. It can be seen from the image that the atomic matching between NiAl and Cr(Mo) has been achieved locally by a slight distortion of the NiAl lattice in the vicinity of the dislocation core. An attempt was made to determine Burgers vector of the misfit dislocation. A Burgers circuit was drawn across the interface encircling the dislocation core. From the Burgers circuit, it can be seen ¯ that the Burgers vector of the misfit dislocation is 1/3[1 1 2]. ¯ It should be noted that 1/3[1 1 2] is a Burgers vector of par-
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Fig. 3. (a) BF TEM image showing fine NiAl precipitates within Cr(Mo) lamella. Coherency is indicated by the strain contrast, g = [1 1 0]. (b) Corresponding HREM image viewed along the common [1 1 1] zone axes of Cr(Mo) and fine NiAl precipitate.
Fig. 4. (a) BF TEM image showing fine Cr(Mo, Ni, Al) precipitates within NiAl matrix. (b) Corresponding HREM image viewed along the common [1 1 1] zone axes of NiAl and Cr(Mo, Ni, Al) precipitate.
tial dislocation in B2 -ordered NiAl. The smallest possible complete dislocation Burgers vector in B2 -ordered NiAl is 0 0 1, since it has, according to the calculation of Ball and Smallman [15], the lowest energy. Therefore, the actual ¯ Burgers vector of the misfit dislocations should be [0 0 1]. The line direction of the misfit dislocation, i.e. 1 1 1, is not ¯ indicating that the misfit dislocaperpendicular to 0 0 1, ¯ tions at the (110) interface are of mixed type. The projection ¯ to 1 1 2 ¯ is 1/31 1 2, ¯ which accommodates the of 0 0 1 lattice mismatch between NiAl and Cr(Mo). Fig. 3(a) is a bright field (BF) TEM image, showing fine precipitates within the lamellar Cr(Mo) phases. The average size of these fine precipitates was measured to be about 25 nm. The EDXS results of these fine precipitates were as following (at.%): 41.29 Al, 56.35 Ni, 1.93 Cr, 0.43 Mo, indicating that these fine precipitates were NiAl but rich in Ni. The zero contrast lines at middle of each NiAl precipitate implies that the fine NiAl precipitates should be coherent with the matrix. Further, a corresponding HREM image viewed along the common [1 1 1] zone axes of the lamel-
lar Cr(Mo) and NiAl precipitate confirms that the fine NiAl precipitates are indeed coherent with the lamellar Cr(Mo), as shown in Fig. 3(b). The small lattice mismatch between the two phases was therefore fully accommodated by elastic strain in both sides of the interface. The dark contrast along the NiAl–Cr(Mo) interface implies that lattice strains are indeed present at the NiAl–Cr(Mo) interface regions. The Cr(Mo) phase containing fine NiAl precipitates indicates some solubility of Ni and Al in Cr(Mo) at high temperature. In addition to lamellar Cr(Mo) phases, fine precipitates with a average value of 30 nm in size were also observed in the NiAl matrix (Fig. 4(a)). EDXS results of these fine precipitates were as followings (at.%): 87.16 Cr, 4.89 Mo, 5.15 Al, 2.80 Ni, indicating that these fine precipitates are Cr with small amount of Mo, Ni and Al. Fig. 4(b) is a corresponding HREM image viewed along the common [1 1 1] zone axes of NiAl and fine Cr(Mo, Ni, Al) precipitate. The fine Cr(Mo, Ni, Al) precipitate is coherent with NiAl matrix. The modulation waves observed in NiAl matrix are due to the B2 -L10 martensitic phase transformation of NiAl [2,16,17].
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state are eliminated, implying that the NiAl–Cr(Mo) interfaces are at their equilibrium state after HIP and aging treatment. 3.4. Misfit dislocation network at the NiAl–Cr(Mo) interface
Fig. 5. BF TEM image showing coarsened NiAl precipitates entangling with mobile dislocations in lamellar Cr(Mo) phase after HIP and aging treatment. Misfit dislocations are visible at the NiAl–Cr(Mo) interface.
3.3. TEM and HREM observation of NiAl and Cr(Mo) phases after HIP and aging treatment The morphology of NiAl and Cr(Mo) phases were almost the same as that of as-processed alloy after HIP and aging treatment except a few notable microstructural changes. The coarsening of NiAl precipitates was observed in Cr(Mo) lamellae, as shown in Fig. 5. Semi-coherent interfaces between the NiAl precipitates and the Cr(Mo) lamellae were therefore formed. The Misfit dislocations are visible at there interfaces. Mobile dislocations which are rarely seen in lamellar Cr(Mo) in as-processed state are entangled with NiAl precipitates, indicating a small plastic deformation of lamellar Cr(Mo) during HIP. Another notable change, as shown in Fig. 6, is that the interface between NiAl and Cr(Mo) become smooth and straight. The cusps which array periodically at the NiAl–Cr(Mo) interfaces in as-processed
Fig. 6. HREM image viewed along the common [1 1 1] zone axes of NiAl and lamellar Cr(Mo) after HIP and aging treatment.
Interfacial coherency is of great interest since the presence of misfit dislocations might be indicative of the relative bonding strength at the interfaces. The geometrical model proposed by Bollmann [18] is a useful method to predict misfit dislocation structure at low-indexed interface between two phases. The O-theory postulates that the structure of a crystal interface is determined by two sets of parameters. The first set is related to the structure and relative orientation of the two crystals neighboring the interface, which determined the O-lattice that is present. The second set of parameters is given by the path of the boundary, i.e. the boundary plane. The dislocation structure of the interface is then given by the intersection of the boundary plane with the O-lattice cell walls, i.e. the Wigner–Seitz cell, with the translation vector attributed to the respective cell wall being the Burgers vector of the dislocation. It was calculated that the O-lattice points formed a bcc lattice with lattice parameter of 15.89 nm by following calculation steps introduced in ref. [12]. The misfit dislocation network at the (0 0 1) interface was obtained by the intersection of the boundary plane with the Wigner–Seitz cell walls of the bcc O-lattice. The resulting dislocation network at the (0 0 1) interface is shown in Fig. 7. The misfit dislocation network at the (0 0 1) interface is a square array of edge dislocations aligned along 1 0 0 direction with Burgers vector of the type 0 1 0. The spacing between misfit dislocations was calculated to be 15.89 nm. Fig. 8 is a BF TEM image taken with a small tilt away from the [0 0 1] zone axes of NiAl and Cr(Mo, Ni, Al) precipitate. From the image, a square array of the misfit dislocations at the (0 0 1) interface is clearly visible. The
Fig. 7. Schematic drawing of prediction of misfit dislocation network at the NiAl–Cr(Mo) (0 0 1) interfaces based on the geometric O-lattice model.
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Fig. 8. BF TEM image showing a square array of misfit dislocation at the NiAl–Cr(Mo, Ni, Al) (0 0 1) interface. The image was taken with a small tilt away from the [0 0 1] zone axes of NiAl and Cr(Mo, Ni, Al).
line direction of the misfit dislocation is 1 0 0. The spacing between misfit dislocations is measured to be 21 nm. This value is larger than the geometrical spacing, i.e. 15.89 nm. According to previous report [7], the lattice constant of Cr(Mo) increased linearly with increasing of Mo content. As described in Section 3.2, the content of Mo in Cr(Mo, Ni, Al) particles is lower than that in lamellar Cr(Mo). It is thus expected that the lattice constant of Cr(Mo, Ni, Al) particles should be shorter than that of lamellar Cr(Mo), resulting in a smaller lattice mismatch between NiAl and Cr(Mo, Ni, Al) precipitates than that between NiAl and lamellar Cr(Mo). Since the calculated dislocation spacing has an inverse relation to the lattice mismatch parameter, it is expected that the calculated misfit dislocation spacing at the NiAl–Cr(Mo, Ni, Al) interface should be very close to the measured value, i.e. 21 nm. The observed configuration, line direction and spacing of misfit dislocations at the (0 0 1) interfaces between NiAl and Cr(Mo, Ni, Al) are in generally good agreement with the theoretical prediction based on the geometric O-lattice model. The O-lattice model has some relevance with misfit dislocation structure formed in highly symmetric NiAl–Cr(Mo, Ni, Al) interface. It was reported that the misfit dislocation network should provide a strengthening for the alloy. The strengthening has an inverse relation to the spacing between misfit dislocations. In the present case, it is expected that the strengthening from the misfit dislocation network should be very small since the lattice mismatch is very small [12]. Additional strengthening should result from the formation of imperfect dislocations as mobile dislocations in NiAl crystal cross misfit dislocation network into lamellar Cr(Mo) phases, as shown in Fig. 9. As more and more dislocations cross from NiAl to Cr(Mo), each should leaves behind an imperfect dislocation. These dislocations soon make it difficult for new dislocations to cross [19].
Fig. 9. BF TEM image showing mobile dislocations in NiAl matrix cross NiAl–Cr(Mo) interface into Cr(Mo) phase after HIP and heating treatment.
The slow cooling of NiAl–Cr(Mo)–Hf alloy from processing temperature where the mismatch is smaller to room temperature where the mismatch is larger provide an increase in the number of misfit dislocations at the interface to accommodate the increasing mismatch. Upon cooling, the interface stress resulting from changing lattice mismatch caused the matrix to deform. Since the matrix mobile dislocations had the same Burgers vector as misfit dislocation, the mobile dislocations generated in the matrix moved into the interface to form misfit dislocation network as observed here, resulting in a reduction in strain energy of the interface.
4. Conclusions The microstructure of a directionally solidified NiAl– Cr(Mo) eutectic alloyed with refractory metal was characterized by lamellar Cr(Mo) phases embedded within NiAl matrix with growth defects near the grain boundary regions. The common growth direction of NiAl and lamellar Cr(Mo) was 1 1 1. The interfaces between NiAl and lamellar Cr(Mo) did not have any transition layers. Mixed type misfit dislo¯ were observed at the cation with Burgers vector of 0 0 1 ¯ NiAl–Cr(Mo) (1 1 0) interface. Coherent Cr(Mo, Ni, Al) precipitates and NiAl precipitates were observed in NiAl matrix and lamellar Cr(Mo) phases, respectively. After HIP and aging treatment, the interfaces between NiAl and lamellar Cr(Mo) became smooth and atomically flat. The coherent NiAl precipitates in the as-processed state lost coherency with the lamellar Cr(Mo) phases. A square array of misfit dislocation at the NiAl–Cr(Mo, Ni, Al) (0 0 1) interface was directly observed. The configuration of misfit dislocation network at the NiAl–Cr(Mo, Ni, Al) (0 0 1) interface showed a generally good agreement with the prediction based on the geometric O-lattice model.
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Acknowledgements This work was supported by National Natural Science Foundation of China on grant nos. 59831020, 59871055, 59895156, and 59895152.
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