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The formation of aluminides in intermetallic nickel aluminide-based nanocomposites
Journal of Alloys and Compounds 392 (2005) 214–219
The formation of aluminides in intermetallic nickel aluminide-based nanocomposites R. Ismail, I.I. Yaacob∗ Department of Mechanical Engineering, Faculty of Engineering, Materials Engineering Programme, University of Malaya, 50603 Kuala Lumpur, Malaysia Received 7 July 2004; received in revised form 17 September 2004; accepted 20 September 2004
Abstract Metal matrix composites based on the intermetallic alloy Ni3 Al and alumina were fabricated using powder metallurgy technique. Nanosize alumina as dispersed phase was mixed with Ni and Al powder in a planetary ball mill for 4 h at 350 rpm to achieve a small mechanical alloying effect. The mixture was then compacted using a hydraulic press at 400 MPa for 15 min. Sintering was performed under an inert condition in a tube furnace at 660 ◦ C with 5 h soaking time. The formation of aluminides, Ni–Al and Ni3 Al in intermetallic nickel aluminide-based composites were observed by measuring the saturation magnetization (Ms ) values of the sintered materials. The sample’s Ms value was very low at 0.9222 emu/g indicating the presence of a small amount of elemental Ni. The occurrence of reaction synthesis during sintering was detected by the presence of a ‘large’ exothermic peak of the differential thermal analysis curve, which occurred at below 600 ◦ C. The heat released during reaction synthesis was probed by differential scanning calorimetry. The heat of reaction was found to be 4.53 × 104 J/mol. Addition of 5% alumina increased the hardness of the composite by about a factor of two with respect to that of Ni3 Al intermetallic. Xray diffraction (XRD) measurements showed peaks corresponding to Ni–Al and Ni3 Al. Optical and SEM investigations revealed growth of lath-like microstructure in the fabricated composite due to a high lattice mismatch between ␥ and ␥ phase. © 2004 Elsevier B.V. All rights reserved. Keywords: Intermetallics; Powder metallurgy; Thermal analysis; Magnetic measurements; Nanocomposites
1. Introduction The intermetallic compound Ni3 Al is well known for its oxidation resistance and high temperature-strength. The latter property is derived from precipitation hardening of the coherent, ordered Ni3 Al ␥ phase within the disordered, solid solution ␥ phase. The compatibility of the ␥ precipitates FCC crystal structure with a lattice constant of approximately 0.1% mismatch with the ␥ phase (Ni matrix phase) results in homogenous nucleation of the precipitates. The precipitates have low surface energy and extraordinary long-time stability. Coherency at elevated temperature between ␥ and ␥ is maintained by tetragonal distortion along the z-axis of the precipitates [1].
∗
Corresponding author. Tel.: +60 379674489/5681; fax: +60 379675317. E-mail address: [email protected] (I.I. Yaacob).
0925-8388/$ – see front matter © 2004 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2004.09.032
This unique ␥ or aluminide intermetallic phase contributes to remarkable strengthening of the intermetallic by forcing dislocations to either bypass the larger aluminide precipitates by Orowan bowing or dislocations cutting of the smaller precipitates in the ␥ –␥ alloy [2]. Moreover, the strength of ␥ increases as temperature increases which is known as anomalous yield behaviour [3,4]. The fact that the intermetallic can be intrinsically strong and can retain this strength at elevated temperatures, makes it especially attractive for high temperature uses. Its low density (7.5 g/cm3 ) is an added advantage to the specific strength where major stresses are determined by the mass of an object. Low selfdiffusion in an ordered intermetallic increases the microstructural stability and decreases creep rates [5]. The intrinsic grain boundary brittleness in Ni3 Al is associated with ordering energy, electronegativity difference, valency difference and atomic size difference between Ni atoms and Al atoms. However, microalloying with boron sharply increases its ductil-
R. Ismail, I.I. Yaacob / Journal of Alloys and Compounds 392 (2005) 214–219
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All the materials were used as received. About 100-mesh nickel powder (99.9% purity) was purchased from Strem Chemicals. Aluminum powder (99.7% purity) and nanosized alumina (average size of 39 nm) were obtained from Aldrich chemicals. Special etchant was prepared using 3 g of molybdic acid powder from Ajax Chemicals, 100 ml of distilled H2 O, 100 ml of hydochloric acid (assay 37%) from Reidel de Haen Chemicals and 100 ml of nitric acid (assay 65%) from Merck. Elemental nickel and aluminum powders at 3:1 mole ratio were pre-mixed with different weight fractions of nanosized alumina (as shown Table 1) using a planetary ball mill (PM400, MA Type, Retsch). The mixtures were milled for 4 h at 175 rpm in a zirconia jar with zirconia balls. The milled samples were then compacted using a hydraulic hand press (Gatesby Specac) in a 20 mm die at 400 MPa for 15 min. Sintering of the green compacts was performed under argon gas atmosphere using a tube furnace (Carbolite, Type CTF 17/75/300, maximum T: 1700 ◦ C). The tube furnace containing the samples was initially purged with argon gas (1 l/min)
and then vacuumed. This process was repeated several times. The samples were heated under flowing argon at a heating rate of 20 ◦ C/min until 660 ◦ C, at which they were isothermally heated for a duration of 5 h [15]. The reaction synthesis was monitored by observing the differential thermal analysis (DTA) curves when the samples were sintered under argon gas atmosphere in a small furnace (TGA/SDTA 851e , Mettler Toledo). A green compact sample weighing about 40–70 mg was placed directly on the DTA sample holder. This was done to ensure an optimum sensitivity during sintering. However, a silica disc was placed underneath the sample to avoid any possible contaminations and degradation of the thermocouples during heating. Changes in the sample, which led to absorption or evolution of heat, were detected relative to the inert reference. The amount of heat released during the reactive sintering process can be measured by heating the samples from 30 to 670 ◦ C under inert argon gas atmosphere using a differential scanning calorimeter (Mettler Toledo DSC820 with FRS5 sensor (Heat flux type)). A small piece of the green compact weighing about 10 mg was placed in a 40 l platinum crucible which was then closed with a Pt lid with a small hole. The sample crucible was placed side by side with an empty Pt reference crucible. The saturation magnetization values for the sintered samples were measured using an alternating gradient magnetometer (AGM) (Micromag 2900) with maximum applied fields of ±10 kOe. A small amount of the sample was mounted on an extension rod attached to a piezoelectric element and placed in a region between two electomagnets. The resulting deflection of the rod was measured by a piezoelectric sensing element mounted on the probe arm. The magnetic response of the sample was plotted in terms of the sample’s magnetization against the magnitude of the applied field. The phases present in the samples before and after sintering were examined by an X-ray diffractometer (XRD) using Cu K␣ radiation (Philips X-Pert MPD PW3040). The diffraction pattern was obtained by scanning the sample from 15◦ to 90◦ 2θ angle at a step size of 0.02◦ and a count time of 1.5 s at each step. The hardness of the sintered samples was determined by Vickers indentation technique using 10 kgf load. For morphological characterizations, the sintered sample was mounted on a resin and polished until mirror finish. The sample’s surface was etched with a specially prepared etchant described earlier. Excess etchant was washed away with running water. The sample was then dried with a blower. Surface morphologies were observed under an optical microscope and a scanning electron microscope (Philips XL40 SEM).
Table 1 Compositions of intermetallic composites (IMCs)
3. Results and discussion
ity and effectively suppresses intergranular fracture [6]. The ductilizing effect of boron can be explained by two possible mechanisms: (i) boron-enhanced grain boundary cohesive strength [7] and (ii) boron-facilitated slip transfer across the grain boundary (or slip nucleation at boundaries) [8]. To improve monolithic Ni3 Al and NiAl properties such as low ambient tensile ductility and insufficient hightemperature strength and creep resistance limit, ceramic particulates are added as reinforcement fillers. The addition of lower density ceramic materials (i.e. oxides, carbides and nitrides) also decreases the composite’s density and increases its specific properties. However, most of these studies are still in the feasibility stage aiming primarily at determining basic mechanical properties and chemical stability [9–11]. In this paper, nanosized alumina powder was used as the reinforcement. Powder metallurgy technique was used for fabricating the composite by which the alumina, nickel and aluminum powders were mixed, compacted and then sintered. Exothermic reaction between the Ni and Al metal powders during the sintering process enabled the formation of intermetallics at temperatures as low as 500 ◦ C [12–14]. The extremely small size of aluminides made them difficult to be observed under an optical or electron microscope. By measuring the saturation magnetization (Ms ) values of the sintered sample, we were able to quantify the amount of aluminides formed in each specimen.
2. Experimental
Samples
Code
Ni3 Al matrix Ni3 Al + 5 wt.% Al2 O3 Ni3 Al + 15 wt.% Al2 O3
IM IMC-5 IMC-15
The saturation magnetization (Ms ) values listed in Table 2 indicate the presence of unreacted elemental nickel in the samples. The formation of aluminide or ␥ precipitate during
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Table 2 Saturation magnetization (Ms ) values for intermetallic, IMCs and pure Ni powder Sintered samples
Saturation magnetization (Ms ) value (emu/g)
IM IMC-5 IMC-15 Ni powder
0.922 3.363 15.38 48.20
sintering reduces the amount of elemental nickel in the samples hence lowering the Ms value. The Ms value of the sintered matrix (IM) is 0.922 emu/g and is comparable with the Ms value of commercial Ni3 Al (Alfa Aesar, lot # F18l52, powder form), which is 0.3927 emu/g. However, the Ms values for the IMC sample are higher. The higher Ms for the IMCs is due to the presence of dispersed alumina which causes incomplete conversion of the metal powder mixture to ␥ precipitate or aluminide. The Ms value for IMC-15 is about 31% of the Ms of pure Ni powder indicating that IMC-15 (containing about 73.95 wt.% Ni) contains about 43 wt.% unreacted elemental nickel. The incomplete conversion of elemental Ni and Al into aluminide in the sintered materials can be further observed form the DTA curves of the nickel aluminide matrix and composites. Differential thermal analysis (DTA) curves show occurrence of reaction synthesis (Fig. 1). Large exothermic peaks as the result from a sudden release of energy emerge in every sintered sample. Reaction synthesis is expected to occur due to the formation of a transient liquid phase from the lower melting point minor phase (aluminum) encapsulating the nickel powder by capillary force [12,16]. Evaluation of the peak’s position shows that reaction synthesis for the intermetallic nickel aluminide composites containing 5 and 15 wt.% alumina as dispersed phase were observed to occur at lower temperatures compared to the ‘pure’ intermetallic nickel aluminide. The disruption of the ease of flow of the liquid aluminum by the dispersed alumina is expected to reduce the reaction synthesis temperature. The presence of alumina
Fig. 2. DSC curve for Ni3 Al shows large exothermic peak indicating heat evolution as the results of reaction synthesis.
as non-reacting dispersed phase can act as heat sinks and thermal barriers limiting temperature rise and also impede the flow of transient aluminum liquid [17]. Measurement of the heat released during reaction synthesis of Ni3 Al can be made by extrapolating the width of the exothermic peak in a DSC curve as shown in Fig. 2. However, two exothermic peaks can be seen. The larger peak occurs at 597 ◦ C with a heat release of 223.23 J/g. The smaller one is at 640.4 ◦ C with a heat release of 4.27 J/g. The heat of reaction at 597 ◦ C for 3Ni + Al powder compact weighing 0.0098 g is calculated to be −4.53 × 104 J/mol. Identification of the desired ␥ precipitate during the reaction synthesis can be performed by observation of the XRD patterns shown in Fig. 3. Comparisons of d-spacings of the XRD peaks with the ICDD files confirmed the presence of both aluminides in the form of ordered ␥ (Ni3 Al) and  (Ni–Al) phases. The presence of dispersed alumina in the IMCs leads to a lower aluminide formation in the sintered samples which is consistent with the higher Ms values. The presence of a prominent Ni (2 0 0) peak in IMC15 and the emergence of the same peak in IMC5 can be observed.
Fig. 1. DTA curves for intermetallic matrix and IMCs. The specimens were heated from 30 to 660 ◦ C at 20 ◦ C/min and were then isothermally heated at 660 ◦ C for 5 h.
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Fig. 3. X-ray diffraction patterns of pure Ni, intermetallic Ni3 Al (IM) and intermetallic composites (IMCs). The peaks are labeled accordingly.
The hardness values of the sintered samples are shown in Table 3. An average hardness of 90.9 VHN is obtained for the Ni3 Al matrix. This is at the lower end of the cited literature value of 100–255 VHN [18]. When 5 wt.% of nanosized alumina is dispersed in the matrix, the hardness increases up to 186 VHN which is more than two times the hardness of the matrix. As anticipated, the randomly dispersed alumina in the matrix acts as obstacles to dislocation glide in the slip plane [19]. However, at higher alumina weight fraction, dispersoids tend to be detrimental as shown in the decrease of hardness value to 133.5 VHN in IMC-15. The sample is presumed to segregate forming separate dispersoids and matrix phase due to a mismatch in atomic structure and surface energy which leads to low cohesiveness. Formation of lesser Ni3 Al (as indicated by AGM and XRD results), creates a weaker matrix, therefore reducing the hardness further. Additions of dispersed nanosized alumina not only results in a lesser formation of aluminide (␥ phase) due to the disruption of flow of the transient liquid aluminium but it can also create lattice mismatch between the Ni/Ni3 Al phase boundary (␥–␥ phase) [20]. The lattice parameter value of a unit ˚ calculated based on Ni3 Al and Ni diffraction patterns cell, A in sintered Ni, IM and IMC are listed in Table 4. The theoretical value for Ni3 Al based on ab initio calculation using ASW method is 0.353 nm [21]. It is found that the experimental Table 3 Comparison of hardness values for intermetallic and IMCs at different amount of addition of nanosized alumina Sintered samples
Hardness value (VHN)
IM IMC-5 IMC-15
90.9 ± 11.1 186.5 ± 36.5 133.5 ± 8.1
value for Ni3 Al in IM is in a good agreement with the theoretical value. However, the value for Ni3 Al in IMC is higher (0.3639 nm). The presence of lattice mismatch between ␥ and ␥ phase in IM and IMC are at 0.6 and 3.97%, respectively. The increase in lattice mismatch of IMC may due to the significant differences in the thermal expansions and elastic compliances between the refractory materials, alumina and its matrix, the Ni3 Al phase [22]. High lattice mismatch could results in the formation of interfacial dislocations between ␥ and ␥ phases. The ␥ precipitates become semi-coherent and the interfacial energy increases. These interfacial dislocations will also shear through the ordered precipitates. Optical images of the etched surface samples are shown in Fig. 4. There are distinguishable difference in the surface morpohologies between IM and IMCs. At 0.6% lattice mismatch, the IMs morphologies are seen as mixtures of cubic single-phase ␥ with equiaxed grains. However, the IMCs morphologies with 3.97% lattice mismatch are lamellar in nature with plate-like microstructures. The morphology of the aluminide precipitates is directly related to the lattice mismatch between ␥ and ␥ phase. Spherical shape is observed when the lattice mismatch is between 0.9 and 0.2%. It changes to cubic at mismatches between 0.5 and 1.0%, and then becomes plates at mismatches above 1.25% [23]. Table 4 Lattice constants for ␥ precipitates and lattice mismatch between ␥ precipitates and ␥ (Ni) phase in the sintered samples Phases in sintered samples
Lattice constants, ˚ (nm) A
Lattice mismatch ␥–␥
␥ (Ni3 Al in IM) ␥ (Ni3 Al in IMC) ␥ (Ni in Pure Ni)
0.3521 0.3639 0.3500
0.6% 3.97% –
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Fig. 4. Optical microscope images at 400× magnification for (a) IM showing single phase ␥ with equiaxed grains and (b) IMC5 showing lath-like microstructures.
Fig. 5. SEM micrograph using BSE for (a) IM at 843× magnification. Elemental mapping procedure reveals Ni and Al signals congregate around the fused region and (b) IMC5 at 1539× magnification. Alumina is found to be randomly dispersed using elemental mapping procedure (not shown).
Fig. 5 depicts the SEM images of etched IM and IMC observed under BSE technique. Elemental mapping on the IM surface revealed that the fused region around the pores consists mainly of Ni and Al signals indicating formation of the aluminide while the rest of the area shows mostly Ni signal. For IMC, the dispersed alumina appears to be scattered all over the IMC as indicated by the close proximity of the Al and oxygen signals. Void formations are prevalent in all samples. This maybe caused by various factors such as hollow powder particles, coalescence of absorbed gas, or inclusion of air during compression.
4. Conclusions Nickel aluminide was successfully formed by a powder metallurgy technique using nickel and aluminum metal powder as starting materials. The formation of aluminide (Ni3 Al) was identified by measuring the magnetic response of the unreacted nickel in the sintered samples. The Ms value of the sintered matrix (IM) was 0.922 emu/g. A higher Ms for IMCs was measured due to the presence of dispersed alumina leading to incomplete conversion to ␥ precipitate or aluminide. This method of measurement might be important as an initial indication on the formation of nickel-based aluminide. XRD investigations indicated that the matrix is in the form of Ni3 Al and NiAl. The lattice parameter for Ni3 Al in IM
was 0.352 nm. A higher value for Ni3 Al was measured at 0.364 nm for the IMC indicating an increased lattice mismatch between the ␥ and ␥ phase. Reaction synthesis or the exothermic reaction during sintering for intermetallic composites containing dispersed alumina was observed to occur at lower temperature than the reaction synthesis of ‘pure’ intermetallic nickel aluminide. The dispersoids were believed to be heat sinks. They disrupted the ease of flow for the transient liquid Al during reaction synthesis. Intermetallic composite with 5 wt.% of nanosized alumina showed higher hardness. However higher amount of dispersoids (15 wt.%) was detrimental because of the decrease in hardness for IMC-15. The matrix (IM) and IMCs revealed distinguishable differences in the surface morphologies. At 0.6% lattice mismatch, the IM morphologies were a mixture of cubic single-phase ␥ with equiaxed grains. However, the IMCs morphologies with 3.97% lattice mismatch were lamellar in nature with plate-like microstructures.
Acknowledgements The authors would like to thank Mr. Said Sakat for his assistance and also the financial funding from UM/IPPP/UpDiT/PASCA/02/01, UMVot F0173/2001D and UM VotF0162/2003A.
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References [1] G.H. Gessinger, Powder Metallurgy of Superalloys, Butterworth, London, 1984, p. 34. [2] N.S. Stolloff, in: C.T. Sims, N.S. Stollof, W.C. Hagel (Eds.), Superalloys II, Wiley, New York, 1987, pp. 70–80. [3] R.E. Smallman, R.J. Bishop, Metals and Materials: Science, Processes and Applications, Butterworth, Oxford, 1995, p. 327. [4] T. Takasugi, E.P. George, P.D. Pope, O. Izumi, Scripta Metall. 19 (1985) 551. [5] C.T. Liu, C.L. White, J.A. Horton, Acta Metall. 33 (2) (1985) 213–229. [6] C.T. Liu, C.L. White, in: C.C. Koch, C.T. Liu, N.S. Stollof (Eds.), High Temperature Ordered Intermetallic Alloys, MRS Proceedings, vol. 39, 1985, pp. 365–380. [7] E.M. Schulson, T.P. Wiehs, D.V. Viens, I. Baker, Acta Metall. 33 (1985) 1587. [8] S. Tohru, T. Nakajima, S. Ueda, K. Niihara, J. Am. Ceram. Soc. 80 (5) (1997) 1139–1148. [9] A.K. Misra, Metall. Trans. A 22 (1991) 2535–2538. [10] S.M. Barinov, V.Y. Evdokimov, V.Y. Shevchenko, J. Mater. Sci. Lett. 11 (1992) 1347–1348. [11] S.M. Barinov, V.Y. Evdokimov, J. Mater. Sci. Lett. 14 (1995) 820–822.
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[12] S.C. Hanyaloglu, B. Aksakal, I.J. McColm, Mater. Charac. 47 (2001) 9–16. [13] A. Bose, R.H. Rabin, R.M. German, Powder Metall. Int. 20 (1988) 25–30. [14] S.K. Mukherjee, G.P. Khanra, J. Mater. Sci. Lett. 10 (1991) 1222–1224. [15] R. Ismail, I.I. Yaacob, Mater. Sci. Forum (2002) 437–438. [16] K. Matsura, T. Kitamura, M. Kudoh, Mater. Process. Technol. 63 (1997) 298–302. [17] J.P. Lebrat, A. Varma, A.E. Miller, Met. Trans. A 23 (1992) 69– 76. [18] W.A. Glaeser, in: J.H. Westbrook, R.L. Fleischer (Eds.), Intermetallic Compounds, vol. 2, Wiley, Chicester, 1995, p. 598. [19] R. Ismail, I.I. Yaacob, Paper presented at Regional Symposium of Chemical Engineering (RSCE), Kuala Lumpur, Malaysia, 2002. [20] A.F. Giamei, D.L. Anton, Metall. Trans. 16A (1985) 1997. [21] D. Hackenbracht, J. K¨ubler, in: J.H. Westbrook, R.L. Fleischer (Eds.), Intermetallic Compounds, vol. 1, Wiley, Chicester, 1995, p. 69. [22] M.V. Nathal, R.A. Mackay, R.G. Garlick, Mater. Sci. Eng. 75 (1985) 195. [23] E.W. Ross, C.T. Sims, in: C.T. Sims, N.S. Stollof, W.C. Hagel (Eds.), Superalloys II, Wiley, New York, 1987, p. 105.
The formation of aluminides in intermetallic nickel aluminide-based nanocomposites R. Ismail, I.I. Yaacob∗ Department of Mechanical Engineering, Faculty of Engineering, Materials Engineering Programme, University of Malaya, 50603 Kuala Lumpur, Malaysia Received 7 July 2004; received in revised form 17 September 2004; accepted 20 September 2004
Abstract Metal matrix composites based on the intermetallic alloy Ni3 Al and alumina were fabricated using powder metallurgy technique. Nanosize alumina as dispersed phase was mixed with Ni and Al powder in a planetary ball mill for 4 h at 350 rpm to achieve a small mechanical alloying effect. The mixture was then compacted using a hydraulic press at 400 MPa for 15 min. Sintering was performed under an inert condition in a tube furnace at 660 ◦ C with 5 h soaking time. The formation of aluminides, Ni–Al and Ni3 Al in intermetallic nickel aluminide-based composites were observed by measuring the saturation magnetization (Ms ) values of the sintered materials. The sample’s Ms value was very low at 0.9222 emu/g indicating the presence of a small amount of elemental Ni. The occurrence of reaction synthesis during sintering was detected by the presence of a ‘large’ exothermic peak of the differential thermal analysis curve, which occurred at below 600 ◦ C. The heat released during reaction synthesis was probed by differential scanning calorimetry. The heat of reaction was found to be 4.53 × 104 J/mol. Addition of 5% alumina increased the hardness of the composite by about a factor of two with respect to that of Ni3 Al intermetallic. Xray diffraction (XRD) measurements showed peaks corresponding to Ni–Al and Ni3 Al. Optical and SEM investigations revealed growth of lath-like microstructure in the fabricated composite due to a high lattice mismatch between ␥ and ␥ phase. © 2004 Elsevier B.V. All rights reserved. Keywords: Intermetallics; Powder metallurgy; Thermal analysis; Magnetic measurements; Nanocomposites
1. Introduction The intermetallic compound Ni3 Al is well known for its oxidation resistance and high temperature-strength. The latter property is derived from precipitation hardening of the coherent, ordered Ni3 Al ␥ phase within the disordered, solid solution ␥ phase. The compatibility of the ␥ precipitates FCC crystal structure with a lattice constant of approximately 0.1% mismatch with the ␥ phase (Ni matrix phase) results in homogenous nucleation of the precipitates. The precipitates have low surface energy and extraordinary long-time stability. Coherency at elevated temperature between ␥ and ␥ is maintained by tetragonal distortion along the z-axis of the precipitates [1].
∗
Corresponding author. Tel.: +60 379674489/5681; fax: +60 379675317. E-mail address: [email protected] (I.I. Yaacob).
0925-8388/$ – see front matter © 2004 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2004.09.032
This unique ␥ or aluminide intermetallic phase contributes to remarkable strengthening of the intermetallic by forcing dislocations to either bypass the larger aluminide precipitates by Orowan bowing or dislocations cutting of the smaller precipitates in the ␥ –␥ alloy [2]. Moreover, the strength of ␥ increases as temperature increases which is known as anomalous yield behaviour [3,4]. The fact that the intermetallic can be intrinsically strong and can retain this strength at elevated temperatures, makes it especially attractive for high temperature uses. Its low density (7.5 g/cm3 ) is an added advantage to the specific strength where major stresses are determined by the mass of an object. Low selfdiffusion in an ordered intermetallic increases the microstructural stability and decreases creep rates [5]. The intrinsic grain boundary brittleness in Ni3 Al is associated with ordering energy, electronegativity difference, valency difference and atomic size difference between Ni atoms and Al atoms. However, microalloying with boron sharply increases its ductil-
R. Ismail, I.I. Yaacob / Journal of Alloys and Compounds 392 (2005) 214–219
215
All the materials were used as received. About 100-mesh nickel powder (99.9% purity) was purchased from Strem Chemicals. Aluminum powder (99.7% purity) and nanosized alumina (average size of 39 nm) were obtained from Aldrich chemicals. Special etchant was prepared using 3 g of molybdic acid powder from Ajax Chemicals, 100 ml of distilled H2 O, 100 ml of hydochloric acid (assay 37%) from Reidel de Haen Chemicals and 100 ml of nitric acid (assay 65%) from Merck. Elemental nickel and aluminum powders at 3:1 mole ratio were pre-mixed with different weight fractions of nanosized alumina (as shown Table 1) using a planetary ball mill (PM400, MA Type, Retsch). The mixtures were milled for 4 h at 175 rpm in a zirconia jar with zirconia balls. The milled samples were then compacted using a hydraulic hand press (Gatesby Specac) in a 20 mm die at 400 MPa for 15 min. Sintering of the green compacts was performed under argon gas atmosphere using a tube furnace (Carbolite, Type CTF 17/75/300, maximum T: 1700 ◦ C). The tube furnace containing the samples was initially purged with argon gas (1 l/min)
and then vacuumed. This process was repeated several times. The samples were heated under flowing argon at a heating rate of 20 ◦ C/min until 660 ◦ C, at which they were isothermally heated for a duration of 5 h [15]. The reaction synthesis was monitored by observing the differential thermal analysis (DTA) curves when the samples were sintered under argon gas atmosphere in a small furnace (TGA/SDTA 851e , Mettler Toledo). A green compact sample weighing about 40–70 mg was placed directly on the DTA sample holder. This was done to ensure an optimum sensitivity during sintering. However, a silica disc was placed underneath the sample to avoid any possible contaminations and degradation of the thermocouples during heating. Changes in the sample, which led to absorption or evolution of heat, were detected relative to the inert reference. The amount of heat released during the reactive sintering process can be measured by heating the samples from 30 to 670 ◦ C under inert argon gas atmosphere using a differential scanning calorimeter (Mettler Toledo DSC820 with FRS5 sensor (Heat flux type)). A small piece of the green compact weighing about 10 mg was placed in a 40 l platinum crucible which was then closed with a Pt lid with a small hole. The sample crucible was placed side by side with an empty Pt reference crucible. The saturation magnetization values for the sintered samples were measured using an alternating gradient magnetometer (AGM) (Micromag 2900) with maximum applied fields of ±10 kOe. A small amount of the sample was mounted on an extension rod attached to a piezoelectric element and placed in a region between two electomagnets. The resulting deflection of the rod was measured by a piezoelectric sensing element mounted on the probe arm. The magnetic response of the sample was plotted in terms of the sample’s magnetization against the magnitude of the applied field. The phases present in the samples before and after sintering were examined by an X-ray diffractometer (XRD) using Cu K␣ radiation (Philips X-Pert MPD PW3040). The diffraction pattern was obtained by scanning the sample from 15◦ to 90◦ 2θ angle at a step size of 0.02◦ and a count time of 1.5 s at each step. The hardness of the sintered samples was determined by Vickers indentation technique using 10 kgf load. For morphological characterizations, the sintered sample was mounted on a resin and polished until mirror finish. The sample’s surface was etched with a specially prepared etchant described earlier. Excess etchant was washed away with running water. The sample was then dried with a blower. Surface morphologies were observed under an optical microscope and a scanning electron microscope (Philips XL40 SEM).
Table 1 Compositions of intermetallic composites (IMCs)
3. Results and discussion
ity and effectively suppresses intergranular fracture [6]. The ductilizing effect of boron can be explained by two possible mechanisms: (i) boron-enhanced grain boundary cohesive strength [7] and (ii) boron-facilitated slip transfer across the grain boundary (or slip nucleation at boundaries) [8]. To improve monolithic Ni3 Al and NiAl properties such as low ambient tensile ductility and insufficient hightemperature strength and creep resistance limit, ceramic particulates are added as reinforcement fillers. The addition of lower density ceramic materials (i.e. oxides, carbides and nitrides) also decreases the composite’s density and increases its specific properties. However, most of these studies are still in the feasibility stage aiming primarily at determining basic mechanical properties and chemical stability [9–11]. In this paper, nanosized alumina powder was used as the reinforcement. Powder metallurgy technique was used for fabricating the composite by which the alumina, nickel and aluminum powders were mixed, compacted and then sintered. Exothermic reaction between the Ni and Al metal powders during the sintering process enabled the formation of intermetallics at temperatures as low as 500 ◦ C [12–14]. The extremely small size of aluminides made them difficult to be observed under an optical or electron microscope. By measuring the saturation magnetization (Ms ) values of the sintered sample, we were able to quantify the amount of aluminides formed in each specimen.
2. Experimental
Samples
Code
Ni3 Al matrix Ni3 Al + 5 wt.% Al2 O3 Ni3 Al + 15 wt.% Al2 O3
IM IMC-5 IMC-15
The saturation magnetization (Ms ) values listed in Table 2 indicate the presence of unreacted elemental nickel in the samples. The formation of aluminide or ␥ precipitate during
216
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Table 2 Saturation magnetization (Ms ) values for intermetallic, IMCs and pure Ni powder Sintered samples
Saturation magnetization (Ms ) value (emu/g)
IM IMC-5 IMC-15 Ni powder
0.922 3.363 15.38 48.20
sintering reduces the amount of elemental nickel in the samples hence lowering the Ms value. The Ms value of the sintered matrix (IM) is 0.922 emu/g and is comparable with the Ms value of commercial Ni3 Al (Alfa Aesar, lot # F18l52, powder form), which is 0.3927 emu/g. However, the Ms values for the IMC sample are higher. The higher Ms for the IMCs is due to the presence of dispersed alumina which causes incomplete conversion of the metal powder mixture to ␥ precipitate or aluminide. The Ms value for IMC-15 is about 31% of the Ms of pure Ni powder indicating that IMC-15 (containing about 73.95 wt.% Ni) contains about 43 wt.% unreacted elemental nickel. The incomplete conversion of elemental Ni and Al into aluminide in the sintered materials can be further observed form the DTA curves of the nickel aluminide matrix and composites. Differential thermal analysis (DTA) curves show occurrence of reaction synthesis (Fig. 1). Large exothermic peaks as the result from a sudden release of energy emerge in every sintered sample. Reaction synthesis is expected to occur due to the formation of a transient liquid phase from the lower melting point minor phase (aluminum) encapsulating the nickel powder by capillary force [12,16]. Evaluation of the peak’s position shows that reaction synthesis for the intermetallic nickel aluminide composites containing 5 and 15 wt.% alumina as dispersed phase were observed to occur at lower temperatures compared to the ‘pure’ intermetallic nickel aluminide. The disruption of the ease of flow of the liquid aluminum by the dispersed alumina is expected to reduce the reaction synthesis temperature. The presence of alumina
Fig. 2. DSC curve for Ni3 Al shows large exothermic peak indicating heat evolution as the results of reaction synthesis.
as non-reacting dispersed phase can act as heat sinks and thermal barriers limiting temperature rise and also impede the flow of transient aluminum liquid [17]. Measurement of the heat released during reaction synthesis of Ni3 Al can be made by extrapolating the width of the exothermic peak in a DSC curve as shown in Fig. 2. However, two exothermic peaks can be seen. The larger peak occurs at 597 ◦ C with a heat release of 223.23 J/g. The smaller one is at 640.4 ◦ C with a heat release of 4.27 J/g. The heat of reaction at 597 ◦ C for 3Ni + Al powder compact weighing 0.0098 g is calculated to be −4.53 × 104 J/mol. Identification of the desired ␥ precipitate during the reaction synthesis can be performed by observation of the XRD patterns shown in Fig. 3. Comparisons of d-spacings of the XRD peaks with the ICDD files confirmed the presence of both aluminides in the form of ordered ␥ (Ni3 Al) and  (Ni–Al) phases. The presence of dispersed alumina in the IMCs leads to a lower aluminide formation in the sintered samples which is consistent with the higher Ms values. The presence of a prominent Ni (2 0 0) peak in IMC15 and the emergence of the same peak in IMC5 can be observed.
Fig. 1. DTA curves for intermetallic matrix and IMCs. The specimens were heated from 30 to 660 ◦ C at 20 ◦ C/min and were then isothermally heated at 660 ◦ C for 5 h.
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Fig. 3. X-ray diffraction patterns of pure Ni, intermetallic Ni3 Al (IM) and intermetallic composites (IMCs). The peaks are labeled accordingly.
The hardness values of the sintered samples are shown in Table 3. An average hardness of 90.9 VHN is obtained for the Ni3 Al matrix. This is at the lower end of the cited literature value of 100–255 VHN [18]. When 5 wt.% of nanosized alumina is dispersed in the matrix, the hardness increases up to 186 VHN which is more than two times the hardness of the matrix. As anticipated, the randomly dispersed alumina in the matrix acts as obstacles to dislocation glide in the slip plane [19]. However, at higher alumina weight fraction, dispersoids tend to be detrimental as shown in the decrease of hardness value to 133.5 VHN in IMC-15. The sample is presumed to segregate forming separate dispersoids and matrix phase due to a mismatch in atomic structure and surface energy which leads to low cohesiveness. Formation of lesser Ni3 Al (as indicated by AGM and XRD results), creates a weaker matrix, therefore reducing the hardness further. Additions of dispersed nanosized alumina not only results in a lesser formation of aluminide (␥ phase) due to the disruption of flow of the transient liquid aluminium but it can also create lattice mismatch between the Ni/Ni3 Al phase boundary (␥–␥ phase) [20]. The lattice parameter value of a unit ˚ calculated based on Ni3 Al and Ni diffraction patterns cell, A in sintered Ni, IM and IMC are listed in Table 4. The theoretical value for Ni3 Al based on ab initio calculation using ASW method is 0.353 nm [21]. It is found that the experimental Table 3 Comparison of hardness values for intermetallic and IMCs at different amount of addition of nanosized alumina Sintered samples
Hardness value (VHN)
IM IMC-5 IMC-15
90.9 ± 11.1 186.5 ± 36.5 133.5 ± 8.1
value for Ni3 Al in IM is in a good agreement with the theoretical value. However, the value for Ni3 Al in IMC is higher (0.3639 nm). The presence of lattice mismatch between ␥ and ␥ phase in IM and IMC are at 0.6 and 3.97%, respectively. The increase in lattice mismatch of IMC may due to the significant differences in the thermal expansions and elastic compliances between the refractory materials, alumina and its matrix, the Ni3 Al phase [22]. High lattice mismatch could results in the formation of interfacial dislocations between ␥ and ␥ phases. The ␥ precipitates become semi-coherent and the interfacial energy increases. These interfacial dislocations will also shear through the ordered precipitates. Optical images of the etched surface samples are shown in Fig. 4. There are distinguishable difference in the surface morpohologies between IM and IMCs. At 0.6% lattice mismatch, the IMs morphologies are seen as mixtures of cubic single-phase ␥ with equiaxed grains. However, the IMCs morphologies with 3.97% lattice mismatch are lamellar in nature with plate-like microstructures. The morphology of the aluminide precipitates is directly related to the lattice mismatch between ␥ and ␥ phase. Spherical shape is observed when the lattice mismatch is between 0.9 and 0.2%. It changes to cubic at mismatches between 0.5 and 1.0%, and then becomes plates at mismatches above 1.25% [23]. Table 4 Lattice constants for ␥ precipitates and lattice mismatch between ␥ precipitates and ␥ (Ni) phase in the sintered samples Phases in sintered samples
Lattice constants, ˚ (nm) A
Lattice mismatch ␥–␥
␥ (Ni3 Al in IM) ␥ (Ni3 Al in IMC) ␥ (Ni in Pure Ni)
0.3521 0.3639 0.3500
0.6% 3.97% –
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Fig. 4. Optical microscope images at 400× magnification for (a) IM showing single phase ␥ with equiaxed grains and (b) IMC5 showing lath-like microstructures.
Fig. 5. SEM micrograph using BSE for (a) IM at 843× magnification. Elemental mapping procedure reveals Ni and Al signals congregate around the fused region and (b) IMC5 at 1539× magnification. Alumina is found to be randomly dispersed using elemental mapping procedure (not shown).
Fig. 5 depicts the SEM images of etched IM and IMC observed under BSE technique. Elemental mapping on the IM surface revealed that the fused region around the pores consists mainly of Ni and Al signals indicating formation of the aluminide while the rest of the area shows mostly Ni signal. For IMC, the dispersed alumina appears to be scattered all over the IMC as indicated by the close proximity of the Al and oxygen signals. Void formations are prevalent in all samples. This maybe caused by various factors such as hollow powder particles, coalescence of absorbed gas, or inclusion of air during compression.
4. Conclusions Nickel aluminide was successfully formed by a powder metallurgy technique using nickel and aluminum metal powder as starting materials. The formation of aluminide (Ni3 Al) was identified by measuring the magnetic response of the unreacted nickel in the sintered samples. The Ms value of the sintered matrix (IM) was 0.922 emu/g. A higher Ms for IMCs was measured due to the presence of dispersed alumina leading to incomplete conversion to ␥ precipitate or aluminide. This method of measurement might be important as an initial indication on the formation of nickel-based aluminide. XRD investigations indicated that the matrix is in the form of Ni3 Al and NiAl. The lattice parameter for Ni3 Al in IM
was 0.352 nm. A higher value for Ni3 Al was measured at 0.364 nm for the IMC indicating an increased lattice mismatch between the ␥ and ␥ phase. Reaction synthesis or the exothermic reaction during sintering for intermetallic composites containing dispersed alumina was observed to occur at lower temperature than the reaction synthesis of ‘pure’ intermetallic nickel aluminide. The dispersoids were believed to be heat sinks. They disrupted the ease of flow for the transient liquid Al during reaction synthesis. Intermetallic composite with 5 wt.% of nanosized alumina showed higher hardness. However higher amount of dispersoids (15 wt.%) was detrimental because of the decrease in hardness for IMC-15. The matrix (IM) and IMCs revealed distinguishable differences in the surface morphologies. At 0.6% lattice mismatch, the IM morphologies were a mixture of cubic single-phase ␥ with equiaxed grains. However, the IMCs morphologies with 3.97% lattice mismatch were lamellar in nature with plate-like microstructures.
Acknowledgements The authors would like to thank Mr. Said Sakat for his assistance and also the financial funding from UM/IPPP/UpDiT/PASCA/02/01, UMVot F0173/2001D and UM VotF0162/2003A.
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