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Investigation of the characteristics of the nanocrystalline Ni3Al-based alloy fabricated by hot pressing and sintering
Journal of Alloys and Compounds 492 (2010) 196–200
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jallcom
Investigation of the characteristics of the nanocrystalline Ni3 Al-based alloy fabricated by hot pressing and sintering Masoud Nazarian-Samani ∗ , Ali Reza Kamali Advanced Materials Research Center, Department of Materials Science and Engineering, Malek Ashtar University of Technology (MUT), Ferdowsi St., Shahin-Shahr, Isfahan, Iran
a r t i c l e
i n f o
Article history: Received 13 July 2009 Accepted 22 November 2009 Available online 2 December 2009 Keywords: Ni3 Al-based alloy Hot pressing Sintering Mechanical alloying Grain growth
a b s t r a c t The characteristics of specimens produced from a Ni3 Al-based alloy using four different routes were investigated and different trends during consolidation of mechanically alloyed (MA-ed) powders were characterized and analyzed using microstructural observations; measurement of long-range order (LRO) parameter, micro-strain, and grain size; as well as determination of microhardness and density. In addition, the compressive yield strength and ductility of two series of specimens at different temperatures were determined. It was observed that hot pressing (HP) led to the best consolidation conditions. In addition, high pressure cold pressing (CP)/sintering caused the consolidated specimens to improve. In all specimens, the nature of nanocrystalline powders remained unchanged and the disordered MA-ed powders were ordered during consolidation. The yield strength in both series of specimens anomalously increased with increased temperature by up to 600 ◦ C, which later decreased beyond this critical temperature. © 2009 Elsevier B.V. All rights reserved.
1. Introduction Ni3 Al-based alloys have attracted a lot of attention due to their promising combination of properties such as high melting point, low density, high corrosion resistance, and good high-temperature strength. The most attractive property of these alloys is their enhanced yield strength with increasing temperature by up to 600–800 ◦ C [1–6]. The synthesis of nanocrystalline Ni3 Al by MA from elemental powders has been successfully achieved by several researchers [7–12] and subsequent consolidation has been accomplished by means of HP [7,8,12] and sintering processes [11]. While the microstructure of the specimens has been frequently discussed in previous studies, few systematic studies of the physical and mechanical properties of Ni3 Al and its alloys have been reported to date. In the present work, we report the results of a study that assesses the effect of HP and CP followed by sintering processes on the microstructure as well as the physical and mechanical properties of a MA-ed Ni3 Al-based alloy. The effects of pressure applied during CP and sintering time have also been investigated. 2. Experimental procedure The Ni3 Al-based alloy with the nominal composition Ni–8.14Al–7.83Cr– 1.45Mo–0.01B (wt%) was prepared by MA of elemental powders. All the powders
∗ Corresponding author. Tel.: +98 912 7085285; fax: +98 312 5228530. E-mail address: [email protected] (M. Nazarian-Samani). 0925-8388/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2009.11.182
used had an average particle size of 10 m. The powders were wet milled under argon in a Fritsch P6 with stainless steel balls for 15 h, using a ball to powder weight ratio (BPR) of 10:1. Details of the MA experiments have been presented elsewhere [12]. The 15 h MA-ed powder was subjected to four different processing routes, each simply termed specimens I through IV as listed in Table 1. CP and HP tests were performed by a uni-axial cold compaction machine and a HP machine (model ASTRO H-20-4560) in a graphite mold, respectively. All sintering processes were conducted under a high purity argon atmosphere. For metallographic and microhardness investigations, the specimens were carefully ground and polished. An etching solution, made of a mixture of glycerin, HCl and HNO3 with a ratio of 2:2:1, was applied to the surface of the polished samples for different times between 15 and 50 s. Vickers microhardness was measured in a Koopa UV1 machine. All specimens were inspected for their microstructure by a “Vega© Tescan” scanning electron microscopy (SEM) with an energy-dispersive X-ray spectroscopy (EDS). The microstructure of fracture surfaces was examined using a Hitachi-make S3400-N SEM. The X-ray diffraction (XRD) profiles of the specimens were carried out using Cu K␣ radiation ( = 0.15406 nm) and a Seifert 3003TT diffractometer operating at 40 kV and 30 mA. The grain size, d, and micro-strain, e, were calculated by applying the well-known Williamson–Hall method [12,13]. The lattice parameter was determined using at least five higher angle 2 peaks. From the XRD patterns, the LRO parameter, S, was determined by comparing the relative intensities of the superlattice (h k l mixed) and the fundamental (h k l are all even or all odd) peaks of all the specimens with respect to a well-annealed reference material [13]: S2 =
(Isup /Ifun )obs (Isup /Ifun )std
(1)
where (Isup /Ifun )obs and (Isup /Ifun )std are the intensities of the superlattice reflection relative to the fundamental line of the specimens (obs) and the fully annealed reference (std), respectively. In this study, the LRO parameter was estimated from the (1 0 0)/(2 0 0) and (1 1 0)/(2 2 0) pairs of reflections. Compressive tests were performed on cylindrical samples (D: 4 mm, L: 6 mm) at room temperature as well as temperatures 200, 400, 600, 700, and 800 ◦ C using an
M. Nazarian-Samani, A.R. Kamali / Journal of Alloys and Compounds 492 (2010) 196–200 Table 1 Different routes of consolidation experiments. Designation
Process
I II III IV
CP under 700 MPa + sintering at 1000 ◦ C for 1 h CP under 700 MPa + sintering at 1000 ◦ C for 2 h CP under 2 GPa + sintering at 1000 ◦ C for 1 h HP at 1000 ◦ C under 6 MPa for 10 min
Instron model 8503 tensile/compressive testing machine at a strain rate of 10−2 s−1 . All the data were reported as the average of at least three test results. The density of consolidated samples was determined using the immersion method in distilled water, based on the Archimedes Principle.
3. Results and discussion 3.1. XRD studies Fig. 1(a)–(e) shows the comparison of XRD patterns for the 15 h-milled powder [12] and the specimens I–IV. The results indicate that no superlattice reflection peaks were detected in the as-milled powder, demonstrating the formation of the fcc phase. The (1 0 0) superlattice peak appeared in specimen I after sintering for 1 h. Increasing the applied pressure during CP and the sintering time caused the amount of ordered structure to increase and the (1 1 0) and (2 1 0) superlattice peaks to appear. However, these were the only superlattice peaks that were revealed, and the consolidated samples may not, therefore, be fully ordered. Other researchers have reported the (2 1 1) superlattice peak to have also been revealed during consolidation [7,8]. Generally, the superlattice diffraction peaks are weak and their detection is not so easy in the Ni3 Al-based alloys. The addition of alloying elements will further affect their detectability. Although the addition of B is effective in stabilizing L12 structure [14], the fully ordered structure cannot be obtained here due to the presence of Cr and Mo. According to Fig. 1, specimen IV exhibits a greater atomic ordering than specimens I–III. This can be explained by the homogenization effect of HP on the composition, in addition to the annealing of the defects. From Table 2, it can be seen that the LRO parameter calculated from the (1 1 0)/(2 2 0) pair is slightly greater than that of the (1 0 0)/(2 0 0) pair in all specimens, which may be because of their texture, extinction, or other reasons [15]. Fig. 1 also reveals that the peak intensities of all the specimens increased and that the peak widths declined during the consolidation processes. Heating the milled Ni3 Al also resulted in a volume expansion as shown by the increased lattice parameter in all the specimens. It is shown that the sintering time causes the increase of the lattice parameter due to the greater dissolution of Cr, Mo, and Al in
Fig. 1. XRD patterns of as-milled powder [12] (a), and specimens I (b), II (c), III (d), and IV (e).
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the Ni3 Al structure. The atomic radius of Al (0.143 nm) and Mo (0.136 nm) is larger than that of Ni (0.125 nm), whereas the atomic radius of Cr (0.125 nm) is equal to that of Ni [16]. Therefore, diffusion of Al and Mo atoms into the Ni lattice leads to an increase in the Ni lattice parameter and the transfer of X-ray peaks of Ni3 Al to lower angles. This was more pronouncedly observed in specimen IV than in other specimens, perhaps due to the applied pressure during HP process. The above explanations are more vividly presented in Table 2. Here, it is also shown that micro-strain decreased during the consolidation process. As expected, the sintering time led to greater annealing of the structure and to the decreased microstrain. The lowest micro-strain was observed in specimen II due to its long time sintering. Also, the grain size slightly increased in all specimens, with its highest value observed in specimen II. It can be seen that the grain size was still small in all specimens and retained its nanocrystalline size after all consolidation routes. The presence of Mo, Cr, and B dissolved in the Ni3 Al phase led to the production of a nanocrystalline specimen even after a high-temperature consolidation at 1000 ◦ C. Generally, the alloying elements segregate in grain boundaries and, therefore, reduce the grain-boundary mobility and surface free energy of the material. Moreover, grain growth in nanocrystalline powders is prevented due to the presence of open porosities in the microstructure. When the porosities are closed and isolated, grains grow without any dike [17]. This is why the grains in specimens I and II did not grow very much and also why grains in specimen III had their maximum size. Moreover, the grains in specimen IV are smaller than those in specimen III; this is due to simultaneous application of pressure and heat. Application of pressure hinders grain growth at elevated temperatures. This is due to the fact that diffusion is involved in grain growth and that the diffusion coefficient decreases with increasing pressure [18]. It can thus be concluded that pressure may reduce grain-boundary mobility. 3.2. Physical and mechanical properties The phenomena involved in cold compression of powders include rearrangement of particles, elastic deformation of particles at their contacts, plastic deformation in metals, and particles crashing in brittle materials. These phenomena are present in the compression of nanocrystalline powders. However, the behavior of very fine particles, i.e., slip, friction, and their diffusion-assisted deformation, can be considerably different from that of larger particles. Slip and rearrangement mechanisms are limited in nanocrystalline powders due to the presence of high friction forces between particles. Surface tension and van der waals as well as electrostatic and mechanical forces are the origins of this friction. Moreover, irregular grain boundaries facilitate the agglomeration of the particles. These facts prevent the rearrangement of particles. Therefore, the possibility for nanocrystalline powders to achieve a high density green compact during compression is lower than that for micron-sized particles [17]. Moreover, nanocrystalline powders have a high potential for agglomeration during MA [13]. Presence of agglomerated powders increases the possibility for large porosities to form in the material during consolidation. Extensive agglomeration of 15 h-milled powders can be seen in Fig. 2. At the atomic level, sintering induces contacting particles to bond together. The most important atomic events are surface diffusion (at lower temperatures) and grain-boundary diffusion (at higher temperatures). Grain-boundary diffusion predominantly occurs at the contacts between particles. On the microstructural scale, the bonding becomes evident as necks grow between touching particles. Such neck growth increases the strength over the green strength and causes many beneficial property changes [19]. However, the HP of powders has certain major benefits compared to CP/sintering. The pressure used in the HP process leads to plastic
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Table 2 Characteristics of the as-milled powder [12] and the specimens I–IV. Specimen
Microhardness (VHN)
(g/cm3 )
a (Å)
e (%)
d (nm)
As-milled I II III IV
– 216 263 387 479
– 4.28 4.65 6.73 7.79
3.60 3.62 3.69 3.64 3.76
0.59 0.28 0.19 0.31 0.36
10.66 38.34 46.93 34.62 28.11
deformation of powders and, therefore, to their interpenetration and decreased porosity of compacts during the consolidation process. Moreover, applied pressure leads to the rupture of oxide layers on the surface of the powders and, thereby, to their more effective connection. In pressure-assisted sintering of agglomerated powders, large porosities can be removed due to the effects cited above which lead to minimized grain growth. It should be noted that applying the pressure at high temperatures leads to decreased resistance of the material against deformation and to increased destruction of porosities. On the other hand, destruction of large porosities at low temperatures is so difficult even when high pressures are applied. Moreover, shear tensor of the applied pressure leads to the rearrangement of the particles and destruction of large porosities. More connection points, therefore, emerge among the particles. The oxide layer formed on the surface of the particles fractures due to the applied shear stress creating a better particle connection. It can be concluded that using pressure during the sintering process leads to the removal of large porosities and to the prevention of grain growth. The smallest grain size in specimen IV can also be explained along the same lines. The relationship between yield strength and grain size in materials can be evaluated using the Hall–Petch equation [17,20]: ∝ d−(1/2)
(2)
where and d are yield stress and grain size, respectively. According to this equation, plastic deformation of particles is more difficult in smaller particles. The relatively low density obtained in specimen III (Table 2) can be explained by considering the yield strength of the alloy at room temperature (600 MPa) and the pressure used in
Fig. 2. SEM micrograph of as-milled powder indicates agglomeration during MA.
LRO (1 0 0)/(2 0 0)
(1 1 0)/(2 2 0)
0 0 0.08 0.22 0.41
0 0.06 0.11 0.32 0.45
the CP process (2 GPa). The lack of a main plastic deformation in powder particles during consolidation is responsible for this low density in the consolidated specimen. Increasing the sintering time from 1 to 2 h led to an increase in density from 4.28 to 4.65 g/cm3 . However, prolonged sintering time beyond 2 h had no effect on the improvement of density. On the other hand, the CP of powders under a high pressure of 2 GPa and the sintering of the green compacts led to the production of consolidated specimens with a considerable density of 6.73 g/cm3 . It can be seen in Table 2 that application of a pressure as low as 6 MPa during the HP process yields the highest density equal to 7.79 g/cm3 . The greater ductility of the particles at a temperature as high as 1000 ◦ C as well as the enhanced diffusion of atoms under the applied stress may be responsible for the considerable consolidation of HP-ed specimen IV. It should be noted that a nanocrystalline material has a lower density than the same material with a normal grain size. This is due to the high amount of grain boundaries in nanocrystalline materials. Grain boundaries have a lower density than grains do. According to Table 2, microhardness increases in the specimens with increased sintering time, which is related to increased density during sintering, the highest value for microhardness belonging to specimen IV. On the other hand, the HP process obtains better conditions for the production of a bulk sample. It was observed that samples sintered for 1 and 2 h did not have sufficient handling strength. Thus, compressive tests were not further continued for measuring yield strength and ductility, but these tests were only performed for specimens III and IV. The yield strengths of both specimens show a positive temperature dependence in the range from room temperature to 600 ◦ C where the strength peak was reached. Further increase in temperature caused a corresponding decrease in strength (Fig. 3). Specimen IV showed yield strengths almost twice as much as that of specimen III for all temperatures applied. This anomalous behavior is common to Ni3 Al and its alloys, as also pointed out by many researchers [1–6]. HP-ed specimens showed
Fig. 3. Comparison of compressive yield strength, y , and ductility, ε, of specimens III (-♦- and --) and IV (-- and -×-) at different temperatures.
M. Nazarian-Samani, A.R. Kamali / Journal of Alloys and Compounds 492 (2010) 196–200
199
Fig. 4. SEM micrographs of specimens II (a), III (b), and IV (c).
the highest yield strength as compared to some literature data [21]. These results represent the high effectiveness of the HP process used compared to other processes. Fig. 3 also presents the variations in ductility of both specimens III and IV, indicating larger ductility and yield strength in specimen IV. This increased yield strength and improved ductility were mainly caused by the finer grain size and lower porosities and pores in specimen IV compared to those in specimen III. It should be noted that the presence of porosities and the non-effective connection between particles lead to considerably reduced ductility of specimens. Moreover, it is observed that compressive yield strength as a function of the consolidation routes used follows trends quite similar to those of density and microhardness in both specimens. It can, therefore, be concluded that yield strength was significantly dependent on the density and microhardness of the consolidated specimens. 3.3. Microstructure The microstructures of specimens II–IV are shown in Fig. 4(a)–(c). It is observed that the CP-ed powders at a pressure of 700 MPa could not be fully consolidated after 1 and 2 h of sintering, but that powders CP-ed under a high pressure (2 GPa) in specimen III led to the formation of a relatively dense sample. The grain boundaries did not appear in specimen IV even after
etching over long periods (50 s). However, the grain boundaries in specimen III appeared only after a light etching (15–20 s). On the other hand, according to Fig. 4(c), the surface is smooth and no pore exists between the bonded powder particles, which are evidence of the good consolidation quality. Table 3 shows the results of EDS analysis for as-milled powder and the specimens I–IV produced via various consolidation routes. The results indicate that the bulk samples were identical to the as-milled powder in composition, and that Al, Cr, and Mo still dissolved in Ni3 Al after sintering at 1000 ◦ C. This is evidenced by the fact that no peaks of these elements and/or their compounds were observed in the XRD patterns (Fig. 1). This suggests that the Ni3 Albased alloy has excellent thermal stability at high temperatures.
Table 3 EDX analysis of the as-milled powder [12] and specimens I–IV. Specimen
As-milled I II III IV a
B was not detected.
Element (wt%)a Ni
Al
Cr
Mo
82.26 82.09 81.97 81.92 81.96
8.09 8.11 8.13 8.13 8.19
7.98 7.99 8.04 8.07 8.14
1.67 1.81 1.86 1.88 1.71
200
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2. Application of pressure during consolidation hinders grain growth at elevated temperatures. Also, the pressure helps to produce a free-pore or void-bulk sample. 3. Both specimens III and IV exhibit a positive temperature dependence of their compression yield strength; maximum yield strength obtained at 600 ◦ C followed by a decrease at higher temperatures. 4. The nature of the nanocrystalline structure remains intact after consolidation. In addition, the ordering transition occurs during consolidation, which is greater in specimen IV than in other specimens. 5. The removal of MA-ed powder agglomeration should be accomplished before the sintering in order to obtain a specimen with the minimum porosity. References Fig. 5. Fracture surface of specimen III indicating crack propagation.
From Table 3, it is also concluded that the specimens were not obviously contaminated by oxygen and/or other elements during different consolidation routes used in this study. Another role of grain boundaries in specimen IV when compared with specimen III is the faster diffusion shortcuts existing in the grain boundaries. These fast shortcuts make a few initial tiny cracks to quickly close during the deformation process. Therefore, the crack growth and propagation can be avoided to a certain extent. Nevertheless, the high pressure CP-sintering technique unavoidably induced grain growth, which is harmful to their properties, especially their yield strength and ductility. It can be seen from Fig. 5 that cracks easily propagated during the compression test in specimen III, but no crack was observed from fracture surfaces of specimen IV (not shown). Based on the results in this investigation, it can be seen that MA-HP process is an effective method for preparing Ni3 Al-based alloys with improved properties. 4. Conclusions The following conclusions may be drawn from the results obtained in this study: 1. The HP process produces better specimens in terms of density, microhardness, grain size, microstructure, yield strength, and ductility than other processes investigated in the present study.
[1] N.S. Stoloff, Int. Mater. Rev. 34 (1989) 153–183. [2] J.H. Westbrook, R.L. Fleischer, Intermetallic Compounds: Structural Applications, vol. 3, John Wiley, New York, 2000. [3] G. Sauthoff, in: R.W. Chan, P. Hassen, E.J. Kramer (Eds.), Materials Science and Technology: A Comprehensive Treatment, vol. 8, VCH Publishers Inc., New York, 2005, pp. 643–803. [4] S.C. Deevi, V.K. Sikka, Intermetallics 4 (1996) 357–375. [5] S.C. Deevi, V.K. Sikka, C.T. Liu, Prog. Mater. Sci. 42 (1997) 177–192. [6] V.K. Sikka, J.T. Mavity, A. Anderson, Mater. Sci. Eng. A 153 (1992) 712–721. [7] J. Meng, C. Jia, Q. He, J. Alloys Compd. 421 (2006) 200–203. [8] M. Krasnowski, A. Antolak, T. Kulik, J. Alloys Compd. 434–435 (2007) 344– 347. [9] M.H. Enayati, Z. Sadeghian, M. Salehi, A. Saidi, Mater. Sci. Eng. A 375–377 (2004) 809–811. [10] L. Lu, M.O. Lai, S. Zhang, Mater. Des. 15 (1994) 79–86. [11] L. D’Angelo, G. Gonzalez, J. Ochao, J. Alloys Compd. 434–435 (2007) 348–353. [12] M. Nazarian-Samani, A. Shokuhfar, A.R. Kamali, M. Hadi, J. Alloys Compd. (2009), doi:10.1016/j.jallcom.2009.02.140. [13] C. Suryanarayana, Mechanical Alloying and Milling, Marcel Dekker, New York, 2004. [14] M. Nazarian-Samani, A.R. Kamali, J. Alloys Compd. 486 (2009) 315–318. [15] R. Ramesh, R. Vasudevan, B. Pathiraj, B.H. Kolster, J. Mater. Sci. 27 (1992) 270–278. [16] J.F. Shackelford, W. Alexander, Materials Science and Engineering Handbook, CRC Press LLC, 2001. [17] H.E. Exner, E. Artz, in: R.W. Cahn, P. Haasen (Eds.), Physical Metallurgy, vol. 3, Elsevier Science, Amsterdam, 2005, pp. 2627–2662. [18] P.G. Shewmon, Diffusion in Solids, McGraw-Hill, New York, 1963. [19] J.R. Groza, J.F. Shackelford, E.J. Lavernia, M.T. Powers, Materials Processing Handbook, CRC Press, New York, 2007. [20] V.S. Arunachalam, R. Sundaresan, in: R.W. Chan, P. Hassen, E.J. Kramer (Eds.), Materials Science and Technology: A Comprehensive Treatment, vol. 15, VCH Publishers Inc., New York, 2005, pp. 137–192. [21] D. Bozic, N. Ilic, M. Mitkov, M.T. Jovanovic, M. Zdujic, J. Mater. Sci. 31 (1996) 3213–3221.
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jallcom
Investigation of the characteristics of the nanocrystalline Ni3 Al-based alloy fabricated by hot pressing and sintering Masoud Nazarian-Samani ∗ , Ali Reza Kamali Advanced Materials Research Center, Department of Materials Science and Engineering, Malek Ashtar University of Technology (MUT), Ferdowsi St., Shahin-Shahr, Isfahan, Iran
a r t i c l e
i n f o
Article history: Received 13 July 2009 Accepted 22 November 2009 Available online 2 December 2009 Keywords: Ni3 Al-based alloy Hot pressing Sintering Mechanical alloying Grain growth
a b s t r a c t The characteristics of specimens produced from a Ni3 Al-based alloy using four different routes were investigated and different trends during consolidation of mechanically alloyed (MA-ed) powders were characterized and analyzed using microstructural observations; measurement of long-range order (LRO) parameter, micro-strain, and grain size; as well as determination of microhardness and density. In addition, the compressive yield strength and ductility of two series of specimens at different temperatures were determined. It was observed that hot pressing (HP) led to the best consolidation conditions. In addition, high pressure cold pressing (CP)/sintering caused the consolidated specimens to improve. In all specimens, the nature of nanocrystalline powders remained unchanged and the disordered MA-ed powders were ordered during consolidation. The yield strength in both series of specimens anomalously increased with increased temperature by up to 600 ◦ C, which later decreased beyond this critical temperature. © 2009 Elsevier B.V. All rights reserved.
1. Introduction Ni3 Al-based alloys have attracted a lot of attention due to their promising combination of properties such as high melting point, low density, high corrosion resistance, and good high-temperature strength. The most attractive property of these alloys is their enhanced yield strength with increasing temperature by up to 600–800 ◦ C [1–6]. The synthesis of nanocrystalline Ni3 Al by MA from elemental powders has been successfully achieved by several researchers [7–12] and subsequent consolidation has been accomplished by means of HP [7,8,12] and sintering processes [11]. While the microstructure of the specimens has been frequently discussed in previous studies, few systematic studies of the physical and mechanical properties of Ni3 Al and its alloys have been reported to date. In the present work, we report the results of a study that assesses the effect of HP and CP followed by sintering processes on the microstructure as well as the physical and mechanical properties of a MA-ed Ni3 Al-based alloy. The effects of pressure applied during CP and sintering time have also been investigated. 2. Experimental procedure The Ni3 Al-based alloy with the nominal composition Ni–8.14Al–7.83Cr– 1.45Mo–0.01B (wt%) was prepared by MA of elemental powders. All the powders
∗ Corresponding author. Tel.: +98 912 7085285; fax: +98 312 5228530. E-mail address: [email protected] (M. Nazarian-Samani). 0925-8388/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2009.11.182
used had an average particle size of 10 m. The powders were wet milled under argon in a Fritsch P6 with stainless steel balls for 15 h, using a ball to powder weight ratio (BPR) of 10:1. Details of the MA experiments have been presented elsewhere [12]. The 15 h MA-ed powder was subjected to four different processing routes, each simply termed specimens I through IV as listed in Table 1. CP and HP tests were performed by a uni-axial cold compaction machine and a HP machine (model ASTRO H-20-4560) in a graphite mold, respectively. All sintering processes were conducted under a high purity argon atmosphere. For metallographic and microhardness investigations, the specimens were carefully ground and polished. An etching solution, made of a mixture of glycerin, HCl and HNO3 with a ratio of 2:2:1, was applied to the surface of the polished samples for different times between 15 and 50 s. Vickers microhardness was measured in a Koopa UV1 machine. All specimens were inspected for their microstructure by a “Vega© Tescan” scanning electron microscopy (SEM) with an energy-dispersive X-ray spectroscopy (EDS). The microstructure of fracture surfaces was examined using a Hitachi-make S3400-N SEM. The X-ray diffraction (XRD) profiles of the specimens were carried out using Cu K␣ radiation ( = 0.15406 nm) and a Seifert 3003TT diffractometer operating at 40 kV and 30 mA. The grain size, d, and micro-strain, e, were calculated by applying the well-known Williamson–Hall method [12,13]. The lattice parameter was determined using at least five higher angle 2 peaks. From the XRD patterns, the LRO parameter, S, was determined by comparing the relative intensities of the superlattice (h k l mixed) and the fundamental (h k l are all even or all odd) peaks of all the specimens with respect to a well-annealed reference material [13]: S2 =
(Isup /Ifun )obs (Isup /Ifun )std
(1)
where (Isup /Ifun )obs and (Isup /Ifun )std are the intensities of the superlattice reflection relative to the fundamental line of the specimens (obs) and the fully annealed reference (std), respectively. In this study, the LRO parameter was estimated from the (1 0 0)/(2 0 0) and (1 1 0)/(2 2 0) pairs of reflections. Compressive tests were performed on cylindrical samples (D: 4 mm, L: 6 mm) at room temperature as well as temperatures 200, 400, 600, 700, and 800 ◦ C using an
M. Nazarian-Samani, A.R. Kamali / Journal of Alloys and Compounds 492 (2010) 196–200 Table 1 Different routes of consolidation experiments. Designation
Process
I II III IV
CP under 700 MPa + sintering at 1000 ◦ C for 1 h CP under 700 MPa + sintering at 1000 ◦ C for 2 h CP under 2 GPa + sintering at 1000 ◦ C for 1 h HP at 1000 ◦ C under 6 MPa for 10 min
Instron model 8503 tensile/compressive testing machine at a strain rate of 10−2 s−1 . All the data were reported as the average of at least three test results. The density of consolidated samples was determined using the immersion method in distilled water, based on the Archimedes Principle.
3. Results and discussion 3.1. XRD studies Fig. 1(a)–(e) shows the comparison of XRD patterns for the 15 h-milled powder [12] and the specimens I–IV. The results indicate that no superlattice reflection peaks were detected in the as-milled powder, demonstrating the formation of the fcc phase. The (1 0 0) superlattice peak appeared in specimen I after sintering for 1 h. Increasing the applied pressure during CP and the sintering time caused the amount of ordered structure to increase and the (1 1 0) and (2 1 0) superlattice peaks to appear. However, these were the only superlattice peaks that were revealed, and the consolidated samples may not, therefore, be fully ordered. Other researchers have reported the (2 1 1) superlattice peak to have also been revealed during consolidation [7,8]. Generally, the superlattice diffraction peaks are weak and their detection is not so easy in the Ni3 Al-based alloys. The addition of alloying elements will further affect their detectability. Although the addition of B is effective in stabilizing L12 structure [14], the fully ordered structure cannot be obtained here due to the presence of Cr and Mo. According to Fig. 1, specimen IV exhibits a greater atomic ordering than specimens I–III. This can be explained by the homogenization effect of HP on the composition, in addition to the annealing of the defects. From Table 2, it can be seen that the LRO parameter calculated from the (1 1 0)/(2 2 0) pair is slightly greater than that of the (1 0 0)/(2 0 0) pair in all specimens, which may be because of their texture, extinction, or other reasons [15]. Fig. 1 also reveals that the peak intensities of all the specimens increased and that the peak widths declined during the consolidation processes. Heating the milled Ni3 Al also resulted in a volume expansion as shown by the increased lattice parameter in all the specimens. It is shown that the sintering time causes the increase of the lattice parameter due to the greater dissolution of Cr, Mo, and Al in
Fig. 1. XRD patterns of as-milled powder [12] (a), and specimens I (b), II (c), III (d), and IV (e).
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the Ni3 Al structure. The atomic radius of Al (0.143 nm) and Mo (0.136 nm) is larger than that of Ni (0.125 nm), whereas the atomic radius of Cr (0.125 nm) is equal to that of Ni [16]. Therefore, diffusion of Al and Mo atoms into the Ni lattice leads to an increase in the Ni lattice parameter and the transfer of X-ray peaks of Ni3 Al to lower angles. This was more pronouncedly observed in specimen IV than in other specimens, perhaps due to the applied pressure during HP process. The above explanations are more vividly presented in Table 2. Here, it is also shown that micro-strain decreased during the consolidation process. As expected, the sintering time led to greater annealing of the structure and to the decreased microstrain. The lowest micro-strain was observed in specimen II due to its long time sintering. Also, the grain size slightly increased in all specimens, with its highest value observed in specimen II. It can be seen that the grain size was still small in all specimens and retained its nanocrystalline size after all consolidation routes. The presence of Mo, Cr, and B dissolved in the Ni3 Al phase led to the production of a nanocrystalline specimen even after a high-temperature consolidation at 1000 ◦ C. Generally, the alloying elements segregate in grain boundaries and, therefore, reduce the grain-boundary mobility and surface free energy of the material. Moreover, grain growth in nanocrystalline powders is prevented due to the presence of open porosities in the microstructure. When the porosities are closed and isolated, grains grow without any dike [17]. This is why the grains in specimens I and II did not grow very much and also why grains in specimen III had their maximum size. Moreover, the grains in specimen IV are smaller than those in specimen III; this is due to simultaneous application of pressure and heat. Application of pressure hinders grain growth at elevated temperatures. This is due to the fact that diffusion is involved in grain growth and that the diffusion coefficient decreases with increasing pressure [18]. It can thus be concluded that pressure may reduce grain-boundary mobility. 3.2. Physical and mechanical properties The phenomena involved in cold compression of powders include rearrangement of particles, elastic deformation of particles at their contacts, plastic deformation in metals, and particles crashing in brittle materials. These phenomena are present in the compression of nanocrystalline powders. However, the behavior of very fine particles, i.e., slip, friction, and their diffusion-assisted deformation, can be considerably different from that of larger particles. Slip and rearrangement mechanisms are limited in nanocrystalline powders due to the presence of high friction forces between particles. Surface tension and van der waals as well as electrostatic and mechanical forces are the origins of this friction. Moreover, irregular grain boundaries facilitate the agglomeration of the particles. These facts prevent the rearrangement of particles. Therefore, the possibility for nanocrystalline powders to achieve a high density green compact during compression is lower than that for micron-sized particles [17]. Moreover, nanocrystalline powders have a high potential for agglomeration during MA [13]. Presence of agglomerated powders increases the possibility for large porosities to form in the material during consolidation. Extensive agglomeration of 15 h-milled powders can be seen in Fig. 2. At the atomic level, sintering induces contacting particles to bond together. The most important atomic events are surface diffusion (at lower temperatures) and grain-boundary diffusion (at higher temperatures). Grain-boundary diffusion predominantly occurs at the contacts between particles. On the microstructural scale, the bonding becomes evident as necks grow between touching particles. Such neck growth increases the strength over the green strength and causes many beneficial property changes [19]. However, the HP of powders has certain major benefits compared to CP/sintering. The pressure used in the HP process leads to plastic
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Table 2 Characteristics of the as-milled powder [12] and the specimens I–IV. Specimen
Microhardness (VHN)
(g/cm3 )
a (Å)
e (%)
d (nm)
As-milled I II III IV
– 216 263 387 479
– 4.28 4.65 6.73 7.79
3.60 3.62 3.69 3.64 3.76
0.59 0.28 0.19 0.31 0.36
10.66 38.34 46.93 34.62 28.11
deformation of powders and, therefore, to their interpenetration and decreased porosity of compacts during the consolidation process. Moreover, applied pressure leads to the rupture of oxide layers on the surface of the powders and, thereby, to their more effective connection. In pressure-assisted sintering of agglomerated powders, large porosities can be removed due to the effects cited above which lead to minimized grain growth. It should be noted that applying the pressure at high temperatures leads to decreased resistance of the material against deformation and to increased destruction of porosities. On the other hand, destruction of large porosities at low temperatures is so difficult even when high pressures are applied. Moreover, shear tensor of the applied pressure leads to the rearrangement of the particles and destruction of large porosities. More connection points, therefore, emerge among the particles. The oxide layer formed on the surface of the particles fractures due to the applied shear stress creating a better particle connection. It can be concluded that using pressure during the sintering process leads to the removal of large porosities and to the prevention of grain growth. The smallest grain size in specimen IV can also be explained along the same lines. The relationship between yield strength and grain size in materials can be evaluated using the Hall–Petch equation [17,20]: ∝ d−(1/2)
(2)
where and d are yield stress and grain size, respectively. According to this equation, plastic deformation of particles is more difficult in smaller particles. The relatively low density obtained in specimen III (Table 2) can be explained by considering the yield strength of the alloy at room temperature (600 MPa) and the pressure used in
Fig. 2. SEM micrograph of as-milled powder indicates agglomeration during MA.
LRO (1 0 0)/(2 0 0)
(1 1 0)/(2 2 0)
0 0 0.08 0.22 0.41
0 0.06 0.11 0.32 0.45
the CP process (2 GPa). The lack of a main plastic deformation in powder particles during consolidation is responsible for this low density in the consolidated specimen. Increasing the sintering time from 1 to 2 h led to an increase in density from 4.28 to 4.65 g/cm3 . However, prolonged sintering time beyond 2 h had no effect on the improvement of density. On the other hand, the CP of powders under a high pressure of 2 GPa and the sintering of the green compacts led to the production of consolidated specimens with a considerable density of 6.73 g/cm3 . It can be seen in Table 2 that application of a pressure as low as 6 MPa during the HP process yields the highest density equal to 7.79 g/cm3 . The greater ductility of the particles at a temperature as high as 1000 ◦ C as well as the enhanced diffusion of atoms under the applied stress may be responsible for the considerable consolidation of HP-ed specimen IV. It should be noted that a nanocrystalline material has a lower density than the same material with a normal grain size. This is due to the high amount of grain boundaries in nanocrystalline materials. Grain boundaries have a lower density than grains do. According to Table 2, microhardness increases in the specimens with increased sintering time, which is related to increased density during sintering, the highest value for microhardness belonging to specimen IV. On the other hand, the HP process obtains better conditions for the production of a bulk sample. It was observed that samples sintered for 1 and 2 h did not have sufficient handling strength. Thus, compressive tests were not further continued for measuring yield strength and ductility, but these tests were only performed for specimens III and IV. The yield strengths of both specimens show a positive temperature dependence in the range from room temperature to 600 ◦ C where the strength peak was reached. Further increase in temperature caused a corresponding decrease in strength (Fig. 3). Specimen IV showed yield strengths almost twice as much as that of specimen III for all temperatures applied. This anomalous behavior is common to Ni3 Al and its alloys, as also pointed out by many researchers [1–6]. HP-ed specimens showed
Fig. 3. Comparison of compressive yield strength, y , and ductility, ε, of specimens III (-♦- and --) and IV (-- and -×-) at different temperatures.
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Fig. 4. SEM micrographs of specimens II (a), III (b), and IV (c).
the highest yield strength as compared to some literature data [21]. These results represent the high effectiveness of the HP process used compared to other processes. Fig. 3 also presents the variations in ductility of both specimens III and IV, indicating larger ductility and yield strength in specimen IV. This increased yield strength and improved ductility were mainly caused by the finer grain size and lower porosities and pores in specimen IV compared to those in specimen III. It should be noted that the presence of porosities and the non-effective connection between particles lead to considerably reduced ductility of specimens. Moreover, it is observed that compressive yield strength as a function of the consolidation routes used follows trends quite similar to those of density and microhardness in both specimens. It can, therefore, be concluded that yield strength was significantly dependent on the density and microhardness of the consolidated specimens. 3.3. Microstructure The microstructures of specimens II–IV are shown in Fig. 4(a)–(c). It is observed that the CP-ed powders at a pressure of 700 MPa could not be fully consolidated after 1 and 2 h of sintering, but that powders CP-ed under a high pressure (2 GPa) in specimen III led to the formation of a relatively dense sample. The grain boundaries did not appear in specimen IV even after
etching over long periods (50 s). However, the grain boundaries in specimen III appeared only after a light etching (15–20 s). On the other hand, according to Fig. 4(c), the surface is smooth and no pore exists between the bonded powder particles, which are evidence of the good consolidation quality. Table 3 shows the results of EDS analysis for as-milled powder and the specimens I–IV produced via various consolidation routes. The results indicate that the bulk samples were identical to the as-milled powder in composition, and that Al, Cr, and Mo still dissolved in Ni3 Al after sintering at 1000 ◦ C. This is evidenced by the fact that no peaks of these elements and/or their compounds were observed in the XRD patterns (Fig. 1). This suggests that the Ni3 Albased alloy has excellent thermal stability at high temperatures.
Table 3 EDX analysis of the as-milled powder [12] and specimens I–IV. Specimen
As-milled I II III IV a
B was not detected.
Element (wt%)a Ni
Al
Cr
Mo
82.26 82.09 81.97 81.92 81.96
8.09 8.11 8.13 8.13 8.19
7.98 7.99 8.04 8.07 8.14
1.67 1.81 1.86 1.88 1.71
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2. Application of pressure during consolidation hinders grain growth at elevated temperatures. Also, the pressure helps to produce a free-pore or void-bulk sample. 3. Both specimens III and IV exhibit a positive temperature dependence of their compression yield strength; maximum yield strength obtained at 600 ◦ C followed by a decrease at higher temperatures. 4. The nature of the nanocrystalline structure remains intact after consolidation. In addition, the ordering transition occurs during consolidation, which is greater in specimen IV than in other specimens. 5. The removal of MA-ed powder agglomeration should be accomplished before the sintering in order to obtain a specimen with the minimum porosity. References Fig. 5. Fracture surface of specimen III indicating crack propagation.
From Table 3, it is also concluded that the specimens were not obviously contaminated by oxygen and/or other elements during different consolidation routes used in this study. Another role of grain boundaries in specimen IV when compared with specimen III is the faster diffusion shortcuts existing in the grain boundaries. These fast shortcuts make a few initial tiny cracks to quickly close during the deformation process. Therefore, the crack growth and propagation can be avoided to a certain extent. Nevertheless, the high pressure CP-sintering technique unavoidably induced grain growth, which is harmful to their properties, especially their yield strength and ductility. It can be seen from Fig. 5 that cracks easily propagated during the compression test in specimen III, but no crack was observed from fracture surfaces of specimen IV (not shown). Based on the results in this investigation, it can be seen that MA-HP process is an effective method for preparing Ni3 Al-based alloys with improved properties. 4. Conclusions The following conclusions may be drawn from the results obtained in this study: 1. The HP process produces better specimens in terms of density, microhardness, grain size, microstructure, yield strength, and ductility than other processes investigated in the present study.
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