Journal of Non-Crystalline Solids 453 (2016) 1–7
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Effect of sintering temperature on phase transformation during consolidation of mechanically alloyed Al86Ni6Y6Co2 amorphous powders by spark plasma sintering Ram S. Maurya, Ashutosh Sahu, Tapas Laha ⁎ Department of Metallurgical & Materials Engineering, Indian Institute of Technology Kharagpur, Kharagpur 721302, India
a r t i c l e
i n f o
Article history: Received 29 July 2016 Received in revised form 7 September 2016 Accepted 21 September 2016 Available online xxxx Keywords: Al based amorphous alloy Mechanical alloying Spark plasma sintering Phase transformation Nanocrystalline phase Microhardness
a b s t r a c t Mechanically alloyed Al86Ni6Y6Co2 amorphous powders (170 h) were consolidated via spark plasma sintering in the temperature range of 300–500 °C. Various phase transformations caused during mechanical alloying of the Al86Ni6Y6Co2 powders and the effect of sintering temperature on the consolidation and devitriﬁcation behavior during the sintering of the amorphous powders were investigated. XRD and TEM study conﬁrmed a decrease in amorphous phase content with increased amount of nanocrystalline FCC-Al and nano-sized intermetallic precipitates as the sintering temperature was increased. High current density at irregular particle contact surfaces during spark plasma sintering resulted in intense Joule heating with rapid temperature rise, which caused localized devitriﬁcation along with improved consolidation. Amorphous alloy sintered at 500 °C yielded higher hardness (298 ± 9 Hv) in comparison to the lower temperature sintered alloys, attributed to better densiﬁcation and uniformly distributed nanocrystalline phases in the amorphous matrix. © 2016 Elsevier B.V. All rights reserved.
1. Introduction Bulk metallic glasses (BMGs) have generated signiﬁcant interest, owing to their outstanding mechanical (tensile and fracture strength, and wear resistance), chemical (corrosion resistance) and physical (soft magnetic) properties in comparison to their crystalline counterparts, attributed to the lack of microstructural defects (viz. dislocation, grain boundaries and anti-phase boundaries) in metallic glasses [1,2]. Considering properties and performance of various BMGs, Al based ones possess comparatively higher speciﬁc strength and better corrosion resistance [3–6]. Various experimental synthesis work (both by rapid quenching [3–4] and solid state processing [5–8]) supported with efﬁcient cluster packing (ECP) model based computational simulation manifested Al (≥85 at.%) – TM (Ni, Co) – RE (Y, La, Ce) systems as good glass formers and thermodynamically stable, attributed to the large atomic size differences and negative heat of mixing of the alloying elements [9–10]. However, Al BMG synthesis by rapid solidiﬁcation requires extremely high cooling rate (105–106 K/s) to suppress crystallization. On the other hand, solid state route, viz. mechanical alloying (MA) offers advantages of (i) broad glass forming composition range, (ii) no restriction of the phase diagram, (iii) some relaxation to the empirical rules viz. negative heat of mixing and atomic size criteria and (iv) capability of synthesizing large amount of amorphous powders . ⁎ Corresponding author. E-mail address: [email protected]
http://dx.doi.org/10.1016/j.jnoncrysol.2016.09.018 0022-3093/© 2016 Elsevier B.V. All rights reserved.
Mechanically synthesized amorphous powders could effectively be consolidated by spark plasma sintering (SPS), which has already demonstrated its potential in sintering metallic amorphous powders [5–7, 12]. The SPS process, which is based on the general principle of electric ﬁeld assisted consolidation, where pulsed electrical discharge combined with rapid Joule heating and a high uniaxial pressure application limits the crystallization events and assists in retaining maximal amorphous phase along with promoting formation of various nanocrystalline phases which could further strengthen the base alloy [7,13–14]. There have been several efforts in synthesizing Al based amorphous powders via gas atomization followed by consolidation via SPS [5–7,13,15,16]. Kim et al. mechanically amorphized Al85Y8Ni5Co2 powders and sintered the powders by SPS, and reported incomplete sintering of the amorphous powders, attributed to the high viscosity of amorphous phase under the given sintering temperature of 250 °C . Li et al. reported abnormal crystallization of nanometric Al5Co2 phase  and formation of anomalously large Al crystals  during SPS of gas atomized Al86Ni6Y4.5Co2La1.5 amorphous powders. Recently, SPS of gas atomized partially amorphous Al86Ni7Y4.5La1.5Co1 powders yielded high strength alloy attributed to the precipitation of FCC-Al, Al4NiY and La7Ni3 intermetallic compounds . Thus, consolidation of amorphous powders via SPS can result in dispersion strengthening of the amorphous matrix which could enhance the fracture toughness by delaying any failure by the virtue of restricted shear band propagation . There are several studies carried out to understand the effect of SPS temperature on the consolidation behavior of Al based [6,18], Cu based
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 and Zr based  amorphous powders. Sasakai et al.  and Kim et al.  reported that an increase in the fractions of hard nanocrystals in Al- and Cu-based amorphous alloys with increasing sintering temperature was the reason behind improved hardness and relative density. Shim et al. reported a dramatic increase in the compressive strength and corrosion resistance of Zr65Al10Ni10Cu15 glassy alloys with increasing sintering temperature under the condition that the temperature was kept below the glass transition temperature to retain maximum amorphous phase fraction . Recently, Deng et al. showed that an alloy produced by sintering of gas atomized partially Al86Ni7Y4.5La1.5Co1 amorphous powders at 250 °C to 350 °C exhibited an increase in the fracture strength, whereas further increases in temperature to 400 °C reduced the fracture strength due to increase in the fraction of soft FCC-Al and La7Ni3 intermetallic compounds . The goal of the present work was to (i) investigate the phase transformation during mechanical alloying of Al86Ni6Y6Co2 powders and (ii) explore the effect of sintering temperature on the consolidation behavior and devitriﬁcation during spark plasma sintering of the Al86Ni6Y6Co2 amorphous powders.
particles (1–2 μm) as displayed in Fig. 1c (100 h) and d (150 h). During the mechanical alloying, the powder particles went through a process of repeated deformation, cold welding, work hardening, fracturing and eventually size reduction attributed to the continuous impact energy transfer phenomena . It is well known that Al based powders are ductile in nature and have a high tendency towards cold welding in the initial stage of milling. With the increase in milling time, fracture tendency dominated over cold welding (assisted by the presence of the PCAs), which led to formation of much smaller size particles (1–2 μm) after 100 h of milling. With further progress in milling (150 h), the tendency of the smaller particles towards welding and the bigger particles towards fracturing kept the particle size within a constant range . However, it can be observed that the 150 h milled powders (Fig. 1d) were more distorted with lesser interlayer spacing indicating larger degree of deformation in these powders.
3.2. Phase transformation in mechanically alloyed powders 2. Materials and methods Al86Ni6Y6Co2 alloy powder mixture was prepared by blending elemental aluminum (99.5%, −44 μm), nickel (99.996%, −125 μm), yttrium (99.9%, −420 μm) and cobalt (99.5%, −44 μm) powders procured from Alfa Aesar, MA, USA. Mechanical alloying was performed in a planetary ball mill (PM 200, Retsch GmbH, Germany), taking the powder mixture in hardened steel vials (capacity of 125 ml) along with hardened steel balls (diameter of 10 mm). Ball to powder weight (BPR) ratio and disc RPM were kept as 15:1 and 300, respectively. Automatically controlled rotating disc in the planetary mill was given 15 min break after every 30 min of alternate clockwise and anticlockwise rotations. Toluene was used as process controlling agent (PCA) to limit the agglomeration and the contamination tendency, whereas, a small amount of stearic acid (0.08 wt%) was added to restrict cold welding between the powders. Powder handling was carried out in a glovebox (O2 and H2O b 5 ppm) to avoid air contamination. The mechanically synthesized amorphous powders, milled for 170 h were consolidated by spark plasma sintering (SPS 625, Fuji Electronic Industrial Co. Ltd., Japan) using a 10 mm diameter tungsten carbide die-punch set. A high sintering pressure of 500 MPa and lesser holding time of 15 min was applied to suppress any crystallization. The sintering temperature was varied from 300 °C to 500 °C at a heating rate of 100 °C/min. Characterization techniques employed on the milled powders and consolidated bulk alloys were: (i) X ray diffraction (Bruker D8 Advance diffractometer, Germany) using Cu Kα (λ = 1.54 Å) radiation for understanding the phase transformation/evolution, (ii) differential scanning calorimetry (DSC Q20, TA Instruments, USA) at a heating rate of 20 °C/ min to investigate the thermal transition behavior, (iii) scanning electron microscopy (SUPRA 40, Carl Zeiss AG, Germany) to reveal microstructure and morphology, (iv) transmission electron microscopy (JEM-2100 LaB6, 200 kV, JEOL USA, Inc.) to observe the transformation of amorphous phase into various crystalline phases and (v) Vickers microhardness test (UHL VMHT - 001, Walter Uhl, Germany) at 500 gf for 20 s to ﬁnd out the variation in microhardness. 3. Results and discussion 3.1. Morphology of mechanically alloyed powders Morphology of 1 h, 50 h, 100 h, and 150 h milled powders is shown in the SEM images in Fig. 1. The powder particles in the form of irregular, conically elongated and spherical particles of various sizes (1–18 μm) could be observed in the initial stage (1 h) of milling as shown in Fig. 1a. The powders encountered severe plastic deformation with increase in the milling time, which decreased the particle size (1–8 μm) after 50 h of milling (Fig. 1b). Continued milling resulted in much smaller
Fig. 2 shows the XRD patterns of Al86Ni6Y6Co2 powders milled up to 1 h, 50 h, 100 h, 150 h and 170 h. As the milling time was increased, the intensity of the XRD peaks were reduced and the peaks became broadened and diffused due to increase in lattice microstrain, lattice parameter change, and crystal defects such as vacancies, dislocations and subgrains. Various defects generated during milling increases the strain energy of the powders and promoted the solid state diffusion reaction [17,22–23]. This led to evolution of some new peaks at 2θ = 29.65°, 32.15°, 34.54° and 50° in the 50 h milled powders as shown by arrows in the corresponding XRD spectrum. The evolution of these peaks indicated the formation of Al-rich complex phases attributed to the solid state diffusion during milling [17,22]. However, the XRD peaks became broader in case of 100 h milled powders indicating a higher degree of microstructural breakdown followed by nanocrystallization and amorphization, attributed to the intensiﬁed inter-diffusion of elements [22,24]. Earlier, Kim et al. also reported formation of amorphous phase along with various Al rich complex phases in 50 h and 100 h milled Al85Y8Ni5Co2 powders . It is a common phenomenon to observe various phase formation and their disappearance during the process of milling of multicomponent powders. Thus, with the progressive milling, new peaks which were observed in 50 h milled powders disappeared in 100 h milled powders showing decrease in amount of Al-rich complex phases as show in the XRD pattern (Fig. 2). Fig. 3 displays the TEM images and SAD patterns of 100 h and 170 h milled Al86Ni6Y6Co2 powders. The TEM image (Fig. 3a) of the 100 h milled powders shows nanocrystalline phases (marked by arrows) distributed in the amorphous matrix. Corresponding SAD pattern, displayed in the inset of Fig. 3a, exhibits diffused background with hazy rings and various bright spots, evidencing presence of amorphous and nanocrystalline phases (viz. FCC-Al and Al rich intermetallics), respectively in the 100 h milled powders. High resolution image (Fig. 3b) of these powders clearly revealed the presence of crystalline (periodic lattice arrangement) and amorphous phases (distorted atomic structure). The powders exhibited atomistically featureless microstructure after 170 h of milling (Fig. 3c), indicating nearly fully amorphous structure. This was also corroborated by the corresponding fully diffused SAD pattern shown in the inset of Fig. 3c. The HRTEM image of the 170 h milled Al86Ni6Y6Co2 powders (Fig. 3d) manifested distorted lattice structure endorsing fully amorphous phase in these powders. Thus, the alloy powders showed decrease in degree of crystallinity with increasing milling time and nearly fully amorphous structure was obtained after 170 h of milling. A similar trend was earlier observed by Zheng et al., who reported a large numbers of intermetallic precipitate formation in the amorphous matrix for Al87Ni8.5Ce3Fe1Cu0.5 alloy powders, and the amount of intermetallic phases decreased with increasing milling time (40 h to 170 h) .
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Fig. 1. SEM micrographs showing variation in the size and morphology of (a) 1 h, (b) 50 h, (c) 100 h and (d) 150 h milled Al86NI6Y6Co2 powders.
3.3. Thermal behavior of the mechanically alloyed powders The thermal behavior of the various hours milled powders was studied to understand the devitriﬁcation behavior of the milled powders. The thermal transformation behavior of 50 h, 100 h, 150 h and 170 h milled Al86Ni6Y6Co2 powders were investigated by differential scanning calorimetry (DSC) and the thermograms are presented in Fig. 4. The DSC curve of the 50 h milled powders exhibited a broad exothermic peak in the vicinity of 530 °C indicating the growth of the nanocrystalline FCCAl grains, which were present in the milled powders. The thermogram also exhibited two very small peaks at 263 °C and 347 °C (indicated by arrows in the inset ﬁgure), which were associated with the structural relaxation and defect annihilation created during ball milling. Thermal transition of the 100 h milled powders was almost similar to that of the 50 h milled powders with some differences. The
Fig. 2. XRD patterns showing the phase transformation and progress in amorphization in the Al86Ni6Y6Co2 powders with increasing milling time up to 170 h.
transformation peak related to the growth of nanocrystalline FCC-Al was noticed at a little earlier (at 515 °C). Furthermore, the two diminutive peaks at 193 °C and 249 °C, respectively related to the structural relaxation and defect annihilation, appeared earlier attributed to the higher defect density in the longer time milled powders. The difference in transition behavior of 50 h and 100 h milled powders could also be anticipated by comparing the corresponding XRD patterns shown earlier in Fig. 2. The XRD peaks of the 100 milled powders were more broadened and of lesser intensity than were observed for the 50 h milled powders. The 100 h milled powders exhibited more defects and contained higher amount of amorphous structure in comparison to those in the 50 h milled powders. Thus, the 100 h milled powder exhibited the structural relaxation events at comparatively lower temperatures. Recently, Zheng et al.  reported similar nature of DSC transition behavior in Al87Ni8.5Ce3Fe1Cu0.5 milled powders. The DSC curves of the 150 h and 170 h milled powders showed almost similar behavior with three exothermic events of crystallization. In case of 150 h milled powders, the ﬁrst crystallization exothermic peak at 191 °C (Tx1) was associated with the formation of nanocrystalline Al phase from the parent amorphous phase. The second exothermic peak (Tx2) at 380 °C corresponded to the devitriﬁcation event in the residual amorphous phase leading to the formation of various Al-rich intermetallic phases and growth of nanocrystalline Al phase formed during the ﬁrst transition at Tx1. The third exothermic event in the vicinity of 593 °C indicated complete transition of any residual amorphous phase into precipitation of various Al rich phases (viz. FCC-Al, Al2Y, Al3Ni5, Al3Ni2, Al9Co2, AlCo and AlNi) and substantial growth of nanocrystalline Al. It has been reported by various researchers that the phase transformations during DSC of multi-component metallic amorphous alloys occur through the sequence of nanocrystallization of base element from the amorphous phase to the growth of the nanocrystalline grains and formation of new intermetallic phases [23,25–27]. In the present study also, the same sequential phase transformation was observed during the DSC experiment of fully amorphous Al85Y6Ni6Co2 powders (milled for 150 h and 170 h). The various exothermic peak at Tx1, Tx2 and Tx3 were associated with (i) the formation of nanocrystalline Al
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Fig. 3. (a, c) TEM images with respective SAD patterns and (b, d) corresponding high resolution images of the 100 and 170 h milled Al86Ni6Y6Co2 powders exhibiting partial and complete amorphization, respectively.
phase from the parent amorphous phase, (ii) the devitriﬁcation event in the residual amorphous phase leading to the formation of various Alrich intermetallic phases and growth of nanocrystalline Al phase formed during the ﬁrst transition and (iii) the complete transition of any residual amorphous phase into precipitation of various Al rich phases. Borner et al.  reported similar type of phase formation (viz. formation of fcc-Al, Al3Y, Al3Ni5, Al9Co2, AlCo, and Al3Ni2) during the DSC
analysis of Al85Y8Ni5Co2 amorphous powders. Various other authors also performed DSC experiment of Al-based amorphous powders and reported similar transformation behavior [23,26–27]. In the present study, the nature of the DSC curve of the 170 h milled powder was similar to that of 150 h milled powders. However, the ﬁrst (Tx1: 175 °C), second (Tx2: 375 °C) and third (Tx3: 585 °C) transition peaks shifted a little to lower temperature values, which could be attributed to the more amount of defects and amorphous phase in the 170 h milled powders. The phase transformation during DSC of the amorphous powders can also be related to the XRD and TEM results of the spark plasma sintered (300–500 °C) bulk alloys, which has been discussed in the following section. 3.4. Phase evolution in sintered bulk amorphous alloys Fig. 5 shows XRD patterns of bulk alloys produced by consolidating amorphous powders in the temperature range of 300–500 °C. Diminishing XRD hump and sharpening of peaks with increasing sintering temperature indicated lesser fraction of retained amorphous phase and higher degree of crystallinity in the higher temperature sintered alloys. Phase fractions in various temperature sintered samples were estimated by calculating the areas of XRD crystalline peaks and amorphous hump, and ﬁnding their ratio using the following equation  vAmor ¼
Fig. 4. DSC thermograms of the 50 h, 100 h, 150 h and 170 h milled Al86Ni6Y6Co2 powders exhibiting various phase transitions.
AAmor ; ACryst þ AAmor
where vAmor is the volume fraction of amorphous phase, and AAmor and ACryst are the area of amorphous hump and the crystalline peaks,
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Fig. 5. XRD patterns of the Al86Ni6Y6Co2 amorphous alloys, spark plasma sintered at various temperatures, showing the evolution of multiple crystalline phases overlaying the broad amorphous humps.
respectively. The estimated amorphous phase fractions were 76%, 65%, 57% and 46% in 300 °C, 350 °C, 400 °C and 500 °C sintered alloys, respectively (Table 1). Increase in the crystalline phase fraction could also be envisaged from the appearance of few higher angle XRD peaks in the 400 °C-(2θ ~ 69.74°, 78.27°) sintered alloy, which became prominent in 500 °C-sintered alloy (2θ ~ 69.29°, 77.05°). These higher angle crystalline peaks were not observed in 300 °C- and 350 °C-sintered amorphous alloys. Multiple crystalline peaks overlaying the broad XRD hump in different temperature sintered alloys correspond to the evolution of various crystalline phases viz. FCC-Al, Al4Ni3, Al0.9Ni1.1, Al3Y, Al2Y Al5Co2, Al4.85Co5.15 and Al13Co4. TEM study (Fig. 6) was performed on the 500 °C-sintered alloy to conﬁrm the presence of various crystalline phases in the amorphous matrix. The micrograph shown in Fig. 6a exhibits the distribution of various nanocrystalline phases (size 5–80 nm) in the amorphous matrix. This observation is supported by corresponding diffused SAD pattern (Fig. 6b), exhibiting dots and discrete rings. High resolution images (Fig. 6c and d) clearly show that the nanocrystalline phases are well adhered to the amorphous matrix. Corresponding fast Fourier transformation (FFT) patterns shown in the inset of Fig. 6c and d, exhibiting diffused background with bright dots also conﬁrm the presence of nanocrystalline phase in the amorphous matrix. Spark plasma sintering results a high current density at the particle contact surface and intense Joule heating during sintering; and thus rapid temperature rise at the particle contacts causing surface melting and facilitating various diffusion processes (viz. line, lattice and surface diffusion) [28,29]. The presence of high defect density in mechanically alloyed powders  decreases the activation energy of mass transfer, which also assists in new phase formation and devitriﬁcation . These phenomena simultaneously played decisive role in the formation of various nanocrystalline phases during sintering of various temperature-sintered alloys in the present study as conﬁrmed by the XRD patterns (Fig. 5) and TEM images (Fig. 6). Recently, Deng et al.  studied Table 1 The processing parameters, physical and mechanical properties of the Al86Ni6Y6Co2 amorphous alloys consolidated via spark plasma sintering. Sintering temperature
300 350 400 500
°C °C °C °C
Sintering Microhardness Relative Amorphous time (Hv) density (%) volume (%) 15 min
148 172 265 298
± 17 ± 16 ± 12 ±9
73 77 84 92
76 65 57 46
Fig. 6. (a) TEM image of the spark plasma sintered Al86Ni6Y6Co2 amorphous alloy, showing the distributed nanocrystalline precipitates in the amorphous matrix, (b) corresponding SAD pattern exhibiting dotted rings with diffused background indicating distribution of nanocrystalline phase in the amorphous matrix, and (c, d) high resolution images of the precipitates shown in (a) and their corresponding fast Fourier transformation image (inset) respectively, conﬁrming the presence of nanocrystalline intermetallic precipitates.
phase transformation in Al86Ni7Y4.5La1.5Co1 nanocrystalline alloys synthesized via spark plasma sintering of gas atomized amorphous powders. They reported that the alloys consisted of multiphase microstructure (viz. FCC-Al, Al4NiY and La7Ni3 intermetallic compounds) attributed to the increase in local temperature at the powder surface leading to crystallization. In another work, Qian et al.  also detected Al5Co2 nanometric phase in gas atomized and spark plasma sintered Al86Ni6Y4.5Co2La1.5 amorphous alloy.
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It should also be mentioned here that in the present work, sintering was performed at high pressure (500 MPa) and high pressure sintering assist in retaining higher amount of amorphous phase with formation of minor amounts of nanocrystalline phases [7,16]. It has also been reported in the literature that high pressure could contribute towards major change in atom spacing, chemical bonding and Gibbs free energy of phase transformation, which collectively play decisive role in governing the devitriﬁcation phenomena in glassy alloys [12,32].
3.5. Effect of morphology and phase transformation on the microhardness of the sintered alloys The morphology of the different temperature sintered amorphous alloys is shown in Fig. 7. The samples sintered at 300 °C (Fig. 7a) and 350 °C (Fig. 7b) showed angular and sub-angular morphology of powder particles with insufﬁcient bonding and clear inter-particle boundaries indicating poor sintering attributed to the high viscosity of the amorphous phase at such low temperature of sintering . When the sintering temperature was increased to 400 °C, improvement in interparticle bonding was observed (Fig. 7c), although a high volume fraction of micro-porosities (16 vol%) were also present. The volume percent of the micro-pores decreased to 8% in the 500 °C-sintered alloy and a very good improvement in the inter-particle bonding could also be seen (Fig. 7d). Higher temperature favors faster mass transfer diffusion kinetics, increases atom mobility and decreases viscosity of amorphous phase leading to improvement in the inter-particle bonding and relative density [7,25]. Estimated relative density via Archimedes' principle was found to be 73%, 77%, 84%, 92% for 300 °C, 350 °C, 400 °C and 500 °C sintered alloys, respectively.
Vickers microhardness test was conducted on the sintered alloys and the values are presented in Table 1, along with sintering parameters and relative density. Considerable improvement in the microhardness was noticed with an increase in sintering temperature. Amorphous alloy sintered at 300 °C (148 ± 17 Hv) and 350 °C (172 ± 16 Hv) exhibited lower values of hardness attributed to incomplete sintering. Improvement in microhardness (265 ± 12 Hv) was observed for the amorphous alloy sintered at 400 °C and the alloy sintered at 500 °C exhibited the highest value of microhardness (298 ± 9 Hv). The improvement in hardness in this sample is attributed to (i) the improvement in relative density and inter-particle bonding, (ii) retention of the high fraction of amorphous phase, (iii) distribution of the hard nanocrystalline phases and (iv) good bonding between the amorphous matrix and the nanocrystalline phases. Similarly, Deng et al.  reported an increase in microhardness from 193 Hv to 259 Hv of Al86Ni6Y4.5Co2La1.5 amorphous alloy, when the sintering temperature was increased from 300 °C to 350 °C during spark plasma sintering of gas atomized amorphous powders, attributed to the retention of high fraction of amorphous phase along with improved densiﬁcation. In another study, Sasaki et al. also reported an increase in microhardness of spark plasma sintered Al85Ni10La5 alloy with increase in sintering temperature from 200 °C to 450 °C . In the present study, the sintering kinetics in the amorphous alloy played signiﬁcant role in controlling the phase transition and consequently the densiﬁcation behavior, and affected the microhardness of the sintered alloys. The effect of spark plasma sintering temperature and pressure on the structural phase transformation, densiﬁcation and consequent effect on mechanical properties of glassy alloys have been studied by various researchers [6,18–20,33] including present authors [7,16]. It could be envisaged from the present study that higher temperature sintering favored faster devitriﬁcation,
Fig. 7. SEM images of Al86NI6Y6Co2 amorphous alloys consolidated at (a) 300 °C, (b) 350 °C, (c) 400 °C and (d) 500 °C, showing improvement in inter-particle bonding with increasing sintering temperature.
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whereas high pressure (500 MPa) sintering promoted the nanocrystalline phase formation by favoring the short range atomic ordering and suppressing the long range one. 4. Conclusions Microstructural and phase evolution study was performed on mechanical alloyed (up to 170 h) Al86Ni6Y6Co2 amorphous powders and spark plasma sintered (300–500 °C) bulk alloys. Higher sintering temperature favored formation of various nanocrystalline intermetallic phases viz. Al4Ni3, Al0.9Ni1.1, Al3Y, Al2Y Al5Co2, Al4.85Co5.15, Al13Co4 at the expense of the parent amorphous phase. A decrease in amorphous phase retention from 76 vol% to 46 vol% was observed from XRD analysis carried out on the sintered alloys when the sintering temperature was increased from 300 °C to 500 °C. Applied high pressure (500 MPa) during sintering assisted in suppressing the long-range diffusion and promoted minor amount of nanocrystalline phase formation by short-range atomic rearrangement. The high current density at irregular particle contact surfaces during spark plasma sintering resulted in intense Joule heating in the vicinity of inter-particle locale causing high localized temperature generation at the particle contacts leading to rapid mass transfer diffusion promoting various phase evolution processes. The bulk alloy sintered at 500 °C exhibited highest hardness of 298 ± 9 Hv, owing to improved inter-particle bonding and uniformly distributed nanocrystalline precipitates in the amorphous matrix. Acknowledgement This work was supported by the Science & Engineering Research Board, Dept. of Sci. & Tech., Govt. of India under grant number SB/S3/ ME/0044/2013; and Sponsored Research and Industrial Consultancy, Indian Institute of Technology Kharagpur, India under grant number GAF. References  W.H. Wang, C. Dong, C.H. Shek, Bulk metallic glasses, Mater. Sci. Eng. R 44 (2004) 45–89.  C.A. Schuh, T.C. Hufnagel, U. Ramamurty, Mechanical behavior of amorphous alloys, Acta Mater. 55 (2007) 4067–4109.  B.J. Yang, J.H. Yao, J. Zhang, H.W. Yang, J.Q. Wang, E. Ma, Al-rich bulk metallic glasses with plasticity and ultrahigh speciﬁc strength, Scr. Mater. 61 (2009) 423–426.  B.J. Yang, J.H. Yao, Y.S. Chao, J.Q. Wang, E. Ma, Developing aluminum-based bulk metallic glasses, Philos. Mag. 90 (2010) 3215–3231.  X.P. Li, M. Yan, H. Imai, K. Kondoh, J.Q. Wang, G.B. Schaffer, M. Qian, Fabrication of 10 mm diameter fully dense Al86Ni6Y4.5Co2La1.5 bulk metallic glass with high fracture strength, Mat. Sci. Eng. A 568 (2013) 155–159.  S.S. Deng, D.J. Wanga, Q. Luo, Y.J. Huang, J. Shen, Spark plasma sintering of gas atomized AlNiYLaCo amorphous powders, Adv. Powder Technol. 26 (2015) 1696–1701.  R.S. Maurya, A. Sahu, T. Laha, Effect of consolidation pressure on phase evolution during sintering of mechanically alloyed Al86Ni8Y6 amorphous powders via spark plasma sintering, Mat. Sci. Eng. A 649 (2016) 48–56.  R.S. Maurya, T. Laha, Effect of rare earth and transition metal elements on the glass forming ability of mechanical alloyed Al–TM–RE based amorphous alloys, J. Mater. Sci. Technol. 31 (2015) 1118–1124.
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