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Nanocrystallization in Mg83Ni17−xYx (x=0, 7.5) amorphous alloys
Journal of Alloys and Compounds 345 (2002) 123–129
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Nanocrystallization in Mg 83 Ni 172xY x (x50, 7.5) amorphous alloys T. Spassov b
a,b ,
b b *, P. Solsona b , S. Surinach ˜ , M.D. Baro´
a Department of Chemistry, University of Sofia ‘ St.Kl.Ohridski’, 1 J.Bourchier str., 1126 Sofia, Bulgaria ´ ` ` , Facultat de Ciencies , Universitat Autonoma de Barcelona, LMT, Edifici Cc, 08193 Bellaterra /Barcelona, Spain Departament de Fısica
Received 17 January 2002; accepted 11 February 2002
Abstract Thermal stability and crystallization of two rapidly quenched Mg-based amorphous alloys (Mg 83 Ni 17 and Mg 83 Ni 9.5 Y 7.5 ) were studied. The influence of yttrium on the crystallization as well as on the microstructure after crystallization was analysed. The binary alloy crystallizes at about 180 8C to a metastable Mg 6 Ni phase with a grain size of about 30 nm. It is shown that the transformation proceeds by nucleation and three dimensional linear growth until the end of the process. At higher temperatures, above 300 8C, the metastable phase transforms to the equilibrium aMg and Mg 2 Ni phases with a very small thermal effect and slow kinetics, being hard to detect by DSC. The ternary alloy (Ni in Mg 83 Ni 17 is partially replaced by Y) reveals rather different crystallization behavior. Nanocrystallization takes place in the range 175–225 8C as a two-step process. The resulting microstructure is extremely fine with an average nanocrystalline size of about 5–6 nm. Kinetic analysis of the first nanocrystallization reaction shows that most probably nucleation with a very high rate leads to high initial density of nuclei, which grow with a diffusion-controlled rate. Impingement of the diffusion fields of the growing nanocrystals occurs at later stages of the transformation and additionally decreases the kinetics. The nanostructure is stable and transforms to the equilibrium Mg 24 Y 5 , Mg 2 Ni and Ni 2 Y 3 phases with coarser microstructure at higher temperatures (250–290 8C). The formation of Mg–Y and Ni–Y intermetallics during crystallization is expected to impede the grain growth of the Mg 2 Ni phase and results in finer microstructure of the fully transformed ternary alloy compared to the binary Mg–Ni alloy. 2002 Elsevier Science B.V. All rights reserved. Keywords: Amorphous materials; Rare earth alloys; Rapid solidification; Kinetics; X-ray diffraction
1. Introduction Due to the increasing interest in magnesium-based alloys as appropriate materials for hydrogen storage and promising structural materials, a number of studies have been devoted during recent years to rapidly quenched Mg–(Cu,Ni)–RE (RE5rare earth or yttrium) alloys [1–8]. The very good glass forming ability of this system was thoroughly investigated [1] as well. In a series of our previous papers the crystallization, microstructure and the hydrogen storage properties for various rapidly quenched Mg–Ni–RE alloys were studied [2–5]. Magnesium alloys with hypoeutectic and hypereutectic compositions (according to the binary Mg–Ni phase diagram) were investigated. A dependence of the hydrogen storage characteristics on the composition and *Corresponding author. E-mail address: [email protected] (T. Spassov).
microstructure of the alloys as well as the appropriate heat treatment conditions for preparing nanocrystalline or nano/ amorphous microstructures with improved hydrogen storage properties were found. It was shown that the nanocrystallization of melt-spun amorphous Mg–Ni–RE alloys is one of the promising approaches for producing Mg-based alloys with excellent hydrogenation / dehydrogenation kinetics [2–6]. Formation of a metastable nanocrystalline Mg 6 Ni phase was found to have a positive effect on the hydrogen sorption kinetics, too [4]. The present work continues a systematic study on the thermal stability and crystallization of rapidly quenched Mg-based alloys for hydrogen storage. The Mg 83 (Ni,Y) 17 alloys in this work are close to the eutectic composition of the binary Mg–Ni system. The main aim is to investigate the possibility for formation of stable nanostructures varying the composition and heat treatment conditions. The influence of Y on the crystallization and microstructure of the crystallized alloys is studied as well.
0925-8388 / 02 / $ – see front matter 2002 Elsevier Science B.V. All rights reserved. PII: S0925-8388( 02 )00286-4
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2. Experimental The pre-alloys were prepared by induction melting of high purity Mg, Ni, Y in a furnace under pure argon. From the master alloy ingots, ribbons were produced by meltspinning with an approximate quenching rate of 25–30 m / s. The microstructure of the melt-spun materials as well as the crystalline phases in the as-quenched and heat treated alloys were characterized by X-ray (Cu Ka) diffraction, using a Philips 3050 diffractometer. Quantitative phase and microstructural information for the crystallized alloys was obtained by applying a program (MAUD) analysing the X-ray diffraction spectra [9,10]. The chemical composition of the alloys was examined by SEM with an energy dispersive X-ray analysis (EDX). Thermal stability and crystallization kinetics of the rapidly solidified alloys were studied by means of DSC (Perkin-Elmer DSC7) under argon atmosphere.
3. Results Both magnesium alloys studied (Mg 83 Ni 17 and Mg 83 Ni 9.5 Y 7.5 ) were melt-spun at approximately the same quenching rate and in the as-cast state they show XRD patterns typical for amorphous / nanocrystalline microstructures, Fig. 1. The effective size of the nanocrystals estimated by Scherrer’s equation from the full width at the half maximum of the peak is about 3–4 nm, practically the same for both rapidly quenched alloys. The position of the main diffraction halo however is different for the two alloys. The maximum of the main diffraction halo of Mg 83 Ni 17 corresponds to the main peaks of a metastable phase, which was recently identified by us as f.c.c. Mg 6 Ni [4], and for Mg 83 Ni 9.5 Y 7.5 it corresponds to Mg–Y and Ni–Y intermetallic phases. The as-quenched Mg 83 Ni 9.5 Y 7.5 alloy shows some better defined diffraction peaks superimposed on the main nano- / amorphous diffraction halo, probably due to the existence of quenched-in
Fig. 1. XRD patterns (Cu Ka) of the as-quenched Mg 83 Ni 172xY x alloys.
(nano)crystals at the surface of the ribbon, because after polishing or grinding the ribbon they disappear from the diffraction pattern. Fig. 2 shows DSC continuous heating analysis of both as-quenched alloys. While the Mg 83 Ni 17 reveals a sharp exothermic peak at about 180 8C, the alloy with yttrium shows a rather broad two-peak exothermic effect with relatively small enthalpy in the temperature range of 175– 225 8C. The temperatures and the enthalpies of the transformations are presented in Table 1. At higher temperatures (250–290 8C for Mg 83 Ni 9.5 Y 7.5 and above 300 8C for Mg 83 Ni 17 ) both alloys undergo phase transformations characterised by very broad exothermal effects. The enthalpy changes of the high temperature reactions are very small, being even difficult to detect by DSC (especially in the case of Mg 83 Ni 17 ). X-ray analysis after each exothermic effect was carried out in order to study the crystalline products of the different reactions during heating of the as-quenched alloys, Fig. 3. After the large single peak at about 180 8C the Mg 83 Ni 17 alloy is almost completely crystalline and consists of the metastable Mg 6 Ni phase. Fig. 3a shows an X-ray diffractogram of the product of the first crystallization reaction after isothermal annealing for 30 min at 165 8C. All diffraction peaks of the later phase are welldefined revealing that the crystals formed are not very small, in contrast to the Mg 6 Ni phase obtained during the primary crystallization in Mg 78 Ni 18 Y 4 [5]. From the XRD peak shapes using the peak fitting program MAUD [9,10] an average grain size of about 30 nm was determined. During further heating (after the high temperature broad peak at about 300–350 8C) the metastable phase transforms to the equilibrium aMg and Mg 2 Ni phases. It has to be pointed out that similar to the rapidly quenched Mg 87 Ni 12 Y 1 and Mg 78 Ni 18 Y 4 alloys studied by us previously [4,5] the intermediate Mg 6 Ni phase in Mg 83 Ni 17 is quite stable, to temperatures higher than 100 8C after its formation. The XRD reveals a preferential orientation of
Fig. 2. DSC scans of the as-quenched Mg 83 Ni 172xY x alloys.
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Table 1 DSC data for the first crystallization reactions in Mg 83 Ni 172xY x (x50 and 7.5) Alloy
Mg 83 Ni 17 Mg 83 Ni 9.5 Y 7.5
I DSC peak
II DSC peak
T x (8C)
T p (8C)
DHcr (J / g)
T x (8C)
T p (8C)
DHcr (J / g)
176.5 175.5
179 187.5
68 27 a
– 187 a
– 208.5
– 15 a
T x , crystallization temperature; T p , temperature corresponding to the maximum of the DSC peak; DHcr , enthalpy of crystallization reaction. a Obtained by combining isothermal pre-annealing with subsequent isochronal DSC scans.
the Mg 6 Ni and aMg crystallites with the (822) and (002) axis, respectively, perpendicular to the ribbon surface. Mg 83 Ni 9.5 Y 7.5 crystallizes in several steps, as the en-
thalpy effects of all reactions are relatively small. In contrast to Mg 83 Ni 17 , after the first crystallization reaction the microstructure in Mg 83 Ni 9.5 Y 7.5 remains nanocrystal-
Fig. 3. XRD patterns of Mg 83 Ni 17 (a) and Mg 83 Ni 9.5 Y 7.5 (b) alloys after different heat treatments.
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line / amorphous, Fig. 3b. Heating above the second DSC peak (90 min at 180 8C) does not change the microstructure essentially, too. The diffraction peak maximum remains at the same position and corresponds to the main peaks of Mg–Y and Ni–Y intermetallics (Mg 24 Y 5 and Ni 2 Y 3 ), rather than to those of Mg 2 Ni and Mg 6 Ni. It is, however, difficult to detect the nanocrystalline phase(s) by XRD, due to the very small crystal size even after the second broad DSC peak, Fig. 3b (180 8C / 90 min). The width of the diffraction peaks is reduced very little with the annealing as the mean nanocrystal size changes from about 3–4 nm in the beginning to 5–6 nm at the end of the second DSC peak. The extremely fine nanostructure in this alloy is stable up to about 250 8C. Above this temperature the Mg 24 Y 5 , Ni 2 Y 3 and Mg 2 Ni phases are formed and the microstructure becomes coarser, but remains still nanocrystalline. XRD of the sample annealed at a high temperature (e.g. 10 min at 400 8C) reveals still some nanocrystalline component in the alloy, Fig. 3b.
4. Discussion Non-isothermal and isothermal kinetic analyses were performed to study the crystallization mechanism of these glasses as well as to highlight the influence of Y on the crystallization process and on the microstructure of the devitrified alloys. DSC isothermal traces of the binary Mg 83 Ni 17 alloy at different temperatures in the range 150–165 8C are presented in Fig. 4. After a well-defined incubation time to the isothermal DSC curves reveal clearly an exothermic peak
with smooth onset and sharper end. This shape is generally an indication for a crystallization process including nucleation (homogeneous or heterogeneous) and finishing with hard impingement. In agreement with the shape of the DSC isothermal peaks the kinetics of the Mg 6 Ni formation in Mg 83 Ni 17 could be satisfactorily described by the JMKA model [11,12] with Avrami exponent n54 (Fig. 4). From the analysis of the dependence of the Avrami exponent on the degree of transformation, n(a ), it can be concluded that the nucleation process takes place during the whole transformation, since n is constantly equal to 4 during the whole transformation. The temperature dependences of the kinetic parameters determined by fitting the experimental transformation curves with the JMKA model are shown in Fig. 5. The activation energy of the incubation period, Q t o 5 150616 kJ / mol, has to be associated with atomic transport during the crystallization (e.g. diffusion). The activation energy obtained from the temperature dependence of the time for half-reaction, Q t 0.5 5 180620 kJ / mol, is an effective value for the overall crystallization process. The very strong temperature dependence of the kinetic constant k is associated with the nucleation and subsequent growth proceeding during the crystallization and the activation energy of k includes the activation energies of both processes according to the expression: Q k 5 Q n 1 3Q g , where Q n is the activation energy of nucleation and Q g for growth. If we assume that the activation energy for linear growth consists of mainly an energy barrier for material transport we obtain an estimate for the activation energy for nucleation of about 350 kJ / mol. The activation energy of this transformation was also estimated according to the Kissinger method [13] to
Fig. 4. Isothermal DSC scans and corresponding transformation curves (inset) of the first crystallization reaction in Mg 83 Ni 17 (dash line, calculated transformation curves according to the JMKA model with n54).
T. Spassov et al. / Journal of Alloys and Compounds 345 (2002) 123–129
Fig. 5. Arrhenius temperature dependence of the incubation time (to ), the time for half-reaction (t 0.5 ) and the kinetic constant (k) in the JMKA equation for the first crystallization reaction in Mg 83 Ni 17 .
be 160610 kJ / mol, Fig. 6, which agrees well with those obtained from the isothermal kinetics. As it was mentioned already, the alloy containing Y devitrifies in several steps. By varying the temperature of isothermal annealing it was not possible to detect the onset of the crystallization DSC peak; after a very sharp initial increase the calorimetric signal continuously decreases until reaching a saturation at long times of annealing, Fig. 7. This is a clear indication that after very fast nucleation (homogeneous or heterogeneous on high density of active sites for nucleation) leading to the formation of a high initial density of nuclei, diffusion-controlled growth takes place. The growth rate as well as the final average nanocrystalline size depends on the annealing temperature. The higher the temperature, the larger the change in enthalpy is observed for the same time of annealing (from 10 J / g at 155 8C to 25 J / g at 180 8C for 30 min
Fig. 6. Kissinger plot for the first crystallization reactions in Mg 83 Ni 17 and Mg 83 Ni 9.5 Y 7.5 .
Fig. 7. DSC isotherms Mg 83 Ni 9.5 Y 7.5 .
of
the
127
first
crystallization
reaction
in
annealing). After sufficiently long time of isothermal heat treatment the first reaction proceeds completely and during subsequent annealing at the same temperature a new decrease in the calorimetric signal was not observed, indicating that the growth has stopped at a certain degree of transformation. XRD after the first crystallization reaction reveals a nanocrystalline microstructure at all temperatures and times of isothermal annealing in the range 155–180 8C (Fig. 3b). Continuous heating DSC following different times of isothermal annealing was carried out in order to determine the fraction transformed at the first crystallization reaction during the isothermal treatment, Fig. 8. For this analysis it is assumed that the sum of the heat released during the isothermal pre-annealing and further continuous heating crystallization is constant. Fig. 8 (inset) shows the fraction transformed as a function of the annealing time at 160 8C. From the isothermal transformation curve it can be seen that after a very short nucleation stage crystal growth takes place. Kinetic analysis of the transformation curve reveals that the growth of nanocrystals is diffusion controlled with soft impingement at the advance stage of the transformation [14–16]. Study on the development of the effective ¯ with the annealing time at constant average grain size, d, temperature of 160 8C showed that d¯ , œDt (D is the diffusion coefficient of the slow diffusing element), a result which confirmed the mechanism proposed. The formation of a high density of nanocrystals, either due to the presence of a high concentration of quenched-in nuclei or due to heterogeneous nucleation on the high density of active sites for nucleation, is probably the main reason for the very slow growth observed, because of diffusion field impingement of the growing nanocrystals (so-called soft impingement). The DSC experiments confirm this mechanism, too; at continuous heating conditions a very small thermal peak can be detected and in the isothermal mode we observe practically only a decreasing calorimetric signal.
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Fig. 8. DSC curves of Mg 83 Ni 9.5 Y 7.5 previously annealed for different times at 160 8C (10 K / min heating rate).
It is important to mention that even after 100 min annealing at 170 8C the second reaction does not start, in spite of the fact that in continuous heating experiments the exothermic peaks of the first two reactions overlap. It was proved at several temperatures in the range 155–180 8C that only after full completion of the first crystallization reaction does the second reaction start. This result shows that the second reaction is most probably the crystallization of amorphous phase, rather than a phase transformation of the nanocrystalline phase, formed during the first reaction. The amount of the amorphous phase left untransformed after the first crystallization reaction is relatively small, because the corresponding thermal effect of the second reaction is only about 15 J / g. As it could be expected, the product of this transformation is also nanocrystalline, due to the short distances between the already existing nanocrystals. At the end of the second transformation the average size of the nanocrystals is about 5–6 nm, i.e. the first nanocrystals have not become larger during further annealing, applied for the second crystallization reaction. In order to obtain information about the activation energy of the nanocrystallization reactions in the Mg–Ni– Y alloy, a Kissinger analysis was applied, Fig. 6. The activation energy of the first reaction, 17565 kJ / mol, has to be associated with growth of existing quenched-in nuclei or heterogeneous nucleation on a high density of potent sites and subsequent growth. The activation energy for the second nanoreaction was difficult to determine accurately due to the overlapping of the DSC peaks, but a rough estimation shows that it is larger than that of the first reaction. From the present study and previous results [2–5] it can be concluded that Y has a strong influence on the
crystallization of rapidly quenched magnesium-based Mg– Ni amorphous alloys, as well as on their microstructure after crystallization. Yttrium additions (above 3–4 at.%) lead to drastic reduction of the average grain size after partial and complete crystallization of the alloy, consisting of mainly the equilibrium aMg and Mg 2 Ni phases and / or the metastable Mg 6 Ni, depending on the general composition of the alloys. Partial Y dissolution in the aMg and / or Mg 2 Ni was also observed [2,5]. At higher Y concentrations (CY .5–6 at.%) Ni–Y and Mg–Y intermetallic phases are formed, too. The nanocrystalline or nano- / amorphous microstructures show high thermal stability. For some Mg–Ni–Y alloys (Mg 78 Ni 18 Y 4 and Mg 83 Ni 9.5 Y 7.5 ) the extremely fine nanostructure (d¯ , 10 nm) is stable up to 300–350 8C. Generally, the presence of Y in the alloy slows down its crystallization kinetics, due to transport difficulties during the growth of the major aMg and Mg–Ni phases, caused by the larger and therefore slowly diffusing yttrium atoms. Additional delay of the growth process is a result of diffusion field impingement of the high density of growing precipitates in the advanced stage of transformation. The formation of Ni–Y and Mg–Y intermetallics (at higher Y concentration) as minor phases prevents the grain growth of the major Mg and Mg 2 Ni phases and leads to much finer microstructure in the completely crystallized ternary Mg–Ni–Y alloys compared to the binary Mg–Ni alloy.
5. Conclusion The devitrification of two rapidly quenched Mg 83 Ni 172xY x (x50 and 7.5 at.%) amorphous or nano-
T. Spassov et al. / Journal of Alloys and Compounds 345 (2002) 123–129
/ amorphous alloys was studied. The influence of yttrium on the crystallization as well as on the microstructure after crystallization was analysed. In the binary Mg 83 Ni 17 alloy after an initial crystallization of a metastable Mg 6 Ni phase (with average crystal size of about 30 nm) the thermodynamically stable aMg and Mg 2 Ni phases are formed. The kinetics of the metastable phase formation were found to obey a JMKA model with nucleation and three-dimensional linear growth with hard impingement of the growing crystals at the end of the transformation. The activation energies of the primary crystallization process were obtained from the temperature dependence of the incubation time, the time for half reaction and the kinetic constant in the JMKA equation as well as by Kissinger analysis. In the alloy with yttrium, Mg 83 Ni 9.5 Y 7.5 , the crystallization starts with the formation of an extremely fine nanocrystalline microstructure (5–6 nm), as the process is characterized by two broad overlapping exothermal peaks with relatively small enthalpy (about 42 J / g). The nanocrystalline phases are most probably Mg–Y and Ni–Y intermetallics. During further annealing larger crystals of the equilibrium Mg 24 Y 5 , Mg 2 Ni and Ni 2 Y 3 phases are formed. It can be concluded that the main reason for the fine nanocrystalline microstructure in Mg–Ni–Y is most probably the high density of nuclei existing in the as-quenched ribbon or formed on active heterogeneous sites at the beginning of the reaction and the subsequent diffusioncontrolled growth (Y is the slow diffusing element). Impingement of the diffusion fields of the growing nanocrystals at later stages of the transformation leads to additional delay of the crystallization (growth) process.
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Acknowledgements One of the authors (T.S.) is very grateful to the Ministry of Education, Culture and Sports of Spain for the financial support. The work was also supported by the project 2001-SGR-00189.
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Nanocrystallization in Mg 83 Ni 172xY x (x50, 7.5) amorphous alloys T. Spassov b
a,b ,
b b *, P. Solsona b , S. Surinach ˜ , M.D. Baro´
a Department of Chemistry, University of Sofia ‘ St.Kl.Ohridski’, 1 J.Bourchier str., 1126 Sofia, Bulgaria ´ ` ` , Facultat de Ciencies , Universitat Autonoma de Barcelona, LMT, Edifici Cc, 08193 Bellaterra /Barcelona, Spain Departament de Fısica
Received 17 January 2002; accepted 11 February 2002
Abstract Thermal stability and crystallization of two rapidly quenched Mg-based amorphous alloys (Mg 83 Ni 17 and Mg 83 Ni 9.5 Y 7.5 ) were studied. The influence of yttrium on the crystallization as well as on the microstructure after crystallization was analysed. The binary alloy crystallizes at about 180 8C to a metastable Mg 6 Ni phase with a grain size of about 30 nm. It is shown that the transformation proceeds by nucleation and three dimensional linear growth until the end of the process. At higher temperatures, above 300 8C, the metastable phase transforms to the equilibrium aMg and Mg 2 Ni phases with a very small thermal effect and slow kinetics, being hard to detect by DSC. The ternary alloy (Ni in Mg 83 Ni 17 is partially replaced by Y) reveals rather different crystallization behavior. Nanocrystallization takes place in the range 175–225 8C as a two-step process. The resulting microstructure is extremely fine with an average nanocrystalline size of about 5–6 nm. Kinetic analysis of the first nanocrystallization reaction shows that most probably nucleation with a very high rate leads to high initial density of nuclei, which grow with a diffusion-controlled rate. Impingement of the diffusion fields of the growing nanocrystals occurs at later stages of the transformation and additionally decreases the kinetics. The nanostructure is stable and transforms to the equilibrium Mg 24 Y 5 , Mg 2 Ni and Ni 2 Y 3 phases with coarser microstructure at higher temperatures (250–290 8C). The formation of Mg–Y and Ni–Y intermetallics during crystallization is expected to impede the grain growth of the Mg 2 Ni phase and results in finer microstructure of the fully transformed ternary alloy compared to the binary Mg–Ni alloy. 2002 Elsevier Science B.V. All rights reserved. Keywords: Amorphous materials; Rare earth alloys; Rapid solidification; Kinetics; X-ray diffraction
1. Introduction Due to the increasing interest in magnesium-based alloys as appropriate materials for hydrogen storage and promising structural materials, a number of studies have been devoted during recent years to rapidly quenched Mg–(Cu,Ni)–RE (RE5rare earth or yttrium) alloys [1–8]. The very good glass forming ability of this system was thoroughly investigated [1] as well. In a series of our previous papers the crystallization, microstructure and the hydrogen storage properties for various rapidly quenched Mg–Ni–RE alloys were studied [2–5]. Magnesium alloys with hypoeutectic and hypereutectic compositions (according to the binary Mg–Ni phase diagram) were investigated. A dependence of the hydrogen storage characteristics on the composition and *Corresponding author. E-mail address: [email protected] (T. Spassov).
microstructure of the alloys as well as the appropriate heat treatment conditions for preparing nanocrystalline or nano/ amorphous microstructures with improved hydrogen storage properties were found. It was shown that the nanocrystallization of melt-spun amorphous Mg–Ni–RE alloys is one of the promising approaches for producing Mg-based alloys with excellent hydrogenation / dehydrogenation kinetics [2–6]. Formation of a metastable nanocrystalline Mg 6 Ni phase was found to have a positive effect on the hydrogen sorption kinetics, too [4]. The present work continues a systematic study on the thermal stability and crystallization of rapidly quenched Mg-based alloys for hydrogen storage. The Mg 83 (Ni,Y) 17 alloys in this work are close to the eutectic composition of the binary Mg–Ni system. The main aim is to investigate the possibility for formation of stable nanostructures varying the composition and heat treatment conditions. The influence of Y on the crystallization and microstructure of the crystallized alloys is studied as well.
0925-8388 / 02 / $ – see front matter 2002 Elsevier Science B.V. All rights reserved. PII: S0925-8388( 02 )00286-4
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2. Experimental The pre-alloys were prepared by induction melting of high purity Mg, Ni, Y in a furnace under pure argon. From the master alloy ingots, ribbons were produced by meltspinning with an approximate quenching rate of 25–30 m / s. The microstructure of the melt-spun materials as well as the crystalline phases in the as-quenched and heat treated alloys were characterized by X-ray (Cu Ka) diffraction, using a Philips 3050 diffractometer. Quantitative phase and microstructural information for the crystallized alloys was obtained by applying a program (MAUD) analysing the X-ray diffraction spectra [9,10]. The chemical composition of the alloys was examined by SEM with an energy dispersive X-ray analysis (EDX). Thermal stability and crystallization kinetics of the rapidly solidified alloys were studied by means of DSC (Perkin-Elmer DSC7) under argon atmosphere.
3. Results Both magnesium alloys studied (Mg 83 Ni 17 and Mg 83 Ni 9.5 Y 7.5 ) were melt-spun at approximately the same quenching rate and in the as-cast state they show XRD patterns typical for amorphous / nanocrystalline microstructures, Fig. 1. The effective size of the nanocrystals estimated by Scherrer’s equation from the full width at the half maximum of the peak is about 3–4 nm, practically the same for both rapidly quenched alloys. The position of the main diffraction halo however is different for the two alloys. The maximum of the main diffraction halo of Mg 83 Ni 17 corresponds to the main peaks of a metastable phase, which was recently identified by us as f.c.c. Mg 6 Ni [4], and for Mg 83 Ni 9.5 Y 7.5 it corresponds to Mg–Y and Ni–Y intermetallic phases. The as-quenched Mg 83 Ni 9.5 Y 7.5 alloy shows some better defined diffraction peaks superimposed on the main nano- / amorphous diffraction halo, probably due to the existence of quenched-in
Fig. 1. XRD patterns (Cu Ka) of the as-quenched Mg 83 Ni 172xY x alloys.
(nano)crystals at the surface of the ribbon, because after polishing or grinding the ribbon they disappear from the diffraction pattern. Fig. 2 shows DSC continuous heating analysis of both as-quenched alloys. While the Mg 83 Ni 17 reveals a sharp exothermic peak at about 180 8C, the alloy with yttrium shows a rather broad two-peak exothermic effect with relatively small enthalpy in the temperature range of 175– 225 8C. The temperatures and the enthalpies of the transformations are presented in Table 1. At higher temperatures (250–290 8C for Mg 83 Ni 9.5 Y 7.5 and above 300 8C for Mg 83 Ni 17 ) both alloys undergo phase transformations characterised by very broad exothermal effects. The enthalpy changes of the high temperature reactions are very small, being even difficult to detect by DSC (especially in the case of Mg 83 Ni 17 ). X-ray analysis after each exothermic effect was carried out in order to study the crystalline products of the different reactions during heating of the as-quenched alloys, Fig. 3. After the large single peak at about 180 8C the Mg 83 Ni 17 alloy is almost completely crystalline and consists of the metastable Mg 6 Ni phase. Fig. 3a shows an X-ray diffractogram of the product of the first crystallization reaction after isothermal annealing for 30 min at 165 8C. All diffraction peaks of the later phase are welldefined revealing that the crystals formed are not very small, in contrast to the Mg 6 Ni phase obtained during the primary crystallization in Mg 78 Ni 18 Y 4 [5]. From the XRD peak shapes using the peak fitting program MAUD [9,10] an average grain size of about 30 nm was determined. During further heating (after the high temperature broad peak at about 300–350 8C) the metastable phase transforms to the equilibrium aMg and Mg 2 Ni phases. It has to be pointed out that similar to the rapidly quenched Mg 87 Ni 12 Y 1 and Mg 78 Ni 18 Y 4 alloys studied by us previously [4,5] the intermediate Mg 6 Ni phase in Mg 83 Ni 17 is quite stable, to temperatures higher than 100 8C after its formation. The XRD reveals a preferential orientation of
Fig. 2. DSC scans of the as-quenched Mg 83 Ni 172xY x alloys.
T. Spassov et al. / Journal of Alloys and Compounds 345 (2002) 123–129
125
Table 1 DSC data for the first crystallization reactions in Mg 83 Ni 172xY x (x50 and 7.5) Alloy
Mg 83 Ni 17 Mg 83 Ni 9.5 Y 7.5
I DSC peak
II DSC peak
T x (8C)
T p (8C)
DHcr (J / g)
T x (8C)
T p (8C)
DHcr (J / g)
176.5 175.5
179 187.5
68 27 a
– 187 a
– 208.5
– 15 a
T x , crystallization temperature; T p , temperature corresponding to the maximum of the DSC peak; DHcr , enthalpy of crystallization reaction. a Obtained by combining isothermal pre-annealing with subsequent isochronal DSC scans.
the Mg 6 Ni and aMg crystallites with the (822) and (002) axis, respectively, perpendicular to the ribbon surface. Mg 83 Ni 9.5 Y 7.5 crystallizes in several steps, as the en-
thalpy effects of all reactions are relatively small. In contrast to Mg 83 Ni 17 , after the first crystallization reaction the microstructure in Mg 83 Ni 9.5 Y 7.5 remains nanocrystal-
Fig. 3. XRD patterns of Mg 83 Ni 17 (a) and Mg 83 Ni 9.5 Y 7.5 (b) alloys after different heat treatments.
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line / amorphous, Fig. 3b. Heating above the second DSC peak (90 min at 180 8C) does not change the microstructure essentially, too. The diffraction peak maximum remains at the same position and corresponds to the main peaks of Mg–Y and Ni–Y intermetallics (Mg 24 Y 5 and Ni 2 Y 3 ), rather than to those of Mg 2 Ni and Mg 6 Ni. It is, however, difficult to detect the nanocrystalline phase(s) by XRD, due to the very small crystal size even after the second broad DSC peak, Fig. 3b (180 8C / 90 min). The width of the diffraction peaks is reduced very little with the annealing as the mean nanocrystal size changes from about 3–4 nm in the beginning to 5–6 nm at the end of the second DSC peak. The extremely fine nanostructure in this alloy is stable up to about 250 8C. Above this temperature the Mg 24 Y 5 , Ni 2 Y 3 and Mg 2 Ni phases are formed and the microstructure becomes coarser, but remains still nanocrystalline. XRD of the sample annealed at a high temperature (e.g. 10 min at 400 8C) reveals still some nanocrystalline component in the alloy, Fig. 3b.
4. Discussion Non-isothermal and isothermal kinetic analyses were performed to study the crystallization mechanism of these glasses as well as to highlight the influence of Y on the crystallization process and on the microstructure of the devitrified alloys. DSC isothermal traces of the binary Mg 83 Ni 17 alloy at different temperatures in the range 150–165 8C are presented in Fig. 4. After a well-defined incubation time to the isothermal DSC curves reveal clearly an exothermic peak
with smooth onset and sharper end. This shape is generally an indication for a crystallization process including nucleation (homogeneous or heterogeneous) and finishing with hard impingement. In agreement with the shape of the DSC isothermal peaks the kinetics of the Mg 6 Ni formation in Mg 83 Ni 17 could be satisfactorily described by the JMKA model [11,12] with Avrami exponent n54 (Fig. 4). From the analysis of the dependence of the Avrami exponent on the degree of transformation, n(a ), it can be concluded that the nucleation process takes place during the whole transformation, since n is constantly equal to 4 during the whole transformation. The temperature dependences of the kinetic parameters determined by fitting the experimental transformation curves with the JMKA model are shown in Fig. 5. The activation energy of the incubation period, Q t o 5 150616 kJ / mol, has to be associated with atomic transport during the crystallization (e.g. diffusion). The activation energy obtained from the temperature dependence of the time for half-reaction, Q t 0.5 5 180620 kJ / mol, is an effective value for the overall crystallization process. The very strong temperature dependence of the kinetic constant k is associated with the nucleation and subsequent growth proceeding during the crystallization and the activation energy of k includes the activation energies of both processes according to the expression: Q k 5 Q n 1 3Q g , where Q n is the activation energy of nucleation and Q g for growth. If we assume that the activation energy for linear growth consists of mainly an energy barrier for material transport we obtain an estimate for the activation energy for nucleation of about 350 kJ / mol. The activation energy of this transformation was also estimated according to the Kissinger method [13] to
Fig. 4. Isothermal DSC scans and corresponding transformation curves (inset) of the first crystallization reaction in Mg 83 Ni 17 (dash line, calculated transformation curves according to the JMKA model with n54).
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Fig. 5. Arrhenius temperature dependence of the incubation time (to ), the time for half-reaction (t 0.5 ) and the kinetic constant (k) in the JMKA equation for the first crystallization reaction in Mg 83 Ni 17 .
be 160610 kJ / mol, Fig. 6, which agrees well with those obtained from the isothermal kinetics. As it was mentioned already, the alloy containing Y devitrifies in several steps. By varying the temperature of isothermal annealing it was not possible to detect the onset of the crystallization DSC peak; after a very sharp initial increase the calorimetric signal continuously decreases until reaching a saturation at long times of annealing, Fig. 7. This is a clear indication that after very fast nucleation (homogeneous or heterogeneous on high density of active sites for nucleation) leading to the formation of a high initial density of nuclei, diffusion-controlled growth takes place. The growth rate as well as the final average nanocrystalline size depends on the annealing temperature. The higher the temperature, the larger the change in enthalpy is observed for the same time of annealing (from 10 J / g at 155 8C to 25 J / g at 180 8C for 30 min
Fig. 6. Kissinger plot for the first crystallization reactions in Mg 83 Ni 17 and Mg 83 Ni 9.5 Y 7.5 .
Fig. 7. DSC isotherms Mg 83 Ni 9.5 Y 7.5 .
of
the
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first
crystallization
reaction
in
annealing). After sufficiently long time of isothermal heat treatment the first reaction proceeds completely and during subsequent annealing at the same temperature a new decrease in the calorimetric signal was not observed, indicating that the growth has stopped at a certain degree of transformation. XRD after the first crystallization reaction reveals a nanocrystalline microstructure at all temperatures and times of isothermal annealing in the range 155–180 8C (Fig. 3b). Continuous heating DSC following different times of isothermal annealing was carried out in order to determine the fraction transformed at the first crystallization reaction during the isothermal treatment, Fig. 8. For this analysis it is assumed that the sum of the heat released during the isothermal pre-annealing and further continuous heating crystallization is constant. Fig. 8 (inset) shows the fraction transformed as a function of the annealing time at 160 8C. From the isothermal transformation curve it can be seen that after a very short nucleation stage crystal growth takes place. Kinetic analysis of the transformation curve reveals that the growth of nanocrystals is diffusion controlled with soft impingement at the advance stage of the transformation [14–16]. Study on the development of the effective ¯ with the annealing time at constant average grain size, d, temperature of 160 8C showed that d¯ , œDt (D is the diffusion coefficient of the slow diffusing element), a result which confirmed the mechanism proposed. The formation of a high density of nanocrystals, either due to the presence of a high concentration of quenched-in nuclei or due to heterogeneous nucleation on the high density of active sites for nucleation, is probably the main reason for the very slow growth observed, because of diffusion field impingement of the growing nanocrystals (so-called soft impingement). The DSC experiments confirm this mechanism, too; at continuous heating conditions a very small thermal peak can be detected and in the isothermal mode we observe practically only a decreasing calorimetric signal.
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Fig. 8. DSC curves of Mg 83 Ni 9.5 Y 7.5 previously annealed for different times at 160 8C (10 K / min heating rate).
It is important to mention that even after 100 min annealing at 170 8C the second reaction does not start, in spite of the fact that in continuous heating experiments the exothermic peaks of the first two reactions overlap. It was proved at several temperatures in the range 155–180 8C that only after full completion of the first crystallization reaction does the second reaction start. This result shows that the second reaction is most probably the crystallization of amorphous phase, rather than a phase transformation of the nanocrystalline phase, formed during the first reaction. The amount of the amorphous phase left untransformed after the first crystallization reaction is relatively small, because the corresponding thermal effect of the second reaction is only about 15 J / g. As it could be expected, the product of this transformation is also nanocrystalline, due to the short distances between the already existing nanocrystals. At the end of the second transformation the average size of the nanocrystals is about 5–6 nm, i.e. the first nanocrystals have not become larger during further annealing, applied for the second crystallization reaction. In order to obtain information about the activation energy of the nanocrystallization reactions in the Mg–Ni– Y alloy, a Kissinger analysis was applied, Fig. 6. The activation energy of the first reaction, 17565 kJ / mol, has to be associated with growth of existing quenched-in nuclei or heterogeneous nucleation on a high density of potent sites and subsequent growth. The activation energy for the second nanoreaction was difficult to determine accurately due to the overlapping of the DSC peaks, but a rough estimation shows that it is larger than that of the first reaction. From the present study and previous results [2–5] it can be concluded that Y has a strong influence on the
crystallization of rapidly quenched magnesium-based Mg– Ni amorphous alloys, as well as on their microstructure after crystallization. Yttrium additions (above 3–4 at.%) lead to drastic reduction of the average grain size after partial and complete crystallization of the alloy, consisting of mainly the equilibrium aMg and Mg 2 Ni phases and / or the metastable Mg 6 Ni, depending on the general composition of the alloys. Partial Y dissolution in the aMg and / or Mg 2 Ni was also observed [2,5]. At higher Y concentrations (CY .5–6 at.%) Ni–Y and Mg–Y intermetallic phases are formed, too. The nanocrystalline or nano- / amorphous microstructures show high thermal stability. For some Mg–Ni–Y alloys (Mg 78 Ni 18 Y 4 and Mg 83 Ni 9.5 Y 7.5 ) the extremely fine nanostructure (d¯ , 10 nm) is stable up to 300–350 8C. Generally, the presence of Y in the alloy slows down its crystallization kinetics, due to transport difficulties during the growth of the major aMg and Mg–Ni phases, caused by the larger and therefore slowly diffusing yttrium atoms. Additional delay of the growth process is a result of diffusion field impingement of the high density of growing precipitates in the advanced stage of transformation. The formation of Ni–Y and Mg–Y intermetallics (at higher Y concentration) as minor phases prevents the grain growth of the major Mg and Mg 2 Ni phases and leads to much finer microstructure in the completely crystallized ternary Mg–Ni–Y alloys compared to the binary Mg–Ni alloy.
5. Conclusion The devitrification of two rapidly quenched Mg 83 Ni 172xY x (x50 and 7.5 at.%) amorphous or nano-
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/ amorphous alloys was studied. The influence of yttrium on the crystallization as well as on the microstructure after crystallization was analysed. In the binary Mg 83 Ni 17 alloy after an initial crystallization of a metastable Mg 6 Ni phase (with average crystal size of about 30 nm) the thermodynamically stable aMg and Mg 2 Ni phases are formed. The kinetics of the metastable phase formation were found to obey a JMKA model with nucleation and three-dimensional linear growth with hard impingement of the growing crystals at the end of the transformation. The activation energies of the primary crystallization process were obtained from the temperature dependence of the incubation time, the time for half reaction and the kinetic constant in the JMKA equation as well as by Kissinger analysis. In the alloy with yttrium, Mg 83 Ni 9.5 Y 7.5 , the crystallization starts with the formation of an extremely fine nanocrystalline microstructure (5–6 nm), as the process is characterized by two broad overlapping exothermal peaks with relatively small enthalpy (about 42 J / g). The nanocrystalline phases are most probably Mg–Y and Ni–Y intermetallics. During further annealing larger crystals of the equilibrium Mg 24 Y 5 , Mg 2 Ni and Ni 2 Y 3 phases are formed. It can be concluded that the main reason for the fine nanocrystalline microstructure in Mg–Ni–Y is most probably the high density of nuclei existing in the as-quenched ribbon or formed on active heterogeneous sites at the beginning of the reaction and the subsequent diffusioncontrolled growth (Y is the slow diffusing element). Impingement of the diffusion fields of the growing nanocrystals at later stages of the transformation leads to additional delay of the crystallization (growth) process.
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Acknowledgements One of the authors (T.S.) is very grateful to the Ministry of Education, Culture and Sports of Spain for the financial support. The work was also supported by the project 2001-SGR-00189.
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