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On the synthesis of Zr-based bulk amorphous alloys from glass-forming compounds and elemental powders
Journal of Alloys and Compounds 367 (2004) 191–198
On the synthesis of Zr-based bulk amorphous alloys from glass-forming compounds and elemental powders N.P. Djakonova a , T.A. Sviridova a,∗ , E.A. Zakharova a , V.V. Molokanov b , M.I. Petrzhik b b
a Moscow State Institute of Steel and Alloys (Technological University), Leninsky pr. 4, Moscow 117936, Russia Institute of Metallurgy and Materials Science, Russian Academy of Sciences, Leninsky pr. 49, Moscow 117911, Russia
Abstract The possibility of producing Zr-based bulk amorphous alloys (Zr65 Cu17.5 Ni10 Al7.5 and Zr57 Ti5 Cu20 Ni8 Al10 ) by mechanical alloying (MA) from either a mixture of elemental powders (EP), or a mixture of glass-forming compounds (GFC), has been investigated. A preliminary study of the crystal structure stability of the GFC-ingredients during mechanical milling was performed. Amorphization was found to be significantly enhanced during MA by using a GFC mixture. A considerable difference in the phase transformations observed on heating the amorphous phase (AP) was noted for alloys obtained from GFC, as compared with those produced by melt quenching (MQ) or milling of a mixture of elemental powders. It is shown that the amorphous phase obtained by MA from GFC has a wider undercooled liquid region between the glass transition temperature Tg and the crystallization temperature Tx than the amorphous phase produced by the conventional MQ technique. © 2003 Published by Elsevier B.V. Keywords: Amorphous materials; Amorphization; Liquid quenching; Mechanical alloying; X-ray diffraction
1. Introduction Recently discovered multicomponent bulk amorphous alloys have good glass-forming ability and can easily be cast into bulk glassy specimens with dimensions of several centimeters at a cooling rate of 1/100 K/s [1–3]. These alloys exhibit an extended undercooled liquid region between the glass transition temperature Tg and the crystallization temperature Tx , allowing easy deformation and shaping of the material by hot-pressing, due to the viscous flow in the undercooled liquid [4,5]. However, melt quenching (MQ) does not guarantee the formation of an amorphous single-phase state. For example, in the Fe base of an Fe61 Co7 Zr10 Mo5 W2 B15 alloy, the refractory phase ZrB2 is often observed [6,7]. The formation of ZrB2 particles leads to a deviation from the near-eutectic composition, which is favorable for obtaining the amorphous state. Thus, mechanical alloying (MA) seems to be a promising method for the preparation of this alloy. We have recently established that the process of amorphous phase (AP) formation during MA depends consider∗
Corresponding author. E-mail address: [email protected] (T.A. Sviridova).
0925-8388/$ – see front matter © 2003 Published by Elsevier B.V. doi:10.1016/j.jallcom.2003.08.036
ably on the characteristics of the starting materials [8]. If the MA synthesis is carried out for a mixture of specially selected intermetallic glass-forming compounds (GFC), then the solid-state amorphization reaction proceeds much faster than for the mixture of elemental powders (EP). Moreover, the characteristics of the amorphous phase—Tg , Tx and T = Tx − Tg —are frequently higher for MA, as compared to those of the corresponding MQ alloy. The purpose of this paper is to analyze and compare stability characteristics of the well-known multicomponent bulk amorphous alloys Zr65 Cu17.5 Ni10 Al7.5 and Zr57 Ti5 Cu20 Ni8 Al10 , produced by conventional MQ and MA techniques from glass-forming compounds and elemental powder mixtures.
2. Experimental details Milling was carried out for 15/480 min in a high-energy planetary ball mill of AGO-2U type. Low-carbon steel vials and balls of a ball-bearing steel with diameters of 6 and 8 mm were used. The ball-to-powder ratio was 10:1, the plate rotational speed 680 rpm. The vials were charged and sealed in Ar atmosphere. Both elementary powders (with 99.5% or better purity) and intermetallic compounds were used as
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initial components for MA. Ingots of the compounds were melted in high-purity argon atmosphere in an arc furnace with a tungsten electrode and a water-cooled hearth and were then ground into powders with a particle size smaller than 0.1 mm. X-ray diffraction patterns were taken with Co K␣ radiation using a monochromated diffraction beam and a step-scanning diffractometer DRON-4. For the analysis of the samples with amorphous phase, a modification of the Rietveld full-profile fitting method was employed, with the experimental X-ray pattern of the pure amorphous phase reference rescaled into theoretical units of intensity [9,10]. The thermal stability and crystallization of the amorphous phase were investigated on heating in a differential scanning calorimeter (DSC-111 SETARAM microcalorimeter). The melting and solidification temperatures were determined by differential thermal analysis (DTA-7 HT thermal analyzer). The samples were annealed in a sealed quartz ampoule evacuated down to 10−3 Torr. In order to prevent oxidation, the powders were wrapped in a double-layer titanium foil with zirconium powder, which also served to keep track of any oxidation.
nealed Zr65 Cu17.5 Ni10 Al7.5 and Zr57 Ti5 Cu20 Ni8 Al10 alloys that show the best glass-forming ability. The ratio between the total volume fractions of binary and ternary crystal phases is close to 3:2 in these alloys [11]. A similar ratio of ZrNiTi ternary Laves phase to binary phases was found to be favorable for a sharp increase of the amorphous layer thickness in rapid-quenched alloys of Zr–Ni–Ti [12]. In order to study this relation, Zr65 Cu17.5 Ni10 Al7.5 and Zr57 Ti5 Cu20 Ni8 Al10 alloys were prepared by the MA method from a mixture of phases taken in the proportions determined from the results of the X-ray quantitative analysis of the annealed samples. These proportions are listed in Table 1. It should be noted that the actual alloy compositions deviate slightly from those of the above-mentioned multicomponent bulk amorphous alloys. For example, it is impossible to obtain the composition Zr65 Cu17.5 Ni10 Al7.5 from a mixture of the phases Zr6 NiAl2 and Zr2 Cu. Therefore, in Table 1 the actual composition obtained for such mixtures is also indicated. 3.1. Stability of different phases during mechanical milling Before the mechanical alloying of the mixtures, a series of experiments was carried out in order to study the influence of mechanical milling on each phase. The Zr2 Ni phase (structure type CuAl2 or C16) was also milled. After mechanical milling all phases showed complete or partial amorphization. The dependence of the amorphous phase content on the milling time for each intermetallic is presented in Fig. 1.
3. Results and discussion It has been found that the following compounds: Zr2 Cu (structure type MoSi2 ), Zr2 Ni (structure type CuAl2 ), ZrNiTi (Laves phase with structure type MgZn2 ) and Zr6 NiAl2 (structure type Zr6 CoAl2 ), are formed in the an-
Table 1 Alloy compositions and proportions of glass-forming compounds in the initial mixtures Phase composition
Volume fraction (%)
Actual composition
Zr65 Cu17.5 Ni10 Al7.5
Zr6 NiAl2 Zr2 Cu
35 65
Zr66.7 Cu22 Ni3.8 Al7.5
Zr57 Ti5 Cu20 Ni8 Al10
ZrNiTi Zr6 NiAl2 Zr2 Cu
25 23 52
Zr59 Ti7.7 Cu17.2 Ni10.5 Al5.6
amorphous phase content, %
Target composition
100 90 80 70 60 50
Zr2Ni
40 30
Zr6NiAl2
20 10 0
ZrNiTi
Zr2Cu
0
60
120
180
240
300
360
420
480
t, min Fig. 1. Dependence of the amorphous phase content on the milling time for some intermetallic compounds.
N.P. Djakonova et al. / Journal of Alloys and Compounds 367 (2004) 191–198
193
amorphous phase content, %
100 80 60 40
mixture of phases mixture of pure elements
20 0 0
60
120
180
240 t, min
300
360
420
480
Fig. 2. Dependence of the amorphous phase content in the alloy Zr57 Ti5 Cu20 Ni8 Al10 on the milling time for different starting materials: mixtures of glass-forming compounds and mixtures of elemental powders.
The highest rate of transition into an AP was observed for the Zr2 Ni phase, and the lowest rate for the Zr6 NiAl2 phase, which was not completely amorphized even after 8 h of milling. The tendency of these phases for amorphization makes it possible to relate them to the so-called glass-forming compounds [13]. The glass-forming compounds are believed to be phases that have a complex crystal structure, frequently with a large fraction of atoms in local icosahedral coordination. The chemical composition of these phases is close to that of the compounds that are prone to transitions into an amorphous state by MQ and satisfy a number of thermodynamic and empirical glass formation criteria [14–18]. Glass-forming compounds frequently have crystal structures with topological and chemical short-range order similar to that of the AP. When the melt is rapidly cooled, the existence of clusters with GFC-like atomic configurations is favorable to melt undercooling and the formation of metal glasses [19]. The ability of GFC to amorphize readily by mechanical milling can be due to several reasons. First of all, a diffusion mass-transfer to form the amorphous phase is not required, as the components are already mixed at the atomic level. Secondly, if
in the crystal state the phase has the chemical and topological short-range order characteristic of the AP, then the amorphization only needs the long-range order to be destroyed, which can be done by introducing crystal lattice defects. Let us consider two of the glass formation criteria: the free energy of the compound being close to that of the amorphous phase [20] and the atomic size misfit factor [14]. If we neglect the entropy contribution, then we can accept as the first criterion the difference in formation enthalpy be-
2 1
30 (a)
35
40
45
50
55
60
2θ, CoK
Table 2 Formation enthalpy of amorphous and crystalline phases and atomic radii ratioa
1 2
Phase
Hform (kJ/mol)
Ham (kJ/mol)
Ham − Hform (kJ/mol)
RA /RB
Zr2 Ni Zr2 Cu Zr6 NiAl2 ZrNiTi
−48.6 −23 −57.5 −34.3
−44.9 −21.2 −52.4 −30.1
3.7 1.8 5.1 4.2
1.28 1.25 1.17 1.18
a Formation enthalpy of binary compounds and amorphous phases is estimated by Miedema’s method [21]. For ternary phases the evaluation is done by enthalpy averaging over two hypothetical binary compounds: Zr6 NiAl2 = (2/3)(Zr 6 Al3 ) + (1/3)(Zr 6 Ni3 ), ZrNiTi = 1/2(ZrNi2 ) + 1/2(ZrTi2 ). The atomic radii ratios in the ternary phases are assumed to be RZr /((1/3)RNi + (2/3)RAl ) and RZr /((1/2)RNi + (1/2)RTi ) for the compounds Zr6 NiAl2 and ZrNiTi, respectively.
30 (b)
35
40
45
50
55
60
2θ, CoK
Fig. 3. X-ray diffraction patterns of (a) a Zr65 Cu17.5 Ni10 Al7.5 alloy prepared by MA of EP (1) and MQ (2) and (b) a Zr57 Ti5 Cu20 Ni8 Al10 alloy prepared by MA from GFC (1) and EP (2) mixtures.
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Table 3 Results of the calorimetric analysis for Zr-based alloys (heating rate 20 K/min) T = Tx − Tg (K)
Alloy
Method of preparation
Tg (K)
Tx (K)
Zr65 Cu17.5 Ni10 Al7.5
MA of GFC mixture MA of elemental powder mixture MQ
650 638 649
746 694 706
96 56 57
Zr57 Ti5 Cu20 Ni8 Al10
MA of GFC mixture MA of elemental powder mixture
672 647
794 711
122 64
tween the amorphous phase and the compound, Ham −Hform . Table 2 lists the corresponding data for the phases under study. It is obvious that the smaller the difference Ham − Hform , the easier the transformation of the compound into an amorphous state. As far as this criterion is concerned, the Zr6 NiAl2 phase should be the most stable with respect to amorphization, just as was observed experimentally. The atomic size factor also explains the higher relative stability of the Zr6 NiAl2 phase.
as starting materials a mixture of elemental powders and a mixture of GFC were compared. Fig. 2 illustrates the dependencies of the AP content on milling time in the Zr57 Ti5 Cu20 Ni8 Al10 alloy for the different starting materials. Similar dependencies were observed for the Zr65 Cu17.5 Ni10 Al7.5 alloy. It is quite obvious that the GFC mixture amorphizes significantly faster than the EP mixture. In the case of GFC, the amorphization was practically complete after 2 h of MA, while in the EP mixture the same AP fraction was obtained only after 8 h of MA. A comparison of the amorphous phases produced by different methods was carried out. Fig. 3 shows diffraction patterns of the alloys under study. As one can see, the width
3.2. Mechanical alloying of GFC mixtures Mechanical alloying of a GFC produces an amorphous phase. For each alloy, the kinetics for amorphization using
*
*
Zr
Zr6NiAl2
Zr2Cu
Zr2Ni (E93)
Zr2(Ni,Cu) (C16)
*
673 K
*
*
*
*
753 K
893 K
1073 K
30
40
50
60
70
80
90
100
2θ, CoK α Fig. 4. X-ray diffraction patterns of a Zr65 Cu17.5 Ni10 Al7.5 alloy produced by MA of an EP mixture after isothermal annealing at different temperatures.
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⊗ ZrCu2Al
+ ZrO2 (cub) Zr2Cu
+
+
195
Zr2Ni (E93)
Zr6NiAl2
Zr2(Ni,Cu) (C16)
ZrO2 (mon)
673 K
+
+
753 K
893 K
⊗
⊗
943 K
⊗ ⊗
30
40
50
60
70
80
90
100
2θ, CoK α Fig. 5. X-ray diffraction patterns of a Zr65 Cu17.5 Ni10 Al7.5 alloy produced by MQ after isothermal annealing at different temperatures.
phase with Ti2Ni structure type 673 K
793 K
893 K
1073 K
30
40
50
60
70
80
90
100
2θ, CoKα Fig. 6. X-ray diffraction patterns of a Zr57 Ti5 Cu20 Ni8 Al10 alloy produced by MA of a GFC mixture after isothermal annealing at different temperatures.
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* *
ZrNiTi
Zr
Zr6NiAl2
Zr2Cu
Zr2Ni (E93)
673 K
* *
*
*
*
*
753 K
793 K
893 K
1073 K
30
40
50
60 70 2θ, CoK α
80
90
100
Fig. 7. X-ray diffraction patterns of a Zr57 Ti5 Cu20 Ni8 Al10 alloy produced by MA of an EP mixture after isothermal annealing at different temperatures.
Table 4 Phase composition of Zr65 Cu17.5 Ni10 Al7.5 alloys after isothermal annealing at different temperatures Annealing temperature (K) 673 753
Annealing time (min)
30 30
Mixture of EP Phases
Volume fraction (%)
Phases
Volume fraction (%)
am Zr
99 1
am ZrO2
Traces
Zr2 Cu Zr6 NiAl2 Zr2 Ni (E93 )
50 20 25
Zr2 Cu Zr6 NiAl2 Zr2 Ni (E93 ) Zr2 (Cu, Ni) ZrO2
30 10 20 35 5
Zr2 Cu Zr6 NiAl2
70 30
X
Traces
Zr2 Cu Zr6 NiAl2 ZrO2 ZrCu2 Al (D03 ) X
40 20 35 5 Traces
–
–
Zr Zr2 Cu Zr6 NiAl2 Zr2 Ni (E93 ) Zr2 (Ni, Cu) (C16) Zr
893
30
943
120
–
In calorimeter (heating rate 20 K/min)
Zr2 Cu Zr6 NiAl2 Zr2 Ni (E93 ) Zr2 (Ni, Cu) (C16)
1073
MQ ribbon
5 65 20 5 10 Traces –
60 20 10 10
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of the amorphous peak is almost the same for the MQ samples and for the samples produced by MA from EP mixtures (the full width at half maximum, FWHM, is about 10◦ in 2θ), while the width of the amorphous peak obtained for the GFC mixture is noticeably larger (FWHM is about 15◦ in 2θ). This most likely means that the amorphous phase produced by the GFC mixture milling differs in its atomic structure from those obtained by either MQ or mechanical alloying of EP mixtures. Differences between the amorphous phases are also revealed in the values of the Tg and Tx temperatures that are listed in Table 3. For the amorphous phase produced from the mixture of phases T = Tx − Tg ≈ 100 ◦ C, whereas for the other two preparation methods T is approximately the same and equals ∼60 ◦ C. It is worth noticing that the entire difference between the T values is related to the higher crystallization temperature, since the temperatures of the glass transition are similar for all preparation methods. 3.3. Isothermal annealing We also compared the phase transition processes that occur on annealing the amorphous phases produced by different methods. The corresponding diffraction patterns of the Zr65 Cu17.5 Ni10 Al7.5 and Zr57 Ti5 Cu20 Ni8 Al10 alloys are presented in Figs. 4–7 and the phase composition of the alloys is listed in Tables 4 and 5. The Zr65 Cu17.5 Ni10 Al7.5 alloy prepared by MA from the EP mixture remains amorphous up to T = 673 K. Its crystallization occurs in a single step with the formation of three phases at 753 K (see Fig. 4). Besides the equilibrium Zr2 Cu and Zr6 NiAl2 phases, the sample reveals a small quantity of Zr2 Ni phase (structure type Ti2 Ni or E93 ), which is ab-
197
sent in the phase diagram. A further rise of the annealing temperature leads to a decrease of this phase fraction since the phase is metastable. As a result of its decomposition the equilibrium Zr2 Ni phase (structure type CuAl2 or C16), possibly enriched by Cu, appears in the sample. The crystallization of the MQ ribbon of the same composition occurs in a similar manner (Fig. 5). However, the complete disappearance of the metastable Zr2 Ni phase (type E93 ) is observed at temperatures above 893 K, where a negligible amount of another phase is detected. This phase was impossible to identify due to its small quantity. For the Zr57 Ti5 Cu20 Ni8 Al10 alloy, the amorphous states obtained by the MA method from EP and GFC mixtures were compared. The amorphous phase from the GFC mixture has a much higher thermal stability: its crystallization occurs at a temperature above 793 K, with the formation of a single Ti2 Ni type phase with lattice parameter a = 1.205 nm (Fig. 6). It is worth noting that this phase remained even after heating the calorimeter up to 1073 K with a heating rate of 20 K/min. The AP produced from the EP mixture is stable up to 673 K and during its crystallization a mixture of the equilibrium Zr2 Cu and Zr6 NiAl2 phases is formed at first and then (at 793 K) two more phases appear: the equilibrium ZrNiTi phase and a Ti2 Ni-type phase (Fig. 7). The latter is stable up to an annealing temperature of about 900 K and after heating the calorimeter. Thus, the comparison of amorphous phases obtained by different methods has revealed similarities between those obtained by MQ and by MA from EP mixtures, but significant differences with respect to the AP produced by MA from GFC mixtures. The latter has a much larger width of the amorphous peak and a noticeably different crystallization temperature, and different phase transitions occur upon isothermal annealing.
Table 5 Phase composition of Zr57 Ti5 Cu20 Ni8 Al10 alloys after isothermal annealing at different temperatures Annealing temperature (K)
Annealing time (min)
Mixture of GFC Phases
Volume fraction (%)
Phases
Volume fraction (%)
673
30
am
100
am
100
753
30
am
100
Zr2 Cu Zr6 NiAl2
70 30
793
30
am
100
Zr2 Cu Zr6 NiAl2 Zr2 Ni (E93 )
60 20 20
893
30
Zr2 Ni (E93 )
100
Zr2 Cu Zr6 NiAl2 Zr2 Ni (E93 ) ZrNiTi (C14)
50 15 30 5
Zr2 Cu
60
Zr6 NiAl2 Zr2 Ni (E93 ) ZrNiTi (C14)
5 30 5
1073
Mixture of EP
In calorimeter (heating rate 20 K/min) Zr2 Ni (E93 ) ZrO2
97 3
198
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Some of the peculiarities described above for the Zr-based amorphous phase obtained from the GFC mixture were also revealed in the Fe-based alloy Fe61 Co7 Zr10 Mo5 W2 B15 . Thus the amorphization rate of the GFC mixture on MA is much higher than that of the EP mixture, for which the production of the AP is impossible. The amorphous peak of the AP produced by MA of the GFC mixture is noticeably wider than that of the AP obtained by MQ. However, the differential thermal analysis did not reveal any detectable difference in either the crystallization or the melting process of the two AP formed by various methods [8,22].
4. Conclusions The feasibility of amorphous phase formation in multicomponent Zr65 Cu17.5 Ni10 Al7.5 and Zr57 Ti5 Cu20 Ni8 Al10 bulk amorphous alloys by the MA method starting from mixtures of elemental powders and mixtures of glass-forming compounds has been studied. The GFC were selected among the intermetallic compounds that satisfy a certain number of thermodynamic and empirical criteria for glass formation. The molar fractions of the GFC were chosen to fit the composition of the specified alloy in the equilibrium state. Prior to that, the crystal structure stability of the selected GFC during mechanical milling was investigated and compared with some crystallochemical characteristics of the phases. It was found that using GFC instead of a mixture of elemental powders significantly speeds up the transformation into the amorphous state during mechanical alloying. The phase transformations occurring during the crystallization of the amorphous phase produced by mechanical alloying of the GFC mixture differ considerably from those taking place in the amorphous phases obtained by melt quenching and mechanical alloying of the mixture of elemental powders. Moreover, some of the thermal characteristics of these amorphous phases differ as well. The values of both Tx and T for the amorphous powder produced by mechanical alloying of the GFC mixture exceed those of the amorphous phase obtained by MQ and by MA of mixtures of elemental powders, whereas Tg is nearly the same in all al-
loys of the same composition, irrespective of the production method. Acknowledgements This work was partially supported by the Russian Foundation for Basic Research (RFBR 01-03-32986). References [1] A. Inoue, T. Zhang, T. Masumoto, Mater. Trans. JIM 31 (1990) 425. [2] A. Inoue, Acta Mater. 48 (2000) 279. [3] A. Inoue, T. Zhang, A. Takeuchi, Mater. Sci. Forum 269–272 (1998) 855. [4] H. Kato, A. Kawamura, A. Inoue, Mater. Trans. JIM 37 (1996) 70. [5] M. Siedel, J. Eckert, H.-D. Bauer, L. Schultz, Mater. Sci. Forum 225–227 (1996) 119. [6] M.I. Petrzhik, T.A. Sviridova, V.V. Molokanov, J. Shen, J. Sun, in: Abs. Conf. Bulk Metallic Glasses II, Bulk Amorphous Alloys, Taiwan, 2002. [7] V.V. Molokanov, M.I. Petrzhik, K.S. Filippov, T.A. Sviridova, A. Castellero, M. Baricco, L. Battezzati, Materialoved. 1 (2002) 42. [8] N.P. Djakonova, T.A. Sviridova, E.A. Zakharova, V.V. Molokanov, M.I. Petrzhik, J. Metastable Nanocryst. Mater. 15–16 (2003) 673. [9] E.V. Shelekhov, in: Proceedings of the National Conference on Applications of X-Ray, Synchrotron Radiation, Neutrons and Electrons for Material Investigations, vol. 3, 1997, p. 316 (in Russian). [10] E.V. Shelekhov, T.A. Sviridova, Yu.A. Skakov, N.P. Djakonova, J. Metastable Nanocryst. Mater. 8 (2000) 615. [11] V.V. Molokanov, M.I. Petrzhik, T.N. Mikhailova, T.A. Sviridova, N.P. Djakonova, J. Non-Cryst. Solids 250–252 (1999) 560. [12] V.N. Chebotnikov, V.V. Molokanov, E.B. Rubina, Yu.K. Kovneristyi, Fiz. Met. Metalloved. 68 (1989) 964. [13] V.V. Molokanov, V.N. Chebotnikov, Key Eng. Mater. 40–41 (1990) 319. [14] P. Ramachandrarao, Z. Metallkd. 71 (1980) 172. [15] S.R. Nagel, J. Tauc, Solid State Commun. 22 (1977) 129. [16] F. Sommer, in: Proceedings of the Fourth International Conference on Rapidly Quenched Metals, 1982, p. 209. [17] A.R. Yavari, P. Hicter, P. Desre, J. Chim. Phys. 79 (1982) 572. [18] A.R. Yavari, J.L. Uriarte, A. Inoue, Mater. Sci. Forum 269–272 (1998) 533. [19] M.I. Petrzhik, V.V. Molokanov, Izv. Rus. Akad. Nauk, Ser. Fiz. 65 (2001) 1384. [20] L. Schultz, Mater. Sci. Eng. A 134 (1991) 1004. [21] H. Bakker, Enthalpies in Alloys: Miedema’s Semi-Empirical, Materials Science Foundation, vol. 1, 1998. [22] N.P. Djakonova, V.V. Molokanov, M.I. Petrzhik, T.A. Sviridova, E.A. Zakharova, Perspect. Mater. 5 (2002) 46.
On the synthesis of Zr-based bulk amorphous alloys from glass-forming compounds and elemental powders N.P. Djakonova a , T.A. Sviridova a,∗ , E.A. Zakharova a , V.V. Molokanov b , M.I. Petrzhik b b
a Moscow State Institute of Steel and Alloys (Technological University), Leninsky pr. 4, Moscow 117936, Russia Institute of Metallurgy and Materials Science, Russian Academy of Sciences, Leninsky pr. 49, Moscow 117911, Russia
Abstract The possibility of producing Zr-based bulk amorphous alloys (Zr65 Cu17.5 Ni10 Al7.5 and Zr57 Ti5 Cu20 Ni8 Al10 ) by mechanical alloying (MA) from either a mixture of elemental powders (EP), or a mixture of glass-forming compounds (GFC), has been investigated. A preliminary study of the crystal structure stability of the GFC-ingredients during mechanical milling was performed. Amorphization was found to be significantly enhanced during MA by using a GFC mixture. A considerable difference in the phase transformations observed on heating the amorphous phase (AP) was noted for alloys obtained from GFC, as compared with those produced by melt quenching (MQ) or milling of a mixture of elemental powders. It is shown that the amorphous phase obtained by MA from GFC has a wider undercooled liquid region between the glass transition temperature Tg and the crystallization temperature Tx than the amorphous phase produced by the conventional MQ technique. © 2003 Published by Elsevier B.V. Keywords: Amorphous materials; Amorphization; Liquid quenching; Mechanical alloying; X-ray diffraction
1. Introduction Recently discovered multicomponent bulk amorphous alloys have good glass-forming ability and can easily be cast into bulk glassy specimens with dimensions of several centimeters at a cooling rate of 1/100 K/s [1–3]. These alloys exhibit an extended undercooled liquid region between the glass transition temperature Tg and the crystallization temperature Tx , allowing easy deformation and shaping of the material by hot-pressing, due to the viscous flow in the undercooled liquid [4,5]. However, melt quenching (MQ) does not guarantee the formation of an amorphous single-phase state. For example, in the Fe base of an Fe61 Co7 Zr10 Mo5 W2 B15 alloy, the refractory phase ZrB2 is often observed [6,7]. The formation of ZrB2 particles leads to a deviation from the near-eutectic composition, which is favorable for obtaining the amorphous state. Thus, mechanical alloying (MA) seems to be a promising method for the preparation of this alloy. We have recently established that the process of amorphous phase (AP) formation during MA depends consider∗
Corresponding author. E-mail address: [email protected] (T.A. Sviridova).
0925-8388/$ – see front matter © 2003 Published by Elsevier B.V. doi:10.1016/j.jallcom.2003.08.036
ably on the characteristics of the starting materials [8]. If the MA synthesis is carried out for a mixture of specially selected intermetallic glass-forming compounds (GFC), then the solid-state amorphization reaction proceeds much faster than for the mixture of elemental powders (EP). Moreover, the characteristics of the amorphous phase—Tg , Tx and T = Tx − Tg —are frequently higher for MA, as compared to those of the corresponding MQ alloy. The purpose of this paper is to analyze and compare stability characteristics of the well-known multicomponent bulk amorphous alloys Zr65 Cu17.5 Ni10 Al7.5 and Zr57 Ti5 Cu20 Ni8 Al10 , produced by conventional MQ and MA techniques from glass-forming compounds and elemental powder mixtures.
2. Experimental details Milling was carried out for 15/480 min in a high-energy planetary ball mill of AGO-2U type. Low-carbon steel vials and balls of a ball-bearing steel with diameters of 6 and 8 mm were used. The ball-to-powder ratio was 10:1, the plate rotational speed 680 rpm. The vials were charged and sealed in Ar atmosphere. Both elementary powders (with 99.5% or better purity) and intermetallic compounds were used as
192
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initial components for MA. Ingots of the compounds were melted in high-purity argon atmosphere in an arc furnace with a tungsten electrode and a water-cooled hearth and were then ground into powders with a particle size smaller than 0.1 mm. X-ray diffraction patterns were taken with Co K␣ radiation using a monochromated diffraction beam and a step-scanning diffractometer DRON-4. For the analysis of the samples with amorphous phase, a modification of the Rietveld full-profile fitting method was employed, with the experimental X-ray pattern of the pure amorphous phase reference rescaled into theoretical units of intensity [9,10]. The thermal stability and crystallization of the amorphous phase were investigated on heating in a differential scanning calorimeter (DSC-111 SETARAM microcalorimeter). The melting and solidification temperatures were determined by differential thermal analysis (DTA-7 HT thermal analyzer). The samples were annealed in a sealed quartz ampoule evacuated down to 10−3 Torr. In order to prevent oxidation, the powders were wrapped in a double-layer titanium foil with zirconium powder, which also served to keep track of any oxidation.
nealed Zr65 Cu17.5 Ni10 Al7.5 and Zr57 Ti5 Cu20 Ni8 Al10 alloys that show the best glass-forming ability. The ratio between the total volume fractions of binary and ternary crystal phases is close to 3:2 in these alloys [11]. A similar ratio of ZrNiTi ternary Laves phase to binary phases was found to be favorable for a sharp increase of the amorphous layer thickness in rapid-quenched alloys of Zr–Ni–Ti [12]. In order to study this relation, Zr65 Cu17.5 Ni10 Al7.5 and Zr57 Ti5 Cu20 Ni8 Al10 alloys were prepared by the MA method from a mixture of phases taken in the proportions determined from the results of the X-ray quantitative analysis of the annealed samples. These proportions are listed in Table 1. It should be noted that the actual alloy compositions deviate slightly from those of the above-mentioned multicomponent bulk amorphous alloys. For example, it is impossible to obtain the composition Zr65 Cu17.5 Ni10 Al7.5 from a mixture of the phases Zr6 NiAl2 and Zr2 Cu. Therefore, in Table 1 the actual composition obtained for such mixtures is also indicated. 3.1. Stability of different phases during mechanical milling Before the mechanical alloying of the mixtures, a series of experiments was carried out in order to study the influence of mechanical milling on each phase. The Zr2 Ni phase (structure type CuAl2 or C16) was also milled. After mechanical milling all phases showed complete or partial amorphization. The dependence of the amorphous phase content on the milling time for each intermetallic is presented in Fig. 1.
3. Results and discussion It has been found that the following compounds: Zr2 Cu (structure type MoSi2 ), Zr2 Ni (structure type CuAl2 ), ZrNiTi (Laves phase with structure type MgZn2 ) and Zr6 NiAl2 (structure type Zr6 CoAl2 ), are formed in the an-
Table 1 Alloy compositions and proportions of glass-forming compounds in the initial mixtures Phase composition
Volume fraction (%)
Actual composition
Zr65 Cu17.5 Ni10 Al7.5
Zr6 NiAl2 Zr2 Cu
35 65
Zr66.7 Cu22 Ni3.8 Al7.5
Zr57 Ti5 Cu20 Ni8 Al10
ZrNiTi Zr6 NiAl2 Zr2 Cu
25 23 52
Zr59 Ti7.7 Cu17.2 Ni10.5 Al5.6
amorphous phase content, %
Target composition
100 90 80 70 60 50
Zr2Ni
40 30
Zr6NiAl2
20 10 0
ZrNiTi
Zr2Cu
0
60
120
180
240
300
360
420
480
t, min Fig. 1. Dependence of the amorphous phase content on the milling time for some intermetallic compounds.
N.P. Djakonova et al. / Journal of Alloys and Compounds 367 (2004) 191–198
193
amorphous phase content, %
100 80 60 40
mixture of phases mixture of pure elements
20 0 0
60
120
180
240 t, min
300
360
420
480
Fig. 2. Dependence of the amorphous phase content in the alloy Zr57 Ti5 Cu20 Ni8 Al10 on the milling time for different starting materials: mixtures of glass-forming compounds and mixtures of elemental powders.
The highest rate of transition into an AP was observed for the Zr2 Ni phase, and the lowest rate for the Zr6 NiAl2 phase, which was not completely amorphized even after 8 h of milling. The tendency of these phases for amorphization makes it possible to relate them to the so-called glass-forming compounds [13]. The glass-forming compounds are believed to be phases that have a complex crystal structure, frequently with a large fraction of atoms in local icosahedral coordination. The chemical composition of these phases is close to that of the compounds that are prone to transitions into an amorphous state by MQ and satisfy a number of thermodynamic and empirical glass formation criteria [14–18]. Glass-forming compounds frequently have crystal structures with topological and chemical short-range order similar to that of the AP. When the melt is rapidly cooled, the existence of clusters with GFC-like atomic configurations is favorable to melt undercooling and the formation of metal glasses [19]. The ability of GFC to amorphize readily by mechanical milling can be due to several reasons. First of all, a diffusion mass-transfer to form the amorphous phase is not required, as the components are already mixed at the atomic level. Secondly, if
in the crystal state the phase has the chemical and topological short-range order characteristic of the AP, then the amorphization only needs the long-range order to be destroyed, which can be done by introducing crystal lattice defects. Let us consider two of the glass formation criteria: the free energy of the compound being close to that of the amorphous phase [20] and the atomic size misfit factor [14]. If we neglect the entropy contribution, then we can accept as the first criterion the difference in formation enthalpy be-
2 1
30 (a)
35
40
45
50
55
60
2θ, CoK
Table 2 Formation enthalpy of amorphous and crystalline phases and atomic radii ratioa
1 2
Phase
Hform (kJ/mol)
Ham (kJ/mol)
Ham − Hform (kJ/mol)
RA /RB
Zr2 Ni Zr2 Cu Zr6 NiAl2 ZrNiTi
−48.6 −23 −57.5 −34.3
−44.9 −21.2 −52.4 −30.1
3.7 1.8 5.1 4.2
1.28 1.25 1.17 1.18
a Formation enthalpy of binary compounds and amorphous phases is estimated by Miedema’s method [21]. For ternary phases the evaluation is done by enthalpy averaging over two hypothetical binary compounds: Zr6 NiAl2 = (2/3)(Zr 6 Al3 ) + (1/3)(Zr 6 Ni3 ), ZrNiTi = 1/2(ZrNi2 ) + 1/2(ZrTi2 ). The atomic radii ratios in the ternary phases are assumed to be RZr /((1/3)RNi + (2/3)RAl ) and RZr /((1/2)RNi + (1/2)RTi ) for the compounds Zr6 NiAl2 and ZrNiTi, respectively.
30 (b)
35
40
45
50
55
60
2θ, CoK
Fig. 3. X-ray diffraction patterns of (a) a Zr65 Cu17.5 Ni10 Al7.5 alloy prepared by MA of EP (1) and MQ (2) and (b) a Zr57 Ti5 Cu20 Ni8 Al10 alloy prepared by MA from GFC (1) and EP (2) mixtures.
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Table 3 Results of the calorimetric analysis for Zr-based alloys (heating rate 20 K/min) T = Tx − Tg (K)
Alloy
Method of preparation
Tg (K)
Tx (K)
Zr65 Cu17.5 Ni10 Al7.5
MA of GFC mixture MA of elemental powder mixture MQ
650 638 649
746 694 706
96 56 57
Zr57 Ti5 Cu20 Ni8 Al10
MA of GFC mixture MA of elemental powder mixture
672 647
794 711
122 64
tween the amorphous phase and the compound, Ham −Hform . Table 2 lists the corresponding data for the phases under study. It is obvious that the smaller the difference Ham − Hform , the easier the transformation of the compound into an amorphous state. As far as this criterion is concerned, the Zr6 NiAl2 phase should be the most stable with respect to amorphization, just as was observed experimentally. The atomic size factor also explains the higher relative stability of the Zr6 NiAl2 phase.
as starting materials a mixture of elemental powders and a mixture of GFC were compared. Fig. 2 illustrates the dependencies of the AP content on milling time in the Zr57 Ti5 Cu20 Ni8 Al10 alloy for the different starting materials. Similar dependencies were observed for the Zr65 Cu17.5 Ni10 Al7.5 alloy. It is quite obvious that the GFC mixture amorphizes significantly faster than the EP mixture. In the case of GFC, the amorphization was practically complete after 2 h of MA, while in the EP mixture the same AP fraction was obtained only after 8 h of MA. A comparison of the amorphous phases produced by different methods was carried out. Fig. 3 shows diffraction patterns of the alloys under study. As one can see, the width
3.2. Mechanical alloying of GFC mixtures Mechanical alloying of a GFC produces an amorphous phase. For each alloy, the kinetics for amorphization using
*
*
Zr
Zr6NiAl2
Zr2Cu
Zr2Ni (E93)
Zr2(Ni,Cu) (C16)
*
673 K
*
*
*
*
753 K
893 K
1073 K
30
40
50
60
70
80
90
100
2θ, CoK α Fig. 4. X-ray diffraction patterns of a Zr65 Cu17.5 Ni10 Al7.5 alloy produced by MA of an EP mixture after isothermal annealing at different temperatures.
N.P. Djakonova et al. / Journal of Alloys and Compounds 367 (2004) 191–198
⊗ ZrCu2Al
+ ZrO2 (cub) Zr2Cu
+
+
195
Zr2Ni (E93)
Zr6NiAl2
Zr2(Ni,Cu) (C16)
ZrO2 (mon)
673 K
+
+
753 K
893 K
⊗
⊗
943 K
⊗ ⊗
30
40
50
60
70
80
90
100
2θ, CoK α Fig. 5. X-ray diffraction patterns of a Zr65 Cu17.5 Ni10 Al7.5 alloy produced by MQ after isothermal annealing at different temperatures.
phase with Ti2Ni structure type 673 K
793 K
893 K
1073 K
30
40
50
60
70
80
90
100
2θ, CoKα Fig. 6. X-ray diffraction patterns of a Zr57 Ti5 Cu20 Ni8 Al10 alloy produced by MA of a GFC mixture after isothermal annealing at different temperatures.
196
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* *
ZrNiTi
Zr
Zr6NiAl2
Zr2Cu
Zr2Ni (E93)
673 K
* *
*
*
*
*
753 K
793 K
893 K
1073 K
30
40
50
60 70 2θ, CoK α
80
90
100
Fig. 7. X-ray diffraction patterns of a Zr57 Ti5 Cu20 Ni8 Al10 alloy produced by MA of an EP mixture after isothermal annealing at different temperatures.
Table 4 Phase composition of Zr65 Cu17.5 Ni10 Al7.5 alloys after isothermal annealing at different temperatures Annealing temperature (K) 673 753
Annealing time (min)
30 30
Mixture of EP Phases
Volume fraction (%)
Phases
Volume fraction (%)
am Zr
99 1
am ZrO2
Traces
Zr2 Cu Zr6 NiAl2 Zr2 Ni (E93 )
50 20 25
Zr2 Cu Zr6 NiAl2 Zr2 Ni (E93 ) Zr2 (Cu, Ni) ZrO2
30 10 20 35 5
Zr2 Cu Zr6 NiAl2
70 30
X
Traces
Zr2 Cu Zr6 NiAl2 ZrO2 ZrCu2 Al (D03 ) X
40 20 35 5 Traces
–
–
Zr Zr2 Cu Zr6 NiAl2 Zr2 Ni (E93 ) Zr2 (Ni, Cu) (C16) Zr
893
30
943
120
–
In calorimeter (heating rate 20 K/min)
Zr2 Cu Zr6 NiAl2 Zr2 Ni (E93 ) Zr2 (Ni, Cu) (C16)
1073
MQ ribbon
5 65 20 5 10 Traces –
60 20 10 10
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of the amorphous peak is almost the same for the MQ samples and for the samples produced by MA from EP mixtures (the full width at half maximum, FWHM, is about 10◦ in 2θ), while the width of the amorphous peak obtained for the GFC mixture is noticeably larger (FWHM is about 15◦ in 2θ). This most likely means that the amorphous phase produced by the GFC mixture milling differs in its atomic structure from those obtained by either MQ or mechanical alloying of EP mixtures. Differences between the amorphous phases are also revealed in the values of the Tg and Tx temperatures that are listed in Table 3. For the amorphous phase produced from the mixture of phases T = Tx − Tg ≈ 100 ◦ C, whereas for the other two preparation methods T is approximately the same and equals ∼60 ◦ C. It is worth noticing that the entire difference between the T values is related to the higher crystallization temperature, since the temperatures of the glass transition are similar for all preparation methods. 3.3. Isothermal annealing We also compared the phase transition processes that occur on annealing the amorphous phases produced by different methods. The corresponding diffraction patterns of the Zr65 Cu17.5 Ni10 Al7.5 and Zr57 Ti5 Cu20 Ni8 Al10 alloys are presented in Figs. 4–7 and the phase composition of the alloys is listed in Tables 4 and 5. The Zr65 Cu17.5 Ni10 Al7.5 alloy prepared by MA from the EP mixture remains amorphous up to T = 673 K. Its crystallization occurs in a single step with the formation of three phases at 753 K (see Fig. 4). Besides the equilibrium Zr2 Cu and Zr6 NiAl2 phases, the sample reveals a small quantity of Zr2 Ni phase (structure type Ti2 Ni or E93 ), which is ab-
197
sent in the phase diagram. A further rise of the annealing temperature leads to a decrease of this phase fraction since the phase is metastable. As a result of its decomposition the equilibrium Zr2 Ni phase (structure type CuAl2 or C16), possibly enriched by Cu, appears in the sample. The crystallization of the MQ ribbon of the same composition occurs in a similar manner (Fig. 5). However, the complete disappearance of the metastable Zr2 Ni phase (type E93 ) is observed at temperatures above 893 K, where a negligible amount of another phase is detected. This phase was impossible to identify due to its small quantity. For the Zr57 Ti5 Cu20 Ni8 Al10 alloy, the amorphous states obtained by the MA method from EP and GFC mixtures were compared. The amorphous phase from the GFC mixture has a much higher thermal stability: its crystallization occurs at a temperature above 793 K, with the formation of a single Ti2 Ni type phase with lattice parameter a = 1.205 nm (Fig. 6). It is worth noting that this phase remained even after heating the calorimeter up to 1073 K with a heating rate of 20 K/min. The AP produced from the EP mixture is stable up to 673 K and during its crystallization a mixture of the equilibrium Zr2 Cu and Zr6 NiAl2 phases is formed at first and then (at 793 K) two more phases appear: the equilibrium ZrNiTi phase and a Ti2 Ni-type phase (Fig. 7). The latter is stable up to an annealing temperature of about 900 K and after heating the calorimeter. Thus, the comparison of amorphous phases obtained by different methods has revealed similarities between those obtained by MQ and by MA from EP mixtures, but significant differences with respect to the AP produced by MA from GFC mixtures. The latter has a much larger width of the amorphous peak and a noticeably different crystallization temperature, and different phase transitions occur upon isothermal annealing.
Table 5 Phase composition of Zr57 Ti5 Cu20 Ni8 Al10 alloys after isothermal annealing at different temperatures Annealing temperature (K)
Annealing time (min)
Mixture of GFC Phases
Volume fraction (%)
Phases
Volume fraction (%)
673
30
am
100
am
100
753
30
am
100
Zr2 Cu Zr6 NiAl2
70 30
793
30
am
100
Zr2 Cu Zr6 NiAl2 Zr2 Ni (E93 )
60 20 20
893
30
Zr2 Ni (E93 )
100
Zr2 Cu Zr6 NiAl2 Zr2 Ni (E93 ) ZrNiTi (C14)
50 15 30 5
Zr2 Cu
60
Zr6 NiAl2 Zr2 Ni (E93 ) ZrNiTi (C14)
5 30 5
1073
Mixture of EP
In calorimeter (heating rate 20 K/min) Zr2 Ni (E93 ) ZrO2
97 3
198
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Some of the peculiarities described above for the Zr-based amorphous phase obtained from the GFC mixture were also revealed in the Fe-based alloy Fe61 Co7 Zr10 Mo5 W2 B15 . Thus the amorphization rate of the GFC mixture on MA is much higher than that of the EP mixture, for which the production of the AP is impossible. The amorphous peak of the AP produced by MA of the GFC mixture is noticeably wider than that of the AP obtained by MQ. However, the differential thermal analysis did not reveal any detectable difference in either the crystallization or the melting process of the two AP formed by various methods [8,22].
4. Conclusions The feasibility of amorphous phase formation in multicomponent Zr65 Cu17.5 Ni10 Al7.5 and Zr57 Ti5 Cu20 Ni8 Al10 bulk amorphous alloys by the MA method starting from mixtures of elemental powders and mixtures of glass-forming compounds has been studied. The GFC were selected among the intermetallic compounds that satisfy a certain number of thermodynamic and empirical criteria for glass formation. The molar fractions of the GFC were chosen to fit the composition of the specified alloy in the equilibrium state. Prior to that, the crystal structure stability of the selected GFC during mechanical milling was investigated and compared with some crystallochemical characteristics of the phases. It was found that using GFC instead of a mixture of elemental powders significantly speeds up the transformation into the amorphous state during mechanical alloying. The phase transformations occurring during the crystallization of the amorphous phase produced by mechanical alloying of the GFC mixture differ considerably from those taking place in the amorphous phases obtained by melt quenching and mechanical alloying of the mixture of elemental powders. Moreover, some of the thermal characteristics of these amorphous phases differ as well. The values of both Tx and T for the amorphous powder produced by mechanical alloying of the GFC mixture exceed those of the amorphous phase obtained by MQ and by MA of mixtures of elemental powders, whereas Tg is nearly the same in all al-
loys of the same composition, irrespective of the production method. Acknowledgements This work was partially supported by the Russian Foundation for Basic Research (RFBR 01-03-32986). References [1] A. Inoue, T. Zhang, T. Masumoto, Mater. Trans. JIM 31 (1990) 425. [2] A. Inoue, Acta Mater. 48 (2000) 279. [3] A. Inoue, T. Zhang, A. Takeuchi, Mater. Sci. Forum 269–272 (1998) 855. [4] H. Kato, A. Kawamura, A. Inoue, Mater. Trans. JIM 37 (1996) 70. [5] M. Siedel, J. Eckert, H.-D. Bauer, L. Schultz, Mater. Sci. Forum 225–227 (1996) 119. [6] M.I. Petrzhik, T.A. Sviridova, V.V. Molokanov, J. Shen, J. Sun, in: Abs. Conf. Bulk Metallic Glasses II, Bulk Amorphous Alloys, Taiwan, 2002. [7] V.V. Molokanov, M.I. Petrzhik, K.S. Filippov, T.A. Sviridova, A. Castellero, M. Baricco, L. Battezzati, Materialoved. 1 (2002) 42. [8] N.P. Djakonova, T.A. Sviridova, E.A. Zakharova, V.V. Molokanov, M.I. Petrzhik, J. Metastable Nanocryst. Mater. 15–16 (2003) 673. [9] E.V. Shelekhov, in: Proceedings of the National Conference on Applications of X-Ray, Synchrotron Radiation, Neutrons and Electrons for Material Investigations, vol. 3, 1997, p. 316 (in Russian). [10] E.V. Shelekhov, T.A. Sviridova, Yu.A. Skakov, N.P. Djakonova, J. Metastable Nanocryst. Mater. 8 (2000) 615. [11] V.V. Molokanov, M.I. Petrzhik, T.N. Mikhailova, T.A. Sviridova, N.P. Djakonova, J. Non-Cryst. Solids 250–252 (1999) 560. [12] V.N. Chebotnikov, V.V. Molokanov, E.B. Rubina, Yu.K. Kovneristyi, Fiz. Met. Metalloved. 68 (1989) 964. [13] V.V. Molokanov, V.N. Chebotnikov, Key Eng. Mater. 40–41 (1990) 319. [14] P. Ramachandrarao, Z. Metallkd. 71 (1980) 172. [15] S.R. Nagel, J. Tauc, Solid State Commun. 22 (1977) 129. [16] F. Sommer, in: Proceedings of the Fourth International Conference on Rapidly Quenched Metals, 1982, p. 209. [17] A.R. Yavari, P. Hicter, P. Desre, J. Chim. Phys. 79 (1982) 572. [18] A.R. Yavari, J.L. Uriarte, A. Inoue, Mater. Sci. Forum 269–272 (1998) 533. [19] M.I. Petrzhik, V.V. Molokanov, Izv. Rus. Akad. Nauk, Ser. Fiz. 65 (2001) 1384. [20] L. Schultz, Mater. Sci. Eng. A 134 (1991) 1004. [21] H. Bakker, Enthalpies in Alloys: Miedema’s Semi-Empirical, Materials Science Foundation, vol. 1, 1998. [22] N.P. Djakonova, V.V. Molokanov, M.I. Petrzhik, T.A. Sviridova, E.A. Zakharova, Perspect. Mater. 5 (2002) 46.