Multistage Phase Decomposition Behavior and Kinetics for Zr₆₅Ni₃₀Pd₅ Ternary Amorphous Alloy

Abstract

We have examined heating-induced phase decomposition and its kinetics for Zr₆₅Ni₃₀Pd₅ amorphous alloy in an argon atmosphere. The amorphous phase decomposes through three stages in the order of amorphous → icosahedral quasicrystal (I-Q) + amorphous → approximant cubic Zr₂(Ni,Pd) → tetragonal Zr₂Ni + tetragonal Zr₂Pd phases. The approximant crystalline (APC) Zr₂(Ni,Pd) phase includes icosahedral-like local atomic arrangements and the structural similarity between the I-Q and APC phases seems to enable the multistage phase decomposition through the precipitation of the APC phase. The primary precipitation of the I-Q phase from the amorphous phase occurs through nucleation and growth mechanisms through two-dimensional (2D) growth mode at lower heating rate and 3D growth mode at a high heating rate.

Keywords:
Metallic glass; Approximant crystals; Amorphous Alloy; Quasicrystalline

1. Introduction

The three transition metals Zr, Ni and Pd have high negative heats of mixing with hydrogen and hence the ternary Zr-Ni-Pd amorphous alloys are expected to exhibit a high hydrogen absorption ability.^1-9 On the other hand, the reactivity with oxygen for the ternary alloys is much larger for Zr, while that for Pd is negligible. Thus, the ternary amorphous alloys have attracted an increasing interest as unique materials in the reactivity with gas elements as well as in the dealloying to fabricate a nanoporous Pd(Ni) structure.

Although there are some papers on the oxidation-induced crystallization behavior for the Zr-Ni-Pd ternary amorphous alloys,^1-4 little is known about the crystallization behavior of the ternary amorphous alloys in an inert gas atmosphere. This situation is significantly different from the academic background that there are a number of papers on the crystallization behavior of Zr-Pd binary,^10-17 and Zr-Ni-Pd ternary amorphous alloys.^18-24 The difference is presumably because an icosahedral quasicrystalline (I-Q) phase is formed as a primary precipitation phase in the crystallization process for Zr_65-70Pd_25-30 binary and Zr₇₀Ni₁₀Pd₂₀ ternary amorphous alloys.^10-24 The crystallization of the binary amorphous alloys occurs through two stages in the process of amorphous → I-Q + residual amorphous → equilibrium tetragonal Zr₂Pd phase. The addition of late transition metals (LTM) such as Fe, Co, Ni and Cu to Zr-Pd binary alloys is expected to enhance the amorphous forming ability through the multicomponent effect leading to the atomic size mismatches of Zr > Pd > LTM. Furthermore, the significant atomic size mismatches are thought to be effective for the increase in the thermal stability of metastable phases of amorphous and I-Q phases. The increase in the stability through the sluggish decomposition reaction of their metastable phases may enable the appearance of the approximant crystalline (APC) phase which lies in an intermediate stage in the aspects of atomic configuration and internal energy between I-Q and equilibrium crystalline phases. The searching for a new amorphous alloy composition where the three metastable phases of amorphous, I-Q and APC phases appear in a simple ternary alloy is very meaning for basic science and engineering aspects to develop useful metastable materials. There have been no data on any binary and ternary amorphous alloys with the crystallization process where the three metastable phases appear. The finding is expected to deep the interpretation on the role of the atomic size mismatch effect in the formation tendency of their three metastable phases, though the satisfaction of only the atomic size factor cannot enhance the thermal stability of supercooled liquid which enables the formation of bulk metallic glasses.

The precipitation of I-Q phase from glassy phase has been reported in a number of alloy systems such as Zr-Cu,^19,23 Zr-M (M=Pd or Pt),^10-17 Zr-Ni-M (M=Pd, Au or Pt),^18-24 Zr-Al-Ni-Cu^20,22 and Zr-Cu-Ni-Pd²¹ glassy alloys, while the transition from I-Q phase to APC phase remains unknown even for their typical I-Q formation glassy alloys where the three component rule for stabilization of supercooled liquid as well as the formation of bulk metallic glasses is satisfied. Very recently, we have noticed that the sequent transition of the three metastable phases, namely, amorphous → I-Q → APQ → equilibrium crystalline phases occurs in a very simple Zr₆₅Ni₃₀Pd₅ ternary amorphous alloy.

In addition, the multicomponent Zr-Ni-Pd alloy gives a unique opportunity to investigate the possibility of synthesizing a glass-type alloy with glass transition and supercooled liquid region. This is because the alloy component deviates slightly from the three component rule for stabilization of supercooled liquid and bulk glass formation.^20,21 , i.e., the satisfaction of significant atomic size mismatches of Zr > Pd > Ni, while the deviation of the bonding state resulting from the coexistence of the nearly zero heat of mixing for Ni-Pd atomic pair and the negative heats of mixing for Zr-Ni and Zr-Pd pairs. This paper aims to examine the thermal stability, phase decomposition behavior and kinetics of a Zr₆₅Ni₃₀Pd₅ amorphous alloy and to investigate the feature of the phase decomposition for the ternary amorphous alloy in comparison with the I-Q precipitation behavior for Zr₇₀Pd₃₀ and Zr₆₅Pd₃₅ binary amorphous alloys,^18-22as well as for Zr-Ni-Pd ternary amorphous alloys.^18-24

2. Experimental Details

An alloy ingot with the nominal atomic composition of Zr₆₅Ni₃₀Pd₅ was chosen because the ternary alloy is expected to increase the thermal stability of metastable phases through the atomic size mismatch effect as compared with binary Zr-Ni and Zr-Pd alloys. The ingot was prepared by arc melting the mixture of pure Zr, Ni and Pd metals with purity above 99.9 mass% in an argon atmosphere.

Alloy ribbon with a cross section of about 0.02 x 1.2 mm² was prepared by the single roller melt spinning method with a high-frequency induction furnace in an argon environment. The diameter of the copper wheel was 200 mm and the wheel rotation speed was 3000 rpm.

Amorphicity and annealed structures were identified and examined by X-ray diffraction (XRD) using Rigaku Smart-Lab X-ray diffractometer with CuKα radiation. Thermal stability, crystallization behavior and kinetics were examined by differential scanning calorimetry (DSC) (Perkin Elmer DSC 8500) at heating rates of 0.083 - 1.33 K/s. Annealing treatment was made for 2 hrs at 423-523 K in an argon atmosphere using one zone programmable furnace from MTI Corporation, model OTF1200X.

3. Results

Figure 1 shows DSC curves of the Zr₆₅Ni₃₀Pd₅ amorphous alloy in the as-spun state and the pre-heated states at 423 K, 473 K and 523 K with a scanning rate of 0.333 K/s. No appreciable glass transition is observed in the temperature range just below the onset temperature of crystallization, indicating the formation of an amorphous type phase. The as-spun and pre-annealed samples exhibit three exothermic peaks, indicating that the amorphous phase completes its crystallization through three stages. The three-stage exothermic behavior remains unchanged even after the pre-annealing for 2 hrs at the temperatures below 523 K. The three-stage exothermic peak behavior is different from the two-stage exothermic behavior for Zr₇₀Ni₃₀, Zr₇₀Pd₃₀ and Zr₇₀Ni₁₀Pd₂₀ amorphous alloy ribbons.^18-24 The two-stage crystallization for these amorphous alloys can be interpreted to originate from the decrease of thermal stability through easy atomic diffusivity and atomic rearrangements for the alloys without the multicomponent effect such as atomic size mismatch.

Figure 1
DSC curves for Zr₆₅Ni₃₀Pd₅ amorphous alloy in as-spun (303 K) and pre-annealed (423, 473 and 523 K) states.

With the aim of clarifying the precipitation phases corresponding to each exothermic peak, Fig. 2 shows the XRD patterns of the as-spun and the annealed samples heated at 0.333 K/s to 700 K (just above the first exothermic peak), 763 K (above the second peak) and 785 K (above the third peak). These XRD patterns can be recognized as an amorphous phase in the as-spun state, I-Q and residual amorphous phases at 700 K, approximant cubic Zr₂Ni phase with a lattice parameter of 1.12 nm at 763 K, and two mixed compounds of tetragonal Zr₂Ni with a=0.654 nm and c=0.534 nm and tetragonal Zr₂Pd with a=0.3299 nm and c=1.088 nm. Based on the DSC and XRD data, it is concluded that the crystallization of the Zr₆₅Ni₃₀Pd₅ amorphous alloy occurs through the multistage process of amorphous → I-Q + amorphous → APC Zr₂Ni(Pd) → Zr₂Ni + Zr₂Pd. It is noticed that the metastable I-Q and APC cubic Zr₂(Ni,Pd) phases precipitate only for the ternary amorphous alloy and this three-stage phase decomposition process is significantly different from the two-stage crystallization mode for binary Zr-Ni and Zr-Pd or even for the ternary Zr₇₀Ni₁₀Pd₂₀ amorphous alloys. The Ni-Pd atomic pair has a nearly zero heat of mixing and the Ni-Pd equilibrium phase diagram is a typical solid soluble type.^10-24 The homogeneously mixed state is presumed to be the origin of the suppression of the directly separated precipitation mode to Zr₂Ni and Zr₂Pd phases. The approximant crystalline (APC) Zr₂(Ni,Pd) phase has a complex cubic structure with a lattice parameter of 1.12 nm including an icosahedral-like local atomic configuration. The similarity of the local atomic configuration in the I-Q and APC structures seems to cause the preferential precipitation of metastable APC Zr₂(Ni,Pd) phase from the I-Q phase.

Figure 2
X-ray diffraction patterns of different Zr₆₅Ni₃₀Pd₅ samples; (a) as-spun and annealed for 900 s at (b) 700 K (c) 763 K and (d) 785 K. The identified phases are also presented.

It was thus shown that the addition of Ni causes a significant influence of crystallization mode in comparison with the Zr_65-70Pd_30-35 binary amorphous alloys. Here, we further examined the kinetic behavior of the I-Q precipitation from the ternary Zr-Ni-Pd amorphous phase, in comparison with the previous data on the Zr-Pd amorphous alloy. ^10-18 To gain some information about the conversion kinetics from the amorphous phase to I-Q phase, the quasicrystallized volume fraction, α, should be determined. Here the Johnson-Mehl-Avrami model can be used to express, α(T) as follows:^25-29

(1) $\mathit{\alpha }\left(T\right)=1-\exp \left[-{\left(\mathit{k}\left(\mathit{T}\right)\mathit{t}\right)}^{\mathit{n}}\right]$

where t is the time, n is the kinetic exponent or Avrami exponent. The importance of n is because the mechanism and dimensionality of the quasicrystallization can be deduced from the value of n. k(T) is constant (reaction rate) and can be described by the Arrhenius temperature dependence:

(2) $\mathit{k}\left(\mathit{T}\right)={\mathit{k}}_0\exp \left[-\frac{{\mathit{E}}_{\mathit{A}}}{\mathit{RT}}\right]$

where k₀ is called as the frequency factor, E_A is the activation energy throughout the crystallization development, R is the ideal gas constant.

The reaction rate constant can be linked with the phase transformation rate dα/dt through the following relation:³⁰

(3) $\frac{\mathit{d}\mathit{\alpha }}{\mathit{dt}}=\mathit{k}\left(\mathit{T}\right)\mathit{\zeta }\left(\mathit{\alpha }\right)$

where ζ(α) is conversion function of the quasicrystallized volume fraction α. The function ζ(α) can be described as follows:^29,30

(4) $\mathit{\zeta }\left(\mathit{\alpha }\right)=\left(1-\mathit{\alpha }\right){\left[-\ln \left(1-\mathit{\alpha }\right)\right]}^{1-\frac{1}{\mathit{n}}}$

The above equation is derived without taking into account the isothermal limitation, and hence it should be used also to describe non-isothermal quasicrystallization kinetics.

Figure 3 shows DSC curves of Zr₆₅Ni₃₀Pd₅ amorphous alloy measured at different heating rates (non-isothermal). It is clear that the position of each peak is shifted toward upper temperatures, as the heating rate is increased. In the non-isothermal process, the quasicrystallized volume fraction, α, is deduced from the thermal analysis curves as a function of the measured temperature. The volume fraction quasicrystallized, α, at a given temperature T is calculated by:³¹

(5) $\mathit{\alpha }\left(\mathit{T}\right)=\frac{A_{\mathit{T}}}{A}$

Figure 3
Non-isothermal DSC curves obtained at different heating rates.

where A is the total area under the exothermic peak between the temperatures T₀ and T_∞ that correspond to the beginning and end of crystallization development, respectively. A_T is the partial area under the exothermic peak from T₀ up to the temperature T. For simplicity, only the first exothermic peak due to the formation of I-Q phase is considered. The calculated quasicrystallization volume fractions are shown in Fig. 4. All the curves exhibit a sigmoidal reliance with temperature which is normal for all amorphous materials during the non-isothermal crystallization process.

Figure 4
Non-isothermal crystallization volume fraction, α as a function of temperature.

The values of kinetic parameter n and the rate constants k(T) are deduced from nonlinear fits of the recorded data, using Eqs. 1 and 2. As summarized in Table 1, the average values of n are 1.93 at the heating rate of 0.083 K/s, 2.25 at 0.167 K/s and 2.63 at 0.333 K/s, respectively. Since n acquires only integer values from 1 to 4, here n values at 0.083 K/s and 0.167 K/s are close to 2 corresponding to 2-dimensional growth of quasicrystals, while the n value at 0.333 K/s becomes close to 3, implying nearly three-dimensional quasicrystallization growth.

Table 1
Non-isothermal kinetic parameters for Zr₆₅Ni₃₀Pd₅ at different heating rates.

	0.083 K/s	0.167 K/s	0.333 K/s
n	1.93 ± 0.086	2.25 ± 0.118	2.63 ± 0.177
EA(kJ/mol)	1652.5 ± 65.387	1704.1 ± 85.846	1784.2 ± 94.88
ΔH (J/g)	-44.95	-55.50	-63.37

With the aim of evaluating the activation energy for the precipitation of I-Q phase, plots of ln k(T) against 1/T are shown in Fig. 5. Good linear relations are recognized at 0.083, 0.167 and 0.333 K/s and the activation energies evaluated from the linear slope are about 1.65, 1.70 and 1.78 kJ/mol, respectively, as shown in Table 1.

Figure 5
Plots of ln k(T) versus 1/T (under non-isothermal process).

Figure 6 shows DSC curves of the Zr₆₅Ni₃₀Pd₅ amorphous alloy under an isothermal crystallization condition obtained at various annealing temperatures. It can be easily observed that the exothermic peak-time decreases with increasing annealing temperature. This is due to the fact that the atoms have larger mobility at superior annealing temperature that guides to easy formation flux for mass quasicrystallization.³² These results agree well with the previous reports for Pd-, Zr- and Cu-based bulk glassy alloys.^33-36

Figure 6
Isothermal DSC curves obtained at different annealing temperatures.

The correspondence involving the crystallization volume fraction and annealing time through isothermal development is revealed in Fig. 7. It is clearly shown that all the curves display a characteristic s-shaped type at all annealing temperatures and that with increasing annealing temperature the phase transformation process goes more rapidly. It is observed that the amorphous phase transfers into two kinds of metastable phases of I-Q and approximant cubic Zr₂(Ni,Pd) before stable tetragonal Zr₂Ni and Zr₂Pd phases.

Figure 7
Crystallization volume fraction α as a function of annealing time at different annealing temperatures.

By combining Eqs. 3 and 4, and taking the logarithm of the resulting equation, the following equation is obtained:

(6) $\ln \left(\frac{\mathit{d}\mathit{\alpha }}{\mathit{dt}}\right)-\ln \left(1-\mathit{\alpha }\right)=\ln \left(\mathit{k}\left(\mathit{T}\right)\right)+\left(\left(1-\frac{1}{\mathit{n}}\right)\ln \left(-\ln \left(1-\mathit{\alpha }\right)\right)\right)$

Scheming of $\left(\frac{\mathit{d}\mathit{\alpha }}{\mathit{dt}}\right)$ versus ln(-ln(1-α)) at different annealing temperatures is displayed in Fig. 8. The kinetic parameter n and the reaction rate (T) can be estimated from the fitted curves. The obtained values of n are given in Table 2.

Figure 8
Plots of ln dt S da X-ln(1-α) versus ln(-ln(1-α)) at different annealing temperatures.

Table 2
Isothermal kinetic parameters for Zr₆₅Ni₃₀Pd₅ at different temperatures.

	673 K	668 K	661 K	658 K
n	2.43 ± 0.01	2.13 ± 0.02	2.78 ± 0.08	2.79 ± 0.07
EA (kJ/mol)	1191.2 ± 69.36	1135.2 ± 70.16	1058.2 ± 86.08	1017.8 ± 92.34
ΔH (J/g)	-30.06	-22.42	-19.25	-18.50

As summarized in Table 2, n diverges from 2.1 to 2.8 with increasing annealing temperature. Similar disparity has been observed in other bulk glassy alloys.^37-41 This disparity is somewhat narrow. According to the theory of diffusion controlled growth,³⁷ n may take the following values: n=1.5, 1.5<n<2.5, n=2.5 and n>2.5 which correspond to the following cases: i- the growth of particles occurs with a nucleation rate close to zero, ii- decreasing nucleation rate, iii- constant nucleation rate, and iv- increasing nucleation rate, respectively.³⁹ Here the average kinetic exponent n is 2.53 which means that the growth of I-Q particles in the Zr₆₅Ni₃₀Pd₅ amorphous alloy happens with a constant nucleation rate. The similar value of n has also been reported for Zr_55.9Cu_18.6Ta₈Al_7.5Ni₁₀.⁴⁰

Using the Arrhenius equation (2), the activation energy for the quasicrystallization development in an isothermal process can be obtained. Plots of ln k(T) versus 1/T are displayed in Fig. 9. All fitted curves are straight lines and coincide. The calculated activation energies are summarized in Table 2. Here, the calculated activation energies under isothermal conditions are smaller than the values calculated under non-isothermal conditions (during a continuous heating). This occurrence is in contrary to other results observed in different bulk glassy alloys.^41-43 Nucleation and growth are two phenomena which take place in the crystallization process. In the non-isothermal condition nucleation is predominant, since crystallization initiates at low temperatures. While in the isothermal condition, equally nucleation and growth are dominant with the same attribution during crystallization. Thus, in the Zr₆₅Ni₃₀Pd₅ amorphous alloy, the energy barrier in non-isothermal annealing condition is higher than that of isothermal conditions.

Figure 9
Plots of ln k(T) versus 1/T (under isothermal process).

4. Discussion

It was shown in the present study that the decomposition of Zr₆₅Ni₃₀Pd₅ amorphous phase upon continuous heating or isothermal annealing occurs in the order of amorphous, I-Q, cubic Zr₂Ni(Pd) and then Zr₂Ni + Zr₂Pd phases, instead of the two-stage process of amorphous, I-Q and then tetragonal Zr₂Ni or Z₂Pd without APC cubic Zr₂(Ni,Pd) phase. Considering that the Zr-Ni and Zr-Pd amorphous alloys crystallize through two stages,^13-17 the appearance of the three metastable amorphous, I-Q and APC phases is presumed to reflect the increase in the thermal stability of each metastable phase which is due to the sluggish phase decomposition reaction resulting from the difficulty of atomic rearrangements caused by the multicomponent effect leading to the distinct atomic size mismatches. Figure 10 shows a schematic illustration of continuous heating decomposition behavior of the Zr₆₅Ni₃₀Pd₅ amorphous phase. It is noticed that the I-Q phase does not decompose directly to stable crystalline mixed phases. The formation of the metastable APC cubic Zr₂Ni(Pd) phase seems to reflect the ease of the phase transition of I-Q phase resulting from the structural similarity in local atomic configurations between the I-Q and the Zr₂(Pd,Ni) phases and the APC cubic phase has a close relation with I-Q phase. The internal energy of these metastable phases lowers in the order of amorphous > I-Q > cubic Zr₂Ni(Pd) > tetragonal Zr₂Ni + tetragonal Zr₂Pd. The formation of the mostly single APC cubic Zr₂Ni(Pd,Ni) phase from the I-Q phase is believed to be the first evidence in the decomposition behavior of the I-Q phase. It is interesting to point out that this alloy gives a unique opportunity to clarify the transformation behavior from I-Q to APC cubic Zr₂Ni(Pd) phase in the absence of any other phase state as well as the coexistent and interface structures between I-Q and APC Zr₂Ni(Pd) phases. The clarification is expected to shed on the light to the mutual phase relationships among amorphous, I-Q and APC Zr₂Ni(Pd), which remain still unknown, and is under investigation.

Figure 10
Schematic illustration of continuous-cooling-transformation (CCT) and continuous-heating-transformation (CHT) curves of Zr₆₅Ni₃₀Pd₅ amorphous alloy.

5. Conclusions

The heating-induced decomposition behavior and its kinetics of Zr₆₅Ni₃₀Pd₅ amorphous alloy, which is attractive as a hydrogen absorption material, were examined in comparison with Zr_65-70Pd_30-35 binary amorphous alloys as well as Zr₇₀Ni₁₀Pd₂₀ ternary amorphous alloy. The amorphous phase decomposes through three stages of amorphous, I-Q + amorphous, APC cubic Zr₂Ni(Pd) and then tetragonal Zr₂Ni + tetragonal Zr₂Pd, being different from the two-stage decomposition of amorphous, I-Q + amorphous and then tetragonal Zr₂Pd for Zr_65-70Pd_30-35 binary amorphous alloys. The formations of I-Q and APC cubic Zr₂Ni(Pd) are presumably due to the development of icosahedral-like local atomic configurations in Zr-Ni-Pd amorphous alloy. The formation kinetics of the I-Q phase from the amorphous phase was also examined in continuous heating and isothermal annealing. The Avrami exponent n value increases from 1.93 at 0.083 K/s to 2.63 at 0.333 K/s in the continuous heating mode and is nearly constant in the isothermal annealing mode. Based on these data, the decomposition from the amorphous to I-Q phase occurs through two-dimensional growth mode with a constant nucleation mechanism at lower heating rate and three-dimensional growth mode at a higher heating rate.

6. Acknowledgments

This project was funded by the National Plan for Science, Technology and Innovation (MAARIFAH) - King Abdulaziz City for Science and Technology - the Kingdom of Saudi Arabia - award number (12-NAN3196-03). The authors also, acknowledge with thanks Science and Technology Unit, King Abdulaziz University for technical support.

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Publication Dates

• Publication in this collection
  20 Jan 2020
• Date of issue
  2019

History

• Received
  23 Oct 2019
• Reviewed
  03 Nov 2019
• Accepted
  11 Nov 2019