Anomalous formation of micrometer-thick amorphous oxide surficial layers during high-temperature oxidation of ZrAl₂

1. Introduction

The oxide layer formed on the surfaces of metals or metallic alloys during thermal oxidation has strong influence on their surficial properties [[1], [2], [3]], such as corrosion resistivity, friction, wear, catalytic properties and long-term stability/reliability [[4], [5], [6]]. An amorphous oxide surficial layer can greatly improve the surface-related properties of metals or metallic alloys as compared to its crystalline counterpart, due to its more homogenous composition and structure (i.e. no grain boundaries or other defects serving as fast transport paths) [[7], [8], [9], [10], [11], [12], [13]]. Thus, metals or metallic alloys, which can form amorphous oxide surficial layers upon oxidation, can have outstanding properties for practical applications. However, normally amorphous oxide surficial layers can only be formed at relatively low oxidation temperatures with rather thin thicknesses (<10 nm), and tend to crystallize at elevated temperature upon thickening [14,15]. This has strongly limited the applications of the metals or metallic alloys.

Zr-Al intermetallic alloys have drawn extensive attention in recent years, due to their low neutron-capture cross-sections, high corrosion resistance, and high tensile strength [[16], [17], [18], [19]]. Therefore, Zr-Al alloys have promising applications in wide areas, for example, as fuel cladding materials for nuclear reactor and hydrogen getters in vacuum systems [[20], [21], [22], [23], [24]]. Among various Zr-Al compounds (e.g. ZrAl₃, ZrAl₂, Zr₂Al₃, ZrAl, Zr₅Al₄, Zr₄Al₃, Zr₃Al₂, Zr₅Al₃, Zr₂Al and Zr₃Al), ZrAl₂ possesses the most outstanding physical properties, such as the largest elastic modulus, the strongest average bond strength of atoms [25] and the highest hardness [26].

In this work, the high-temperature thermal oxidation of ZrAl₂ has been investigated by X-ray diffraction, scanning electron microscopy, Auger electron spectroscopy and cross-sectional transmission electron microscopy. The oxidation leads to the growth of anomalously thick (1-4.5 μm) amorphous oxide surficial layers, which can be explained by interface thermodynamics and the occurrence of high interface stability associated with a synchronous oxidation of Al and Zr elements. The oxidation kinetics obeys a parabolic law and the activation energy is determined to be 143 kJ/mol. The results enlighten ZrAl₂ as a promising engineering material for industrial applications.

2. Experimental

2.1. Specimen preparation

ZrAl₂ alloy ingots were prepared by melting the mixture of pure Zr (purity ≥99.99 wt.%) and pure Al (purity ≥99.999 wt.%) in a magnetic levitation melting furnace (LGX-5B) under protective pure Ar gas (purity ≥99.995 vol.%) atmosphere. The furnace was firstly evacuated to a base pressure of 1 × 10^-2 Pa and then backfilled with argon to a pressure of 0.4 bar. The mixture of two pure metals were melted at 1750 °C for 10 min and then solidified. Next, the alloy was flipped over and re-melted for four cycles to improve the homogeneity and purity. Finally, the liquid alloy was cooled quickly and solidified into bulk ZrAl₂ specimens.

The ZrAl₂ specimens were cut into pieces of 5 mm × 5 mm × 2 mm, and then polished using a series of polishing papers and finally by Al₂O₃ slurry. After polishing, the ZrAl₂ specimens were ultrasonically cleaned in acetone and ethanol for 1 h.

2.2. Thermal oxidation

The as-cleaned ZrAl₂ specimen was put into a quartz tube. The tube was firstly evacuated and then backfilled with pure oxygen (purity ≥99.999 vol. %), using a rotary vacuum sealing device (MRVS-3002, Partulab Technology). To ensure an oxygen partial pressure of pO₂ = 1 × 10⁵ Pa at the oxidation temperature of 550 °C, the oxygen partial pressure was set at 3.6 × 10⁴ Pa at room temperature (RT) in the tube. The sealed tube was introduced into a pre-heated tube furnace (OTF-1200X) at 550 °C and isothermally oxidized for 24 h. The oxidized specimen was cooled naturally to RT after the oxidation experiment. Similarly, oxidation experiments were also carried out at 650 °C (pO₂ = 3.2 × 10⁴ Pa at RT) and 750 °C (pO₂ = 2.9 × 10⁴ Pa at RT), respectively.

2.3. Thermogravimetric analysis

For thermogravimetric analysis (TGA), the polished ZrAl₂ specimens were cut into smaller pieces of 2 mm × 2 mm × 1 mm. The mass changes of ZrAl₂ specimens during isothermal oxidation in pure oxygen at 550 °C, 650 °C and 750 °C were recorded as a function of the oxidation time in a Mettler Toledo-1 1600 H T instrument.

2.4. X-ray diffraction

The ZrAl₂ specimens before and after thermal oxidation were investigated by X-ray diffraction (XRD) on a Bruker D8 Advance diffractometer with Cu-Kα radiation (40 kV/40 mA, λ = 1.5418 Å). The scanning range of the diffraction angle 2θ was 15 to 80°, with a step size of 0.2°.

2.5. Scanning electron microscopy

The oxidized ZrAl₂ specimens have been investigated by using scanning electron microscopy (SEM) in secondary electron (SE) and backscattered electron (BSE) modes, respectively. The SEM microscope (JEOL JSM-7800 F) was equipped with an energy dispersive spectrometer (EDS, EDAX Octane Plus). The spacial distributions of O, Al and Zr elements for the ZrAl₂ oxidized at 650 °C and 750 °C for 24 h have been investigated by making EDS line scans along cross-section of the specimens.

2.6. Auger electron spectroscopy

The elemental depth-distribution in the ZrAl₂ specimen oxidized at 550 °C for 24 h has been determined by Auger electron spectroscopy (AES) sputter-depth profiling in a PHI 700Xi Scanning Auger Nanoprobe system equipped with a cylindrical mirror analyzer and a field emission electron gun (acceleration voltage: 5 kV). The quantification of Al and Zr in their metallic (Al^met and Zr^met) and oxidic (Al³⁺ and Zr⁴⁺) states, as well as of O, have been carried out according to the procedure described in Ref. [27]. The determined relative sensitivity factors S (with respect to metallic Zr) are: $S_{Al^{3+}}$ = 0.411, $S_{Al^{met}}$= 0.178, SZr4+ = 0.627, $S_{Zr^{met}}$ = 1.000, S_O = 1.312.

2.7. Transmission electron microscopy

The cross-sectional thin lamellae of ZrAl₂ oxidized at 750 °C for 24 h were prepared by using a so-called tripod-polishing method (see Ref. [27] for details). The transmission electron microscopy (TEM) investigated was carried out in a JEOL JEM-2100 F microscope, operated at an acceleration voltage of 200 keV.

3. Results

3.1. Microstructure and composition of the oxide surficial layer

Fig. 1a presents the XRD pattern of the as-cast ZrAl₂ alloy and those collected after oxidation of the ZrAl₂ alloy at 550 °C, 650 °C and 750 °C for 24 h. The magnified XRD pattern of the ZrAl₂ alloy oxidized at 750 °C for 24 h is shown in Fig. 1b. From the sharp diffraction peaks observed in Fig. 1a, it can be confirmed that the ZrAl₂ alloy is single phase and polycrystalline. Observation of the as-cast ZrAl₂ alloy by optical microscopy has indicated that its grain sizes are very large, within the range of tens to hundreds of micrometers. A broad diffraction hump is observed in the range of 2θ ∼ 25-35° in the XRD pattern of the ZrAl₂ oxidized at 750 °C for 24 h (Fig. 1b). These observations imply an amorphous state of the oxide surficial layer formed upon oxidation of ZrAl₂ at 550-750 °C. Fig. 1c-e show the surface SEM morphologies of the oxide layers. It is obvious that the surfaces of all the oxide layers are smooth and uniform, without formation of any cracks, which may suggest excellent oxidation resistance of ZrAl₂ at the studied temperatures.

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Fig. 1.   (a) XRD patterns of the as-cast and the oxidized ZrAl₂ alloys at oxidation temperatures of 550 °C, 650 °C and 750 °C (in pure oxygen for 24 h); (b) magnified XRD pattern for the ZrAl₂ oxidized at 750 °C; (c-e) surface SEM micrographs of the ZrAl₂ alloy oxidized at 550 °C (c), 650 °C (d) and 750 °C (e) for 24 h, respectively.

Fig. 2a-c show the cross-sectional SEM micrographs of ZrAl₂ oxidized at 550 °C, 650 °C and 750 °C for 24 h. The micrographs indicate that the thicknesses of the oxide surficial layers formed at 550 °C, 650 °C and 750 °C are 1.0 ± 0.2 μm, 2.0 ± 0.2 μm and 4.5 ± 0.5 μm, respectively. Consistent with the results of plan-view SEM, these cross-sectional SEM micrographs show that a continuous, homogenous and single-layered oxide layer has been formed on surfaces of the ZrAl₂ alloy at the studied oxidation temperatures, without the presence of any cracks or holes/pores in the oxide surficial layer.

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Fig. 2.   Cross-sectional SEM micrographs of ZrAl₂ oxidized at (a) 550 °C, (b) 650 °C, (c) 750 °C for 24 h in pure oxygen (pO₂ = 1 × 10⁵ Pa). (d) AES elemental concentration-depth profiles of ZrAl₂ oxidized at 550 °C for 24 h. (e, f) Cross-sectional EDS-line profiles of ZrAl₂ oxidized at 650 °C (e) and 750 °C (f) for 24 h.

The elemental distributions of O, Al and Zr in the oxide surficial layer have further been investigated by AES depth profiling and by cross-sectional EDS line scanning (see Fig. 2d-f). The atomic fractions of O, Al and Zr keep constant over the whole depth range of the oxide surficial layer, and are also independent of the oxidation temperature. Both the AES and EDX analyses indicate that the Al³⁺/Zr⁴⁺ atomic ratio in the amorphous oxide surficial layer is about 2, corresponding to a homogenous oxide composition of (Zr_0.33Al_0.67)O_1.66 at 550-750 °C. The Al³⁺/Zr⁴⁺ atomic ratio is practically the same as the Al/Zr atomic ratio in the ZrAl₂ substrate. The AES and EDX analyses further indicate that the solubility of O in the ZrAl₂ substrate is rather low at 550-750 °C.

The cross-sectional TEM image of ZrAl₂ oxidized at 750 °C for 24 h is shown in Fig. 3a, and b (the selected area in the Fig. 3a) shows a cross-sectional HRTEM image of the oxide/substrate interface region, indicating a rather sharp interface of the substrate and the oxide surficial layer. The corresponding selected-area electron diffraction patterns (SADP) of the ZrAl₂ substrate and the oxide surficial layer are presented in Fig. 3c and d, respectively. The SADP of the ZrAl₂ substrate confirms the crystalline structure of the ZrAl₂ substrate. The diffuse halo in the SADP of the oxide surficial layer indicates that the oxide surficial layer is amorphous, in consistence with the corresponding XRD results (Fig. 1).

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Fig. 3.   (a) Cross-sectional TEM image of ZrAl₂ oxidized at 750 °C for 24 h; (b) Cross-sectional HRTEM image of the region marked in (a). Arrow marks the location of the oxide/substrate interface; (c) SAD pattern of the substrate region and (d) SAD pattern of the oxide region.

3.2. Oxidation kinetics

TGA measurements in a pure O₂ atmosphere have been performed to investigate the oxidation kinetics of ZrAl₂ at different temperatures (550 °C, 650 °C and 750 °C). The measured mass gain per unit surface area as a function of the oxidation time at the studied temperatures is shown in Fig. 4a. For a parabolic oxidation kinetics, the relationship between the mass gain per unit surface area w and the oxidation time t is expressed as:

W(t)²=k_pt (1)

where k_p is the parabolic growth-rate constant. Fig. 4b presents the square of mass gain per unit surface area as a function of the oxidation time. It follows that the oxidation kinetics of ZrAl₂ indeed obeys a parabolic law at 550-750 °C. The parabolic rate constants k_p have been determined to be 4.43 × 10^-11, 3.80 × 10^-10 and 2.42 × 10^-9 kg²/(m⁴·s) at 550 °C, 650 °C and 750 °C, respectively. The k_p increases with increasing oxidation temperature. According to the Arrhenius relationship:

k_pT=A×exp-($\frac{Q}{RT}$) (2)

where A is a pre-exponential factor, Q is the activation energy for the rate-determining step in the oxidation process, R is the gas constant and T is the absolute temperature. The Arrhenius plots for k_p at 550 °C, 650 °C and 750 °C are shown in Fig. 4c. On this basics, the activation energy is calculated to be 143 kJ/mol.

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Fig. 4.   (a) Measured mass gain of ZrAl₂ per unit surface area as a function of the oxidation time at 550 °C, 650 °C and 750 °C. (b) The square of mass gain per unit surface area as a function of oxidation time at 550 °C, 650 °C and 750 °C. (c) Arrhenius plot of parabolic rate constants k_p for oxidation of ZrAl₂ at 550 °C, 650 °C and 750 °C.

4. Discussion

Zr and Al have very similar affinity towards oxygen [27]. During thermal oxidation of crystalline ZrAl₂, Zr and Al elements are oxidized more or less simultaneously. As a result, the formed oxide surficial layer [(Zr_0.33Al_0.67)O_1.66] has an Al³⁺/Zr⁴⁺ atomic ratio practically the same as the Al/Zr atomic ratio in the ZrAl₂ alloy, i.e. no preferential oxidation of any metallic elements occurs during the thermal oxidation of ZrAl₂ at 550-750 °C.

During initial oxidation of metals or metallic alloys, the formation of a (very) thin amorphous oxide surficial layer can be thermodynamically preferred, because of relatively lower surface and interface energy associated with the amorphous oxide surficial layer. When the oxide surficial layer grows thicker upon continued oxidation and exceeds a critical thickness, the higher bulk Gibbs energy of amorphous oxide relative to that of the corresponding crystalline oxide starts to overweight the effect of the lower surface and interface energy, and the oxide surficial layer then becomes crystalline [15]. Such a critical thickness is typically in the range of only 1-3 nm [28,29]. In the present study, strikingly, amorphous oxide surficial layers as thick as 4.5 μm (i.e. more than 3 orders of magnitude larger than the critical thickness; see Fig. 2) have been observed to form upon high-temperature thermal oxidation of ZrAl₂. Such thick and uniform amorphous oxide surficial layers can lead to outstanding corrosion resistivity and other surface-related properties [30,31]. The underlying mechanism for the growth of such unusually thick amorphous oxide surficial layer can be understood as follows.

In the pseudo-binary Al₂O₃-ZrO₂ system, there is practically no mutual solubility of crystalline Al₂O₃ and ZrO₂ [32,33]. For the oxidation of the ZrAl₂ alloy, the expected stable oxide phases are therefore phase-separated crystalline γ-Al₂O₃ and tetragonal ZrO₂ [15,27,28], rather than a ternary crystalline (Al, Zr)-oxide solid solution or compound. The crystallization of the initially formed amorphous (Zr_0.33Al_0.67)O_1.66 into crystalline oxides would thus involve significant compositional redistribution, requiring pronounced atomic/ionic diffusion in the amorphous (Zr_0.33Al_0.67)O_1.66 oxide. For ionic oxide as (Zr_0.33Al_0.67)O_1.66, such atomic/ionic diffusion would be rather difficult, due to the strong interaction between cations and anions and the absence of lattice defects in amorphous oxide. As a result, the formation of crystalline oxide nuclei within the formed amorphous (Zr_0.33Al_0.67)O_1.66 oxide would be hindered owing to such diffusion constraints.

It has been suggested previously that the crystalline oxide may firstly nucleate at the amorphous oxide/substrate interface [[34], [35], [36], [37]]. During oxidation, the oxygen diffuses through the amorphous oxide to the oxide/substrate interface, leading to the formation of crystalline oxide at the interface [[38], [39], [40]]. Once the oxide crystallizes at the interface, it may even trigger the transition of the entire oxide layer from amorphous to crystalline [41,42]. As schematically sketched in Fig. 5, the difference in the total Gibbs energy change for the new nucleation of a crystalline oxide layer of thickness h at the oxide/alloy-substrate interface (Case ii in Fig. 5), as compared to that for the continuous growth of existing amorphous oxide (Case i in Fig. 5), ΔGi→iitotal (in J/mol), is given by:

$ΔG^{total}_{i→ii}=ΔG^{bulk}_{am-ox|cry-ox}+(γ^{interface}_{cry-ox|cry-alloy}+γ^{interface}_{am-ox|cry-ox}-γ^{interface}_{am-ox|cry-alloy})/h$ (3)

where $ΔG^{bulk}_{am-ox|cry-ox}$ is the difference in the bulk Gibbs energies of formation of the crystalline and amorphous oxides, $γ^{interface}_{cry-ox|cry-alloy}$, $γ^{interface}_{am-ox|cry-ox}$ and $γ^{interface}_{am-ox|cry-alloy}$ represent the interface energy of the crystalline-oxide/crystalline-alloy interface, the amorphous-oxide/crystalline-oxide interface and the amorphous-oxide/crystalline-alloy interface, respectively (see Fig. 5).

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Fig. 5.   Schematic illustration of (i) the continuous growth of existing amorphous oxide and (ii) the new nucleation of a crystalline oxide layer at the oxide/alloy-substrate interface.

In Eq. (3), on the one hand, the $ΔG^{bulk}_{am-ox|cry-ox}$ is only slightly negative since the Gibbs energies for crystalline Al₂O₃ and ZrO₂ are only slightly smaller than that of amorphous (Zr, Al)-oxide phase, because of the presence of strong mutual miscibility between amorphous Al₂O₃ and amorphous ZrO₂ [[43], [44], [45]]. On the other hand, the interface energy $γ^{interface}_{cry-ox|cry-alloy}$ is known to be highly positive, as a consequence of the high interface strain energy and energy of misfit dislocations at the interface of the crystalline oxide and the crystalline alloy substrate [46]. Furthermore, the term $γ^{interface}_{am-ox|cry-ox}$-$γ^{interface}_{am-ox|cry-alloy}$ is also positive, realizing that a stronger chemical interaction between the amorphous oxide and the crystalline alloy than that between the amorphous oxide and the crystalline oxide [[47], [48], [49], [50]] will definitely result in a lower interface energy $γ^{interface}_{am-ox|cry-alloy}$ than $γ^{interface}_{am-ox|cry-ox}$. Therefore, on the basis of the above discussion, the initial formation of a thin crystalline oxide layer at the amorphous-oxide/alloy-substrate interface would be strongly inhibited by the associated high sum of $(γ^{interface}_{cry-ox|cry-alloy}+γ^{interface}_{am-ox|cry-ox}-γ^{interface}_{am-ox|cry-alloy})/h$. The diffusion of oxygen towards the interface thus only leads to the continued growth of amorphous oxide (Case i in Fig. 5).

Besides, it has been previously found that the compositional instability at the amorphous-oxide/alloy interface is also a key factor influencing the amorphous-to-crystalline transition of the oxide layer [41,42,48,51]. During thermal oxidation of amorphous Cu-Zr [52] and Al-Zr [27] alloys, crystalline tetragonal ZrO₂ can nucleate at the amorphous oxide/substrate interface, as triggered by the continuous enrichment of certain metallic elements (Cu or Al) and the associated oxygen solubility change at the reacting oxide/substrate interface during oxidation [27,42,52]. During oxidation of crystalline ZrAl₂, the Al³⁺/Zr⁴⁺ atomic ratio in the amorphous oxide surficial layer keeps the same as the Al/Zr atomic ratio in the ZrAl₂ substrate (see above discussion). The Al and Zr are synchronously oxidized and oxidation-induced enrichment/redistribution of elements (Al and Zr) does not occur at the oxide/substrate. Thus, the amorphous oxide/ZrAl₂ interface keeps stable throughout long-term oxidation even at high temperatures, and any formation of crystalline oxide phase as mediated by interface instability can be ruled out.

The parabolic rate law of oxidation kinetics at studied temperatures confirmed that the diffusion of oxygen ions [1,53] through the amorphous oxide surficial layer controls the rate of the oxidation process. The activation energy for oxidation of ZrAl₂ is determined to be 143 kJ/mol at the studied temperatures, which is higher than the values for oxidation of pure Al [54], amorphous Zr_0.32Al_0.68 [55] and pure Zr [56] (shown in Table 1), likely as a result of the dense-packed atomic structure of the amorphous oxide. Therefore, micrometer-thick amorphous (Zr_0.33Al_0.67)O_1.66 oxide surficial layer with a dense-packed atomic structure, formed upon thermal oxidation of ZrAl₂, can effectively improve the oxidation and corrosion resistance of ZrAl₂ at high temperatures, thus enlightening ZrAl₂ as an important new high-temperature alloy for potential industrial applications.

5. Conclusions

The thermal oxidation of ZrAl₂ in the temperature range of 550-750 °C has been investigated comprehensively. The thermal oxidation leads to the growth of anomalously thick (up to 4.5 μm) amorphous (Zr_0.33Al_0.67)O_1.66 surficial layers at the ZrAl₂ surface. The oxidation kinetics follows a parabolic law with an activation energy of 143 kJ/mol.

The continuous growth of micrometer-thick amorphous oxide surficial layers at temperatures as high as 750 ℃, rather than a crystalline oxide surficial layer, is ascribed to two reasons: (1) high energy barrier for nucleation of the crystalline oxide layer owing to the high sum of $(γ^{interface}_{cry-ox|cry-alloy}+γ^{interface}_{am-ox|cry-ox}-γ^{interface}_{am-ox|cry-alloy})/h$ and (2) high compositional stability in front of the reacting amorphous-oxide/alloy interface due to the synchronous oxidation of Al and Zr. The finding can be useful in the design of new types of high-temperature metallic alloys with outstanding stability/reliability at high temperatures, associated with thick amorphous oxide formation.

Acknowledgements

This work is financially supported by the National Natural Science Foundation of China (No. 51571148), the National Key Research and Development Program of China (No. 2017YFE0302600 and No. 2017YFB0701801), and the Thousand Talents Program for Distinguished Young Scholars of China.

The authors have declared that no competing interests exist.