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Structural and mechanical modifications induced on Cu47.5Zr47.5Al5 metallic glass by surface laser treatments
Applied Surface Science 290 (2014) 188–193
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Applied Surface Science journal homepage: www.elsevier.com/locate/apsusc
Structural and mechanical modifications induced on Cu47.5 Zr47.5 Al5 metallic glass by surface laser treatments a ˜ J. Fornell a,∗ , E. Pellicer a , E. Garcia-Lecina b , D. Nieto b , S. Surinach , M.D. Baró a , J. Sort c a b c
Departament de Física, Universitat Autònoma de Barcelona, 08193 Bellaterra, Spain Surfaces Division, IK4-CIDETEC, Parque Tecnológico de San Sebastián, 20009 Donostia, Spain Institució Catalana de Recerca i Estudis Avanc¸ats and Departament de Física, Universitat Autònoma de Barcelona, 08193 Bellaterra, Spain
a r t i c l e
i n f o
Article history: Received 9 September 2013 Received in revised form 25 October 2013 Accepted 9 November 2013 Available online 21 November 2013 Keywords: Metallic glass Surface laser treatment Nanoindentation Wettability
a b s t r a c t We have investigated the effects of surface laser treatment (SLT) on the structure, mechanical properties and wettability of Cu47.5 Zr47.5 Al5 metallic glass alloy. SLT has been carried out at three different intensities with the aim of inducing variable surface damage and tuneable changes in the resulting properties. X-ray diffraction characterization and scanning electron microscopy observations reveal that the alloy laser treated at 28.5 A remains amorphous while the alloy treated at 29 A becomes partially crystalline (CuZr B2 phase). When the alloy is treated at 30 A, it is mainly composed of copper and zirconium oxides. Nanoindentation tests, carried out on-top of the as-cast and laser-treated surfaces, reveal that SLT at 28.5 A causes an increase in hardness, which can be attributed to annihilation of free volume (i.e. structural relaxation). Conversely, hardness values of the alloy laser-treated at 29 A are almost the same as those of the as-cast alloy. This could be ascribed to the counterbalance effect between the softer nature of the CuZr B2 phase and the harder nature of the remaining relaxed amorphous phase. Larger hardness values are observed for the alloy laser treated at 30 A as a result of oxide phase formation. © 2013 Elsevier B.V. All rights reserved.
1. Introduction During the last decades bulk metallic glasses (BMGs) have attracted deal of attention due to their unique properties. In many cases the mechanical properties of BMGs are superior to those of their crystalline counterparts. They exhibit outstanding strength, large elastic strain and high energy storage. Such properties, together with their superior corrosion resistance, make them promising candidates in many structural applications [1–3]. Among metallic glasses, the binary Cu–Zr system has triggered considerable interest due to its tuneable microstructure and reasonable plasticity during compression [4,5]. Moreover, apart from their good mechanical properties, Cu-based alloys are relatively lowcost, thus being attractive for engineering applications. Conventionally, surface treatments have been widely used to broaden the applicability and to increase the fatigue life of engineering materials. Protective layers for harsh environments, reflective or anti-reflective coatings to be used in solar cells or biocompatible coatings on implants are just a few examples of the range of surface treatments available. To achieve enhanced mechanical properties is one of the purposes of surface treatments.
∗ Corresponding author. Tel.: +34 93 5811401; fax: +34 93 5812155. E-mail addresses: [email protected] (J. Fornell), [email protected] (E. Garcia-Lecina). 0169-4332/$ – see front matter © 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.apsusc.2013.11.032
In this framework, surface hardening is suitable to improve wear resistance, to increase surface strength for load carrying, to induce suitable residual and compressive stresses and to improve fatigue life and/or impact resistance. Methods to harden surfaces may be divided into three main groups: (i) case hardening, where the surface is chemically changed by thermochemical treatments (i.e. carburizing, chroming, nitriding, etc.), (ii) heat treatment, where microstructural changes are induced by heating the surface (i.e. induction, flame, laser, light, electron beam, etc.) and (iii) mechanical processes, where the structural changes are induced by elastic–plastic cold working on the surface (i.e. shot peening, deep rolling and shot blasting). In metallic glasses, shot peening and ion/electron irradiation are the most commonly used surface treatments techniques to improve mechanical properties [6–11] by inducing residual stresses and increasing the amount of free volume within the amorphous structure [8], eventually causing pronounced increase of plasticity [6]. Recently, researchers have shown an increased interest in laser treatments on polycrystalline materials. This technique can adjust the strengthening effect by accurately selecting the working parameters. In addition it is not a time-consuming technique and is a well-established method in the industry. A considerable amount of literature has been published on laser peening. Such studies reported an increase in hardness and better fatigue performance on peened surfaces [10–13]. Changes induced by shot and laser peening were compared in various studies [10,14]. These
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results reveal that laser peening can induce deeper compressive residual stress fields making metallic materials with better fatigue resistance. Although some research has been carried out on polycrystalline materials, far too little attention has been paid to laser treatments on metallic glasses. In this framework, high power pulsed lasers have been used to induce surface modifications in Cu- and Zr-based metallic glasses [15–18]. In most cases, surface melting takes place as a result of the high energy density of the laser beam. During this melting, the free volume content increases in the melt layer as a result of the rapid heating and cooling process; moreover, crystallization of primary crystalline phases is eventually observed. Most of these treatments have been carried out in He or Ar atmosphere to prevent oxidation. To the best of our knowledge, the effects of SLT performed in continuous mode and using lower laser power have not been investigated. The purpose of the current study was to induce surface modifications on Cu–Zr–Al metallic glass to enhance mechanical and hydrophobicity properties with a novel laser infrared treatment operating in continuous mode in air and at relatively low power. These SLTs permit to harden the alloy’s surface without crystallization. This method allows to precisely tune the microstructure of BMGs: from structurally relaxed to crystalline states, depending on the laser intensity. Furthermore, laser treatments under specific conditions in air atmosphere enable oxide phase formation which is known to be beneficial in terms of hydrophobicity and hardness. The laser treatment parameters were optimized to induce various degrees of surface modification and, subsequently, different properties resulted. Specifically, three different laser intensities were chosen to obtain relaxed, partially crystallized and completely crystallized states at the surface of the Cu–Zr–Al alloys. This method is proven to be very efficient to achieve the desired microstructural states in a much faster way than conventional heat-treatments (the overall duration of the laser treatments was 8 s in each case).
2. Experimental details A master alloy with nominal composition Cu47.5 Zr47.5 Al5 was prepared by arc melting the pure elements under an argon atmosphere. The alloy was re-melted several times in order to ensure homogeneity. Then, cylindrical rods with 2 mm in diameter were obtained from the arc-melt by suction casting into a watercooled copper mould in a purified argon atmosphere. Surface laser treatments (SLT) were carried out using a 25 W Nd:YVO4 laser (PowerLine E-Air 25, Rofin) emitting with a wavelength of 1064 nm (infrared) on 0.8 mm thick disks cut from the as-cast rods and subsequently polished to a mirror-like appearance. The specimens were placed at 53 mm from the focal plane and the treatment was applied at 10 mm/s for 8 s at three different intensities (28.5, 29 and 30 A) in a continuous mode. Working parameters (intensity, beam scanning speed and distance from the surface) were optimized experimentally, following our previous studies on Cu–Zr–Al BMG, to induce variable surface damage and tuneable changes in the resulting properties. The samples were structurally characterized by grazing incidence X-ray diffraction (XRD) on an X’pert PRO MRD PANalytical X-ray diffractometer by setting the incidence angle of the Xray beam at 1◦ . The superficial changes induced by the laser treatment were observed with a scanning electron microscope (SEM) (Zeiss Merlin) equipped with energy dispersive X-ray (EDX) analysis. X-ray photoelectron spectroscopy (XPS) analyses were carried out on a PHI equipment 5500 Multitechnique using Al K␣ radiation (1486.6 eV), after sputtering the surface of the lasertreated samples with Ar ions for 1 min. All spectral positions
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were charge corrected taking C 1s peak at 285 eV. Nanoindentation measurements were performed in load control mode, using an UMIS equipment from Fischer-Cripps Laboratories furnished with a Berkovich pyramidal-shaped diamond tip. The maximum applied load was 5 mN. The thermal drift was always kept below ±0.05 nm/s. From the load–displacement curves, the hardness (H) and the reduced Young’s modulus (Er ) values were determined at the beginning of the unloading segments, following the method of Oliver and Pharr [19,20]. Wetting experiments were carried out in 3.5 wt%NaCl by the sessile drop wettability technique on a CAM 200 from Iberlaser equipment with a liquid droplet of 1 L at room temperature. The roughness of the as-cast alloy and SLT specimens was assessed with a 3D Optical Surface Metrology System (DCM 3D) from Leica which combines confocal and interferometry technology. 3. Results and discussion Fig. 1 shows XRD patterns corresponding to the alloys laser treated at 28.5, 29 and 30 A. For the sample SLT at 28.5 A no diffraction peaks are observed suggesting the amorphous nature of this alloy. In the case of the alloy SLT at 29 A, we can identify Bragg reflections corresponding to B2 CuZr austenite structure (Pm3m space group) overlapping to two broad halos, characteristic of the amorphous structure. The XRD pattern of the alloy SLT at the highest intensity indicates the presence of B2 CuZr, Al2 O3 (R-3c space group), ZrO2 (P2/c space group)and perhaps some traces of other oxides. SEM images of the SLT alloys at 29 and 30 A support the XRD results. On the less treated alloy no traces of crystalline phases were detected throughout the specimen (not shown). Conversely, in the alloy SLT at 29 A some dendrites embedded in an apparently amorphous (i.e. featureless) structure can be observed in the low magnification SEM image (Fig. 2a). EDX quantification from the higher magnification image in Fig. 2b is presented in Fig. 2c. The Zr:Cu ratio on the dendrite is 1.27, quite in agreement with the stoichiometry (1:1) of the CuZr phase identified by XRD. The difference in atomic radii between Al and Cu is less than 15% suggesting that Al could occupy substitutional positions of Cu. Assuming that Al is occupying Cu positions, the Zr:Cu ratio would be 1.15 (almost the theoretical one). As expected, the ratios between Cu, Zr and Al on the matrix are maintained in line with the master alloy. The morphology of the alloy SLT at 30 A is presented in Fig. 2d (low magnification) and Fig. 2e (high magnification). At naked eye bluish colour can be observed at the surface of this alloy indicating oxide formation. Low magnification image shows black dendrites embedded in a matrix (Fig. 2e). EDX from Fig. 2e quantifies high O
Fig. 1. Grazing-incidence X-ray diffraction patterns of the alloys SLT at 28.5, 29 and 30 A.
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Fig. 2. SEM images of SLT specimens ((a) and (b)) at 29 A ((d) and (e)) at 30 A. EDX compositional analysis of the SLT specimens at (c) 29 A and (f) 30 A.
content in the dendrites as well as in the matrix (Fig. 2f) confirming oxide formation in both parts. In addition, strong depletion of Cu is observed in this case. In order to gain a better knowledge of the chemical composition of the surface layers, XPS analyses have been carried out after sputtering the surface with Ar ions for 1 min to avoid quantification of environmental contamination. Fig. 3 shows detailed spectra of the Zr 3d, Cu 2p and Al 2s. The Zr 3d5/2 and 3d3/2 doublet is located at 178.9 and 181.3 eV, respectively, which matches the position of metallic Zr [21]. The Zr 3d5/2 and 3d3/2 doublet located at 183.1 and 185.5 eV matches the Zr4+ valence state [22]. In the alloy SLT at 28.5 A both states are clearly identified; however, as the intensity of the laser increases the peaks corresponding to Zr4+ become more prominent while the peaks of Zr0 state tend to disappear. Likewise, the Cu 2p core splits into 2p3/2 (933 eV) and 2p1/2 (952.9 eV), which are characteristic of Cu2+ [23] and Cu0 [24]. It is interesting to note that these peaks tend to move to slightly lower binding energies as the laser intensity increases, tending to match the binding energies of Cu+ state (932.1 and 951.9 eV) [25]; however, the asymmetry of the peaks may indicate that there is still a contribution coming from the Cu2+ valence state. The Cu 3s peaks can also be detected in the binding energy range of Al 2s. Regarding Al, in the alloy SLT at 28.5 A the contribution of metallic Al is uniquely detected (118.1 eV) [26] while in the heaviest treated alloys the main contribution comes
Table 1 Element composition determined by XPS after sputtering with Ar ions for 1 min. Alloy
O (%)
Zr (%)
Al (%)
Cu (%)
SLT at 28.5 A SLT at 29 A SLT at 30 A
54.7 55.1 60.4
32.2 36.5 35
12.4 8.8 4.3
0.7 0.5 0.3
from the Al3+ valence state. This contribution, located at 116.4 eV, has been associated to ␥-Al2 O3 (I41/amd space group). This peak is usually accompanied with a lower intensity shoulder toward higher binding energy at 119 eV [27]. The shift in binding energies – from 118.5 eV for the alloy SLT at 30 A to 119.7 eV for the alloy SLT at 29 A – may be attributed to distinct crystallographic phases of Al2 O3 . As a matter of fact, a binding energy of 118.9 eV has been reported for corundum ␣-Al2 O3 (R-3c space group) phase [28] whereas a binding energy of 119.2 eV has been associated to ␥-Al2 O3 (I41/amd space group) phase [29]. Table 1 lists the atomic percentages of Zr, Cu, Al and O elements measured by XPS. It is interesting to note that almost no Cu is detected in the specimens, the surface being rich in O, Zr and Al. This is in agreement with EDX observations, and is consistent with the formation enthalpies (Hf ) of Zr, Cu and Al oxides. Values of formation enthalpies are: Hf (ZrO2 ) = −1100.8 kJ/mol, Hf (Al2 O3 ) = −1678.2 kJ/mol, Hf (CuO) = −155.2 kJ/mol and
Fig. 3. Zr 3d, Cu 2p and Al 2s core-level XPS spectra of the SLT Cu–Zr–Al alloys at 28.5, 29 and 30 A after sputtering with Ar ions for 1 min.
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Hf (Cu2 O) = −167.4 kJ/mol). Hence, from the thermodynamic viewpoint, Al2 O3 and ZrO2 should be more prone to grow compared to copper oxides. XPS results reveal that, even if not detected by XRD, a thin oxide layer is formed in all specimens. For instance, the XPS spectra of the alloy SLT at 28.5 A indicates the presence of ZrO2 , metallic Zr, CuO (Cu) and Al. This indicates that kinetics also plays a role on oxide formation. Fig. 4 shows the 3D profilometry images of the as-cast (a) and SLT alloys at 29 A (b) and 30 A (c). The Z-scale bar increases with the laser intensity compared to the as-cast alloy. It is clear that surface damage increases with laser intensity, therefore larger roughness values are obtained (Fig. 4d). No significant increase in Sq (root mean squared) and Sz (maximum height) were observed between the alloy SLT at 28.5 and 29 A but both values become much larger in the heavily treated alloy. It is interesting to note that roughness is homogeneous overall the laser treated specimens. Wettability results assessed in NaCl medium are presented in Fig. 5. The equilibrium contact angle ( eq ) in a non-reactive medium is defined by Young’s equation [30]: cos eq =
(SV − SL ) LV
(1)
where SV , SL and LV are the interfacial tensions between solid–gas, solid–liquid and liquid–gas phases, respectively. Our results indicate that the crystalline alloy exhibits markedly lower contact angle than the as-cast and SLT alloys at low intensities. However, a slight increase in the contact angle value between the as-cast and SLT alloys at 28.5 and 29 A is noticed. Since LV depends mainly on the liquid used for the sessile drop experiments, its value can be regarded as constant. Hence, the change in contact angle has to be attributed to variations in SV and/or SL . Laser treatments, as well as shot peening, high pressure torsion or annealing treatments, carried out under certain conditions, usually result in structural relaxation and the concomitant decrease in free volume concentration, thus reducing the atomic mobility. Moreover, structural
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relaxation may also be accompanied by cluster formation as a result of the high Cu–Zr binding energy which also contributes to decrease the free surface energy. Hence, lower SV values are expected. In turn, when the surface free energy of the solid decreases (i.e. in the relaxed state), the attraction between the liquid molecules and the atoms in the solids is lower than in the as-cast state, disfavouring the spreading of the liquid drop on the solid surface, hence increasing SL [31]. This concept could also be applied to crystalline materials, and consequently, even higher contact angles may be expected in the sample SLT at 30 A. However, lower contact angles are obtained in our case, suggesting the influence of other mechanisms. Actually, contact angles are larger on very rough surfaces than on chemically identical, smooth surfaces [32]. The observed increase in roughness (Fig. 4) could therefore have an influence on the hydrophobicity of the alloy. In addition, the alloy SLT at 30 A has larger O content at the surface. Hence, there is higher chance for O to combine with water molecule to form hydrogen bond, resulting in good wettability. These factors could explain the trends observed in this work. Fig. 6 displays the load (P) vs. penetration depth (h) curves of the as-cast and SLT alloys. Except for the alloy SLT at 30 A, the loading segments of these curves show serrated flow behaviour. These serrations are typically reported in metallic glasses, resulting from the creation and propagation of discrete shear bands [1,33]. Nanoindentation curves are consistent with XRD results: the alloy SLT at 30 A is completely crystalline; therefore no pop-in events are expected in the curve. Conversely, the other alloys, which are fully or mainly amorphous, show the existence of the serrations typically found in metallic glasses. The inset in Fig. 6 lists hardness and reduced Young’s modulus values. The increase in H and Er for the relaxed specimen (SLT at 28.5 A) can be explained because of the stronger bonding between atoms as a result of the annihilation of free volume. This enhancement could also be partly due to the formation of fine clusters. The decrease in hardness of the alloy SLT at 29 A can be explained by the formation of CuZr B2 phase. This
Fig. 4. Three-dimensional topography images of the (a) as-cast (b) SLT at 29 A and (c) SLT at 30 A alloys and (d) roughness parameters table (Sq = root mean squared and Sz = maximum height).
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Fig. 5. Photographs of as-deposited 3.5 wt% NaCl droplets onto the surface of (a) as-cast (b) SLT at 28.5 A (c) SLT at 29 A and (d) SLT at 30 A alloys.
to investigate the effect of surface laser treatments on wetting behaviour and mechanical properties of Cu–Zr–Al metallic glass. The lower contact angle in the crystalline specimen than in the ascast and SLT alloys at low intensities can be explained because of chemical and roughness contributions coming from oxide phase formation. Hardness results are explained in terms of free volume annihilation and phase formation. The alloy SLT at 28.5 A exhibits the largest hardness value. Acknowledgements This work has been supported by the MAT2011-27380-C02-01 and 2009 SGR 1292 projects. MDB acknowledges the ICREAAcademia Award. References Fig. 6. Load (P) vs. penetration depth (h) curves for the as-cast, SLT at 28.5 A, SLT 29 A and SLT 30 A alloys. In the inset the Hardness (H) and the reduced Young’s modulus (Er ) values are listed.
phase has been widely reported in the Cu–Zr system [34–37]; it is softer than its amorphous counterpart but it tends to improve the ductility of Cu-based BMGs via seeding and branching of shear bands. Lastly, the alloy SLT at 30 A exhibits large hardness values evidencing the harder nature of oxide phases. 4. Conclusions In this study Cu–Zr–Al alloy has been laser treated at different intensities. An in-depth structural characterization was employed
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Applied Surface Science journal homepage: www.elsevier.com/locate/apsusc
Structural and mechanical modifications induced on Cu47.5 Zr47.5 Al5 metallic glass by surface laser treatments a ˜ J. Fornell a,∗ , E. Pellicer a , E. Garcia-Lecina b , D. Nieto b , S. Surinach , M.D. Baró a , J. Sort c a b c
Departament de Física, Universitat Autònoma de Barcelona, 08193 Bellaterra, Spain Surfaces Division, IK4-CIDETEC, Parque Tecnológico de San Sebastián, 20009 Donostia, Spain Institució Catalana de Recerca i Estudis Avanc¸ats and Departament de Física, Universitat Autònoma de Barcelona, 08193 Bellaterra, Spain
a r t i c l e
i n f o
Article history: Received 9 September 2013 Received in revised form 25 October 2013 Accepted 9 November 2013 Available online 21 November 2013 Keywords: Metallic glass Surface laser treatment Nanoindentation Wettability
a b s t r a c t We have investigated the effects of surface laser treatment (SLT) on the structure, mechanical properties and wettability of Cu47.5 Zr47.5 Al5 metallic glass alloy. SLT has been carried out at three different intensities with the aim of inducing variable surface damage and tuneable changes in the resulting properties. X-ray diffraction characterization and scanning electron microscopy observations reveal that the alloy laser treated at 28.5 A remains amorphous while the alloy treated at 29 A becomes partially crystalline (CuZr B2 phase). When the alloy is treated at 30 A, it is mainly composed of copper and zirconium oxides. Nanoindentation tests, carried out on-top of the as-cast and laser-treated surfaces, reveal that SLT at 28.5 A causes an increase in hardness, which can be attributed to annihilation of free volume (i.e. structural relaxation). Conversely, hardness values of the alloy laser-treated at 29 A are almost the same as those of the as-cast alloy. This could be ascribed to the counterbalance effect between the softer nature of the CuZr B2 phase and the harder nature of the remaining relaxed amorphous phase. Larger hardness values are observed for the alloy laser treated at 30 A as a result of oxide phase formation. © 2013 Elsevier B.V. All rights reserved.
1. Introduction During the last decades bulk metallic glasses (BMGs) have attracted deal of attention due to their unique properties. In many cases the mechanical properties of BMGs are superior to those of their crystalline counterparts. They exhibit outstanding strength, large elastic strain and high energy storage. Such properties, together with their superior corrosion resistance, make them promising candidates in many structural applications [1–3]. Among metallic glasses, the binary Cu–Zr system has triggered considerable interest due to its tuneable microstructure and reasonable plasticity during compression [4,5]. Moreover, apart from their good mechanical properties, Cu-based alloys are relatively lowcost, thus being attractive for engineering applications. Conventionally, surface treatments have been widely used to broaden the applicability and to increase the fatigue life of engineering materials. Protective layers for harsh environments, reflective or anti-reflective coatings to be used in solar cells or biocompatible coatings on implants are just a few examples of the range of surface treatments available. To achieve enhanced mechanical properties is one of the purposes of surface treatments.
∗ Corresponding author. Tel.: +34 93 5811401; fax: +34 93 5812155. E-mail addresses: [email protected] (J. Fornell), [email protected] (E. Garcia-Lecina). 0169-4332/$ – see front matter © 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.apsusc.2013.11.032
In this framework, surface hardening is suitable to improve wear resistance, to increase surface strength for load carrying, to induce suitable residual and compressive stresses and to improve fatigue life and/or impact resistance. Methods to harden surfaces may be divided into three main groups: (i) case hardening, where the surface is chemically changed by thermochemical treatments (i.e. carburizing, chroming, nitriding, etc.), (ii) heat treatment, where microstructural changes are induced by heating the surface (i.e. induction, flame, laser, light, electron beam, etc.) and (iii) mechanical processes, where the structural changes are induced by elastic–plastic cold working on the surface (i.e. shot peening, deep rolling and shot blasting). In metallic glasses, shot peening and ion/electron irradiation are the most commonly used surface treatments techniques to improve mechanical properties [6–11] by inducing residual stresses and increasing the amount of free volume within the amorphous structure [8], eventually causing pronounced increase of plasticity [6]. Recently, researchers have shown an increased interest in laser treatments on polycrystalline materials. This technique can adjust the strengthening effect by accurately selecting the working parameters. In addition it is not a time-consuming technique and is a well-established method in the industry. A considerable amount of literature has been published on laser peening. Such studies reported an increase in hardness and better fatigue performance on peened surfaces [10–13]. Changes induced by shot and laser peening were compared in various studies [10,14]. These
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results reveal that laser peening can induce deeper compressive residual stress fields making metallic materials with better fatigue resistance. Although some research has been carried out on polycrystalline materials, far too little attention has been paid to laser treatments on metallic glasses. In this framework, high power pulsed lasers have been used to induce surface modifications in Cu- and Zr-based metallic glasses [15–18]. In most cases, surface melting takes place as a result of the high energy density of the laser beam. During this melting, the free volume content increases in the melt layer as a result of the rapid heating and cooling process; moreover, crystallization of primary crystalline phases is eventually observed. Most of these treatments have been carried out in He or Ar atmosphere to prevent oxidation. To the best of our knowledge, the effects of SLT performed in continuous mode and using lower laser power have not been investigated. The purpose of the current study was to induce surface modifications on Cu–Zr–Al metallic glass to enhance mechanical and hydrophobicity properties with a novel laser infrared treatment operating in continuous mode in air and at relatively low power. These SLTs permit to harden the alloy’s surface without crystallization. This method allows to precisely tune the microstructure of BMGs: from structurally relaxed to crystalline states, depending on the laser intensity. Furthermore, laser treatments under specific conditions in air atmosphere enable oxide phase formation which is known to be beneficial in terms of hydrophobicity and hardness. The laser treatment parameters were optimized to induce various degrees of surface modification and, subsequently, different properties resulted. Specifically, three different laser intensities were chosen to obtain relaxed, partially crystallized and completely crystallized states at the surface of the Cu–Zr–Al alloys. This method is proven to be very efficient to achieve the desired microstructural states in a much faster way than conventional heat-treatments (the overall duration of the laser treatments was 8 s in each case).
2. Experimental details A master alloy with nominal composition Cu47.5 Zr47.5 Al5 was prepared by arc melting the pure elements under an argon atmosphere. The alloy was re-melted several times in order to ensure homogeneity. Then, cylindrical rods with 2 mm in diameter were obtained from the arc-melt by suction casting into a watercooled copper mould in a purified argon atmosphere. Surface laser treatments (SLT) were carried out using a 25 W Nd:YVO4 laser (PowerLine E-Air 25, Rofin) emitting with a wavelength of 1064 nm (infrared) on 0.8 mm thick disks cut from the as-cast rods and subsequently polished to a mirror-like appearance. The specimens were placed at 53 mm from the focal plane and the treatment was applied at 10 mm/s for 8 s at three different intensities (28.5, 29 and 30 A) in a continuous mode. Working parameters (intensity, beam scanning speed and distance from the surface) were optimized experimentally, following our previous studies on Cu–Zr–Al BMG, to induce variable surface damage and tuneable changes in the resulting properties. The samples were structurally characterized by grazing incidence X-ray diffraction (XRD) on an X’pert PRO MRD PANalytical X-ray diffractometer by setting the incidence angle of the Xray beam at 1◦ . The superficial changes induced by the laser treatment were observed with a scanning electron microscope (SEM) (Zeiss Merlin) equipped with energy dispersive X-ray (EDX) analysis. X-ray photoelectron spectroscopy (XPS) analyses were carried out on a PHI equipment 5500 Multitechnique using Al K␣ radiation (1486.6 eV), after sputtering the surface of the lasertreated samples with Ar ions for 1 min. All spectral positions
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were charge corrected taking C 1s peak at 285 eV. Nanoindentation measurements were performed in load control mode, using an UMIS equipment from Fischer-Cripps Laboratories furnished with a Berkovich pyramidal-shaped diamond tip. The maximum applied load was 5 mN. The thermal drift was always kept below ±0.05 nm/s. From the load–displacement curves, the hardness (H) and the reduced Young’s modulus (Er ) values were determined at the beginning of the unloading segments, following the method of Oliver and Pharr [19,20]. Wetting experiments were carried out in 3.5 wt%NaCl by the sessile drop wettability technique on a CAM 200 from Iberlaser equipment with a liquid droplet of 1 L at room temperature. The roughness of the as-cast alloy and SLT specimens was assessed with a 3D Optical Surface Metrology System (DCM 3D) from Leica which combines confocal and interferometry technology. 3. Results and discussion Fig. 1 shows XRD patterns corresponding to the alloys laser treated at 28.5, 29 and 30 A. For the sample SLT at 28.5 A no diffraction peaks are observed suggesting the amorphous nature of this alloy. In the case of the alloy SLT at 29 A, we can identify Bragg reflections corresponding to B2 CuZr austenite structure (Pm3m space group) overlapping to two broad halos, characteristic of the amorphous structure. The XRD pattern of the alloy SLT at the highest intensity indicates the presence of B2 CuZr, Al2 O3 (R-3c space group), ZrO2 (P2/c space group)and perhaps some traces of other oxides. SEM images of the SLT alloys at 29 and 30 A support the XRD results. On the less treated alloy no traces of crystalline phases were detected throughout the specimen (not shown). Conversely, in the alloy SLT at 29 A some dendrites embedded in an apparently amorphous (i.e. featureless) structure can be observed in the low magnification SEM image (Fig. 2a). EDX quantification from the higher magnification image in Fig. 2b is presented in Fig. 2c. The Zr:Cu ratio on the dendrite is 1.27, quite in agreement with the stoichiometry (1:1) of the CuZr phase identified by XRD. The difference in atomic radii between Al and Cu is less than 15% suggesting that Al could occupy substitutional positions of Cu. Assuming that Al is occupying Cu positions, the Zr:Cu ratio would be 1.15 (almost the theoretical one). As expected, the ratios between Cu, Zr and Al on the matrix are maintained in line with the master alloy. The morphology of the alloy SLT at 30 A is presented in Fig. 2d (low magnification) and Fig. 2e (high magnification). At naked eye bluish colour can be observed at the surface of this alloy indicating oxide formation. Low magnification image shows black dendrites embedded in a matrix (Fig. 2e). EDX from Fig. 2e quantifies high O
Fig. 1. Grazing-incidence X-ray diffraction patterns of the alloys SLT at 28.5, 29 and 30 A.
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Fig. 2. SEM images of SLT specimens ((a) and (b)) at 29 A ((d) and (e)) at 30 A. EDX compositional analysis of the SLT specimens at (c) 29 A and (f) 30 A.
content in the dendrites as well as in the matrix (Fig. 2f) confirming oxide formation in both parts. In addition, strong depletion of Cu is observed in this case. In order to gain a better knowledge of the chemical composition of the surface layers, XPS analyses have been carried out after sputtering the surface with Ar ions for 1 min to avoid quantification of environmental contamination. Fig. 3 shows detailed spectra of the Zr 3d, Cu 2p and Al 2s. The Zr 3d5/2 and 3d3/2 doublet is located at 178.9 and 181.3 eV, respectively, which matches the position of metallic Zr [21]. The Zr 3d5/2 and 3d3/2 doublet located at 183.1 and 185.5 eV matches the Zr4+ valence state [22]. In the alloy SLT at 28.5 A both states are clearly identified; however, as the intensity of the laser increases the peaks corresponding to Zr4+ become more prominent while the peaks of Zr0 state tend to disappear. Likewise, the Cu 2p core splits into 2p3/2 (933 eV) and 2p1/2 (952.9 eV), which are characteristic of Cu2+ [23] and Cu0 [24]. It is interesting to note that these peaks tend to move to slightly lower binding energies as the laser intensity increases, tending to match the binding energies of Cu+ state (932.1 and 951.9 eV) [25]; however, the asymmetry of the peaks may indicate that there is still a contribution coming from the Cu2+ valence state. The Cu 3s peaks can also be detected in the binding energy range of Al 2s. Regarding Al, in the alloy SLT at 28.5 A the contribution of metallic Al is uniquely detected (118.1 eV) [26] while in the heaviest treated alloys the main contribution comes
Table 1 Element composition determined by XPS after sputtering with Ar ions for 1 min. Alloy
O (%)
Zr (%)
Al (%)
Cu (%)
SLT at 28.5 A SLT at 29 A SLT at 30 A
54.7 55.1 60.4
32.2 36.5 35
12.4 8.8 4.3
0.7 0.5 0.3
from the Al3+ valence state. This contribution, located at 116.4 eV, has been associated to ␥-Al2 O3 (I41/amd space group). This peak is usually accompanied with a lower intensity shoulder toward higher binding energy at 119 eV [27]. The shift in binding energies – from 118.5 eV for the alloy SLT at 30 A to 119.7 eV for the alloy SLT at 29 A – may be attributed to distinct crystallographic phases of Al2 O3 . As a matter of fact, a binding energy of 118.9 eV has been reported for corundum ␣-Al2 O3 (R-3c space group) phase [28] whereas a binding energy of 119.2 eV has been associated to ␥-Al2 O3 (I41/amd space group) phase [29]. Table 1 lists the atomic percentages of Zr, Cu, Al and O elements measured by XPS. It is interesting to note that almost no Cu is detected in the specimens, the surface being rich in O, Zr and Al. This is in agreement with EDX observations, and is consistent with the formation enthalpies (Hf ) of Zr, Cu and Al oxides. Values of formation enthalpies are: Hf (ZrO2 ) = −1100.8 kJ/mol, Hf (Al2 O3 ) = −1678.2 kJ/mol, Hf (CuO) = −155.2 kJ/mol and
Fig. 3. Zr 3d, Cu 2p and Al 2s core-level XPS spectra of the SLT Cu–Zr–Al alloys at 28.5, 29 and 30 A after sputtering with Ar ions for 1 min.
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Hf (Cu2 O) = −167.4 kJ/mol). Hence, from the thermodynamic viewpoint, Al2 O3 and ZrO2 should be more prone to grow compared to copper oxides. XPS results reveal that, even if not detected by XRD, a thin oxide layer is formed in all specimens. For instance, the XPS spectra of the alloy SLT at 28.5 A indicates the presence of ZrO2 , metallic Zr, CuO (Cu) and Al. This indicates that kinetics also plays a role on oxide formation. Fig. 4 shows the 3D profilometry images of the as-cast (a) and SLT alloys at 29 A (b) and 30 A (c). The Z-scale bar increases with the laser intensity compared to the as-cast alloy. It is clear that surface damage increases with laser intensity, therefore larger roughness values are obtained (Fig. 4d). No significant increase in Sq (root mean squared) and Sz (maximum height) were observed between the alloy SLT at 28.5 and 29 A but both values become much larger in the heavily treated alloy. It is interesting to note that roughness is homogeneous overall the laser treated specimens. Wettability results assessed in NaCl medium are presented in Fig. 5. The equilibrium contact angle ( eq ) in a non-reactive medium is defined by Young’s equation [30]: cos eq =
(SV − SL ) LV
(1)
where SV , SL and LV are the interfacial tensions between solid–gas, solid–liquid and liquid–gas phases, respectively. Our results indicate that the crystalline alloy exhibits markedly lower contact angle than the as-cast and SLT alloys at low intensities. However, a slight increase in the contact angle value between the as-cast and SLT alloys at 28.5 and 29 A is noticed. Since LV depends mainly on the liquid used for the sessile drop experiments, its value can be regarded as constant. Hence, the change in contact angle has to be attributed to variations in SV and/or SL . Laser treatments, as well as shot peening, high pressure torsion or annealing treatments, carried out under certain conditions, usually result in structural relaxation and the concomitant decrease in free volume concentration, thus reducing the atomic mobility. Moreover, structural
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relaxation may also be accompanied by cluster formation as a result of the high Cu–Zr binding energy which also contributes to decrease the free surface energy. Hence, lower SV values are expected. In turn, when the surface free energy of the solid decreases (i.e. in the relaxed state), the attraction between the liquid molecules and the atoms in the solids is lower than in the as-cast state, disfavouring the spreading of the liquid drop on the solid surface, hence increasing SL [31]. This concept could also be applied to crystalline materials, and consequently, even higher contact angles may be expected in the sample SLT at 30 A. However, lower contact angles are obtained in our case, suggesting the influence of other mechanisms. Actually, contact angles are larger on very rough surfaces than on chemically identical, smooth surfaces [32]. The observed increase in roughness (Fig. 4) could therefore have an influence on the hydrophobicity of the alloy. In addition, the alloy SLT at 30 A has larger O content at the surface. Hence, there is higher chance for O to combine with water molecule to form hydrogen bond, resulting in good wettability. These factors could explain the trends observed in this work. Fig. 6 displays the load (P) vs. penetration depth (h) curves of the as-cast and SLT alloys. Except for the alloy SLT at 30 A, the loading segments of these curves show serrated flow behaviour. These serrations are typically reported in metallic glasses, resulting from the creation and propagation of discrete shear bands [1,33]. Nanoindentation curves are consistent with XRD results: the alloy SLT at 30 A is completely crystalline; therefore no pop-in events are expected in the curve. Conversely, the other alloys, which are fully or mainly amorphous, show the existence of the serrations typically found in metallic glasses. The inset in Fig. 6 lists hardness and reduced Young’s modulus values. The increase in H and Er for the relaxed specimen (SLT at 28.5 A) can be explained because of the stronger bonding between atoms as a result of the annihilation of free volume. This enhancement could also be partly due to the formation of fine clusters. The decrease in hardness of the alloy SLT at 29 A can be explained by the formation of CuZr B2 phase. This
Fig. 4. Three-dimensional topography images of the (a) as-cast (b) SLT at 29 A and (c) SLT at 30 A alloys and (d) roughness parameters table (Sq = root mean squared and Sz = maximum height).
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Fig. 5. Photographs of as-deposited 3.5 wt% NaCl droplets onto the surface of (a) as-cast (b) SLT at 28.5 A (c) SLT at 29 A and (d) SLT at 30 A alloys.
to investigate the effect of surface laser treatments on wetting behaviour and mechanical properties of Cu–Zr–Al metallic glass. The lower contact angle in the crystalline specimen than in the ascast and SLT alloys at low intensities can be explained because of chemical and roughness contributions coming from oxide phase formation. Hardness results are explained in terms of free volume annihilation and phase formation. The alloy SLT at 28.5 A exhibits the largest hardness value. Acknowledgements This work has been supported by the MAT2011-27380-C02-01 and 2009 SGR 1292 projects. MDB acknowledges the ICREAAcademia Award. References Fig. 6. Load (P) vs. penetration depth (h) curves for the as-cast, SLT at 28.5 A, SLT 29 A and SLT 30 A alloys. In the inset the Hardness (H) and the reduced Young’s modulus (Er ) values are listed.
phase has been widely reported in the Cu–Zr system [34–37]; it is softer than its amorphous counterpart but it tends to improve the ductility of Cu-based BMGs via seeding and branching of shear bands. Lastly, the alloy SLT at 30 A exhibits large hardness values evidencing the harder nature of oxide phases. 4. Conclusions In this study Cu–Zr–Al alloy has been laser treated at different intensities. An in-depth structural characterization was employed
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