Investigation of the mechanical and microstructural evolution of a Cu based bulk metallic glass during ion irradiation

Intermetallics 116 (2020) 106655

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Investigation of the mechanical and microstructural evolution of a Cu based bulk metallic glass during ion irradiation J. Brechtl a, *, S. Agarwal b, M.L. Crespillo c, J. Salasin c, T. Yang c, H. Bei d, S.J. Zinkle a, b, c, d, ** a

Bredesen Center for Interdisciplinary Research and Graduate Education, University of Tennessee, Knoxville, TN, 37996, USA Department of Nuclear Engineering, University of Tennessee, Knoxville, TN, 37996, USA c Department of Materials Science and Engineering, University of Tennessee, Knoxville, TN, 37996, USA d Materials Science and Technology Division, Oak Ridge National Laboratory, Oak Ridge, TN, 37831, USA b

A R T I C L E I N F O

A B S T R A C T

Keywords: Metallic glasses Irradiation effects Annealing Microstructure Nanoindentation X-ray diffraction

Ion irradiation and annealing experiments were performed on Cu60Zr20Hf10Ti10 bulk metallic glass (BMG) specimens to investigate their irradiation- and temperature-induced microstructural and mechanical property evolution. For the ion irradiations, samples were exposed to 9 MeV Ni3þ ions to a midrange (~1.2 μm depth) dose of 10 displacements per atom (dpa) at temperatures ranging from room temperature to 360 � C (the corre­ sponding peak dose at ~2.8 μm depth was ~25 dpa). Bulk X-ray diffraction (XRD) and transmission electron microscopy (TEM) characterization revealed that the alloy did not crystallize during irradiation up to 290 � C but did partially crystallize at 360 � C. XRD analysis revealed that the crystallization which occurred in the sample irradiated at 360 � C was caused by thermal effects instead of irradiation displacement damage. Subsequent Rietveld refinement analysis of the XRD measurements revealed the presence of two distinct crystal phases, namely a CuTiZr hexagonal structure belonging to the P63/mmc space group and a CuTi tetragonal structure belonging to the P4/mmm space group. Nanoindentation experiments revealed that no pronounced hardness changes occurred in the specimens irradiated at room temperature and 290� C, although significant hardening was observed in the sample irradiated at 360 � C. The significant increase in the hardness at 360� C was ascribed to thermally induced partial crystallization of the alloy instead of the ion irradiation. In general, the results of the nanoindentation experiments and XRD characterization suggest that although the Cu BMG exhibits good stability during irradiation at temperatures up to 290 � C it is not suitable for irradiation environments where the tem­ perature is 360 � C for extended periods of time. The Lam and Chong extrapolation method, which has been used to study the indentation size effect (ISE) in amorphous alloys, was employed to quantify how irradiation and temperature affect this type of behavior in the BMG. However, the poor linear fitting of the indentation hardness data by this model indicate that a new ISE model is likely needed to quantify indentation hardening in BMGs.

1. Introduction In the 1980’s, Inoue et al. discovered methods to produce bulk metallic glasses (BMGs) with thicknesses exceeding 1 cm [1]. With respect to their material properties, metallic glasses can exhibit high strength [2–6] excellent hardness [7,8], wear and corrosion resistance [9–13], and high toughness [14]. Furthermore, BMGs possess a topo­ logically disordered structure that does not contain typical crystalline defects such as grain boundaries or lattice sites [15]. This lack of lattice sites required for the formation of Frenkel pair defects should make BMGs potentially viable for use as structural components in nuclear

fusion systems [16,17]. However, the suitability of BMGs in these types of environments will depend on their resistance to crystallization and other radiation-induced degradation processes after exposure to a range of irradiation doses and temperatures [18]. In the following paragraphs, a brief literature review of the irradiation studies on Cu based BMGs will be discussed. Cu–Zr-Hf-Ti BMGs, such as the one studied for the present investi­ gation, were found to have exceptional properties, such as a compressive fracture strength >2,000 MPa [19], fracture toughness of 67.6 MPa�m1/2 (room temperature) [20,21], a relatively high glass transition temperature (Tg) of 480 � C [22,23], and low cost as compared

* Corresponding author. ** Corresponding author. Bredesen Center for Interdisciplinary Research and Graduate Education, University of Tennessee, Knoxville, TN, 37996, USA. E-mail addresses: [email protected] (J. Brechtl), [email protected] (S.J. Zinkle). https://doi.org/10.1016/j.intermet.2019.106655 Received 13 February 2019; Accepted 2 November 2019 Available online 16 November 2019 0966-9795/© 2019 Elsevier Ltd. All rights reserved.

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to Zr-based and Pd-based BMGs [22,24]. For a comparison, typical steels have a fracture strength and toughness of 1,000 MPa and ~100 MPa�m1/2, respectively [25,26]. Furthermore, these BMGs exhibit deformation in the form of shear bands instead of dislocations [27,28], which surprisingly contribute little to the overall plastic strain during deformation [30]. Finally, the plastic strain in BMGs is thought to be related to the phenomenon of “fertile sites” [31] or “liquid-like” zones [32] that occupy the glassy matrix. Carter et al. examined the radiation response of Cu50Zr45Ti5 metallic glass ribbons that were irradiated by 1 MeV Cuþ ions to a fluence of 1 � 1016 ions/cm2 (17 dpa) at room temperature [33]. Bright-field TEM imaging revealed that partial crystallization that occurred during irra­ diation, and was due to irradiation effects. Furthermore, X-ray diffrac­ tion and selected electron area diffraction (SAED) revealed that the nanocrystals were composed of Cu10Zr7 and CuZr2 phases with di­ ameters which ranged in size from 2 to 14 nm. From the results, it was hypothesized that the nanocrystal formation during irradiation is a consequence of the enhanced atomic mobility that is caused by the introduction of excess free volume via energetic atomic mixing. An investigation by Mei et al. compared the irradiation resistance of (Cu47Zr45Al8)98⋅5Y1.5 BMG with polycrystalline W during He irradiation [34]. Here, they exposed the materials to 500 keV He2þ at room tem­ perature to fluences of 2 � 1017 ions/cm2 [4 dpa, 10 atomic percent (at. %) He], 1 � 1018 ions/cm2 (20 dpa, 35 at.% He), and 2 � 1018 ions/cm2 (40 dpa, 50 at.% He). Scanning electron microscopy (SEM) showed that surface peeling and delamination, as caused by He bubble formation and expansion, occurred in the W alloy during irradiation whereas the BMG only exhibited slight damage at depths within the ion range. Further­ more, XRD patterns revealed that the BMG specimens did not crystallize under any irradiation condition. Therefore, this study provides impor­ tant evidence that Cu based BMGs have superior radiation resistance as compared to polycrystalline W. Subsequently, Wang et al. extended the previous study and included Zr and Co based BMGs, in addition to Cu based BMGs [35]. Here the samples were exposed to the same irradiation conditions as described in Ref. [34]. The XRD patterns indicated that all the BMGs retained their amorphous structure under the given irradiation fluences, further showing that these materials have irradiation resistance. Also, the highest degree of damage due to He embrittlement was seen in the W alloy, followed by the Co BMG. In contrast, there was no visible damage observed in the Zr and Cu based BMGs. As compared to the Zr and Cu based BMGs, the higher embrittlement in the Co based BMG and W alloy was thought to be linked to the relatively lower atomic spacing, which can less accommodate the trapped He bubbles. Therefore, from Refs. [34,35], two important points can be concluded. Firstly, the results showed that all the BMGs displayed greater resistance to He damage as compared to the W alloy. Secondly, the Zr and Cu based BMGs demonstrated better resistance than the Co based BMG. However, there are very few papers available in the literature that have explored the irradiation induced mechanical property changes of Cu based BMGs. One of the important works that is worth highlighting is by Zhang et al., where they investigated the effects of ion fluence on the nanoindentation hardness of (Cu47Zr45Al8)98⋅5Y1.5 BMG [36]. The irra­ diations were performed with 3 MeV Ar12þ at room temperature to maximum ion doses of 7.5 � 10 2 dpa (1 � 1014 ions/cm2), 7.5 � 10 1 dpa (1 � 1015 ions/cm2), and 7.5 dpa (1 � 1016 ions/cm2), which cor­ responded to a range of ~1.2 μm. Similar to Ref. [34], the samples did not crystallize during irradiation to the above fluences. Subsequent nanoindentation experiments showed that within the irradiated region of the alloy (<1000 nm), the hardness decreased (or softening increased) with increasing ion fluence. The decrease in the hardness was thought to be caused by an irradiation induced atomic disordering in the glass [37]. Although the investigations, as discussed above, gives some insight on the irradiation response of Cu based BMGs, there is still little known. Due to this issue, a literature survey was conducted on non-Cu based

BMGs, where investigations have revealed that both irradiation induced softening and hardening can occur [37–41]. For instance, irradiation-induced hardening that was accompanied by crystallization was observed in Ref. [38], whereas softening was reported in Refs. [37, 39–41]. More recent studies have examined the combined effects of irradiation dose and temperature on the evolution of mechanical prop­ erties in Zr52⋅5Cu17⋅9Ni14⋅6Al10Ti5 (BAM-11) BMG [42–45,99]. Perez-Bergquist et al. found that in samples irradiated by 3 MeV Niþ ions to a midrange dose of 0.1 dpa at 200 � C significant softening occurred [42]. A later study found that the hardness also decreased after being annealed at 300 � C for 48 h and subsequently irradiated (3 MeV Niþ) at room temperature to a midrange dose of 10 dpa [43]. In contrast to the softening observed in Refs. [42,43], hardening was reported in the BAM-11 BMG after irradiation by 9 MeV Ni3þ to a midrange dose of 10 dpa at 360 � C [44,45]. Importantly, this hardening during irradiation was found to be associated with the partial crystallization of the specimen. Despite these recent efforts, the response of BMGs (especially Cu based BMGs) during irradiation at different temperatures is still largely unknown [42]. To broaden the fundamental understanding of the thermal and irradiation response of this amorphous alloy system, the present work focuses on the ion bombardment of a Cu based BMG at different temperatures. Here, Cu60Zr20Hf10Ti10 BMG was exposed to 9 MeV Ni3þ ion irradiation at temperatures ranging from room tem­ perature up to ~75% of Tg. 2. Experimental procedures 2.1. Sample preparation The arc melting procedure under an argon atmosphere, was done at Oak Ridge National Laboratory to fabricate the Cu60Zr20Hf10Ti10 alloy using a mixture of base metals Zr, Cu, Hf, and Ti. Here, the purity for each metal was 99.99%. The provided rod had a cylindrical geometry with a length of 7 mm and diameter of 5 mm. To confirm that the ma­ terial was fully amorphous, sections of the drop cast rod were evaluated via X-ray diffraction (XRD) and differential scanning calorimetry. The specimens with nominal dimensions of 3 mm diameter with 0.4–0.5 mm thickness were made using a combination of diamond saw and me­ chanical grinding. The 0.4–0.5 mm thickness, and particularly the 3 mm diameter, were intentionally chosen so that the specimens could fit into the stage that is typically used to irradiate relatively thin TEM disks. Furthermore, to remove the surface defects the specimens were ground via a Minimet 1000 polishing/grinding system. As a final step, the samples were polished to a mirror finish using colloidal silica particles. 2.2. Ion irradiations The ion irradiation experiments were conducted at the Ion Beam Materials Laboratory at the University of Tennessee, Knoxville (UTK), which is equipped with a 3 MV tandem accelerator and three highenergy beamlines [46]. The specimens were irradiated by 9 MeV Ni3þ ions at three different temperatures of 25, 290, and 360 � C. The ion energy was selected to produce a reasonably deep irradiated region without causing “swift heavy ion” effects that emerge at higher ion energies. The accumulated ion fluence during the irradiation was calculated from the beam spot size that is normally regulated by the double slits on the beamline close to the target chamber [46]. It is important to note that two samples for each temperature condition were irradiated simultaneously using the sample holder. A tungsten heating element, which is mounted behind the Mo plate, was used to increase the temperature during irradiation. To mitigate issues associated with beam heating and charge accumulation on the samples, a low Ni3þ beam current density of ~480 nA/cm2 was used [47]. During the ion implantations, the ion beam was scanned both vertically and horizontally for lateral homogeneity of the beam. The 2

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beam homogeneity was within 10% throughout the irradiated area. The method used to check the homogeneity can be found elsewhere [100]. Adjustable slits were used to maintain the irradiation area of 4 � 9 mm2 on the sample surface. For accurate sample temperature control, two K-type (chromel-alumel) thermocouples were used. One thermocouple was spot welded to the surface of an adjacent dummy sample that was composed of an FeCrAl alloy. Furthermore, the second thermocouple was attached to the surface of the top plate outside of the beam region to provide an estimate of the baseline temperature that is not affected by ion beam heating. During the entire experiment, the temperature did not fluctuate significantly, indicating that substantial beam heating did not occur. More information regarding the experimental apparatus used in the irradiation can be found in Ref. [47]. The depth dependent dpa, ion range, and the corresponding Ni concentration was estimated using SRIM 2013 (see Fig. 1). These values were calculated using quick Kinchin-Pease calculations of SRIM 2013 [48,49], as suggested by Stoller et al. using threshold displacement en­ ergies of 40 eV, where a density of 8.6 g cm 3 was assumed. The pro­ jected ion range was calculated to be ~3.0 μm, with implanted Ni concentrations of ~0.36 at. % at this depth. For these implantations, the calculated damage and displacements per atom (dpa) for the midrange and peak doses, respectively, were 10 (1.2 μm) and 25 dpa (2.8 μm). To achieve the prescribed doses, an ion flux in the range of ~9.5 � 1011 cm 2s 1 was used. The time for each irradiation was approximately 7 h, which produced a fluence of 2.4 � 1016 ions/cm2. To minimize the inaccuracies associated with surface effects and diffusional broadening of the implanted ions [19,50,51] the microstructure and mechanical properties were examined at an intermediate depth of ~1.2 μm, which corresponds to a midrange dose of 10 dpa (dotted line in Fig. 1). Finally, the ionization stopping power values at the surface and the midrange region were found to be 5.4 keV/nm and 2.0 keV/nm, respectively.

ranging from 0.5 to 1� (~1.2 μm penetration depth) were used. For further details regarding the calculation of the prescribed angles, please see Ref. [52]. TEM characterization was performed at Oak Ridge National Labo­ ratory to obtain micrographs and selected area electron diffraction patterns using a JEOL JEM-2100F TEM/STEM at an accelerating voltage of 200 kV. The cross-section TEM foils were fabricated using an FEI Quanta Dual-beam focused ion beam (FIB)/SEM system. Here, 2 kV Gaþ ions were used for the thinning steps where the energy was gradually reduced from 30 keV to 3 keV. An ion current of 30 pA with a glancing angle of about 4� was used during the final thinning step to minimize ion beam milling damage. A Fischione Nanomill (model 1040) 4 keV Arþ ion polisher (at low incident angles) was used to polish the specimens after the thinning of the specimen was completed. Room temperature nanoindentation characterizations were per­ formed to observe whether any irradiation induced changes in the hardness occurred. These tests were performed using a Tencor-KLA G200 nanoindenter with a Berkovich diamond (3-sided pyramidal ge­ ometry) tip. To calibrate both the area function of the tip and the ma­ chine stiffness of the nanoindenter, indentions were performed on a standard fused silica sample [42]. The experiment was run in the continuous stiffness measurement mode with a constant loading rate of 400 μNs 1. To ensure adequate statistics, ~25 indents were made where hardness was measured as a function of depth from the point of contact of the nanoindenter with the surface to a depth of ~2,500 nm. To mitigate effects associated with the large data scatter that is associated with surface roughness, data from the surface to ~100 nm from the specimen surface was omitted. After testing, hardness was calculated using the Oliver and Pharr method [53,54]. 2.4. Indentation size effect analysis To study the influence of irradiation and annealing on the indenta­ tion size effect (ISE) of the Cu BMG, the hardness was plotted according to the following equation [55]: � � h* H ¼ H0 1 þ pffiffi (1) h

2.3. Characterization techniques X-ray diffraction (XRD) characterization was carried out at the Joint Institute for Advanced Materials (JIAM) Diffraction Facility at the Uni­ versity of Tennessee. The grazing incidence XRD was performed using a PANalytical Empyrean diffractometer equipped with a Xe proportional detector, 0.09� parallel plate collimator, and 0.04 Rad soller slits. The Xray consisted of a Cu beam with a K-alpha wavelength of 1.54 Å, in addition to an accelerating voltage and current of 45 kV and 40 mA, respectively. To determine whether any crystallization or phase decomposition occurred in the irradiation region, glancing angles

where H is the nanoindentation hardness, H0 is the extrapolated (bulk) hardness value that arises from the statistically stored clusters in the absence of the geometrically-necessary clusters that are associated with strain gradients [55,56], h is the indentation depth, and h* is a term which characterizes the depth dependence of the hardness. This model was derived by Lam and Chong to examine the ISE in metallic glasses [55], in which the depth dependent terms typically depend on material parameters such as H0 and the number of defects in the amorphous matrix. Importantly, h* also depends on other parameters of the metallic glass, such as the angle between the tip and surface of the sample, the local shear strain required to transform a single cluster, the number of clusters associated with the strain, the Helmholtz free energy associated with the shear transformation of clusters occurring, and the experi­ mental temperature. It should also be mentioned that in the present work, the concept of cluster defects will be extended to include other defect types such as liquid-like sites and soft-zones [57,58]. In Ref. [55], there was an apparent discrepancy in the linear fitting of the data that was not accounted for. Here, it was assumed that the fitted data, according to Eq. (1), fell on a straight line. However, there was an apparent deviation away from linearity for 1/h0.5 < 0.8 μm-0.5, which is roughly equivalent to indentation depths beyond ~1.6 μm. 3. Results

Fig. 1. Left) Irradiation damage (dpa) profile (corresponding to a fluence of 2.4 � 1016 ions/cm2) versus depth for the Cu60Zr20Hf10Ti10 BMG. Right) at. % Ni versus depth. Both graphs were obtained from SRIM 2013 simulation using quick Kinchin-Pease calculations (see text for more details). Dashed line points out the depth and dpa that was chosen for this study.

The XRD patterns of the as-cast and irradiated Cu BMG samples is presented in Fig. 2. For the samples bombarded at 25 and 290 � C, a broad peak centered around 32 to 46� can be observed, which indicates 3

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that the specimens remained amorphous during irradiation. For the sample irradiated at 360 � C, however, the XRD spectrum shows numerous sharp peaks which indicates the presence of crystalline pha­ ses. Glancing XRD was also performed on the unirradiated side of the specimen that was irradiated at 360 � C, in which the resulting pattern is displayed in Fig. 3. As can be observed, the pattern for the unirradiated side of the specimen is quite similar to the irradiated side, implying that the partial crystallization was thermally induced. Fig. 3 also compares the bulk XRD measurements for the sample irradiated at 360 � C (irra­ diated and unirradiated sides) for 7 h with an unirradiated specimen that was annealed at 520 � C for 2.5 h. The bulk X-ray had a penetrating distance of ~10-50 μm. This comparison was intended to examine the XRD behavior at ~120 � C below vs. ~40 � C above the reported Cu BMG glass transition temperature of 480 � C [22]. As can be seen, the patterns for the 360 � C irradiated sample are quite similar to the one corre­ sponding to the 520 � C thermally annealed specimen. However, for the pattern of the 520 � C annealed specimen, the peaks centered at ~37, 39, and 40� exhibited a greater intensity as compared to the same peaks for the sample irradiated at 360 � C. This indicates that a greater degree of crystallization occurred in the sampled heated at 520 � C for 2.5 h as compared to the 360 � C (~7 h) irradiated condition. Fig. 4 presents a medium-magnification TEM bright field image of the specimen that underwent irradiation to a midrange dose of 10 dpa at 360 � C. The corresponding diffraction pattern is displayed in the upper left corner of the figure. Evidence for polycrystallization can be clearly observed in the matrix of the irradiated specimen. Furthermore, the crystallites that formed throughout the irradiated region consisted of various shapes and sizes. Subsequent powder diffraction in conjunction with qualitative and quantitative phase analysis was performed to analyze the crystal phases in the alloy (see Fig. 5). The inset of the figure shows the same data for scattering angles between 30 and 50� . Here, qualitative phase identifi­ cation (ID) was determined in Highscore Plus [101] using the Powder Diffraction File-4þ (PDF-4þ) database while quantitative phase analysis was performed via Rietveld refinement via the GLAS Science Algorithm Software II (GSAS II) software [59]. The results of the Rietveld refine­ ment analysis can be seen in Table 1, which shows that the crystalline phases are dominated by some form of CuZrTi solid solution phase at 81.7 wt % with the secondary phase being CuTi at 18.3 wt %. It was also determined that the crystallites were most likely composed of tetragonal CuTi (space group P4/mmm) [60] and hexagonal CuTiZr (space group P63/mmc) [61]. In addition, the results indicated that Hf did not combine to form any compounds, and therefore most likely remained in the matrix.

Fig. 3. Glancing X-ray diffraction patterns of the irradiated and unirradiated sides of the Cu60Zr20Hf10Ti10 BMG specimen irradiated by Ni3þ at 360 � C (9 MeV, 10 dpa) in addition to a bulk scan of the sample annealed at 520 � C for 2.5 h.

Fig. 4. TEM bright field imaging and inset which contains the corresponding diffraction pattern of the irradiated region in the Cu60Zr20Hf10Ti10 BMG sample irradiated with 9 MeV Ni3þ ion beam to a dose of 10 dpa at 360 � C.

Fig. 5. Rietveld refinement analysis of X-ray diffraction pattern of partially crystallized Cu60Zr20Hf10Ti10 BMG after irradiation by 9 MeV Ni3þ to midrange dose of 10 dpa at 360 � C.

The depth dependent nanoindentation hardness for the as-cast and the ion irradiated samples is presented in Fig. 6(a). For each condition there was a noticeable indentation size effect, as the hardness values decreased with respect to the indentation depth. In addition, the

Fig. 2. Glancing X-ray diffraction patterns for Cu60Zr20Hf10Ti10 BMG samples in the as-cast and Ni3þ irradiated condition (9 MeV, 10 dpa) at different temperatures. 4

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tip is sensitive to substrate regions that are 5–10 times beyond its penetration depth [42,62]. Here, it is important to note that the indent depth of 450 nm coincides with the transition point between the ion irradiated and unirradiated regions. To demarcate these two regions in the irradiated specimen, a vertical line was inserted into Fig. 6(a). Fig. 6(b) displays the depth dependent (direction normal to the irradiated surface) hardness data for the irradiated region, which cor­ responds to depths of 100–450 nm beneath the surface. In the irradiated region, the hardness for the room temperature condition did not noticeably increase as compared to the values for the as-cast condition. The hardness values for the unirradiated region exhibited a large in­ crease for the 360 � C exposure condition, where the sample partially crystallized. Also, the hardness values for the specimen irradiated at room temperature were higher (greater than a standard deviation) for depths beyond 900 nm as compared to the as-cast condition [see Fig. 6 (a)]. This discrepancy in the hardness is probably linked to slight vari­ ations in the homogeneity of the sample. The corresponding hardness data for the indentation depths of 200 and 450 nm are presented in Table 2. As can be seen, the hardness values were higher at a depth of 200 nm as compared to that of 450 nm. Furthermore, there was no significant change in the irradiated vs. un­ irradiated hardness at 25 and 290 � C, whereas there was a significant increase in hardness for the sample irradiated at 360 � C due to partial crystallization (in both the irradiated and unirradiated regions). To compare the effects of temperature and irradiation damage on the nanoindentation properties in the 360 � C partially crystallized sample, the nanoindentation hardness was measured in the irradiated and un­ irradiated regions of the surface exposed to irradiation. The unirradiated region of the surface corresponded to the polished front surface shielded from the ion beam. Fig. 7 compares the nanoindentation hardness of the two sections, and as can be observed, the hardness values were quite similar for all indentation depths. Fig. 8(a)–(c) display the ISE hardness data (H vs. h 0.5) as well as the corresponding linear fit from the Lam and Chong model [Eq. (1)], as applied to the as-cast and irradiated (9 MeV Ni3þ) specimens. Similar to Fig. 6(a), a vertical line was inserted at ~1.5 μm 0.5 in Fig. 8(a) to separate the nonirradiation and irradiation-dominant regimes. As can be observed, the fitted values based on this model for the as-cast and the irradiated specimens did not follow the predicted linear behavior versus 1/h0.5 throughout the entire indentation region, indicating that the quantitative values derived from this analysis are somewhat unreliable. The corresponding fitting parameters for the irradiated region (100–450 nm) are presented in Table 3. As can be seen in the table, the fitted H0 hardness values did not vary significantly for the unirradiated (as-cast) and irradiated samples at 25 and 290 � C. For the near-surface irradiated region, the extrapolated bulk hardness was ~7.1–7.6 GPa for the as-cast materials and for the samples irradiated up to 290 � C. The detailed fitted hardness values imply a slight radiation-induced soft­ ening. The fitted reduction in hardness was most pronounced for the specimen irradiated at 290 � C ( 6.6%), while the sample bombarded at 25 � C softened by about 1.3%. On the other hand, H0 was 9.3 GPa for the specimen irradiated at 360 � C, which corresponded to an increase in the hardness of 20% as compared to the as-cast state due to the partial crystallization of the material. The hardness parameters, as derived from the Lam and Chong model [Eq. (1)] for indentation depths 500–~2,500 nm (unirradiated region),

Table 1 Results of the Rietveld refinement which list the goodness of fit (χ2), software used for the analysis, background fitting function, lattice parameters (a, b, c), the crystal system, the space group, the crystallite volume, and the weight percent for the multi-phase Cu60Zr20Hf10Ti10 BMG. CuZrTi Refinement: Goodness of fit, (χ2) Software Variables Background Lattice parameters Crystal Data: Crystal system Space group

a ¼ ba [Å] ca, [Å] Volume, [Å3] Weight Percent a

CuTi

4.66 GSAS II [59] Chebyschev-Background a ¼ b, c

a ¼ b, c

Hexagonal ð194Þ P63 =mmc

Tetragonal ð123Þ P4=mmm

5.167 (2) 8.265 (2) 191.12 (2) 81.7 (7) %

3.182 (3) 2.853 (3) 28.904 (9) 18.3 (5) %

Estimated standard deviations are 3σ.

hardness values corresponding to a midrange dose of 10 dpa and at the projected ion range were examined, which respectively coincide to distances of 1.2 and 2.8 μm beneath the surface. However, these values actually correspond to depths of 200 and 450 nm since the nanoindenter

Table 2 Summary of nanoindentation results on virgin and irradiated Cu60Zr20Hf10Ti10 BMG samples (9 MeV Ni3þ, dose of 10 dpa) at depths of 200 and 450 nm. Fig. 6. Nanoindentation hardness vs. indentation depth for Cu60Zr20Hf10Ti10 BMG as-cast and irradiated samples (9 MeV Ni3þ to 10 dpa) at different tem­ peratures ranging from ambient to 360 � C for (a) depths ranging from 100 to 2,500 nm and (b) the irradiated region (100–450 nm). Each data point is based on the mean values for ~25 indents.

Average Hardness (GPa)

5

Depth (nm)

As-cast

25 � C

290 � C

360 � C

200 450

8.4 � 0.3 7.6 � 0.1

8.4 � 0.2 7.6 � 0.1

8.4 � 0.3 7.4 � 0.2

10.6 � 0.4 9.5 � 0.2

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Table 3 Comparison of the results for the parameters h, H0, and the percent change in H0 [Eq. (1)] for the as-cast and the 9 MeV Ni3þ irradiated Cu60Zr20Hf10Ti10 BMG for the irradiated region (100–450 nm). Condition

h* (nm)

H0 (GPa)

Change H0 (%)

As-cast 9 MeV Ni3þ 25 � C 9 MeV Ni3þ 290 � C 9 MeV Ni3þ 360 � C

33.8 38.3 63.2 47.4

7.6 7.5 7.1 9.3

– 1.3 6.6 22

are presented in Table 4. With regards to the extrapolated bulk hardness, H0, the values ranged from ~4.7 to 5.0 GPa for the as-cast and irradiated samples (25 and 290 � C). Comparing the results for the irradiated and unirradiated regions for the specimen irradiated at 360 � C, H0 was 9.3 and 6.0 GPa, respectively. This difference in the hardness was due to non-linearity of the data, which produced different extrapolation values. For all conditions, as listed in Table 4, the characteristic depth values, i. e. h*, were one order of magnitude greater than those from the nearsurface (ion irradiated) region, as listed in Table 3.

Fig. 7. A comparison of the nanoindentation hardness in both the irradiated region and the unirradiated substrate (polished front surface shielded from the ion beam) in the Cu60Zr20Hf10Ti10 BMG alloy after irradiation by 9 MeV Ni3þ to 10 dpa at 360 � C.

4. Discussion The glancing XRD characterization shows that the Cu60Zr20Hf10Ti10 BMG remained amorphous during the ion irradiation up to 290 � C. This resistance to crystallization during irradiation at temperatures well below Tg has also been observed in other studies [34,37,39,42–45, 63–67]. For instance, Brechtl et al. [44,45] found that Zr52⋅5Cu17⋅9Ni14⋅6Al10Ti5 (BAM-11) BMG did not crystallize when irra­ diated by 9 MeV Ni3þ ions to a midrange dose of 10 dpa at room tem­ perature and 290 � C. In another study, Perez-Bergquist et al. found that the same alloy did not crystallize when irradiated by 3 MeV Niþ ions to a midrange dose of 1 dpa at 200 � C [42]. Interestingly, Yuka et al. observed that 200 MeV Xe irradiation (dose of 0.25 dpa) of Zr50Cu40Al10 BMG at room temperature lead to a decrease in the free volume content of the material without inducing crystallization [65,67]. The crystalli­ zation resistance, as exhibited by the BMGs during irradiation at lower temperatures, could be attributed to various factors. These factors include a high required driving force for irradiation-induced crystalli­ zation and low irradiation enhanced diffusivity [68]. On the other hand, the Cu BMG sample partially crystallized when irradiated at 360 � C for 7 h. This result indicates that the material is not suitable for irradiation environments where it will be exposed to tem­ peratures greater than ~0.75 Tg for prolonged periods of time. Glancing XRD that was performed on the unirradiated side of the specimen (see Fig. 3) indicates that the partial crystallization that occurred during irradiation at the highest temperature was a consequence of thermal instead of irradiation effects. Bulk XRD measurements were also per­ formed on a separate specimen that was annealed at 520 � C for 2.5 h, which are displayed in the same figure. The temperature of 520 � C was chosen since it is above the reported crystallization temperature (Tx) of the alloy [22]. The pronounced similarity between all three patterns further suggests that crystallization occurred in the material due to thermal, instead of irradiation effects. Other studies have reported the irradiation induced crystallization of Cu based BMGs. For instance, Xie et al. studied the nanocrystallization Table 4 Comparison of the results for the parameters h*, H0, and the percent change in H0 [Eq. (1)] for the as-cast and the 9 MeV Ni3þ irradiated Cu60Zr20Hf10Ti10 BMG for the unirradiated region (500–~2,500 nm).

0.5

Fig. 8. H vs. 1/h (with extrapolated hardness Ho) for Cu60Zr20Hf10Ti10 BMG in the as-cast vs. irradiated samples with 9 MeV Ni3þ to 10 dpa at different temperatures ranging from ambient to 360 � C for the (a) depths ranging from 150 to 2,500 nm, (b) the unirradiated region (450 nm–2,500 nm), and (c) the irradiated region (150–450 nm).

6

Condition

h* (nm)

H0 (GPa)

Change H0 (%)

As-cast 9 MeV Ni3þ 25 � C 9 MeV Ni3þ 290 � C 9 MeV Ni3þ 360 � C

7.2 � 102 3.7 � 102 7.8 � 102 8.0 � 102

5.0 5.8 4.7 6.0

– 16.0 6.0 20

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Intermetallics 116 (2020) 106655

behavior of a Cu50Zr45Ti5 metallic glass melt spun ribbons [69]. Here, the specimens were irradiated by 200 keV electrons to fluences ranging from 2.4 � 1021 cm 2 to 2.9 � 1022 cm 2 at room temperature. The nanocrystallites that formed during electron irradiation (room temper­ ature) were identified as monoclinic CuZr phase. However, a later study found that Cu10Zr7 phases formed in the same alloy after 200 keV electron irradiation (room temperature) to a fluence of 4.4 � 1022 cm 2 [70]. Another investigation studied the Ar-ion-milling induced nano­ crystallization of Cu50Zr45Ti5 metallic glass [71]. Here, the samples were irradiated by Ar ions with accelerating voltages of 2–4 kV at room temperature. Subsequent TEM characterization revealed the presence of Cu10Zr7 nanocrystals. Another study, which examined the crystalliza­ tion of Cu50Zr45Ti5 metallic glass during irradiation by 1 MeV Cu ions (room temperature), observed that Cu10Zr7 and CuZr2 phases formed in the matrix [33]. Another investigation, which examined the combined effects of pressure and temperature on Cu60Zr20Hf10Ti10 BMG, observed that Cu10Zr7 formed during exposure to 700 � C, 5 GPa during in-situ XRD experiments [22]. Based on the above discussion, therefore, it can be inferred that CuZr polymorphs typically form in Cu based BMGs during crystallization. However, the results of the present study appear to contradict these previous results. As discussed earlier, the Rietveld refinement (see Table 1) found that the crystalline structures that formed in the sample irradiated at 360 � C consisted of tetragonal CuTi and hexagonal CuZrTi phases. However, it should be stated that in Ref. [22] there were some crystalline peaks that were not identified in the XRD pattern and could have been representative of either the CuTi or CuZrTi phases. Carter et al. hypothesized that crystallites which formed in the Cu50Zr45Ti5 BMG during 1 MeV Cu irradiation (room temperature) was a consequence of the enhanced atomic mobility resulting from the intro­ duction of excess free volume during quenching [33]. In this scenario, the free volume is spread over a relatively large region, resulting in increased atomic mobility that enhances the short-range order which leads to crystallite nucleation. It should be noted that the alloy studied in our work did not crystallize during irradiation by a more energetic ion (9 MeV) at temperatures up to 290 � C. This lack of induced crystalliza­ tion may be a consequence of the higher number of elements with differing radii. As compared to the Cu50Zr45Ti5 BMG, the greater degree of atomic size mismatch in the Cu60Zr20Hf10Ti10 BMG favors a more dense random packing structure that increases the difficulty of atomic rearrangement which leads to crystallization [72–74]. Kraposhin et al. reported that irradiation can induce a different crystallization mechanism in metallic glass, as compared to thermal annealing [75]. Here, Co86⋅7Fe3⋅6Si2⋅7Mn3⋅5B3.5 metallic glass ribbons were observed to crystallize after irradiation by 30 keV Arþ ions to a dose of 840 dpa at 200 � C. However, the same material did not crys­ tallize during thermal annealing until the temperature reached 500 � C. The crystallization of the metallic glass during irradiation at 200 � C may have been caused by an increased temperature in the sample induced by beam heating [76]. Furthermore, the temperature during irradiation may have been higher than reported if the chamber temperature was not being accurately monitored. The fact that the authors did not provide details on how the specimen temperature was monitored during irra­ diation raises questions as to the accuracy of the reported temperature. In terms of the microstructure, XRD revealed that during thermally induced (no irradiation) crystallization, the Co-based solid solution and CoB crystalline phases formed during devitrification. In addition to these two phases, an unknown cubic aperiodical phase was observed after irradiation. This phase was thought to be isostructural to the α-Mn aperiodical cubic phase. Additionally, thermal annealing after irradia­ tion (250� and 300 � C for 1 h) was found to recover the amorphous phase. Interestingly, this newly recovered amorphous phase was thought to contain non-crystalline long-range order structures in the matrix. These structures were thought to consist of polytopes with a noninteger order symmetry axis [77]. Table 5 summarizes the results of previous irradiation studies on

BMGs. As can be observed, some of the metallic glasses crystallized during irradiation (or exposure to high temperatures during irradiation), while others remained amorphous. Several interesting or puzzling re­ sults can be seen in Table 5. For instance, Ti40Zr10Cu38Pd12 BMG was found to crystallize during Ar (200 keV followed by 75 keV) ion irradi­ ation at 347 � C, whereas Zr55Cu28Al10Ni7 BMG did not [78]. Neither of the two BMGs experienced crystallization during room temperature irradiation [78]. This result is qualitatively consistent with expectations since the Tg, Tx, and ΔTx of the Zr based BMG (429 � C, 498 � C, and 69 � C [79,80]) is larger than that of the Ti based BMG (412 � C, 447 � C, and 35 � C [81]), indicating that the former BMG has greater glass formability and thermal stability [82]. It should be noted that in nearly all of the investigations listed in Table 5 there was no mention as to how the temperature was monitored during irradiation, which raises serious concerns as to the accuracy of the reported irradiation temperatures. Nevertheless, the results of these studies have important implications pertaining to the irradiation response of amorphous alloys. Considering only the few ion irradiation studies that utilized thermocouples to monitor the irradiation temperature ([42,44,45] and this study) along with the studies from Refs. [78,83], it appears that recrystallization has been reported only for irradiation temperatures approaching Tg for the given BMG. Secondly, in some metallic glasses total irradiation dose appears to be a greater contributing factor than ion mass or energy as to whether an amorphous alloy will crystallize. Finally, it appears that the composition (chemical complexity) of the metallic glass also appears to play an important role in determining its radiation resistance. More specifically, the enhanced irradiation resistance may arise from both the mass disparity and the “confusion principle”, where a greater number of elements comprising the BMG enhances its glass forming ability (and Tg), making it more stable [84,85]. In addition to the constituent ele­ ments comprising the alloy, oxygen impurities in the matrix may also promote crystallization of the material during irradiation [86]. Here, the impurities can aid in the heterogeneous nucleation of intermetallic phases, effectively decreasing the temperature or heating time at which crystallization can occur. On the other hand, beam heating may have been a contributing factor to the crystallization of the specimen during irradiation. It has been reported that in commercial ion implanters, ion beam heating can result in temperature increases approaching hundreds of degrees [76]. In the worse-case scenario, this could allow the radiation environment to exceed temperatures in which thermal-induced crystallization would occur. As presented in Fig. 6(a)–(b), only slight hardening was observed in the Cu BMG after irradiation by 9 MeV Ni3þ to midrange doses of ~10 dpa (25 dpa at peak damage region) at 25� and 290 � C. On the other hand, the sample irradiated at 360 � C (~7 h) exhibited a significant in­ crease in hardness, indicating that the sample crystallized during irra­ diation. Since the nanoindentation hardness (Fig. 6) and XRD evaluation (Fig. 3) indicate similar crystallization and depth-dependent hardness changes for the unirradiated and irradiated regions, these changes appear to be associated with thermal annealing effects rather than irradiation damage. Crystallization induced hardening has also been observed in He irradiated Cu50Zr45Ti5 metallic glass [87]. Here, the crystallization occurred in the samples after irradiation by 140 keV He (room temperature) to a dose of 4 dpa (11 at.% He). Post-irradiation TEM characterization revealed the presence of either He bubbles and/or voids in the matrix. Moreover, microindentation tests revealed that the sample exhibited greater hardness as compared to the as-cast state. This increase in the hardness was the highest at the interrogated depth of 600 nm below the surface, which coincided with either the maximal nuclear stopping power or the presence of implanted He at that depth. The poor linear fit that can be observed in Fig. 8(a) for both the unirradiated and irradiated specimens, which suggests that the model as derived by Lam and Chong [Eq. (1)] [88], is not valid for analyzing the ISE in the Cu BMG. This deviation away from linearity is especially 7

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Intermetallics 116 (2020) 106655

Table 5 Summary of ion irradiation studies in amorphous alloys R.T.: Room temperature. Alloy

Specimen Type

Ion Species

Ti40Zr25Be30Cr5

Bulk

C, Cl

Ti40Zr10Cu38Pd12

Bulk

Ar

Fe81B13⋅5Si3⋅5C2 Fe80Si7⋅43B12.57 Fe78B13Si9 Fe79B16Si5 Fe40Ni40P14B6 Ni52⋅5Nb10Zr15Ti15Pt7.5 Ni52⋅5Nb10Zr15Ti15Pt7.5 Co61⋅2B26⋅2Si7⋅8Ta4.8 Co86⋅7Fe3⋅6Si2⋅7Mn3⋅5B3.5 Zr55Cu28Al10Ni7

Ribbon Ribbon Ribbon Ribbon Ribbon Ribbon Ribbon Bulk Ribbon Bulk

He H He He He Ni Xe He Ar Ar

Zr64Cu17⋅8Ni10⋅7Al7.5 Zr52⋅5Cu17⋅9Ni14⋅6Al10Ti5

Bulk Bulk

Zr52⋅5Cu17⋅9Ni14⋅6Al10Ti5

Ion Energy (MeV)

Irradiation Dose Rate (dpa/s)

Irradiation Dose (dpa)

Crystallized? (Y/N)

Temperature Measurement

Ref.

R.T.

N

–

[39]

Y N Y N Y N Y Y N N Y N N N N

– – – – – – – – – – – –

[78] [90] [91] [92] [93] [94] [95] [83] [35] [75] [78]

– Thermocouple

[35] [42]

N N Ya N Y Y

Thermocouple

[44, 45]

– – –

[65] [96] [66, 97] [98] [36] [33] [34, 35] This work

Irr. Temp (0C)

4.2 � 1012

0.008

– 6.2 � 1013 – 5.0 � 1013 3.1 � 1013 1.0 � 1012 – 4.3 � 1012 3.1 � 1014 4.2 � 1012

– 0.0003 – 0.005 0.0009 0.0018 – 0.0009 0.17 0.006

0.05 5 2 1 28 18 0.00014 40 840 22 (total)

He Ni

0.2, 0.075 (consecutive) 2.8 0.25 2.8 0.005 0.04 1 18 0.5 0.03 0.2, 0.075 (consecutive) 0.5 3

0.1(C) 5.7 (Cl) 29 (total)

4.3 � 1012 2.1 � 1012

0.00009 0.005

40 1

Bulk

Ni

9

9.2 � 1011

0.0004

10

Zr50Cu40Al10 Zr55Cu30Al10Ni5 Zr55Cu30Al10Ni5

Bulk Ribbon Bulk

Xe Cu Co

200 1 0.04

– 6.3 � 1010 1.5 � 1014

– 0.0001 0.34

0.25 16 160

347 R.T. R.T. R.T. R.T. R.T. 400 R.T. R.T. R.T. 200 347 R.T. 125 R.T 200 R. T. 290 360 R.T. <50 <140

Zr50Cu35Al7Pd5Nb3 (Cu47Zr45Al8)98⋅5Y1.5 Cu50Zr45Ti5 (Cu47Zr45Al8)98⋅5Y1.5

Ribbon Bulk Ribbon Bulk

Kr Ar Cu He

1 3 1 0.5

– – 6.2 � 1011 5.8 � 1013

– – 0.001 0.0012

0.1 7.5 16 40

420 R.T. R.T. R.T.

Y N Y N

– – – –

Bulk

Ni

9

9.2 � 1011

0.0004

10

R. T. 290 360

N N Ya

Thermocouple

Cu60Zr20Hf10Ti10

a

25

Ion Flux (cm 2s 1) –

–

Sample crystallized via thermal effects instead of irradiation damage.

noticeable in the unirradiated region of the specimens, where the de­ viation becomes more significant at shallower depths. For the as-cast specimen, this behavior in the shallower region is unexpected since a variation in hardness vs. depth should not be observed. It is thought that the inadequate fit may arise from either an inherent limitation of the model or surface artifacts created during sample preparation. As discussed previously, there was a discrepancy in the linear fitting (for H vs. h 0.5) of the Lam and Chong model that was not accounted for in Ref. [55]. This discrepancy was a result of the apparent deviation away from linearity for indentation depths beyond ~1.6 μm. The above result suggests that this model may not be applicable for metallic glasses. A second model, which is based on the Nix-Gao extrapolation method [89], was also applied to the data. However, this model gave a poorer fit as compared to the previous technique [Eq. (1)]. From a mechanics standpoint, it is not surprising that the Nix-Gao model does not provide an adequate fit to the data, since it is based on proximity to geometri­ cally necessary dislocations to induce plastic deformation whereas dis­ locations are not present in amorphous materials. From the above discussion, it can therefore be inferred that a new model may be needed to more accurately quantify the ISE of the Cu60Zr20Hf10Ti10 BMG.

crystallization that occurs during prolonged (~7 h) exposure at 360 � C. However, as revealed by the XRD patterns, there does not appear to be any pronounced radiation-enhanced crystallization of the irradiated samples at 25–290 � C. TEM characterization revealed that the crystal­ lites in the sample irradiated at 360 � C consisted of various morphol­ ogies. Subsequent Rietveld refinement indicated that crystallized portion of the sample irradiated at 360 � C were comprised of hexagonal Cu–Zr–Ti (P63/mmc) and tetragonal Cu–Ti (P4/mmc) phases. In terms of the nanoindentation hardness, only modest changes were observed after irradiation by 9 MeV Ni3þ ions to a midrange dose of ~10 dpa (peak dose of 25 dpa) at temperatures of 25� and 290 � C. In contrast, the hardness was significantly higher for the specimen irradiated at 360 � C for ~7 h, and is a consequence of the partial crystallization that occurred due to thermal annealing effects. Importantly, the similar hardness values observed in both the irradiated and unirradiated regions of the near-surface region further supports the conclusion that the partial crystallization occurred due to thermal effects instead of irradiation damage. It should also be noted that all specimens (including unirradi­ ated as-cast controls) exhibited a pronounced indentation size effect on the measured nanoindentation hardness. The experiments revealed that in the irradiated region of the alloy the depth dependent hardness did not significantly change during irradiation at room temperature to 290 � C as compared to the as-cast state. The Lam and Chong extrapo­ lation method, that has been previously used to examine the ISE in metallic glass, was employed on Cu BMG but resulted in poor quanti­ tative fit to the experimental nanoindentation hardness data. This failure suggests that a new model is needed to properly analyze the ISE phe­ nomenon in amorphous alloy systems.

5. Conclusions An investigation into the response of Cu60Zr20Hf10Ti10 BMG exposed to 9 MeV Ni3þ ion irradiation (~10 dpa midrange dose) at different temperatures led to several important conclusions. The results of the experiments suggest that the maximum operating temperature for the Cu BMG is ~300 � C, because of the thermally induced partial

8

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Intermetallics 116 (2020) 106655

Acknowledgements

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