Glass-Forming Ability, Mechanical Properties, and Energetic Characteristics of ZrCuNiAlNbHfY Bulk Metallic Glasses

The effects of partially substituting Al for Cu in Zr59.62Cu18.4-xNi12Al6+xNb3Hf0.78Y0.2 (x = 0, 2, 4, 6, 8 at.%) bulk metallic glasses (BMGs) on their glass-forming ability (GFA), quasi-static and dynamic mechanical properties, and energy characteristics were investigated. The results showed that an appropriate substitution of Al for Cu can improve GFA and reach a critical casting size up to 10 mm. Additionally, with Al replacement of Cu, the change in the distribution and content of free volume inside the BMGs was the main reason for the quasi-static compression plasticity. In contrast, the BMGs exhibited no plasticity during dynamic compression and high-speed impact, owing to the short loading time and thermal softening effect. In terms of energy characteristics, all alloys have a high combustion enthalpy. And on the surface of the fragments collected from impact, the active elements Zr, Al, and Nb reacted because of the adiabatic temperature rise. Further, x = 4 at.% Zr-based BMG with its superior overall performance could penetrate a 6 mm Q235 plate at a speed of 1038 m/s, combining excellent mechanical properties and energy characteristics. This study contributes to the development of Zr-based BMGs as novel energetic structural materials.


Introduction
Bulk metallic glasses (BMGs) are characterized by high strength, high hardness, large elastic strain limit, and excellent corrosion resistance [1][2][3] due to their unique long-range disordered atomic arrangement structure.And BMGs have a higher free energy and energy-release efficiency than crystalline materials [4,5].Furthermore, since Zr combines a low activation threshold and high calorific value for oxidation reactions, Zr-based BMGs exhibit a dual destructive effect of penetration and deflagration through energy release on targets, which facilitates military applications such as fragmentation warheads and shaped charges [6,7].However, because of strain softening and the lack of inherent propagation barriers, shear bands may rapidly crack, leading to catastrophic failure at room temperature [8,9].Besides plasticity, the practical applications of BMGs are also limited by their poor GFA.Moreover, current energy research has mainly focused on the classic components Zr 55 Ni 5 Al 10 Cu 30 and Zr 41.25 Ti 13.75 Cu 12.5 Ni 10 Be 22.5 [7,[10][11][12][13], and more BMG systems that simultaneously meet the dimensional, mechanical, and energy requirements are required.
According to recent reports, plasticity can be improved by employing nanocrystals, phase-separated structures, and the introduction of large amounts of free volume [14][15][16].The first two strategies are typically limited to quasi-static compression, which can dramatically decrease strength and plasticity with an increasing strain rate [15,17].For instance, the increase in the in situ crystalline phase content of Zr 55 Ni 5 Al 10 Cu 30 results in enhanced brittleness of fragments under high-speed impact, which promotes adequate fragmentation The Zr 59.62 Cu 18.4-x Ni 12 Al 6+x Nb 3 Hf 0.78 Y 0.2 (x = 0, 2, 4, 6, 8 at.%) alloy ingots were prepared in a vacuum arc furnace (WK-II A, Beijing Physcience Optoelectronics Co., Ltd., Beijing, China) using a Ti-gettered argon atmosphere.The purity of raw material was higher than 99.9 wt% (Zhongnuo New Material Technology Co., Ltd., Beijing, China).The intermediate Zr-Nb alloy was first pre-melted and then mixed with other metals and remelted four times to ensure compositional homogeneity.The surface oxide layer of the ingot was polished.Then, a series of cylindrical rods with diameters (Φ) of 3, 5, 7, 8, 10, and 12 mm and a length (l) of 150 mm was prepared using water-cooled copper-mold suction-casting equipment (ML-MG35, Peshing New Metal Co., Ltd., Changzhou, China).And the melt-spun ribbons were prepared by a melt-spinning machine (VF-RQT50, Makabe Giken Co., Ltd., Sendai, Japan) at a speed of 20 m/s.

Microstructure Characterization
The structures and morphologies of the samples were analyzed at room temperature using X-ray diffraction (XRD, Bruker D8 Advance, Karlsruhe, Germany), scanning electron microscopy (SEM, JSM 7200F, Tokyo, Japan), and high-resolution transmission electron microscopy (HRTEM, FEI Talos F200X, Waltham, MA, USA).The HRTEM sample was cut from the rods with Φ = 3 mm and thinned down to a thickness of 30 µm using SiC papers, and then prepared using a precision ion polishing system (PIPS, Gatan-691, Pleasanton, CA, USA).The thermal parameters were obtained using differential scanning calorimetry (DSC, Netzsch 404F3, Selb, Germany) in a flow argon atmosphere with a heating rate of 20 K/min, and DSC specimens, which had a weight of 10-15 mg, were taken from the same parts of the Φ = 3 mm rods.

Performance Test
The as-cast samples with Φ = 3 mm were machined into a length-to-width ratio of 2:1 for quasi-static compression and 1:1 for dynamic compression.All samples were carefully polished to ensure parallel ends.Quasi-static compression tests were performed using a universal testing machine (MTS E43.504, Eden Prairie, MN, USA) at a strain rate of 1 × 10 −3 s −1 , and dynamic compression tests were performed using the Split Hopkinson Pressure Bar (SHPB, ATL1500, Archimedes Industry Technology Co., LTD., Beijing, China) at a strain rate of 3 × 10 3 s −1 .Each compression test was repeated three times to ensure the reliability of the data.
A microcomputer automatic oxygen bomb calorimeter (TRHW-7000E, Tianrun Electronic Technology Co., Hebi, China) was used in 3 MPa pure oxygen to measure the combustion enthalpy at room temperature, and the weight of the test samples prepared from the melt-spinning method for each test was about 0.2 g.Ballistic gun (Nanjing University of Science and Technology, Nanjing, China) tests were used to investigate the damage performance of the BMG samples with Φ = 10 mm and l = 10 mm.The targets were two separate plates placed 1000 mm apart, including a 6 mm thick Q235 plate and a 1 mm thick aftereffect Al plate, to observe the fragmentation distribution of the fragments.The penetration process was recorded using a high-speed camera (Phantom, Wayne, NJ, USA).The fragments generated under high-speed impact were collected in a closed steel chamber.heating rate of 20 K/min, and DSC specimens, which had a weight of 10-15 mg, were tak from the same parts of the Φ = 3 mm rods.

Performance Test
The as-cast samples with Φ = 3 mm were machined into a length-to-width ratio of for quasi-static compression and 1:1 for dynamic compression.All samples were carefu polished to ensure parallel ends.Quasi-static compression tests were performed using universal testing machine (MTS E43.504, Eden Prairie, MN, USA) at a strain rate of 1 × 1 s −1 , and dynamic compression tests were performed using the Split Hopkinson Pressu Bar (SHPB, ATL1500, Archimedes Industry Technology Co., LTD., Beijing, China) a strain rate of 3 × 10 3 s −1 .Each compression test was repeated three times to ensure t reliability of the data.
A microcomputer automatic oxygen bomb calorimeter (TRHW-7000E, Tianr Electronic Technology Co., Hebi, China) was used in 3 MPa pure oxygen to measure t combustion enthalpy at room temperature, and the weight of the test samples prepar from the melt-spinning method for each test was about 0.2 g.Ballistic gun (Nanji University of Science and Technology, Nanjing, China) tests were used to investigate t damage performance of the BMG samples with Φ = 10 mm and l = 10 mm.The targets we two separate plates placed 1000 mm apart, including a 6 mm thick Q235 plate and a 1 m thick aftereffect Al plate, to observe the fragmentation distribution of the fragments.T penetration process was recorded using a high-speed camera (Phantom, Wayne, NJ, US The fragments generated under high-speed impact were collected in a closed steel chamb

GFA and Thermal Ability
Figure 1 shows the XRD patterns of as-cast Zr59.62Cu18.4-xNi12Al6+xNb3Hf0.78Y0.2(x = 0 4, 6, 8 at.%) alloy rods with different diameters.The XRD patterns under critical diame (Dc) conditions exhibit a broad diffraction peak, indicating that the cast samples a completely amorphous within the resolution range of the XRD technique.When replaces Cu, the Dc of the BMGs increases at first and then decreases.Dc reaches maximum of more than 10 mm at x = 4 at.% and approaches 10 mm at x = 6 at.%, which more than twice that of x = 8 at.%.When the diameter exceeds the Dc, many crystalli phases precipitate.The HRTEM image and corresponding selected area electron diffraction pattern (SAED) of x = 0, x = 4, and x = 8 at.%BMGs with Φ = 3 mm are shown in Figure 2. The results show typical amorphous structures with isotropic maze patterns and the SAED in the inset displays a typical amorphous halo ring with no diffraction spots, indicating the glassy nature of the rod.
The HRTEM image and corresponding selected area electron diffraction pattern (SAED) of x = 0, x = 4, and x = 8 at.%BMGs with Φ = 3 mm are shown in Figure 2. The results show typical amorphous structures with isotropic maze patterns and the SAED in the inset displays a typical amorphous halo ring with no diffraction spots, indicating the glassy nature of the rod.The DSC curves of Zr59.62Cu18.4-xNi12Al6+xNb3Hf0.78Y0.2(x = 0, 2, 4, 6, 8 at.%)BMGs with Φ = 3 mm at a heating rate of 20 K/min are shown in Figure 3.In the temperature range of 400-1300 K, some evident glass transition characteristics and the production of exothermic peaks corresponding with crystallization are observed.A two-stage crystallization exothermic peak appears at x = 8 at.%, suggesting that Al substitution changes the internal structure of BMGs and affects the crystallization process.In Zr-Cu-Ni-Al-Nb and many other multicomponent BMGs, the formation of crystal phases can be thermodynamically divided into several stages [20,24,26,27].The first stage is the precipitation of primary crystal phases from supercooled liquids, which mainly refers to the precipitation of local ordering clusters.The other stages correspond to the formation of stable crystal phases.When the instability of the precipitated local ordering clusters gradually increases upon the Al content exceeding (x = 8 at.%), the corresponding precipitation crystallization temperature decreases, resulting in a single crystallization event divided into two steps, which is similar to the research of Li and Tan, et al. [24,34].3.In the temperature range of 400-1300 K, some evident glass transition characteristics and the production of exothermic peaks corresponding with crystallization are observed.A two-stage crystallization exothermic peak appears at x = 8 at.%, suggesting that Al substitution changes the internal structure of BMGs and affects the crystallization process.In Zr-Cu-Ni-Al-Nb and many other multicomponent BMGs, the formation of crystal phases can be thermodynamically divided into several stages [20,24,26,27].The first stage is the precipitation of primary crystal phases from supercooled liquids, which mainly refers to the precipitation of local ordering clusters.The other stages correspond to the formation of stable crystal phases.When the instability of the precipitated local ordering clusters gradually increases upon the Al content exceeding (x = 8 at.%), the corresponding precipitation crystallization temperature decreases, resulting in a single crystallization event divided into two steps, which is similar to the research of Li and Tan, et al. [24,34].The HRTEM image and corresponding selected area electron diffraction pattern (SAED) of x = 0, x = 4, and x = 8 at.%BMGs with Φ = 3 mm are shown in Figure 2. The results show typical amorphous structures with isotropic maze patterns and the SAED in the inset displays a typical amorphous halo ring with no diffraction spots, indicating the glassy nature of the rod.The DSC curves of Zr59.62Cu18.4-xNi12Al6+xNb3Hf0.78Y0.2(x = 0, 2, 4, 6, 8 at.%)BMGs with Φ = 3 mm at a heating rate of 20 K/min are shown in Figure 3.In the temperature range of 400-1300 K, some evident glass transition characteristics and the production of exothermic peaks corresponding with crystallization are observed.A two-stage crystallization exothermic peak appears at x = 8 at.%, suggesting that Al substitution changes the internal structure of BMGs and affects the crystallization process.In Zr-Cu-Ni-Al-Nb and many other multicomponent BMGs, the formation of crystal phases can be thermodynamically divided into several stages [20,24,26,27].The first stage is the precipitation of primary crystal phases from supercooled liquids, which mainly refers to the precipitation of local ordering clusters.The other stages correspond to the formation of stable crystal phases.When the instability of the precipitated local ordering clusters gradually increases upon the Al content exceeding (x = 8 at.%), the corresponding precipitation crystallization temperature decreases, resulting in a single crystallization event divided into two steps, which is similar to the research of Li and Tan, et al. [24,34].The detailed thermal parameters are listed in Table 1.When Cu is gradually substituted by Al, both the glass transition temperature (T g ) and the onset temperature of crystallization (T x ) initially increase before decreasing.The supercooled liquid region (∆T = T x − T g ) obtains a maximum value of 59 K at x = 4 at.%. Inoue et al. [2] reported that a larger ∆T could indicate a more stable supercooled liquid.Two other parameters reflecting GFA, namely reduced glass transition temperature (T rg = T g /T l ) and the γ parameter (γ = T x /(T g + T l )), also exhibit the largest values of 0.610 and 0.411, respectively, indicating that x = 4 at.%BMG has the best GFA and stability in the system.Moreover, as the Al content increases, the liquids temperature (T l ) decreases and then increases, and the x = 4 at.%alloy exhibits the lowest T l of 1130 K, reflecting that this composition is closer to the eutectic point when compared with others.The thermodynamic stability of the molten liquid phase is higher when the T l of deep eutectic composition is lower.In other words, the x = 4 at.%alloy possesses the highest stability of alloy melting against crystallization upon cooling and ultimately achieves the highest GFA.In a multicomponent system, it is known that the mixed enthalpy (∆H i-j ) and atomic size are the two main factors affecting GFA.Cu has a positive ∆H i-j with major elements Ni (∆H Cu-Nb = 3 kJ/mol) and Nb (∆H Cu-Ni = 4 kJ/mol), and excess Cu may weaken the strong bonding structure between atoms and reduce the packing densities of atoms.In contrast, Al has a large negative ∆H i-j with all other elements Zr, Ni, Nb, Y, and Hf (∆H Al-Zr = −44 kJ/mol, ∆H Al-Cu = −1 kJ/mol, ∆H Al-Ni = −22 kJ/mol, ∆H Al-Nb = −18 kJ/mol, ∆H Al-Hf = −38 kJ/mol, ∆H Al-Y = −39 kJ/mol), preferring to form atomic pairs.The system mixing enthalpy (∆H mix ) of the Zr-Cu-Ni-Al-Nb-Y-Hf system can be calculated using the following equation [35]: where x i and x j are the atomic percentages of the elements i and j, respectively, and the values of ∆H i−j are obtained from the literature [36].As summarized in Table 1, the ∆H mix of the system increases when Cu is substituted by Al.An enhanced negative ∆H mix would encourage the chemical short-range order.On the other hand, in terms of atomic size, appropriate differences in atomic size increase the random packing density and improves GFA [37].However, when the Al content continues to increase (x = 8 at.%), some cluster structures are unstable and preferentially precipitate, resulting in a decrease in T g and the degradation of GFA.

Quasi-Static and Dynamic Mechanical Properties of BMGs
The stress-strain curves for the quasi-static and dynamic compression of the BMGs at strain rates of 1 × 10 −3 s −1 and 3 × 10 3 s −1 are shown in Figure 4a,c.Characterization data obtained from the curves are listed in Table 2, including yield strength (σ y ), fracture strength (σ max ), and plastic strain (ε p ).Under quasi-static compressive loading, the σ y value of the x = 0 at.%alloy is 1582 MPa and the ε p value is 1.06%.When the Al content increases from x = 0 at.% to x = 6 at.%, both σ y and ε p of the alloys increase and then decrease with the increasing Al content.The x = 6 at.% alloy exhibits the largest σ y of 1722 MPa and a ε p of 3.45%, which is 3.25 times larger than that of the pristine alloy.While there exist serrated flow behaviors (Figure 4b) in quasi-static compressive deformation, reflecting that a large number of shear bands were formed to adapt to the plastic strain, the dynamic fracture is complete elastic deformation without any macroplasticity.As shown in Figure 4d, when Al replaces Cu, the quasi-static and dynamic compression fracture strengths follow basically the same trend.And the compressive strength decreases as the strain rate increases, and the dynamic fracture strength ranges from 1100 to 1310 MPa.
that a large number of shear bands were formed to adapt to the plastic strain, the dynamic fracture is complete elastic deformation without any macroplasticity.As shown in Figure 4d, when Al replaces Cu, the quasi-static and dynamic compression fracture strengths follow basically the same trend.And the compressive strength decreases as the strain rate increases, and the dynamic fracture strength ranges from 1100 to 1310 MPa.To further clarify the deformation mechanism of BMGs, Figure 5 shows the lateral and fracture surface morphology images after the final fracture.The fracture surface angle is found to be from 39.8° to 41.7°.Under compression, the normal stresses prohibit the shear fracture of the BMG, resulting in a fracture angle of less than 45° [38,39].Some primary and secondary shear bands (SBs) appear on the specimens, and the numbers of SBs and their bifurcations and crisscrossings (marked by circles) at x = 4 and 6 at.% are higher than those at x = 0, 2, and 8 at.%.In the fractography images, the specimens with x  To further clarify the deformation mechanism of BMGs, Figure 5 shows the lateral and fracture surface morphology images after the final fracture.The fracture surface angle is found to be from 39.8 • to 41.7 • .Under compression, the normal stresses prohibit the shear fracture of the BMG, resulting in a fracture angle of less than 45 • [38,39].Some primary and secondary shear bands (SBs) appear on the specimens, and the numbers of SBs and their bifurcations and crisscrossings (marked by circles) at x = 4 and 6 at.% are higher than those at x = 0, 2, and 8 at.%.In the fractography images, the specimens with x = 4 and 6 at.% exhibit more vein-like patterns along the direction of shear stress, caused by the instantaneous release of high elastic energy within the shear bands, while the x = 0 at.%BMG demonstrates the largest proportion of smooth areas.The plasticity is revealed by the interaction between SBs and the competitive relationship between the patterned regions and smooth regions.The high density of shear bands and vein patterns generally reflects good plasticity [14,40,41].
= 4 and 6 at.% exhibit more vein-like patterns along the direction of shear stress, caused by the instantaneous release of high elastic energy within the shear bands, while the x = 0 at.%BMG demonstrates the largest proportion of smooth areas.The plasticity is revealed by the interaction between SBs and the competitive relationship between the patterned regions and smooth regions.The high density of shear bands and vein patterns generally reflects good plasticity [14,40,41].As mentioned above, Al has a large negative ∆H i-j with other elements.In particular, Al and Zr easily form stable Al-Zr atomic pairs because they have the largest negative enthalpy of mixing, which may contribute to their strength in a certain range [36].In this system, the atomic radii of Zr, Cu, Ni, Al, Nb, Hf, and Y are 0.162, 0.128, 0.125, 0.143, 0.143, 0.159, and 0.181 nm, respectively.The proper replacement of small atoms of Cu with Al may lead to enhanced structural inhomogeneity and an increase in free volume.According to the free-volume model, the loosely packed defective regions with a higher free volume serve as the initiation regions for shear band generation during plastic deformation, while the high-density-packing area usually inhibits the propagation of the SBs [20], causing the SBs to proliferate, branch, and deflect.In such interactions, the formation of numerous shear bands is promoted, resulting in greater plasticity.
Contrary to the vein-like patterns with obvious density differences under quasi-static compression, the fracture surfaces under dynamic compression show a complex pattern with no obvious regularity.Figure 6 shows the typical dynamic fracture surfaces of the x = 4 at.%alloy and the x = 8 at.% alloy.In addition to vein-like patterns, there are dendritelike patterns, cracks, and molten zones.The direct linking of the vein-like pattern and the radial vein-like pattern indicates a direct transition from compressive shear stress into tensile shear stress [42], as shown in Figure 6b,e.Moreover, the molten droplets produced during quasi-static compression are approximately 10 µm (Figure 5f), and the size of the molten zones at the edge of the cracks after dynamic compression can reach more than 200 µm (Figure 6c,f), which indicates that the temperature rise caused by adiabatic shearing exceeds the glass transition temperature or even the melting point of the alloys during quasi-static compression [43], and is greater under dynamic compression.
As mentioned above, Al has a large negative ΔHi-j with other elements.In particu Al and Zr easily form stable Al-Zr atomic pairs because they have the largest negat enthalpy of mixing, which may contribute to their strength in a certain range [36].In t system, the atomic radii of Zr, Cu, Ni, Al, Nb, Hf, and Y are 0.162, 0.128, 0.125, 0.143, 0.1 0.159, and 0.181 nm, respectively.The proper replacement of small atoms of Cu with may lead to enhanced structural inhomogeneity and an increase in free volum According to the free-volume model, the loosely packed defective regions with a high free volume serve as the initiation regions for shear band generation during plas deformation, while the high-density-packing area usually inhibits the propagation of SBs [20], causing the SBs to proliferate, branch, and deflect.In such interactions, formation of numerous shear bands is promoted, resulting in greater plasticity.
Contrary to the vein-like patterns with obvious density differences under quasi-sta compression, the fracture surfaces under dynamic compression show a complex patte with no obvious regularity.Figure 6 shows the typical dynamic fracture surfaces of th = 4 at.%alloy and the x = 8 at.% alloy.In addition to vein-like patterns, there are dendr like patterns, cracks, and molten zones.The direct linking of the vein-like pattern and radial vein-like pattern indicates a direct transition from compressive shear stress i tensile shear stress [42], as shown in Figure 6b,e.Moreover, the molten droplets produc during quasi-static compression are approximately 10 µm (Figure 5f), and the size of molten zones at the edge of the cracks after dynamic compression can reach more th 200 µm (Figure 6c,f), which indicates that the temperature rise caused by adiaba shearing exceeds the glass transition temperature or even the melting point of the allo during quasi-static compression [43], and is greater under dynamic compression.Compared with quasi-static compression, the high-density-packing area does n prohibit the propagation of SBs [44], and there is not enough time for shear transformat zones to adapt to the dynamic strain rate, resulting in premature failure of the samp The free-volume softening and thermal softening effects for dynamic compression more remarkable than those for quasi-static compression [45]; thus, the dynamic streng is lower than that under quasi-static loading.It should be emphasized that despite overall brittle morphology, local toughness persists as vein-like patterns (Figure 6b under dynamic compression.Compared with quasi-static compression, the high-density-packing area does not prohibit the propagation of SBs [44], and there is not enough time for shear transformation zones to adapt to the dynamic strain rate, resulting in premature failure of the samples.The free-volume softening and thermal softening effects for dynamic compression are more remarkable than those for quasi-static compression [45]; thus, the dynamic strength is lower than that under quasi-static loading.It should be emphasized that despite the overall brittle morphology, local toughness persists as vein-like patterns (Figure 6b,e) under dynamic compression.

Energetic Characteristics
The average mass combustion enthalpy of the Zr 59.62 Cu 18.4-x Ni 12 Al 6+x Nb 3 Hf 0.78 Y 0.2 (x = 0, 2, 4, 6, 8 at.%) melt-spun ribbon measured by the oxygen bomb calorimeter is above 10 kJ/g (TNT: 4.1 kJ/g).As the Al content increases from 6 at.% to 14 at.%, the combustion enthalpy of the alloys has a slightly increasing trend (Figure 7a), from 10.826 KJ/g at x = 0 at.% to 11.749 KJ/g at x = 8 at.%. Hu et al. [46] have reported that the average mass combustion enthalpy of Zr 55 Cu 30 Ni 5 Al 10 and Zr 41.2 Ti 13.8 Cu 12.5 Ni 10 Be 22.5 is 10.797 KJ/g and 6.496 KJ/g, respectively.Our results show that the Zr 59.62 Cu 18.4-x Ni 12 Al 6+x Nb 3 Hf 0.78 Y 0.2 alloy has a higher combustion enthalpy.And the main combustion products are ZrO 2 and a small amount of CuO, Cu 2 O, and NiO (Figure 7b).

Energetic Characteristics
The average mass combustion enthalpy of the Zr59.62Cu18.4-xNi12Al6+xNb3Hf0.78Y0.2(x = 0, 2, 4, 6, 8 at.%) melt-spun ribbon measured by the oxygen bomb calorimeter is above 10 kJ/g (TNT: 4.1 kJ/g).As the Al content increases from 6 at.% to 14 at.%, the combustion enthalpy of the alloys has a slightly increasing trend (Figure 7a), from 10.826 KJ/g at x = 0 at.% to 11.749 KJ/g at x = 8 at.%. Hu et al. [46] have reported that the average mass combustion enthalpy of Zr55Cu30Ni5Al10 and Zr41.2Ti13.8Cu12.5Ni10Be22.5 is 10.797 KJ/g and 6.496 KJ/g, respectively.Our results show that the Zr59.62Cu18.4-xNi12Al6+xNb3Hf0.78Y0.2alloy has a higher combustion enthalpy.And the main combustion products are ZrO2 and a small amount of CuO, Cu2O, and NiO (Figure 7b).Among all of the alloys, the x = 4 at.%BMG demonstrates excellent comprehensive properties, so it is necessary to verify its target penetration capability.The high-speed photography images of the x = 4 at.%BMG passing through the targets at a speed of 1038 m/s are shown in Figure 8.Under impact, the BMG firstly penetrates the Q235 target, leaving a hole larger than 10 mm and melting traces around it caused by the violent reaction during the penetration process.The point to emphasize is that no burning is observed before the BMG hits the target, implying that it does not crack prematurely under detonation loading.The remaining fragments subsequently penetrate and damage the Al plate with kinetic and chemical energy.The light emitted by impacts arises from burning debris [47].Due to the large specific surface area, small fragments are fully in contact with oxygen and contribute more to the chemical reaction, releasing energy to generate flames and leaving black reaction traces on the Al plate.In the recovered impact debris (dimensions less than 400 µm), only the characteristic peaks of ZrO2 are observed and other oxides are contained insufficiently to be detected (Figure 7b).Among all of the alloys, the x = 4 at.%BMG demonstrates excellent comprehensive properties, so it is necessary to verify its target penetration capability.The high-speed photography images of the x = 4 at.%BMG passing through the targets at a speed of 1038 m/s are shown in Figure 8.Under impact, the BMG firstly penetrates the Q235 target, leaving a hole larger than 10 mm and melting traces around it caused by the violent reaction during the penetration process.The point to emphasize is that no burning is observed before the BMG hits the target, implying that it does not crack prematurely under detonation loading.The remaining fragments subsequently penetrate and damage the Al plate with kinetic and chemical energy.The light emitted by impacts arises from burning debris [47].Due to the large specific surface area, small fragments are fully in contact with oxygen and contribute more to the chemical reaction, releasing energy to generate flames and leaving black reaction traces on the Al plate.In the recovered impact debris (dimensions less than 400 µm), only the characteristic peaks of ZrO 2 are observed and other oxides are contained insufficiently to be detected (Figure 7b).
In order to confirm the oxidation state of the constituent elements, XPS analysis is performed.As shown in Figure 9, on the surface of the combustion products and recovered impact debris, all elements show oxidation peaks, indicating that the elements all participate in the oxidation reaction.Zr, Ni, Al, and Nb elements are fully oxidized, and Cu is relatively inert because of the existence of partially incompletely oxidized Cu.By comparison, the oxidation degree of Cu in the combustion products is higher than that in the impact debris.In order to confirm the oxidation state of the constituent elements, XPS analysis is performed.As shown in Figure 9, on the surface of the combustion products and recovered impact debris, all elements show oxidation peaks, indicating that the elements all participate in the oxidation reaction.Zr, Ni, Al, and Nb elements are fully oxidized, and Cu is relatively inert because of the existence of partially incompletely oxidized Cu.By comparison, the oxidation degree of Cu in the combustion products is higher than that in the impact debris.To further understand the mechanism of energy release, SEM and EDS are used to analyze the combustion products, the oxide layer on the cross-section of the impact debris,  In order to confirm the oxidation state of the constituent elements, XPS an performed.As shown in Figure 9, on the surface of the combustion prod recovered impact debris, all elements show oxidation peaks, indicating that the all participate in the oxidation reaction.Zr, Ni, Al, and Nb elements are fully o and Cu is relatively inert because of the existence of partially incompletely oxid By comparison, the oxidation degree of Cu in the combustion products is higher in the impact debris.To further understand the mechanism of energy release, SEM and EDS are used to analyze the combustion products, the oxide layer on the cross-section of the impact debris, and the hollow spherical particles on the surface of the impact debris, as shown in Figure 10.Obviously, there are clearly distinguished Zr-rich phases and Cu-rich phases in the combustion products (Figure 10a), which are caused by the uneven distribution of oxides produced by complete combustion.Moreover, in the oxide layer (Figure 10b), the active elements Zr, Al, and Nb demonstrate obvious concentrated distribution with O, and the relative distributions of Cu and Ni are different.Segregation does not exist in as-cast amorphous alloys, and because of the higher melting point of Nb than that of Cu, this phenomenon of differing distributions is mainly caused by the preferential combustion of active elements rather than just simple melting, which becomes more pronounced with an increase in the degree of oxidation reaction.According to their Gibbs energies of formation (per mole of oxygen) [48], the order of elements that are more likely to react with O is Zr > Al > Nb > Ni > Cu.Under high-speed impact, the smaller the size of the fragments, the higher the degree of oxidation [49]; so, the separation is more severe in hollow spherical particles (Figure 10c) than in the oxide layer.
and the hollow spherical particles on the surface of the impact debris, as shown in Figure 10.Obviously, there are clearly distinguished Zr-rich phases and Cu-rich phases in the combustion products (Figure 10a), which are caused by the uneven distribution of oxides produced by complete combustion.Moreover, in the oxide layer (Figure 10b), the active elements Zr, Al, and Nb demonstrate obvious concentrated distribution with O, and the relative distributions of Cu and Ni are different.Segregation does not exist in as-cast amorphous alloys, and because of the higher melting point of Nb than that of Cu, this phenomenon of differing distributions is mainly caused by the preferential combustion of active elements rather than just simple melting, which becomes more pronounced with an increase in the degree of oxidation reaction.According to their Gibbs energies of formation (per mole of oxygen) [48], the order of elements that are more likely to react with O is Zr > Al > Nb > Ni > Cu.Under high-speed impact, the smaller the size of the fragments, the higher the degree of oxidation [49]; so, the separation is more severe in hollow spherical particles (Figure 10c) than in the oxide layer.

Conclusions
We have reported the GFA, mechanical performance, and energetic characteristics of Be-free Zr59.62Cu18.4-xNi12Al6+xNb3Hf0.78Y0.2 (x = 0, 2, 4, 6, 8 at.%)BMGs.With increasing Al content, the GFA increases significantly and reaches a maximum Dc of 10 mm at x = 4 at.%.And the x = 6 at.%BMG possesses the highest σmax of 1791 MPa and the best quasi-static compression plasticity of 3.45%.In contrast, owing to the short loading time and thermal softening effect, no plasticity exists in Zr-based BMGs during high-rate compression.Regarding the energy characteristics, the increase in Al content increases the enthalpy of combustion from 10.826 KJ/g at x = 0 at.% to 11.749 KJ/g at x = 8 at.%.And elements Zr, Al, and Nb burn preferentially over Cu and Ni under impact due to different oxygen affinities on the surface of the fragments collected from the impact.The x = 4 at.%Zr-based BMG

Conclusions
We have reported the GFA, mechanical performance, and energetic characteristics of Be-free Zr 59.62 Cu 18.4-x Ni 12 Al 6+x Nb 3 Hf 0.78 Y 0.2 (x = 0, 2, 4, 6, 8 at.%)BMGs.With increasing Al content, the GFA increases significantly and reaches a maximum D c of 10 mm at x = 4 at.%.And the x = 6 at.%BMG possesses the highest σ max of 1791 MPa and the best quasi-static compression plasticity of 3.45%.In contrast, owing to the short loading time and thermal softening effect, no plasticity exists in Zr-based BMGs during high-rate compression.Regarding the energy characteristics, the increase in Al content increases the enthalpy of combustion from 10.826 KJ/g at x = 0 at.% to 11.749 KJ/g at x = 8 at.%.And elements Zr, Al, and Nb burn preferentially over Cu and Ni under impact due to

Figure 1
Figure 1 shows the XRD patterns of as-cast Zr 59.62 Cu 18.4-x Ni 12 Al 6+x Nb 3 Hf 0.78 Y 0.2 (x = 0, 2, 4, 6, 8 at.%) alloy rods with different diameters.The XRD patterns under critical diameter (D c ) conditions exhibit a broad diffraction peak, indicating that the cast samples are completely amorphous within the resolution range of the XRD technique.When Al replaces Cu, the D c of the BMGs increases at first and then decreases.D c reaches a maximum of more than 10 mm at x = 4 at.% and approaches 10 mm at x = 6 at.%, which is more than twice that of x = 8 at.%.When the diameter exceeds the D c , many crystalline phases precipitate.

Figure 2 .
Figure 2. The HRTEM and corresponding SAED (in the inset) images of the prepared Zr 59.62 Cu 18.4-x Ni 12 Al 6+x Nb 3 Hf 0.78 Y 0.2 (x = 0, 4, 8 at.%)BMGs with Φ = 3 mm: (a) x = 0; (b) x = 4; (c) x = 8.The DSC curves of Zr 59.62 Cu 18.4-x Ni 12 Al 6+x Nb 3 Hf 0.78 Y 0.2 (x = 0, 2, 4, 6, 8 at.%)BMGs with Φ = 3 mm at a heating rate of 20 K/min are shown in Figure3.In the temperature range of 400-1300 K, some evident glass transition characteristics and the production of exothermic peaks corresponding with crystallization are observed.A two-stage crystallization exothermic peak appears at x = 8 at.%, suggesting that Al substitution changes the internal structure of BMGs and affects the crystallization process.In Zr-Cu-Ni-Al-Nb and many other multicomponent BMGs, the formation of crystal phases can be thermodynamically divided into several stages[20,24,26,27].The first stage is the precipitation of primary crystal phases from supercooled liquids, which mainly refers to the precipitation of local ordering clusters.The other stages correspond to the formation of stable crystal phases.When the instability of the precipitated local ordering clusters gradually increases upon the Al content exceeding (x = 8 at.%), the corresponding precipitation crystallization temperature decreases, resulting in a single crystallization event divided into two steps, which is similar to the research of Li and Tan, et al.[24,34].

Figure 4 .
Figure 4.The mechanical behavior of BMGs: (a) quasi-static compressive stress-strain curves of Zr 59.62 Cu 18.4-x Ni 12 Al 6+x Nb 3 Hf 0.78 Y 0.2 (x = 0, 2, 4, 6, 8 at.%)BMGs at a strain rate of 1 × 10 −3 s −1 and the serrated flow behaviors in region b of the x = 6 curve; (b) the magnified image of serrated flow behaviors in (a); (c) dynamic compressive stress-strain curves of BMGs at a strain rate of 3 × 10 3 s −1 ; (d) the quasi-static and dynamic fracture strength of the as-cast specimens as a function of Al content.

Figure 6 .
Figure 6.The fracture surfaces of Zr59.62Cu18.4-xNi12Al6+xNb3Hf0.78Y0.2(x = 4, 8 at.%)BMGs a dynamic compression: (a) the full view of x = 4 BMG fracture surface and the classic patterns (b) a molten zone (c); (b) the magnified image of radial vein-like pattern and vein-like pattern of x BMG; (c) the magnified image of molten zone of x = 4 BMG; (d) the full view of x = 8 BMG fract surface and the classic patterns (e) and molten zone (f); (e) the magnified image of radial vein-l pattern and vein-like pattern of x = 8 BMG; (f) the magnified image of molten zone of x = 8 BMG

Figure 6 .
Figure 6.The fracture surfaces of Zr 59.62 Cu 18.4-x Ni 12 Al 6+x Nb 3 Hf 0.78 Y 0.2 (x = 4, 8 at.%)BMGs after dynamic compression: (a) the full view of x = 4 BMG fracture surface and the classic patterns (b) and molten zone (c); (b) the magnified image of radial vein-like pattern and vein-like pattern of x = 4 BMG; (c) the magnified image of molten zone of x = 4 BMG; (d) the full view of x = 8 BMG fracture surface and the classic patterns (e) and molten zone (f); (e) the magnified image of radial vein-like pattern and vein-like pattern of x = 8 BMG; (f) the magnified image of molten zone of x = 8 BMG.

Figure 8 .
Figure 8.The x = 4 at.%Zr-based BMG fragment penetrating two separated plates at a speed of 1038 m/s: (a) the penetration process; (b) the holes formed on the Q235 target; (c) the distribution of debris on the Al target.

Figure 9 .
Figure 9. SEM images of samples and XPS results of Zr, Cu, Ni, Al, and Nb elements on surface of (a-f) combustion products of x = 4 at.%alloy; (g-l) recovered impact debris of x = 4 at.%alloy.

Figure 8 .
Figure 8.The x = 4 at.%Zr-based BMG fragment penetrating two separated plates at a speed of 1038 m/s: (a) the penetration process; (b) the holes formed on the Q235 target; (c) the distribution of debris on the Al target.

Figure 8 .
Figure 8.The x = 4 at.%Zr-based BMG fragment penetrating two separated plates at a spe m/s: (a) the penetration process; (b) the holes formed on the Q235 target; (c) the distribution on the Al target.

Figure 9 .
Figure 9. SEM images of samples and XPS results of Zr, Cu, Ni, Al, and Nb elements on (a-f) combustion products of x = 4 at.%alloy; (g-l) recovered impact debris of x = 4 at.% a

Figure 9 .
Figure 9. SEM images of samples and XPS results of Zr, Cu, Ni, Al, and Nb elements on surface of (a-f) combustion products of x = 4 at.%alloy; (g-l) recovered impact debris of x = 4 at.%alloy.

Figure 10 .
Figure 10.SEM and EDS images: (a) the combustion products of the x = 4 at.%alloy; (b) the oxide layer in the cross-section of the impact debris of the x = 4 at.%alloy; (c) the hollow spherical particles (the marked box area) on the surface of the impact debris of the x = 4 at.%alloy.

Figure 10 .
Figure 10.SEM and EDS images: (a) the combustion products of the x = 4 at.%alloy; (b) the oxide layer in the cross-section of the impact debris of the x = 4 at.%alloy; (c) the hollow spherical particles (the marked box area) on the surface of the impact debris of the x = 4 at.%alloy.