Aluminum- and Zirconium-Based Amorphous Alloys: A Comparative Review of Glass Formation, Crystallization, Mechanical Behavior, Processing and Applications

Abstract

Aluminum- and zirconium-based amorphous alloys occupy complementary positions in metallic glass research. This review compares these two alloy families within a common framework that links thermodynamic and kinetic glass-forming criteria, crystallization pathways, deformation mechanisms, and processing routes. Zr-based alloys generally combine lower critical cooling rates, wide supercooled liquid regions, and bulk castability, whereas Al-rich alloys offer low density and high specific strength but remain marginal glass formers that are commonly produced as ribbons or powders. The supercooled liquid region, ΔT_x = T_x − T_g, is treated here as a kinetic processing window rather than a direct measure of ductility. Room-temperature plastic flow is discussed in terms of shear-transformation zones and shear-band initiation, propagation, multiplication, and arrest. Crystalline second phases may improve plasticity by redistributing stress and promoting the formation of multiple shear bands, but the outcome depends on their phase fraction, size, spacing, morphology, and loading mode. Recent original studies demonstrate the fabrication of dense, fully amorphous Zr-based components by laser powder bed fusion, as well as bulk Al-based glasses produced by copper mold casting or powder consolidation. The comparative analysis provides practical design guidance: combine a negative enthalpy of mixing and atomic-size mismatch with a low liquidus temperature; suppress oxygen contamination and heterogeneous nucleation; and select processing routes according to the attainable cooling rate and the stability of the supercooled liquid.

1. Introduction

Aluminum- and zirconium-based metallic glasses illustrate two distinct routes to amorphization. Zr-based multicomponent alloys can often be cast as bulk metallic glasses (BMGs) because crystallization is sufficiently slow at cooling rates achievable in metallic molds. Al-rich alloys have much lower glass-forming ability (GFA) and usually require melt spinning, atomization, or subsequent consolidation of amorphous powders. Comparing these two families is useful because the same descriptors—atomic-size mismatch, mixing enthalpy, liquidus temperature, reduced glass-transition temperature, and crystallization kinetics—lead to different processing limits and mechanical responses [1,2,3,4,5,6,7].

Table 1 highlights the key contrast addressed in this review. Zr-based alloys can be produced as monolithic BMGs by melt casting, whereas Al-rich alloys generally require substantially higher cooling rates. Therefore, when evaluating the dimensions of Al-based samples, the processing route must also be taken into account: a cast rod with a sub-millimetre diameter and a centimetre-scale compact consolidated from amorphous powder are not equivalent indicators of the GFA of the melt. This distinction is main-tained throughout the discussion of thermal and mechanical properties.

Table 1. Representative formation conditions and achievable dimensions for Al- and Zr-based amorphous alloys [5,8,9,10].

Alloy Family or Representative Composition	Milestone	Product Form/Maximum Dimension	Cooling or Consolidation Condition	Reference
Al–TM–RE alloys	First Al-rich metallic glasses reported in 1988	Melt-spun ribbons, typically tens of micrometers thick	Critical cooling rate typically 105–106 K s−1	[8]
Al–Ni–Y	Cast monolithic glass reported in 2009	Rod, 0.75 mm diameter	Copper-mold casting after melt treatment	[9]
Al86Ni6Y4.5Co2La1.5	Bulk glass by powder consolidation in 2013	Fully dense disk, 10 mm diameter	Spark-plasma sintering; dimension is not set by melt-quench cooling rate	[10]
Zr–Al–(Cu,Ni)	Bulk glass-forming family reported from 1990	Cast sections up to ≈30 mm	Critical cooling rate ≈ 1–10 K s−1	[5]
Zr–Ti–Al–(Cu,Ni)	Bulk multicomponent family	Cast sections up to ≈30 mm	Critical cooling rate ≈ 1–5 K s−1	[5]

For Al-rich glasses, the most successful compositions usually contain a transition metal and a rare-earth element. Increasing chemical complexity, atomic size mismatch, and the high interaction energy between unlike atoms can slow the structural rear-rangements required for crystallization; however, this effect is composition-specific. Oxygen has a particularly strong influence on Zr-based melts: oxide particles and oxygen-enriched clusters can serve as heterogeneous nucleation sites, thereby reducing the attainable fraction of the amorphous phase [11,12]. These observations suggest that both thermodynamic and kinetic factors play an important role in glass formation, as discussed below.

As a criterion for amorphization in work [13] it was proposed to use the ratio I_r/kT_l, where I_r is the reduced ionization potential of the alloy components. For alloys of the Zr-Cu system, a correlation of the maximum of this criterion with the concentration region of alloys that are amorphized is observed, which can be explained by the same nature of the chemical bond in the regions of short-range order for the liquid and amorphous states, and, accordingly, by the presence of dynamically stable atomic groups that do not crystallize. The higher the dynamic stability of clusters, i.e., regions of regular structure that do not correspond to the crystalline state, the more difficult it is for crystallization nuclei to appear in the supercooled melt. Overheating the melt and keeping it at a temperature above the liquidus line contribute to increasing the stability of the liquid state due to the fact that dissociation of the crystalline phase residues occurs in the melt. The cluster model of the amorphous state is based on the fact that a certain number of atoms forms an ordered micro-grouping (cluster), which, due to a certain ordering in the arrangement of atoms, has a lower energy. Such a cluster does not have a physical interface, and the transition from one local grouping to another occurs continuously through more disordered transition regions. It has been established that for the number of atoms less than 50, the energy minimum can be achieved only with polyhedral coordination of atoms. In works [14,15] it is noted that clusters that form an fcc lattice are unstable and transform into structures with icosahedral coordination of atoms or into a densely packed tetrahedral one. In the Brant cluster model, the basic structural unit of the amorphous state is taken to be a 13-atomic icosahedron with a fifth-order symmetry axis, which is forbidden for the case of other crystalline structures. These 13 atoms of the icosahedron have 42 close neighbors, which, under the condition of attraction between atoms, give such a polygon high stability. The similarity between the microcrystalline and cluster models lies in the assumption of the existence of some structural units that differ in topology. In the case of the microcrystalline model, these are small crystals of crystallographic packing with a structure that is characteristic of this material, and in the case of the cluster model, this unit can be densely packed structures such as icosahedrons or Casper polyhedra, which indicate a similarity to the cluster model with close packing [16]. Studies of chemical bonding and the near-field environment of atoms by X-ray diffraction spectroscopy of the Zr₆₀Ni₂₀Ti₂₀ alloy in the amorphous, quasi-crystalline and liquid states in one experiment showed the identity of the near-field order [17]. In the clusters, dense packing of dissimilar atoms and strong chemical bonds were observed due to strong spatial overlap and hybridization of d-electrons of the alloy components. It was shown that in the temperature range up to 470–500 K, the surface of the amorphous alloy is enriched mainly with Zr-Ti clusters. These clusters are the most stable for a given temperature, i.e., they are characteristic only for the amorphous state and are inherited from the melt. An increase in temperature (above 500 K) leads to a sudden change in the chemical composition of the surface layer, Zr-Ni clusters appear on the surface, which are more stable than Zr-Ti due to the stronger chemical bond formed by the d-electrons of nickel atoms with the d-electrons of zirconium atoms. In work [18], the number of atoms in ordered regions was estimated based on the temperature dependence of the shear viscosity of amorphous alloys Zr₆₅Cu₁₅Ni₁₀Al₅ and La₅₅Al₂₅Ni₂₀ and it was shown that for the first alloy this number corresponds to 22 atoms, and for the second—to 14 atoms, i.e., within the second—third coordination spheres.

The arrangement of atoms in the clusters of an amorphous alloy does not correspond to the crystal lattice of the phases formed during slow cooling of an alloy of this composition. To ensure crystallization, it is necessary to overcome a certain energy barrier, which is associated with the increase in disorder in the system and the destruction of clusters that do not crystallize. The activation energy of this “disorder” is a component of the entropy of the state of the system. An increase in the number of components in the alloy increases the entropy of the systems, and taking into account the tendency to amorphization, the added element must meet the above requirements.

The authors of [13] give the ratio of the activation energy of the process of structural reorganization of the melt in the pre-crystallization region, which depends on the bond energy in the cluster that does not crystallize to the crystallization temperature in the following form [13]:

$$I_r=f\left[\exp (-{\displaystyle \frac{b{\sigma }^3}{k{T}_m∆T^2}})\right]\exp \left(-{\displaystyle \frac{E_a}{\mathit{kT}}}\right),$$

(1)

where I_r is the reduced ionization potential of the alloy components; σ is the surface energy at the liquid-crystal interface; T is the current temperature; T_m is the melting point; f is the kinetic coefficient; b is a value that depends on the physical constants of the substance; E_a is the activation energy; ∆T is the supercooling temperature value; k is the Boltzmann constant.

Melts with a large ${\displaystyle \frac{E_a}{\mathit{kT}}}$ ratio should have a high tendency to amorphize. In addition, according to the Maxwell-Boltzmann law of particle energy distribution, the statistical number of particles with high energy, i.e., those with higher activation energies in the short-range order region (ordering region), can be determined by the formula [13]:

$$n_a\approx n\sqrt{{\displaystyle \frac{E_a}{\mathit{kT}}}}\exp (-{\displaystyle \frac{E_a}{\mathit{kT}}}),$$

(2)

where n is the total number of particles in the system. Accordingly, this review is restricted to Al- and Zr-based amorphous alloys. It critically compares (i) thermodynamic and kinetic GFA indicators, (ii) amorphous-to-crystalline transformations, (iii) mechanical properties and deformation by shear-transformation zones and shear bands, (iv) the role of crystalline second phases, and (v) recent processing and application developments. Its additional value is the direct comparison of both alloy families in terms of structure, processing, and properties, which provides practical guidelines for alloy design and processing.

2. Glass-Forming Ability and Thermodynamic Screening

The characteristics of the tendency of metallic systems to amorphization are the following parameters: T_rg = T_g/T_l (ratio of glass transition temperature T_g to liquidus temperature T_l), α = T_x/T_l (ratio of crystallization temperature to liquidus temperature) and γ = T_x/(T_g + T_l). In alloys based on aluminum and zirconium, the parameter T_rg increases with increasing concentration of alloying elements. The tendency to amorphization also correlates with the enthalpy of mixing of amorphous alloys (${∆H}_{\mathit{mix}}^{\mathit{amorph}}$) [6]:

$${∆H}_{\mathit{mix}}^{\mathit{amorph}}={∆H}^{\mathit{chem}}+{\sum }_i{x}_i{H}_i^{\mathit{amorph}},$$

(3)

where x_i is the molar concentration of the i-th element, $H_i^{\mathit{amorph}}$ is the enthalpy of the pure amorphous metal ($H_i^{\mathit{amorph}}=C_p{T}_{m,i}$, C_p = 3.5 J/mol∙K, T_m,i is the melting point of the i-th element). ∆H^chem for binary alloys can be calculated using the following equation [6]:

$${∆H}^{\mathit{chem}}=x_i{x}_j(x_j{∆H}_{i\to j}^{\mathit{sol}}+x_j{∆H}_{j\to i}^{\mathit{sol}}),$$

(4)

where x_i and x_j are the molar concentrations of i- and j- components, respectively. For the amorphous phase, the ΔH^chem of the binary AB system is multicomponent with the factor (1 + 5(x_A^c∙x_B^c)²), which is used to calculate the short-range order in the amorphous phase. In this case, x_A^c = x_AV_A^2/3/(x_AV_A^2/3 + x_BV_B^2/3), where V is the molar volume, and x_B^c = 1 − x_A^c. To calculate the ΔH^chem of multicomponent systems that include more than three components, the following equation is used [6]:

$${∆H}^{\mathit{chem}}={\sum }_{i=1,j=1,i\ne j}^N{∆H}^{\mathit{chem}}{A}_i{A}_j,$$

(5)

where i and j denote the alloy components, A is the name of the element and N is the total number of elements. The authors of [19,20,21,22,23,24,25,26,27,28,29] present the enthalpy of mixing calculated by this method and the thermodynamic properties of some bulk amorphous alloys based on zirconium and aluminum (Table 2).

Table 2. Thermodynamic properties of amorphous alloys based on zirconium and aluminum [19,20,21,22,23,24,25,26,27,28,29].

Alloy	Tg, K	Tx, K	Tm, K	Tl, K	∆Tx, K	Trg	γ	α	ΔHmix, kJ/mol
Zr58Cu22Co4Fe4Al12	673	788	−	1217	115	0.55	0.416	0.65	−47.61
Zr58Cu22Co4Ag4Al12	679	761	−	1183	82	0.57	0.408	0.64	−44.74
Zr58Cu22Co2Ag6Al12	675	751	−	1227	76	0.55	0.395	0.61	−43.09
Zr58Cu22Ag4Fe4Al12	675	757	−	1205	82	0.56	0.403	0.63	−39.84
Zr58Cu22Ag6Fe2Al12	676	748	−	1227	72	0.55	0.393	0.61	−40.47
Zr58Cu22Fe8Al12	677	761	−	1192	84	0.57	0.407	0.64	−43.15
Zr65Cu18Ni9Al8	612	741	1103	1123	129	0.54	0.427	0.66	−
(Zr65Cu18Ni9Al8)98Y2	603	689	1062	1104	86	0.55	0.403	0.62	−
(Zr65Cu18Ni9Al8)98Ti2	616	735	1064	1100	119	0.56	0.429	0.66	−
(Zr65Cu18Ni9Al8)98Nb2	616	737	1110	1116	121	0.55	0.425	0.66	−
(Zr65Cu18Ni9Al8)98Er2	611	730	1077	1107	119	0.55	0.425	0.66	−
(Zr65Cu18Ni9Al8)98W2	616	719	1101	1124	103	0.55	0.413	0.64	−
Zr55Cu30Ni5Al10	−	−	−	−	−	0.59	0.409	−	−
(Zr55Cu30Ni5Al10)99.98Sc0.02	−	742	−	−	−	0.60	0.411	−	−
Zr60.7Al15.5Ni15.5Fe8.3	710	770	1170	1285	60	0.55	0.362	0.60	−
Zr60Cu17.5Ni10Al7.5Si4B1	674	754	−	1087	80	0.62	0.43	0.69	−
Zr51Cu28Al21	694	755	−	−	61	−	−	−	−
Zr65Cu17.5Ni10Al7.5	639	695	−	−	56	−	−	−	−
Zr57Cu20Ni8Al10Ti5	648	712	−	−	64	−	−	−	−
Zr41.2Ti13.8Cu12.5Ni10Be22.5	626	772	−	−	145	−	−	−	−
Al88Ni6La6	−	506	−	915	−	−	−	0.55	−
Al86Ni6La6Cu2	−	509	−	909	−	−	−	0.56	−
Al85Ni9Nd4Co2	533	553	−	−	20	−	−	−	−
Al89.5Ni3.5Y7	−	428	906	1146	−	−	−	0.37	−
Al88Ni4Y8	−	465	904	1189	−	−	−	0.39	−
Al86.5Ni4.5Y9	505	526	904	1222	21	0.41	0.305	0.43	−
Al85Ni5Y10	537	560	906	1249	24	0.43	0.314	0.45	−
Al89.5Co3.5Y7	−	460	911	1188	−	−	−	0.39	−
Al88Co4Y8	−	507	911	1189	−	−	−	0.43	−
Al86.5Co4.5Y9	543	563	911	1231	20	0.44	0.317	0.46	−
Al85Co5Y10	−	587	911	1232	−	−	−	0.48	−
Al88La4Ni8	−	457	−	1082	−	−	−	0.42	−
Al88La5Ni7	−	480	−	1182	−	−	−	0.41	−
Al87La4Ni9	−	479	−	1113	−	−	−	0.43	−
Al87La5Ni8	−	500	−	1088	−	−	−	0.46	−
Al86La5Ni9	499	513	−	1109	14	0.45	0.319	0.46	−
Al86La6Ni8	519	533	−	1105	14	0.47	0.328	0.48	−
Al86La4Ni10	−	473	−	1107	−	−	−	0.43	−
Al85La6Ni9	530	546	−	1096	16	0.48	0.328	0.50	−
Al85La5Ni10	517	534	−	1124	17	0.46	0.325	0.48	−
Al84La5Ni11	539	556	−	1132	17	0.48	0.333	0.49	−
Al84La6Ni10	547	563	−	1119	16	0.49	0.338	0.50

The data indicate three robust trends. First, the melting interval, (T_l–T_m), is generally narrower for the Zr-based alloys listed (approximately 20–40 K) than for the Al-based alloys listed (approximately 120–150 K); however, this comparison is composition-specific and should not be regarded as a universal characteristic. Second, the T_rg, α, and γ values are generally more favorable for the Zr-based systems, consistent with their lower critical cooling rates and larger achievable cast-section thicknesses. Third, ΔT_x = T_x − T_g represents the supercooled liquid region. It is a kinetic indicator of the temperature interval available for thermoplastic forming or controlled devitrification, rather than a direct measure of room-temperature ductility or superplasticity.

When studying the tendency of metallic systems to amorphization, conditions of three interrelated factors are distinguished:

-

Thermodynamic (melting temperature, liquidus temperature, ratio of glass transition temperature to liquidus temperature, enthalpy of mixing, activation energy);

-

Kinetic (melt viscosity, rate of homogeneous or heterogeneous formation of crystallization centers);

-

Technological (degree of supercooling and cooling rate).

For metallic systems characterized by strong interaction of atoms, additional signs of the ability of alloys to amorphize can be distinguished [1,2]: the content in alloys of at least three chemical elements; the difference in the sizes of the radii of the atoms of the components (less than 0.88 or more than 1.25); the correspondence of the alloy composition to deep eutectic according to phase diagrams; the presence of intermetallics in this region; a large negative heat of alloy formation and “asymmetry” in the concentration dependence of the heat of mixing; high interfacial energy at the solid–liquid interface; difficulty of atomic rearrangement; decreased diffusion, and increased viscosity during solidification. Our studies of metal-metalloid systems that are amorphized and characterized by strong interparticle interactions of components [30,31,32,33,34,35] have shown the presence of certain concentration regions with large negative values of partial enthalpies of mixing. These regions, as a rule, correspond to alloy compositions that are prone to amorphization. The integral enthalpies of formation of alloys of these systems are asymmetric and the presence of extrema is observed in the concentration dependences.

The nature of the interaction of the main components of the alloys of metal-metal systems based on zirconium and aluminum and the influence of some elements on the interaction with these melts was analyzed on the basis of binary and ternary systems [21,24,36,37,38,39,40]. The first partial enthalpies of mixing and the radii of chemical elements of binary systems according to these works are presented in Table 3.

Table 3. Size of atoms and the first partial enthalpies of mixing (ΔH_mix) of binary alloys [21,24,36,37,38,39,40].

Chemical Element	Atomic Radius, (10−1 nm)	Enthalpy of Mixing, ∆Hmix, kJ/mol
		Zr(1.58) *	Al(1.43)	Ni(1.24)	Cu(1.28)	Y(1.81)	La(1.87)
Ag	1.44	−20; −41	−4	+15	+2	-	-
Pd	1.37	−91	−46; −84	0	−14	−123	−119
Au	1.44	−74	−22	+7	−9	-	-
Pt	1.39	−100	−44	−5	−12	-	-
Ir	1.35	−76	−30	−2	0	-	-
Re	1.37	−35	−9	+2	+18	-	-
Zn	1.33	−29	+1	−9	+1	-	-
Mo	1.36	−6	−5	−7	+19	-	-
V	1.32	−4	−16	−18	+5	-	-
Nb	1.45	+4	−18; −44	−30	+3	+41	+51
Ta	1.43	+3	−19	−29	+2	-	-
Cr	1.25	−12	−10	−7	+12	-	-
Sn	1.41	−43	+4	−4	+7	-	-
Y	1.81	−19	−38	−31	−22	-	-
Ti	1.47	0	−30	−35	−9	-	-
Fe	1.24	−44	−11	−2	-	−31	−38; +5
Co	1.25	−25	−43	-	-	−44	−38
Ni	1.24	-	−48; −22	-	-	-	−27

* The atomic radius (nm) follows the element in parentheses.

The influence of selected additions on the partial enthalpy of mixing in the Zr-containing Cu–Ti–Zr liquid is summarized in Table 4 [41].

Table 4. Partial enthalpies of mixing for selected additions to Cu–Ti–Zr liquids at 1873 K [41].

Atomic Radius, 10−1 nm	Base-Alloy Composition	ΔHmix, kJ/mol
Al (1.43) *	xTi = 0.18 xZr = 0.42	−84
Si (1.33)	xTi = 0.18 xZr = 0.42	−182
Sn (1.41)	xTi = 0.18 xZr = 0.42	−107
Al (1.43)	xTi = 0.30 xZr = 0.30	−57
Y (1.81)	xTi = 0.30 xZr = 0.30	−69
Al (1.43)	xTi = 0.42 xZr = 0.18	−94

* The atomic radius of the element (nm) is given in parentheses.

Negative pairwise mixing enthalpies favor chemical short-range order and can increase the barrier to atomic rearrangement, but they are not sufficient on their own to guarantee glass formation: very strong negative interactions may instead promote the formation of competing intermetallic compounds, thereby reducing GFA. The most useful screening strategy is therefore a combined rather than a single-parameter approach. Candidate compositions should combine: (i) negative but balanced mixing enthalpies, (ii) substantial atomic-size mismatch, (iii) a low liquidus temperature or proximity to a deep eutectic composition, and (iv) favorable kinetic indicators such as T_rg and γ. Processing constraints must then also be considered: oxygen control is especially important for reactive Zr melts, whereas the stable Al₂O₃ oxide film can complicate the melting, atomization, and powder consolidation of Al-rich alloys [5,6,11,12,26,27,28,29].

The ratio ΔH/T_l and related thermodynamic functions have historically been proposed as qualitative indicators of glass formation [42,43,44]. Their principal value lies in comparative screening within chemically related alloy families. Because experimental enthalpies, liquidus temperatures, nucleation kinetics, and impurity effects are composition-dependent, no single thermodynamic ratio should be regarded as a universal predictor. In this review, thermodynamic descriptors are used only when they help define combined criteria for selecting glass-forming compositions. Thermal analysis provides the experimentally accessible quantities needed to evaluate this screening approach, including (T_g), (T_x), crystallization enthalpy, (T_m), and (T_l). Together with compositional and cooling-rate data, these quantities enable consistent comparisons of GFA and thermal stability among Al- and Zr-based alloys.

3. Amorphous-to-Crystalline Transition

For amorphous alloys, the onset and peak crystallization temperatures depend on the experimental conditions and are not fixed equilibrium properties. Isothermal experiments describe the transformed fraction as a function of time, whereas continuous-heating experiments enable comparative kinetic analysis when performed at multiple heating rates. For isothermal transformations, the crystallized fraction is commonly described by the Johnson–Mehl–Avrami–Kolmogorov (JMAK) relation [45,46]:

X(τ) = 1 − exp(−kτⁿ),

(6)

where X(τ) is the crystalline volume fraction after time τ, k is a temperature-dependent rate constant, and n is the Avrami exponent. The exponent n can provide information about the transformation mechanism when the model assumptions and fitting range are clearly defined. The temperature dependence of k is commonly described by an Arrhenius relation [46]:

$$k=k_0\exp (-{\displaystyle \frac{E_a}{\mathit{RT}}}),$$

(7)

where k₀ is the pre-exponential factor and E_a is the apparent activation energy. For continuous heating, the Kissinger method relates the peak temperature to the heating rate [47]:

$${\displaystyle \frac{v_i}{T_x^2}}=a\exp \left(-{\displaystyle \frac{E_a}{R{T}_x}}\right),$$

(8)

$$\ln \left({\displaystyle \frac{v_i}{T_x^2}}\right)=-{\displaystyle \frac{E_a}{{\mathit{RT}}_x}}+\mathit{lna},$$

(9)

where R is the universal gas constant (J/(mol·K)), a is a temperature-dependent constant, and T_x is the crystallization temperature (K) for the i-th heating at rate v_i (K/min). Using this method, the activation energy of the transition from the amorphous to the crystalline state is found as the tangent of the slope of the linear dependence ln(v_i/T²_x) − f(1/T_x) for a series of experiments with different heating rates. Instead of T_x, T_peak temperature can be used. T_peak is the maximum crystallization temperature, the temperature at which the heat flux during the phase transition reaches a maximum. These methods were used to determine the crystallization temperatures and calculate the activation energy of the transition from the amorphous to the crystalline state for some alloys [37,38,48,49,50,51,52]. The relationship between the fraction of the microcrystalline component in the amorphous matrix and the crystallization temperatures was also established. The obtained data are presented in Table 5 and Table 6.

Table 5. Characteristic crystallization temperatures and activation energy for the transition from amorphous to crystalline state of zirconium- and aluminum-based alloys [37,38,48,49,50,51,52].

Alloy	Tx, K				* Ea, kJ/mol
	Heating Rate, K/min
	10	20	30	40
Zr65Cu17.5Ni10Al7.5	−	−	−	−	314
Zr60Cu17.5Ni10Al7.5Si4B1	−	−	−	−	370
Al83Ni10Si2Ce5	511	516	522	526	191
Al85Ni10Ce5	529	532	535	539	290
Al87Ni5Co2Nd6	553	557	563	568	221
Al87Ni7Nd6	489	494	497	505	166
Al82Fe5Ni5Ce8	667	680	690	−	165

* E_a—activation energy of the transition from amorphous to crystalline state.

Table 6. Characteristic annealing temperatures and crystallite size of aluminum-based amorphous alloys (annealing was carried out for an hour) [36].

Alloy	Annealing Temperature, Tanneal ± 1 K	Crystallite Size, nm
Al87Y5Ni8	532	20 ± 4
Al87Gd5Ni8	510	15 ± 3
Al87Y4Gd1Ni8	503	9 ± 2
Al87Y4Gd1Ni4Fe4	543	15 ± 3
Al87Gd5Ni4Fe4	558	22 ± 4

Table 5 illustrates that T_x shifts to higher temperatures as the heating rate increases, but the magnitude of this shift varies depending on the alloy and experimental conditions rather than following a universal “5 K per 10 K/min” rule. The practical significance of multi-rate measurements is that they enable the estimation of an apparent activation energy and show why T_x values obtained under different DSC conditions should not be compared directly. For Al₈₂Fe₅Ni₅Ce₈, for example, the activation energies calculated from the onset and peak temperatures were statistically similar (164 ± 10.6 and 165 ± 7.9 kJ/mol, respectively) [36].

Thus, the formation and growth of new phases proceeds according to the classical mechanism of nucleation and growth of a new phase. At the same time, during long-term exposure of amorphous alloys at a temperature below the temperature T_x, a gradual ordering of the matrix structure occurs with the formation of a supersaturated solid solution with a crystalline matrix lattice. Literature data on the crystallization of amorphous alloys based on zirconium and aluminum have shown that the process can be divided into several stages [50]: (1) the initial stage of crystallization, during which the ordering of the initial structure occurs; (2) the formation of small crystallites of the base metal in the amorphous matrix; (3) the transformation of the amorphous matrix into a complex ordered metastable phase; (4) the formation of stable phases of the base metal and metal-metal phases (intermetallics).

4. Mechanical Properties and Deformation Mechanisms

Mechanical data for metallic glasses should be interpreted in the context of the testing method and the type of strength reported. Tensile yield strength, ultimate tensile strength, compressive yield strength, compressive fracture strength, and fracture toughness describe different mechanical properties and should not be treated as equivalent. Therefore, Table 7 reports the available data for Zr-based metallic glasses as compressive fracture strength and strain to fracture. For the Al-based alloys, where only indentation results were available, the data are given as Vickers hardness [24,25,51,53,54,55,56,57,58,59,60,61,62,63,64,65].

Table 7. Mechanical properties of amorphous alloys based on zirconium and aluminium [25,51,56,57,59].

Alloy	σcf, MPa	εcf, %	HV	E, GPa
Zr65Cu18Ni9Al8	1764	6.54	600	99.8
(Zr65Cu18Ni9Al8)98Y2	1478	3.97	570	92.0
(Zr65Cu18Ni9Al8)98Ti2	1848	8.79	580	93.2
(Zr65Cu18Ni9Al8)98Nb2	1757	7.30	610	97.2
(Zr65Cu18Ni9Al8)98Er2	1677	7.07	620	93.3
(Zr65Cu18Ni9Al8)98W2	1441	12.20	620	100.3
Zr38Ti17Cu10.5Co12Be22.5	1850	−	−	−
Zr66.4Nb6.4Cu10.5Ni8.7Al8	1922	2.0	−	93
Zr73.5Nb9Cu7Ni1Al9.5	1740	2.1	−	89
Al78Si13Ni7Mm2	−	−	243	−
Al79.4Si11.6Ni7Mm2	−	−	297	−
Al81Si10Ni7Mm2	−	−	317	−

Mm—mischmetall, σ_cf—compressive fracture strength, ε_cf—compressive strain to fracture, HV—Vickers hardness, E—Young’s modulus.

The authors of [51] showed that the mechanical properties of amorphous alloys similar in composition depend little on the method of their preparation under the condition of similar cooling rates. The results presented in Table 8 can serve as confirmation of this fact.

Table 8. Methods for obtaining amorphous materials and their mechanical properties [51].

Alloy	Production Method	d, mm	σcf, MPa	εcf, %	E, GPa
Zr66.4Nb6.4Cu10.5Ni8.7Al8	Die casting	5	1840	2.3	82
Zr66.4Nb6.4Cu10.5Ni8.7Al8	Centrifugal casting	5	1755	1.6	117
Zr66.4Nb6.4Cu10.5Ni8.7Al8	Vacuum suction	5	1922	2.0	93
Zr73.5Nb9Cu7Ni1Al9.5	Vacuum suction	5	1740	1.9	89
Zr73.5Nb9Cu7Ni1Al9.5	Centrifugal casting	5	1771	1.8	84
Zr73.5Nb9Cu7Ni1Al9.5	Levitation melting and cold mold die casting	10	1754	1.7	82

σ_cf—compressive fracture strength, ε_cf—compressive strain to fracture, E—Young’s modulus.

In Refs. [25,60], isothermal aging of bulk amorphous alloys was carried out. The authors of [25]—at different temperatures for 20 min, and the authors of [60]—at a temperature of 673 K for different times. It was established that with increasing temperature and increasing holding time of amorphous alloys, the proportion of the microcrystalline component increases and, as a result, the hardness of the materials increases; the strength increases to a certain limit, and with further holding and increasing temperature, the opposite effect is observed (Table 9 and Table 10).

Table 9. Vickers hardness (HV) of initial and annealed at different temperatures for 20 min amorphous zirconium-based alloys [25].

Alloy	Initial Hardness, HV	Hardness After Annealing, HV		ΔHV, % *
		at 714 K	at 824 K
Zr65Cu18Ni9Al8	385	606	615	59.7
(Zr65Cu18Ni9Al8)98Y2	401	581	677	68.8
(Zr65Cu18Ni9Al8)98Ti2	408	614	667	63.5
(Zr65Cu18Ni9Al8)98Nb2	420	509	557	32.6
(Zr65Cu18Ni9Al8)98Er2	400	517	695	73.8
(Zr65Cu18Ni9Al8)98W2	431	562	544	26.2

* Vickers hardness increase after annealing at temperature 714 K.

Table 10. Mechanical properties of the bulk amorphous alloy Zr₅₇Cu₂₀Ni₈Al₁₀Ti₅ with different fractions of the microcrystalline phase, which was isothermally aged at a temperature of 673 K for different times [60].

Crystalline Phase, vol.%	Deformation Strain, %			σcf, MPa	E, GPa	Hardness, HV
	Reverse	Irreversible	Total
Amorphous state	2.6	1.2	3.8	1680	65	510
0 *	2.6	1.0	3.6	1741	70	530
14	2.4	0.7	3.1	1695	71
50	2.4	0.6	3.0	1835	72	590
54	2.4	0.1	2.5	1895	78
70	2.0	0	2.0	1261	73	640
78	1.9	0	1.9	1169	70	660

* without holding, E—Young’s modulus.

For the Zr₅₇Cu₂₀Ni₈Al₁₀Ti₅ alloy reported in Ref. [60], cylindrical specimens with a diameter of 3 mm and a height of 6 mm were tested in compression after isothermal aging at 673 K. Under these conditions, the highest compressive fracture strength was obtained at about 50 vol.% crystalline phase. Samples with a lower crystalline-phase fraction showed some irreversible compressive deformation, while those with a higher fraction fractured with little or no measurable plastic strain. The approximately 50 vol.% value should therefore not be regarded as a general ductile-to-brittle transition point for metallic glasses. The mechanical response depends on the alloy composition, type of crystalline phase, particle size and spacing, interface strength, specimen geometry, strain rate, and loading conditions. The nearly linear stress–strain curve of a monolithic metallic glass before yielding is mainly associated with elastic deformation. However, this behavior alone does not demonstrate that dislocations are completely absent. Since the amorphous matrix does not have a periodic crystal lattice, plastic deformation is generally described in terms of shear-transformation zones (STZs). These involve the cooperative rearrangement of small groups of atoms, and their interaction can eventually lead to the formation of shear bands [66]. Failure becomes catastrophic when deformation is concentrated in a dominant shear band. Structural heterogeneity or the introduction of a suitable crystalline phase can help distribute the deformation by providing additional sites for strain localization. This may result in the formation of several shear bands, hinder their propagation, and increase energy absorption through plastic deformation or a martensitic transformation [24,67,68]. However, a crystalline second phase does not necessarily improve the mechanical properties. Its effect depends on the phase fraction, particle size and spacing, morphology, strength difference between the phases, and interface cohesion. Depending on these factors, the second phase can either improve the toughness of the glass or promote crack initiation.

5. Recent Advances, Applications, and Comparative Design Guidance

Recent reviews describe the broad development and applications of Zr- and Al-based amorphous alloys [69,70,71,72,73,74,75]. To avoid merely repeating those reviews, this section emphasizes representative original studies [8,9,10,76,77,78,79,80,81,82] and uses them to compare current processing–structure–property relationships. Table 11 summarizes the principal examples.

Table 11. Selected original studies showing recent developments and representative processing routes for Zr-and Al- based amorphous alloys.

Alloys	Processing/Structure	Key Measured Results	Implication for Alloy Design	Reference
Zr59.3Cu28.8Nb1.5Al10.4	LPBF/>99% relative density and fully XRD-amorphous at an energy density of ~20–33 J·mm−3	Up to 6% compressive plastic strain; KQ = 38 MPa·m0.5	Both energy density and oxygen content strongly affect relaxation, GFA, and embrittlement.	[78]
Zr48Cu48Al4	Copper-mold suction-casting method/B2-containing composite plus cryogenic thermal cycling	Yield strength from 1680 MPa to 1832 MPa; plastic strain from 5.6% to 9.1%	Matrix rejuvenation and transformable second phases can act cooperatively	[79]
(Zr56Co28Al16)99.25Nb0.75	Cast BMG; nanoscale structural heterogeneity	Ultimate compressive strength 2318 ± 29 MPa; plastic strain 5.1± 0.1%	Minor alloying can promote shear-band multiplication and deformation-induced nanocrystals	[80]
Al79Ni5Fe5Y11, Al79Ni11Fe5Y5	Melt-spun amorphous/nanocrystalline ribbons; cast ingots crystalline	First crystallization near 387–408 °C	Rapid cooling remains essential even for chemically complex Al-rich alloys	[8]
Al–Ni–Y	Copper-mold-cast monolithic glass, 0.75 mm rod	Compressive yield 1090 MPa; fracture strength 1150 MPa	High specific strength is attainable, but cast section size remains small	[9]
Al86Ni6Y4.5Co2La1.5	SPS consolidation of amorphous powder; fully dense 10 mm diameter disk	Fracture strength 1250 MPa	Powder consolidation can produce larger bulk samples without requiring high melt GFA	[10]
Al86Ni6.75Co2.25Y3.25La1.75	Cu-mold casting, salt-fluxing treatment/fully amorphous BMG	Critical diameter increased from 1.5 to 2.5 mm after fluxing	Melt purification and suppression of oxide induced nucleation improve GFA	[81]
Al74.26Ni7.69Co2.49Y11.30L4.27	Data-driven composition design; experimentally validated amorphous alloy	High GFA and enhanced pitting resistance	Machine-learning guided multi-element optimization enables simultaneous GFA and corrosion resistance improvement	[82]

For Zr-based alloys, additive manufacturing offers a way to overcome some of the geometric limitations of conventional casting. Li et al. produced Zr_59.3Cu_28.8Nb_1.5Al_10.4 by laser powder-bed fusion and obtained a relative density above 99% and a fully amorphous structure according to XRD. This was achieved within an energy-density range of about 20–33 J mm⁻³ using coarse, low-oxygen powder. The best specimens showed about 6% compressive plastic strain and a fracture toughness of approximately 38 MPa m^1/2. At higher energy densities, however, structural relaxation, devitrification, and embrittlement occurred [78]. These results show that oxygen content and powder size, together with energy density and repeated thermal exposure during processing, can strongly affect the final structure and properties.

Recent studies also show that a controlled heterogeneous structure can improve the balance between strength and plasticity. In Zr₄₈Cu₄₈Al₄ composites containing 15–17 vol.% B2 CuZr, cryogenic thermal cycling followed by sub-T_g heating increased the compressive yield strength from 1680 to 1832 MPa and the plastic strain from 5.6% to 9.1%. The improvement was attributed to rejuvenation of the amorphous matrix, partial B2 → B19′ martensitic transformation, and more effective blocking of shear-band propagation [79]. In (Zr₅₆Co₂₈Al₁₆)_99.25Nb_0.75, an ultimate compressive strength of 2318 ± 29 MPa and a plastic strain of 5.1 ± 0.1% were reported. These properties were related to nanoscale structural heterogeneity and interactions between shear bands and deformation-induced nanocrystals [80].

The Al-based case is different. Conventional melt casting has produced a fully amorphous Al–Ni–Y rod only 0.75 mm in diameter, with compressive yield and fracture strengths of 1090 and 1150 MPa, respectively [9]. Powder-based processing can be used to produce relatively large bulk samples even when the alloy has limited glass-forming ability during conventional melt casting. For example, spark-plasma sintering of amorphous Al₈₆Ni₆Y_4.5Co₂La_1.5 powder produced a fully dense glass with a diameter of 10 mm and a reported fracture strength of 1250 MPa [10]. These examples directly demonstrate both the attractive strength of Al-rich glasses and the severe processing constraint imposed by their low GFA.

Recent studies on Al-rich alloys have continued to examine how composition and processing conditions affect the resulting structure and glass-forming ability. Melt-spun Al₇₉Ni₅Fe₅Y₁₁ and Al₇₉Ni₁₁Fe₅Y₅ ribbons consisted mainly of amorphous matrices with nanocrystals detected by neutron diffraction, whereas the ingots cooled under conventional conditions showed a multiphase crystalline structure. The first crystallization events occurred at about 408 °C and 387 °C, respectively [8]. These results show that neither a lower Al content nor greater chemical complexity removes the need for sufficiently rapid cooling. They also highlight the importance of detecting nanoscale crystalline phases, which may not be apparent from a broad XRD halo alone. At the same time, appropriate compositional and melt-processing strategies can substantially improve the glass-forming ability of Al-rich alloys. For example, Al₈₆Ni_6.75Co_2.25Y_3.25La_1.75 has demonstrated exceptional glass-forming ability, with a critical casting diameter of 2.5 mm achieved after salt-fluxing treatment, compared with 1.5 mm without fluxing [81]. The improvement was attributed primarily to the removal of oxygen impurities and the consequent suppression of heterogeneous crystallization. More recently, data-driven alloy design has been applied to Al-based metallic glasses, leading to the experimentally validated composition Al_74.26Ni_7.69Co_2.49Y_11.30La_4.27, which combines high glass-forming ability with enhanced pitting-corrosion resistance [82]. These results demonstrate that the glass-forming ability of Al-rich alloys is governed not only by cooling rate but also by a combination of multicomponent chemical interactions, melt cleanliness, and optimized alloy composition.

For both alloy families, ΔT_x is mainly useful as an indicator of processing potential. A wide supercooled liquid region can be beneficial for thermoplastic forming, welding, powder consolidation, and controlled partial crystallization, provided that the viscosity and crystallization kinetics are suitable for the process. However, ΔT_x alone does not indicate the ductility of the material at room temperature. Mechanical behavior is governed mainly by how STZs and shear bands form and develop, while in composites the second phase must help redistribute stress without becoming a site for premature cracking [66,67,68].

The comparison of suggests different approaches for the two alloy families. For cast Zr-based BMGs, the main priorities are a low liquidus temperature, high GFA, low oxygen content, and avoiding the formation of brittle crystalline phases. For Zr-based composites, greater attention should be paid to the fraction, size, spacing, and morphology of the crystalline phase, as well as to its ability to undergo a beneficial transformation. For Al-rich glasses, alloy design must be combined with a processing method that provides the required cooling rate. An alternative is powder consolidation below the crystallization temperature, with sufficient disruption of the oxide layers at the particle interfaces. In both alloy families, thermal characteristics should be reported together with the DSC conditions used for their measurement. Mechanical data should also specify the loading mode, specimen geometry, strain rate, and the particular strength or fracture parameter being reported.

6. Conclusions

A comparison of Al- and Zr-based amorphous alloys shows that glass formation cannot be predicted from a single thermodynamic parameter. A more useful approach is to consider several factors together, including mixing enthalpy, atomic-size mismatch, liquidus temperature, and kinetic parameters such as T_rg and γ. Oxygen content, heterogeneous nucleation, and the cooling rate achievable during processing also need to be considered. For the alloys reviewed here, Zr-based compositions generally have lower critical cooling rates and wider supercooled liquid regions than Al-rich alloys. The parameter ΔT_x = T_x − T_g is mainly useful for evaluating the processing window for thermoplastic forming or controlled devitrification. It should not be taken as a direct indication of ductility or superplasticity. Since T_x depends on the heating rate, values obtained at different rates should only be compared when the DSC conditions are specified. Mechanical properties should be discussed in terms of the loading mode and the type of strength measured. In monolithic metallic glasses, plastic deformation occurs through the activation of shear-transformation zones (STZs), which can subsequently develop into localized shear bands, rather than through conventional dislocation glide. A crystalline second phase can improve plasticity by promoting the formation of multiple shear bands, deflecting their propagation, or preventing their growth. It can also absorb energy through plastic slip or phase transformation. However, its effect depends on the phase fraction, size, spacing, morphology, and interface cohesion. Therefore, the approximately 50 vol.% crystalline fraction reported for Zr₅₇Cu₂₀Ni₈Al₁₀Ti₅ should be considered specific to this alloy and the testing conditions, rather than as a general transition point. Zr-based BMGs currently offer a good combination of castability and mechanical properties and can also be processed into fully amorphous components by additive manufacturing. Al-rich glasses have high specific-strength potential, but producing bulk amorphous materials remains more difficult. Current approaches therefore rely mainly on submillimetre casting, rapidly solidified powders, and consolidation methods that limit crystallization and help break up oxide layers at particle interfaces.