**3. Results and Discussion**

Figure 1a–d shows temperature dependences of the shear modulus *G*(*T*) for all MGs under investigation. The room-temperature shear moduli *Grt* listed in this Figure are taken from Ref. [26]. It is seen that *<sup>G</sup>*(*T*)-patterns for all MGs are quite similar. Heating of glassy samples leads to

a monotonous decrease of *G* while the slope |*dG*/*dT*| decreases by several times near *Tg*, which constitutes a typical behavior upon high-frequency *G*-measurements [17]. Heating by 50–70 K above *Tg* results in the beginning of crystallization, which leads to an increase of the shear modulus by 22% (*ZrFe*1) to 40% (*ZrNi*10). In the crystalline state (2nd run), the shear modulus *μ*(*T*) demonstrates a featureless decrease with temperature.

**Figure 1.** Temperature dependences of the shear modulus measured at 3 K/min for the glasses under investigation in the initial state (**<sup>a</sup>**–**<sup>c</sup>**), after relaxation (**d**) and after full crystallization (**<sup>a</sup>**–**d**). The calorimetric *Tg*'s are indicated by the arrows. The room-temperature shear moduli are taken from Ref. [26].

Figure 2 shows DSC traces of the MGs under investigation at the same heating rate of 3 K/min. Initial glasses (*ZrAl*5, *ZrNi*10 and *ZrFe*1) demonstrate (*i*) exothermal reaction below *Tg* (not seen in the scale of Figure 2a–c), (*ii*) heat absorption in the supercooled liquid state (i.e., at temperatures *Tg* ≤ *T* < *Tx*), which is bigger compared with that of the glassy and crystalline phases, and (*iii*) large heat release due to the crystallization. In the relaxed MG (*ZrTi*5), the feature (*i*) is absent, as one would expect.

Figure 2 also gives the heat flow <sup>Δ</sup>*W*(*T*)-curves calculated with Equation (1) at *T*˙ = 3 K/min using the corresponding temperature dependences of the shear moduli *G*(*T*) and *μ*(*T*) given in Figure 1. The way of the determination of the shear susceptibility *β* from the experimental data and its values together with the densities for the MGS under investigation are given in Ref [26]. The insets in Figure 2 show experimental and calculated <sup>Δ</sup>*W*(*T*)-curves in the supercooled liquid and crystallization temperature regions on enlarged scales. Figure 2 in general demonstrates that the calculation provides a rather good quantitative reproduction of the experimental <sup>Δ</sup>*W*(*T*)-curves. First of all, this applies to the heat absorption for all glasses in the supercooled liquid state. Temperature position of the crystallization exothermal peak is reproduced within ≈10 K for *ZrNi*10 and ≤6 K for other glasses. The height of this peak is reproduced within 10%–15% accuracy with the exception of *ZrFe*1 for which the calculation gives about 60% of the experimental height (not shown in Figure 2). It should be emphasized that the density of data points provided by our EMAT system in the range of fast crystallization (i.e., upon a rapid change of the shear modulus) is significantly smaller than that for temperatures *T* < *Tx* (see Figure 1) and this constitutes a significant source for the calculation errors when using Equation (1), which contains both derivatives of the shear moduli and their difference. Nonetheless, one can conclude that this equation provides a good description of the heat effects on the basis of shear modulus data. It is interesting to note that the calculation reproduces the fine details of the crystallization kinetics as exemplified by the data around *T* = 700 K shown in the inset of Figure 2d.

**Figure 2.** Experimental and calculated using Equation (1) DSC traces for the MGs under investigation in the initial (**<sup>a</sup>**–**<sup>c</sup>**) and relaxed states (**d**). The insets show the heat flow in the supercooled liquid state and upon crystallization on an enlarged scale. The shear susceptibilities *β* and the densities *ρ* are taken from Ref. [26].

The obtained results confirm the basic idea of the IT sketched above: the origin of the heat absorption and heat release occurring in the glass transition range as well as upon crystallization can be understood as a result of the change of the concentration of interstitial-type defects frozen-in from the melt upon glass production. In particular, structural relaxation below *Tg* provides relatively small decrease of the defect concentration and constitutes the reason of moderate exothermal heat effect (feature (*i*) mentioned above). The defect concentration above *Tg* increases with temperature leading to the heat absorption (feature (*ii*)). Finally, fast crystallization at *T* > *Tx* results in a rapid drop of the defect concentration down to zero providing quick dissipation of the defect elastic energy into heat, which is released as a strong crystallization peak (feature (*iii*)) [20].

As mentioned above, the interstitial defects in crystalline structures are unambiguously geometrically defined as two atoms trying to occupy the same lattice site. However, the situation in metallic non-crystalline structures is more complex. Molecular static simulation performed on a

monoatomic glassy structure revealed localized nano-regions, which display the properties similar to those of dumbbell interstitials in simple metals [16]. These findings imply that interstitial-type defects can indeed exist in a monatomic non-crystalline structure. The issue whether such structural entities can be found in polyatomic glassy structures constitutes a major challenge for further research in this direction.
