**1. Introduction**

Paraffin waxes are defined as the materials consisting of saturated carbonhydrogen chains integrated with branched, straight and ring-like (aromatics) structures [1] that are relatively complex by nature. This chemical configuration endows the amorphous characteristics and inertness to paraffin waxes, resulting in inactive functional groups where the external chemical reactions become impossible. Based on this stance, paraffin waxes are supposed to be the green thermal reservoirs, lying within the sustainable targets of the current era. Therefore, the applications of paraffin waxes ranging from biomedical [2] to thermal storage/release [3–5] are declared relatively safe and environmental-friendly.

In thermal storage/release applications, phase change of the paraffin waxes is the backbone that governs the under-lying mechanism depending on thermal excitation or de-excitation driven by the heat source. Thermal excitation is the phase change process of melting during which thermal energy is absorbed, while thermal de-excitation is the phase change process of solidification during which thermal energy is released. These both processes build up the reversible functionality involving the sensible heat and latent heat storage/release. In case of paraffin waxes (or generally for PCMs), the sensible heat is counted before the phase change process, while latent heat is considered during the phase change process. By definition, phase change [6] process refers to either structural change or state change, as shown in **Figure 1**. In structural change, a single phase of paraffin wax undergoes thermal excitation that brings about the conversion of the internal structures and it is called the solid–solid phase change. Whilst, in the state change, two phases of the paraffin wax undergoes thermal excitation, resulting in conversion of solid phase into liquid phase at the melting temperature, which is known as solid–liquid phase change. The occurrence of either kind of phase change is dependent on the melting temperatures of the paraffin waxes. The paraffin waxes with low-melting temperature such as <40°C (or also called soft paraffin waxes) do not demonstrate structural change, so they only undergo state change. While, the paraffin waxes with high-melting temperature (or also called hard paraffin waxes) provide both structural change and state change.

#### **Figure 1.**

*Classification of phase change and paraffin wax-based phase change materials. A few contents of this figure are partially restructured from references [6, 7].*

#### *Paraffin Wax-Based Thermal Composites DOI: http://dx.doi.org/10.5772/intechopen.97195*

In between structural change and state change, another phase consisting of solid flakes and liquid sites ensues which is called mush phase. Normally, the temperature range of the mush phase lies closer to the onset temperature of the liquid phase. In brief, all these phases are responsible for the total latent heat, while before and after these phases, total sensible heat is considered, depending on the melting temperatures of the involved paraffin waxes. Therefore, it is emphasized that the total thermal energy of the paraffin waxes should carefully include thorough observation of the onset and endset temperatures which can be determined either by DSC curves (**Figure 1**) or by the transient heat diffusion process (temperature versus time analysis). With help of transient heat diffusion process, the thermal storage/ release performance of paraffin waxes can precisely be assessed by looking into the time consumed during the melting (thermal excitation) and solidification (thermal de-excitation) processes, which is called charging time and discharging time, respectively. It is generally deducted that, the charging/discharging time becomes short, if the heat diffusion occurs fast, and vice versa. The heat diffusion is further related to the intrinsic property of materials known as thermal conductivity, which is defined as the ability of material allowing the fast or slow transfer of heat. In short, with high thermal conductivity, the charging/discharging time are reduced on behalf of faster heat transfer. However, heat transfer in paraffin waxes is sluggish, and their charging/discharging time is sufficiently long based on the low thermal conductivity. The possible reason for the low thermal conductivity of paraffin waxes is their amorphous nature wherein tightly-packed and inter-connected thermal paths are unavailable. However, a great care is essential because every application may not need fast heat transfer rates, for example, thermal management of buildings where the objective is to keep the heat either outside the buildings in summer or inside the buildings in winter, which is possible only if paraffin wax serves as a thermal insulator necessitating slow heat transfer rates, but the challenge of liquid drainage needs to be simultaneously addressed. Therefore, depending on the target applications, enhancing the thermal conductivity of paraffin waxes is of great practical interest so that the charging/discharging can be reduced. For example, the long charging/discharging time of paraffin waxes is the major bottleneck that can potentially hampers their functionality for thermal management of batteries where the heat generation rate is prone to high and the objective is to dissipate the heat so that hot-spots can be avoided. Keeping different practical scenarios ahead, thermal reinforcements are essential to improve the thermo-physical bottlenecks of the base materials (paraffin waxes), helping create paraffin wax-based thermal composites (PWTCs) which are discussed henceforth.


**1. Introduction**

In recent years, an increasing number of studies have shown the impact of mild traumatic brain injury (mTBI) on cognitive functions. The fact that mTBIs are considered a major health issue involving long-term health risks raises questions as to how their impact can be best identified and measured in order to be treated and managed optimally.

A TBI is produced by a large transfer of energy generated by a direct impact of the head against a hard surface, or by forces (acceleration, deceleration, rotation) created during the impact [1]. In the case of mTBI, these forces are responsible for stretching of the axons and microbleeds, which lead to a complex neurometabolic cascade [2]. There is now growing and reproducible evidence that following mTBI microstructural damage and neurochemical imbalances occur in a number of brain regions (e.g., frontal, temporal, motor cortex) and in white matter integrity (e.g., corpus callosum) [3–8]. These alterations have been related to the known post-mTBI physical (e.g., headaches, drowsiness, fatigue, dizziness), cognitive (e.g., attention and memory problems, bradyphrenia), and a ffective (e.g., irritability, depression, anxiety) symptoms [1,9,10]. These symptoms are more intense in the first days and weeks following the injury and slowly decrease during the subacute recovery period, which has been described as the first three months post-trauma. However, in a non-negligible proportion of individuals these symptoms are at risk of becoming chronic [11].

Attentional problems such as di fficulties orienting and disengaging attention [12], sustaining it over time, and/or dividing attentional resources across multiple simultaneous demands [13,14] are amongst the most common complaints reported by individuals after a mTBI. On a behavioural level, these deficits usually disappear over a period of seven to 10 days [10,12,15]. However, mTBI can have negative long-term neuropsychological impacts on subtle aspects of complex attention and working memory [16,17], even with normal behavioural performances, where some individuals report the persistence of attentional di fficulties that interfere with the demands of their daily life and social participation. The need for a more sensitive functional measure of the impact of mTBI on distinct steps involved in the deployment of attention thus remains.

Event-related brain potentials (ERPs) can represent a relatively simple, inexpensive, and precise answer to that question because they allow us to assess the integrity (reflected by the amplitude of a component) and e fficiency (reflected by its latency) of specific and complex cognitive processes [1,18]. With excellent temporal resolution, ERPs possess an advantage over reaction time measures because they provide measures of multiple stages of cognitive processing, instead of the summation of the duration of all the intervening mechanisms involved in the generation of a response [18].

Experimental paradigms have historically been designed to study precise cognitive processes along with their event-related components. For instance, oddball paradigms have been used to evoke the P3, a component reflecting a target stimulus upgrade in working memory [19]. Such paradigms have been shown to be sensitive following mTBI [20–24]. However, since mTBI can present with deficits at di fferent levels within the deployment of attention, one needs to be able to rapidly track the steps of the process that are distinctly a ffected.

The objective of this study was to track, with the help of a single task, the deployment of visuospatial attentional mechanisms and to identify underlying deficits at di fferent recovery time points after mTBI. The task we used allowed us to measure the integrity of the di fferent visual and cognitive processes involved in the deployment of visuospatial selective attention. We studied components reflecting earlier visual–attentional processing, such as the N1 (discrimination processes within the focus of visual attention and central attention) [25,26], P1 (facilitation e ffect for stimuli presented at an attended localisation) [27], and P2 (early attentional modulation) [28] components, which we expected would not be di fferent between the mTBI and control groups. We also investigated later attention-related and visuocognitive mechanisms such as the N2pc (deployment and orientation of visuospatial attention) [28], Ptc (target isolation once identified among distractors) [29], SPCN (encoding in visual short-term memory, working memory capacity) [30], P3a (disengaging of attention to re-orient toward novel stimuli) [31], and P3b (tracking of task-relevant stimuli during updating in working memory) [19] components, which were expected to be a ffected, at least for later components.

#### **2. Materials and Methods**


*Article* **Multicomponent Fe-Based Bulk Metallic Glasses with Excellent Corrosion and Wear Resistances**

**Guan Zhang 1,2, Wenlei Sun 1,\*, Lei Xie 3,4, Chengwu Zhang 3,4, Jie Tan 5, Xuan Peng 2, Qiang Li 4,\*, Xu Ma 3,4, Dongmei Zhao 6 and Jiangtong Yu 1**

**Abstract:** In this study, new multicomponent Fe54M5Cr15Mo6Si2B4P10C4 (M = Fe, Co, and Ni, denoted as Fe59, Fe54Co5, and Fe54Ni5, respectively) bulk metallic glasses (BMGs) with excellent corrosion and wear resistances were synthesized using the J-quenching technique and fluxing treatment. The synthesized Fe-based BMGs possessed a large glass-forming ability, and the maximum diameters of the Fe59, Fe54Co5, and Fe54Ni5 glassy alloy rods reached 5.5, 4.5, and 4.0 mm, respectively. The Fe59 BMG had a wide supercooled liquid region of 65 K. Potentiodynamic tests in 3.5 wt.% NaCl solution showed that the corrosion resistances of the synthesized Fe-based BMGs were relatively better than that of the 316L stainless steel. The Fe59 BMG had the highest corrosion resistance, with the lowest self-corrosion current density in the order of 10−8 A·cm−2. Wear tests showed that the synthesized Fe-based BMGs exhibited excellent wear resistances, and the wear rate of the Fe59 BMG was as low as approximately 1.73 × 10−15 m3·N−1·m−1. The rare-earth-element-free Fe-based BMGs, especially the Fe59 BMG, have a low cost, large glass-forming ability, and excellent wear and corrosion resistance, which make them good candidates for wear-and corrosion-resistant coating materials.

**Keywords:** Fe-based bulk metallic glasses; glass forming ability; corrosion resistance; wear resistance

#### **1. Introduction**

Fe-based bulk metallic glasses (BMGs) have gained substantial attention in recent decades owing to their low cost and unique performance, such as high strength and hardness, excellent wear and corrosion resistance, and excellent soft magnetic properties [1–4]. The strength of Fe-based BMGs is generally higher than 2 GPa, which currently exceeds the strength of high-strength steel; in particular, the compressive fracture strength of a Fe33.5Co33.5Nb6B27 BMG reaches 4.84 GPa [5]. Owing to their high strength, the microhardness of the Fe-based BMGs is larger than that of the Zr- and Cu-based BMGs, and their wear resistance is three orders of magnitude higher than that of the Zr- and Cu-based BMGs. Additionally, owing to their unique amorphous structure, Fe-based BMGs exhibit a chemically homogeneous single phase and high atomic reactivity, which results in the formation of an extremely uniform and stable passive film on the surface of the alloys; thus, Fe-based BMGs have higher corrosion resistances compared with their crystalline alloy counterparts. For example, Li et al. reported an Fe59Cr12Mo3.5Ni5P10C4B4Si2.5 BMG with excellent corrosion resistance, with *I*corr and *E*corr values of 2.47 × 10−7 A·cm−2 and −0.22 V, respectively [6].

**Citation:** Zhang, G.; Sun, W.; Xie, L.; Zhang, C.; Tan, J.; Peng, X.; Li, Q.; Ma, X.; Zhao, D.; Yu, J. Multicomponent Fe-Based Bulk Metallic Glasses with Excellent Corrosion and Wear Resistances. *Metals* **2022**, *12*, 564. https://doi.org/10.3390/ met12040564

Academic Editor: George A. Pantazopoulos

Received: 18 February 2022 Accepted: 24 March 2022 Published: 27 March 2022

**Publisher's Note:** MDPI stays neutral with regard to jurisdictional claims in published maps and institutional affiliations.

**Copyright:** © 2022 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (https:// creativecommons.org/licenses/by/ 4.0/).

Owing to their low cost, ultra-high strength, and excellent wear and corrosion resistance, Fe-based BMGs have been considered good candidates as wear-and corrosionresistant coating materials [7]. It has been extensively reported that alloying elements can significantly influence the corrosion and wear properties of Fe-based BMGs. Cr and Mo are the most effective elements for providing a high passivation ability for Fe-based BMGs [8,9]. Cr can form a dense and stable hydrated chromium oxyhydroxide passive film on the surface of alloys to prevent corrosion within the alloys [10]. Generally, the higher the Cr content, the better the corrosion resistance of Fe-based BMGs. However, when the Cr content is greater than 15 at.%, further increasing the Cr content does not significantly improve the corrosion resistance of the Fe-based BMGs but increases the material cost. The effect of Mo on the protective ability of the passivation layer is weaker than that of Cr. However, Mo can promote the enrichment of Cr in the passive film, thereby enhancing the corrosion and pitting resistances of Cr-containing Fe-based BMGs. However, the excess addition of Mo is detrimental to the corrosion resistance of Fe-based BMGs [11]. Recently, it was found that the substitution of Fe with a similar element, such as Ni or Co, can produce positive effects on the glass forming ability (GFA) and the mechanical and magnetic properties of Fe-based BMGs. For example, the substitution of 20 at.% Ni for Fe in Fe80P13C7 BMG resulted in an increase in the critical diameter (*D*c) from 2.0 mm to 2.3 mm for full glass formation and a significant increase in room-temperature compressive plastic strain from 1.1% to 11.2% [12]. The room-temperature compressive plasticity of Fe80P13C7 BMG improved from 1.1% to 3.0%, and the saturation magnetization increased from 1.45 T to 1.55 T due to the replacement of 10 at.% Fe with Co [13]. Based on the above considerations, new Fe54M5Cr15Mo6Si2B4P10C4 (M = Fe, Co, and Ni) BMGs were successfully prepared in this study, and the corrosion and wear properties of the present Fe-based BMGs were systematically investigated.

#### **2. Materials and Methods**

Fe54M5Cr15Mo6Si2B4P10C4 (M = Fe, Ni, and Co) alloy ingots were prepared by torchmelting a mixture of Fe3P pieces (99.5 mass% purity), Cr powder (99.9 mass% purity), Mo powder (99.9 mass% purity), Fe powder (99.9 mass% purity), Si powder (99.99 mass% purity), boron pieces (99.95 mass% purity), graphite powder (99.95 mass% purity), Co powder (99.9 mass% purity), and Ni powder (99.7 mass% purity) under a high-purity Ar atmosphere. All the as-prepared alloy ingots had a mass of 2 g. Subsequently, the ingots were purified through a fluxing treatment at an elevated temperature of approximately 1450 K for 2–3 h under a vacuum of ~50 Pa, in which a mixture of B2O3 and CaO with a mass ratio of 3:1 was used as the fluxing agent. After the fluxing treatment, the alloys were subjected to the J-quenching technique [14–16]. Consequently, Fe54M5Cr15Mo6Si2B4P10C4 (M = Fe, Ni, and Co) alloy rods with diameters ranging from 1.0–6.0 mm and lengths of several centimeters were prepared.

The amorphous nature of the specimens was examined using an X-ray diffractometer (XRD, Bruker D8 Discover, Bruker Inc., Karlsruhe, Germany) with Co Kα radiation. The thermal behavior of the specimens was examined through differential scanning calorimetry (DSC, NETZSCH DSC 404F1, NETZSCH Inc., Selb, Germany) at a heating rate of 0.33 K/s under an Ar atmosphere. The corrosion behaviors of the specimens, as well as that of 316 stainless steel, which was used for comparison, were evaluated through electrochemical measurements (Zahner Zennium X, Zahner Inc., Kronach, Germany). This was conducted in a three-electrode cell using the glassy rod alloy specimens as the working electrode, a saturated calomel electrode as the reference electrode, and a metal platinum electrode as the counter electrode. The working face of the samples was polished to a mirror face using sandpaper, and the non-working parts were sealed with epoxy resin. The potentiodynamic polarization curves were measured in a 3.5% NaCl solution at room temperature with a potential sweep rate of 1 mV/s after immersing the specimens for approximately 30 min to stabilize the open-circuit potential. Electrochemical impedance spectroscopy (EIS) measurements were carried out in a frequency range of 10 kHz to

10 mHz, using a sinusoidal potential perturbation of ±10 mV relative to the OCP. The fitting of EIS data was operated by using the software ZSimpWin (3.60, AMETEK Inc. Middleboro, MA, USA). The surface morphology of the specimens after immersion was observed using scanning electron microscopy (SEM, SU8010). The chemical states of the surface elements of the specimens after polarization were analyzed using X-ray photoelectron spectroscopy (XPS, ESCALAB 250Xi, Thermo Fisher Scientific Inc., Waltham, MA, USA) with a monochromatic Al Kα X-ray source (hν = 1486.6 eV). The binding energies were calibrated using carbon contamination with a C 1s peak value of 284.8 eV. The Vickers hardness (HV) of the specimens was measured at a load of 200 g and dwell time of 15 s using a microhardness test instrument (MT-401MVA, BangYi Co., Shanghai, China). Prior to testing, the specimens were polished to a mirror face using sandpaper and cleaned with alcohol, and the average of ten measurement values was taken as the final hardness value for each specimen. Tribological tests of the specimens were conducted under ambient conditions using a pin-on-disc tribometer (CSM, Graz, Austria) with a pin-on-disk contact geometry; a GCr15 (860HV) plate with a diameter of 60 mm was used as the counter body.  The wear rate was calculated using Equation (1).

$$V = \frac{\Delta m}{PS\rho} = \frac{m\_1 - m\_2}{PS\rho} \tag{1}$$

where *V* is the wear rate, Δ*m* is the wear mass loss, *m*1 and *m*2 are the masses of the sample before and after the test, respectively, *S* is the sliding distance of the sample, *p* is the load, and *ρ* is the density of the test sample, which was measured using density testing equipment (Micromeritics, AccuPyc-1340, Micromeritics Inc., Atlanta, GA, USA). The surface and subsurface morphologies of the worn tracks were observed using scanning electron microscopy (SEM, Hitachi s4800, Hitachi lnc., Tokyo, Japan). The topography and chemical composition of the worn surfaces were examined using energy-dispersive X-ray spectrometry (EDS HORIBA 7593-H, HORIBA lnc., Tokyo, Japan).

#### **3. Results and Discussion**

#### *3.1. Material Preparation and Glass Forming Ability (GFA)*

The XRD (X-ray diffraction) patterns of the as-prepared Fe54M5Cr15Mo6Si2B4P10C4 (M = Fe, Ni, and Co) glassy alloy rods with critical diameters (*D*max) for glass formation are shown in Figure 1a. There was only a broad diffraction peak and no detectable sharp crystalline peaks in the XRD patterns of all the specimens, indicating the fully amorphous structure of the specimens [17]. Among the present Fe-based BMGs, the Fe59 alloy had the largest GFA with a *D*max of 5.5 mm, and the substitution of Fe with 5 at.% Co or Ni degraded the GFA of Fe59 alloy; the *D*max of Fe54Co5 and Fe54Ni5 alloys were 4.5 and 4.0 mm, respectively [17]. Figure 1b shows the DSC curves of Fe54M5Cr15Mo6Si2B4P10C4 (M = Fe, Ni, and Co) BMGs at a heating rate of 0.33 K/s. The DSC (Differential scanning calorimetry) curves of all specimens exhibited a distinct glass transition, followed by a supercooled liquid region and multistep crystallization peaks. The thermal parameters associated with the glass transition temperature (*T*g), crystallization temperature (*T*x), and supercooled liquid region (Δ*T*x = *T*x − *T*g) are summarized in Table 1. It can be seen that the substitution of Co and Ni for Fe led to a decrease in the *T*g and *T*x of the Fe-based BMGs. The *Tg* of amorphous alloys mainly depends on the atomic bonding strength among the constituent elements, which can be evaluated through the mixing enthalpy between the constituent elements. It is suggested that the dominant short-range order unit for Fe-metalloid amorphous alloys can be characterized as solute-centered atomic clusters, in which the metallic atoms are located on the shell and the metalloid atoms are located in the center [18]. The mixing enthalpies between Fe and the metalloid elements Si, B, P, and C are −35, −26, −39.5, and −50 kJ/mol, respectively; those between Co and the metalloid elements are −38, −24, −35.5, and −42 kJ/mol, respectively; and those between Ni and the metalloid elements are −40, −24, −34.5, and −39 kJ/mol, respectively [19]. Except for Si, the mixing enthalpies between Fe and the metalloid elements were more negative than those between Co/Ni and the metalloid elements, which could be the reason for the decrease in the *T*g and *T*x of the Fe59, Fe54Co5, and Fe54Ni5 alloys. It was also observed that the Δ*T*x of the Fe59, Fe54Co5, and Fe54Ni5 BMGs decreased. Inoue suggested that a wider Δ*T*x reflects a better GFA [20], which may account for the compositional dependence of *D*max of the Fe-based BMGs in this study. Additionally, it can be noted that the Δ*T*x of the Fe59 BMG reaches 65 K, indicating an extremely high thermal stability.

**Figure 1.** (**a**) XRD patterns and (**b**) DSC curves of as-prepared Fe54M5Cr15Mo6Si2B4P10C4 (M = Fe, Ni, and Co) glassy alloy rods with the critical diameter (*D*max) for glass formation. Photos of (**c**) Fe59, (**d**) Fe54Co5, and (**e**) Fe54Ni5 glassy alloy rods with the *D*max values.

**Table 1.** Summary of the critical diameters for glass formation (*D*max), thermal properties, and the GFA indicators of Fe54M5Cr15Mo6Si2B4P10C4 (M = Fe, Ni, and Co) glassy alloy rods determined from the DSC curves at a heating rate of 0.33 K/s (*Tg*: glass transition temperature, *Tx*: onset crystallization temperature, Δ*Tx* = *Tx* − *Tg*).

#### *3.2. Corrosion Resistance*

To evaluate the corrosion resistance, electrochemical measurements of the Fe-based BMGs were performed in a 3.5 wt.% NaCl solution at room temperature. Figure 2 shows representative potentiodynamic polarization curves of the Fe59, Fe54Co5, and Fe54Ni5 glassy alloy rods. The self-corrosion current density (*I*corr) and self-corrosion potential (*E*corr) obtained using the Tafel extrapolation method are listed in Table 2.

**Table 2.** Electrochemical parameters derived from potentiodynamic polarization curves and EIS analysis of the Fe54M5Cr15Mo6Si2B4P10C4 (M = Fe, Co, and Ni) BMGs and 316L stainless steel (*E*corr: self-corrosion potential, *I*corr: self-corrosion current density, *R*s: solution resistance, *R*t: charge transfer resistance, *R*u: charge transfer resistance, *Q*t: non-ideal capacitance, *Q*u: non-ideal capacitance).

**Figure 2.** (**a**) Potentiodynamic polarization curves in 3.5 wt% NaCl solution at room temperature for Fe54M5Cr15Mo6Si2B4P10C4 (M = Fe, Ni, and Co) glassy alloy rods and 316L stainless steel; (**b**) Nyquist plots and the corresponding equivalent circuit (inset) for fitting the impedance spectra of Fe54M5Cr15Mo6Si2B4P10C4 (M = Fe, Co, and Ni) glassy alloy rods and 316L in 3.5 wt.% NaCl solution at room temperature.

Figure 2a shows the potentiodynamic polarization curves in 3.5 wt.% NaCl solution exposed to air at room temperature for the Fe-based glassy alloy rod specimens as well as the 316L stainless steel (316L) for comparison. The self-corrosion potential (*E*corr) and self-corrosion current density (*I*corr) of the specimens listed in Table 2 were obtained from the potentiodynamic polarization curves using the Tafel extrapolation method. Compared with the 316L, the Fe-based BMGs exhibited a more positive *E*corr, lower *I*corr, and a more stable and wider passive region, indicating a relatively higher corrosion resistance and superior passive film protection. Among the Fe-based BMGs used in this study, the Fe59 BMG exhibited the highest corrosion resistance, with the lowest *I*corr and *I*pass in the order of 10−8 and 10−7 A/cm2, respectively. Further, the potentiodynamic polarization curve of the Fe59 BMG had a smooth shape, as shown in Figure 2a, indicating that a stable passivation layer was formed on the surface and tiny pit corrosion occurred during the entire measurement period. However, few sharp peaks were found in the curves of the Fe54Co5 and Fe54Ni5 BMGs, as shown in Figure 2a, indicating that pit corrosion occurred during the measurement period; evidence of pit corrosion is shown in Figure 3. It is worth noting that the corrosion resistance of the Fe59 BMG was better than that of most reported Fe-based BMGs, of which *I*corr is usually in the order of 10−6~10−7 A/cm2. For example, it has been reported that a Fe36Cr23Mo10W8C15B6Y2 BMG possesses good corrosion resistance with an *I*corr of 6.16 × 10−7 A·cm−2 and *E*corr of −0.275 V [1]. Compared with this Fe-based BMG, the Fe59 BMG exhibited a lower *I*corr but contained less Cr and Mo. This result may be related to the metalloid elements in the alloys, which can significantly influence the corrosion properties of Fe-based BMGs [4]. For example, Si can improve the passivation ability and corrosion resistance of Fe-based BMGs, which can be attributed to the formation of a dense and stable passive film rich in Si and Cr- oxides [21]. The Fe43Cr16Mo16C10B5P10 BMGs exhibited a wide passive region, indicating a higher corrosion resistance compared to that of the P-free metallic glass [22].

**Figure 3.** SEM micrographs of the electrochemical corroded surfaces of the Fe59 BMG (**a**), Fe54Co5 BMG (**c**), Fe54Ni5 BMG (**e**), and 316L (**g**), respectively, and (**b**,**d**,**f**,**h**) are corresponding partial enlarged views.

Figure 2b shows the Nyquist plots of the Fe-based BMGs and 316L in 3.5 wt.% NaCl solution at room temperature. The Nyquist plots of all samples show a double capacitive loop, implying two time constants in the electrochemical measurements. The semicircle diameter in the Nyquist plots of the Fe59, Fe54Co5, and Fe54Ni5 BMGs and 316L stainless steel decreased, implying that their corrosion resistances also decreased. An appropriate equivalent circuit for fitting the impedance spectra is shown in Figure 2b, where *R*s is the solution resistance; *R*t and *R*u are the charge transfer resistances; and *Q*t and *Q*u represent the possibility of non-ideal capacitance. These fitting parameters are listed in Table 2. The *R*t values of the Fe59, Fe54Co5, Fe54Ni5 BMGs, and 316L stainless steel decreased. It is known that a larger *R*t implies better corrosion resistance [23], and thus, the results of the EIS measurements agree well with those of the electrochemical measurements.

Figure 3 shows the SEM micrographs of the electrochemical corroded surfaces of the Fe-based BMGs and 316L specimens. There were a few pitting pits and a film rupture zone on the surface of the Fe59 BMG specimen, as shown in Figure 3a,b, revealing the excellent corrosion resistance of the Fe59 BMG. It can be seen in Figure 3c,d that there were pitting pits with diameters of approximately 4 μm and of a few nanometers on the surface of the Fe54Co5 BMG specimen. Compared with the Fe59 and Fe54Co5 BMGs, there were more and larger pitting pits and more corrosion products on the surface of the Fe54Ni5 BMG specimen, as shown in Figure 3e,f, which could be attributed to the sharp peaks in the potentiodynamic polarization curve of the Fe54Ni5 BMG shown Figure 2a. The surface of the 316L specimen was completely damaged, covered with a thick layer of polarized product, and had huge pitting pits caused by a severe corrosion, as shown in Figure 3g,h. The observed electrochemical corroded surfaces of the specimens confirmed the results of the electrochemical and EIS measurements.

To better understand the effect of the elements on the corrosion mechanism of the Fe-based BMGs, an XPS (X-ray Photoelectron Spectroscopy) analysis was performed to characterize the chemical composition of the oxide films formed on the surfaces of the specimens after electrochemical measurement in 3.5 wt.% NaCl solution. Figure 4 shows the Fe 2p, Cr 2p, Mo 3d, Co 2p, and Ni 2p spectra of the Fe-based BMGs. The Fe 2p spectra of the specimens comprised the peaks of Fe 2p1/2 and 2p3/2 corresponding to the peaks of the metallic (Fe) and Fe2+, Fehy3+, and Feox3+ oxide (Feox) states of Fe [23–27]. The peak intensities of Fem in the Fe54Co5 and Fe54Ni5 BMGs were relatively higher than that of the Fe59 BMG. The Cr 2p peaks represented the metallic (Cr) and Cr3+ and Cr6+ oxide (Crox) states of Cr [26,28–30]. Cr6+ dissolves in water and causes the corrosion current density to increase; however, it has a self-repairing ability and can promote the formation of denser Cr2O3 to prevent chloride ions from penetrating into the internal damage [31,32]. As shown in Figure 4b, the total area of the Cr 2p peaks in the Fe59 BMG was the largest, suggesting a higher concentration of Cr on its surface than on the Fe54Co5 and Fe54Ni5 BMGs. The Mo 3d spectrum, consisting of the Mo 3d 3/2 and 3d 5/2 peaks, corresponded to the metallic Mo and Mo4+ and Mo6+ oxide states [24,27,29,30]. The Fe54Co5 BMG had the highest intensity of the low-valence Mo4+ peak on the surface, followed by the Fe59 and Fe54Ni5 BMGs. The Fe59 BMG had the highest intensity of the high-valence Mo6+ peak in the surface film, followed by the Fe54Co5 and Fe54Ni5 BMGs. Mo can promote the enrichment of Cr in the passive film and, consequently, enhance the corrosion resistance of amorphous alloys [33–35]. The Co 2p spectrum consisted of the metallic Co (Com), Cohy2+, and Coox2+ peaks, as shown in Figure 4d. The Ni 2p spectrum consisted of the peaks corresponding to the metallic Ni (Ni), Ni2+, and Ni3+, as shown in Figure 4e. It is known that Co and Ni are more effective in providing a high passivation ability for alloys compared with Fe; however, the corrosion resistance of the Fe-based BMGs in this study decreased with the substitution of Co and Ni for Fe. This could be due to the small amount of crystallization in the Fe54Co5 and Fe54Ni5 BMGs owing to the lower GFA, which degraded their corrosion resistance.

**Figure 4.** XPS spectra of Fe 2p (**a**), Cr 2p (**b**), Mo 3d (**c**), Co 2p (**d**), and Ni 2p (**e**) in the passive films on the surface of the Fe54M5Cr15Mo6Si2B4P10C4 (M = Fe, Co, and Ni) glassy alloy rod specimens after electrochemical measurements in a 3.5 wt.% NaCl solution.

The outstanding corrosion resistance of amorphous alloys is primarily influenced by the stability of a uniform passive film. Figure 4b,c show that the Crox and Moox fractions in the Fe59 and Fe54Co5 BMGs in the surface film after the electrochemical test were apparently higher than those in the Fe54Ni5 BMG. The enrichment of Crox and Moox was beneficial for the modification of the passive film quality of the BMGs and thus improved the corrosion resistance. The XPS results for the Fe59, Fe54Co5, and Fe54Ni5 BMGs were consistent with the SEM morphology of the samples after the electrochemical test, as shown in Figure 3, in which the Fe54Ni5 BMG exhibited a more severely corroded surface than the Fe59 and Fe54Co5 BMGs.

#### *3.3. Microhardness and Wear Resistance*

The measured microhardness values (HV0.2) of the Fe-based BMG specimens are listed in Table 3 and shown in Figure 5a. The Fe-based BMGs exhibited extremely high microhardnesses, exceeding 1000 HV0.2. The microhardness values of the Fe59, Fe54Co5, and Fe54Ni5 BMGs decreased in turn. Similar to the *T*g of amorphous alloys, the hardness of the alloys depends on the atomic bonding strength between the constituent elements; therefore, the microhardness has a similar change trend as the *T*g of the synthesized Febased BMGs.

**Table 3.** Density, average microhardness, average COF, and wear rate of the Fe54M5Cr15Mo6Si2B4P10C4 (M = Fe, Ni, and Co) BMGs.

**Figure 5.** (**a**) Microhardness, (**b**) average coefficient of friction (COF), and wear rate of the Fe54M5Cr15Mo6Si2B4P10C4 (M = Fe, Ni, and Co) BMGs.

Friction and wear tests of the Fe-based BMGs were performed using pin-on-disk experiments, which were run against a GCr15 plate with a diameter of 60 mm and a hardness of 860 HV under a constant normal load of 7 N, a rolling speed of 1.8 m/s, and a sliding distance of 5000 m. The obtained average steady-state coefficient of friction (COF) and the wear rate, which is calculated by Equation (1), of the Fe-based BMG specimens are listed in Table 3 and shown in Figure 5b. Among the Fe-based BMGs, the Fe59 BMG had the smallest wear rate of 1.73 × 10−15 m3·N−1·m−1, followed by the Fe54Ni5 BMG with a wear rate of 5.05 × 10−15 m3·N−1·m−1; the Fe54Co5 BMG had the largest wear rate of 6.18 × 10−15 m3·N−1·m−1. Furthermore, the wear rate of the Fe59 BMG was relatively lower than those of the Fe54Co5 and Fe54Ni5 BMGs. The COF values for the Fe59, Fe54Co5, and Fe54Ni5 BMGs under dry friction conditions were 0.716, 0.795, and 0.727, respectively. It was clear that the wear rates of the Fe-based BMG specimens exhibited an obvious

positive association with their COF under dry sliding conditions [36]. It should also be noted that the wear rates of the Fe54Co5 and Fe54Ni5 BMGs were negatively correlated, but not inversely proportional, with the microhardness. This indicates that the wear resistance of the Fe-based BMGs cannot be simply determined by their hardness, as predicted by the Archard equation.

Figure 6 shows the SEM (Scanning electron microscope) images of the worn areas of the Fe59, Fe54Co5, and Fe54Ni5 BMGs. An EDS mapping test was performed on typical regions of the worn surfaces; their chemical compositions are summarized in Figure 6c,f,i. The worn surfaces of Fe59 exhibited various ploughed grooves parallel to the sliding direction, with different widths and depths, which are typical features corresponding to the abrasive-wear mechanism. In addition, as shown in Figure 6a,b, the worn surface of the Fe59 BMG was covered with tribo-patches, in which a high fraction of oxygen was observed through EDS (Energy Disperse Spectroscopy) mapping analysis, as shown in Figure 6c. This result demonstrates the oxidation of the tribo-patches of the Fe59 BMG during the sliding process. The worn surface of Fe54Co5, as shown in Figure 6d,e, exhibited a smooth morphology covered with a uniform oxide film (Figure 6f), including some small tribo-patches, indicating an oxidation wear mechanism. As shown in Figure 6g,h, the worn morphology of the Fe54Ni5 BMG exhibited much wear debris with grain-like wear particles, which resulted from a brittle failure characteristic, and many tribo-patches with diverse widths and lengths, indicating an adhesive wear mechanism.

**Figure 6.** SEM images of wear surface morphology and EDS mapping of the Fe59 (**a**–**c**), Fe54Co5 (**d**–**f**), and Fe54Ni5 (**g**–**i**) glassy alloy rod specimens.

#### **4. Conclusions**

In this study, a new series of rare-earth-element-free Fe54M5Cr15Mo6Si2B4P10C4 (M = Fe, Ni, and Co) BMGs were prepared, and the effects of Fe, Co, and Ni contents on the thermal stability, corrosion, and wear resistances of the Fe-based BMGs were investigated. The main results are as follows.

(1) Fe54M5Cr15Mo6Si2B4P10C4 (M = Fe, Ni and Co) bulk glassy alloy rods were successfully prepared using the J-quenching and fluxing techniques. The *D*max of the Fe59, Fe54Co5, and Fe54Ni5 alloys reached 5.5, 4.5, and 4.0 mm, respectively.

**Author Contributions:** Conceptualization, G.Z., Q.L. and W.S.; methodology, G.Z. and J.T.; validation, G.Z., Q.L. and W.S.; formal analysis, G.Z.; investigation, G.Z. and L.X.; resources, Q.L. and D.Z.; data curation, Q.L.; writing—original draft preparation, G.Z.; writing—review and editing, G.Z., Q.L., W.S., C.Z., X.P., X.M. and J.Y.; visualization, G.Z.; project administration, Q.L. and W.S.; funding acquisition, G.Z., Q.L. and W.S. All authors have read and agreed to the published version of the manuscript.

**Funding:** This research was funded by the University Science Research Planning Project of Xinjiang Uygur Autonomous Region, grant number XJEDU2021Y008, the key University Science Research Planning Project of Xinjiang Uygur Autonomous Region, grant number XJEDU2021I003, the Tianshan Innovation Team Program of Xinjiang Uygur Autonomous Region, grant number 2020D14038, the Xinjiang Autonomous Region Key Laboratory Open Fund, grant number<PHONE_NUMBER> and the "Talent Projects" of Urumqi Key Training Object.

**Institutional Review Board Statement:** Not applicable.

**Informed Consent Statement:** Not applicable.

**Data Availability Statement:** Not applicable.

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**