Review of the development and application of aluminum alloys in the nuclear industry

3.1 Utilization of aluminum alloy in the development of nuclear materials

Figure 8a depicts the utilization of aluminum alloy in the construction of nuclear fuel components such as side plates, end fittings, and cladding [34–36]. On the other hand, Figure 8b presents the composition of the fuel element plate, consisting of aluminum alloy cladding and uranium meat. The primary function of cladding is to act as a protective covering for the uranium meat fuel and the fission products that arise from the irradiation process. Uranium meat represents the central component of the fuel, which undergoes nuclear fission via a chain reaction, leading to the generation of heat and fission products. Typically, fuel cladding is fabricated using aluminum alloys such as AlMg, Al 6061, and various other aluminum alloys, including AlFeNi.

Figure 8

(a) Components constituting nuclear fuel and (b) plates composing fuel elements [34,36].

3.2 AlMg alloy

Magnesium aluminum alloys belong to the classification group 5xxx series, and they have characteristics that cannot be hardened through heat treatment (non-heat-treatable alloys). The solubility of magnesium in aluminum is approximately 14.9% at 451°C, decreasing as temperature decreases. However, aluminum–magnesium alloys do not show the effect of precipitation hardening. This is because the solubility is significant, the sediments are collected at the grain boundary, and the nature of the sediments is coarse and soft. The effect of strengthening and hardening provided by magnesium is through a solid solution strengthening process carried out by substitution; namely, the Mg atom replaces the position of the Al atom [37].

The substitution process initiates a distortion in the metal crystal lattice, generating a stress field surrounding the soluble atoms. The presence of dislocations with stress fields hinders the movement speed, mainly when traversing through the solid-soluble atoms. Consequently, an elevated stress level becomes necessary as the metal exhibits enhanced strength in the company of the solid-soluble atoms. The reinforcement and toughening of AlMg alloys are achievable through a strain-hardening process attained via deformation [37]. Nevertheless, the deformation process introduces irregularities in a lattice, leading to lattice defects manifested as atomic vacancies. These conditions culminate in alterations in mechanical and thermal properties, including heat conductivity.

The decline in heat conductivity is linked to the decrease in electron velocity, which leads to the hindrance or restriction of heat transfer in the presence of atomic vacancies. Heat rolling occurs at temperatures exceeding the corresponding metal’s recrystallization threshold. The metal’s recrystallization point is significantly affected by its level of deformation – a higher degree of deformation results in a lower recrystallization temperature. Deformation exceeding the recrystallization temperature triggers softening phenomena involving recovery mechanisms like recrystallization and grain growth [38].

Several factors, such as the type of metal, temperature, and deformation or strain rate, influence the softening level. When the metal exhibits low stacking fault energy at elevated temperatures, it tends to undergo recrystallization throughout the deformation process. This event is referred to as dynamic recrystallization. Conversely, if recrystallization takes place post-deformation, it is called static recrystallization. Dynamic recrystallization is triggered by the restricted engagement of the recovery mechanism in metals with low stacking fault energy during the softening process and requires a significant amount of driving energy for the commencement of recrystallization [38]. However, aluminum and its alloys demonstrate a high stacking fault energy, which leads to the maintenance of elongated grain shapes even during deformation at elevated temperatures, thereby inhibiting recrystallization. Despite this, the properties of aluminum alloys remain relatively soft because of the slight increase in hardness. This is primarily attributed to the dominant role of softening through dynamic recovery, resulting in low driving energy that is inadequate for prompting dynamic recrystallization [29]. Specifically, aluminum boasts a stacking fault energy of approximately 200 erg/cm², whereas 304 stainless steel registers at 20 erg/cm² [39].

Figure 9 shows the phase diagram of the Al–Mg binary system. The eutectic point of the Al–Mg alloy, as illustrated in Figure 9, is positioned at a compositional level of 35% Mg, corresponding to a temperature of 415°C. The aluminum-rich solid solution encompasses 14.9% Mg at this critical eutectic temperature. Furthermore, with a decline in temperature, the magnesium content within the alloy also diminishes to 11.8% at 400°C, 6% at 300°C, 4% at 200°C, and less than 2% at 100°C. This solubility decrease is followed by removing Mg₂Al₃ from the solid (Al) solution. The formation of Mg₂Al₃ and Mg₅Al₈ phases occurs when the magnesium content exceeds 2% in the alloy [40], exhibiting brittleness and roughness below 327°C yet demonstrating plastic properties at elevated temperatures [40]. These phases tend to precipitate at the grain boundary as an anodic phase, consequently facilitating the onset of intergranular corrosion and stress corrosion cracking in a corrosive milieu [41].

Figure 9

Al–Mg phase equilibrium diagram [42].

MAGNOX reactor is a nuclear reactor developed for power generation, which uses uranium fuel and carbon dioxide as a coolant. This reactor is designed to generate electrical energy through the process of nuclear fission. The MAGNOX reactor uses uranium metal as fuel. Uranium is processed into fuel rods that are coated with a special cladding material. This cladding is made of a mixture of magnesium and oxygen (MgO). The coating serves to protect the uranium fuel from direct reaction with the coolant gas (carbon dioxide) and prevents corrosion of the fuel that may occur due to high temperatures. Magnesium cladding also helps to reduce damage to the fuel structure caused by neutron radiation and high heat during the nuclear fission process. Carbon dioxide (CO₂) gas is used as the coolant in the MAGNOX reactor. The gas flows through the reactor core to carry the heat generated by the fission reaction, which is then pumped out to produce steam that drives the electrical power turbine. To maintain a controlled fission reaction, the MAGNOX reactor uses control rods made from neutron-absorbing materials (such as boron or cadmium). These control rods can be inserted or removed from the reactor core to regulate the rate of the nuclear reaction [43,44].

3.3 Al 6061 alloys

In contrast to AlMg, Al 6061 alloy is one of the 6xxx series AlMgSi alloys that can be hardened by heat treatment (heat-treatable alloy) [45,46]. These AlMgSi alloys exhibit favorable mechanical, thermal, and corrosion resistance characteristics [47,48]. Moreover, they have been widely used in the cladding of nuclear materials in plate-type research reactors, as illustrated in Figure 10. For example, the Multipurpose Reactor, GA Siwabesi (RSG-GAS), uses aluminum alloys AlMg and Al 6061 for structural material components with the following specifications. The fuel elements consist of 21 fuel plates assembled with end fittings and side plates. The fuel plates are part of the cladding and meat (uranium fuel). AlMgSi alloys must undergo various processing stages when used as nuclear material cladding, as depicted in Figure 11.

Figure 10

Diagrammatic representation of nuclear fuel components [49].

Figure 11

The schematic illustration delineates the progression of nuclear fuel plate manufacturing utilizing the picture frame technique (PFT) [49].

The aluminum 6061 alloy consists of magnesium and silicon as its primary alloying constituents. Incorporating these elements into aluminum can lead to solid solution strengthening or second-phase metal compounds, thereby enhancing the strength and hardness of AlMgSi alloys. The effectiveness of strengthening AlMgSi alloys through the secondary phase can be optimized by aiming to generate delicate, uniformly dispersed particles of the AlMgSi second phase. This strengthening method is called precipitation hardening [49,50].

Figure 12 shows the changes in the mechanical properties of material hardness in line with the stages of the fabrication process for Al 6061 alloy plates. The initial availability of Al 6061 alloy plates includes 6061-T6, a product of the aging process, and 6061-O, a product of the annealing process. Aluminum 6061-T6 alloy plates are designated for side plates and end fittings, while aluminum 6061-O alloy is selected for cladding in nuclear fuel enclosures. Al 6061-O alloy plates undergo various fabrication stages involving thermomechanical deformation and heat treatment processes until the final phase. The progression of fabrication stages applied to Al 6061-O alloy until the completion of the procedure will impact the alterations in the hardness properties of the material. The fabrication process (PFT) leads to an enhancement in the hardness of the Al 6061 alloy in its final phase compared to the annealed Al 6061-O state. This phenomenon is attributed to the impact of cold rolling, which induces strain hardening.

Figure 12

Relationship between microhardness and PFT fabrication stages for the Al 6061 alloy [49].

On the other hand, the Al 606-clad alloy plate showcases the lowest level of hardness characteristics. This condition is due to the two-step hot rolling deformation treatment implemented on the Al 6061-clad plates. The hot rolling process is executed at temperatures exceeding the metal’s recrystallization temperature. Consequently, recrystallization takes place alongside a softening mechanism that involves recovery processes, such as recrystallization and grain growth. In contrast, the Al 6061-insp alloy samples undergo an initial heating phase at 440°C for 15 min, followed by deformation through 7–9 stages of hot rolling, with a higher deformation level than the Al 6061-clad alloy specimens. As a result, the hardening effect during hot rolling can be attributed to an increase in dislocation density. The dislocation area is the site for phase β precipitation, impacting the enhancement of the hardness characteristics of the Al 6061-insp alloy.

In addition, the Al 6061-insp alloy is subjected to a subsequent heat treatment for blister testing at 440°C for 60 min. The outcomes of the tests revealed an augmentation in the hardness properties of the Al 6061-insp alloy, slightly surpassing those of the Al 6061-clad alloy. Nevertheless, compared to the final Al 6061 alloy PFT, the hardness properties of the Al 6061-insp alloy are inferior. This discrepancy arises since the final Al 6061 alloy undergoes a two-step cold working deformation process, resulting in reinforcement and strain hardening.

Figure 13 presents the changes in microstructural properties of the Al 6061 alloy (Al–Mg–Si), resulting from a thermomechanical treatment. The findings from the tempering process of the Al 6061-T6 alloy, as illustrated in Figure 13(a), demonstrate a reduction in grain size compared to the other test specimens due to the consecutive implementation of heat treatment procedures involving heating, quenching, and aging at temperatures ranging from 175 to 180°C for periods of 10–20 h. On the other hand, Figure 13(b) explicitly shows the increased flexibility and larger grain size of the annealed Al 6061-O alloy compared to the Al 6061-T6 variant. Furthermore, Figure 13(c) shows the Al 6061-clad alloy sample with significantly larger grains than those observed in the Al 6061-T6 and Al 6061-O samples. This condition can be understood because the 6061-clad alloy goes through a two-step hot rolling deformation process conducted above the material’s recrystallization temperature. Consequently, this process induces softening mechanisms involving recovery, recrystallization, and grain growth. Figure 13(d) shows that the Al 6061-inspected alloy presents the most extensive microstructure grains, ascribed to the initial heating at 440°C for 15 min followed by 7–9 stages of hot rolling. Subsequent reheating at 440°C for 60 min led to significant grain growth. Conversely, the Al 6061-end PFT alloy specimens exhibit a smaller microstructure grain size than the Al 6061-insp alloy samples, as illustrated in Figure 13(e). This result stems from the cold rolling treatment involving 1–2 stages post-inspection. Consequently, there is an increase in dislocation frequency, thereby augmenting the grain boundaries and refining the grains.

Figure 13

Microstructural characteristics of alloy 6061 are depicted in optical micrographs, showcasing the grain morphology: (a) Al 606-T6 tempered, (b) Al 6061-O annealed, (c) Al 6061-clad), (d) Al 6061-Insp, and (e) Al 6061-final PFT [49].

Figure 14 illustrates the DSC curve of the Al 6061-T6 alloy, which shows the correlation between microhardness and the precipitate phase. The appearance of exothermic peaks in the DSC thermogram is closely related to the components present in the solid solution. It subsequently results in the formation of precipitates as the heating process progresses following the precipitation sequence of the Al–Mg–Si alloy. The DSC graphs demonstrate that the specimen undergoes a phase transition at relatively modest temperatures ranging from 300 to 400°C. This indicates that the thermal stability of the Al 6061 alloy is greatly affected by the creation and alteration of the second phase in the alloy.

Figure 14

DSC thermograms of 6061 alloy under different conditions: black line for Al 6061-T6 alloy, red line for Al 6061-O alloy, green line for Al 6061-clad alloy, blue line for Al 6061-Insp alloy, and cyan line for Al 6061-end PFT alloy [49].

The DSC thermograms of the Al 6061-T6 alloy condition demonstrate the dissolution (decomposition) peak of the precipitation phase within the temperature interval of 300–350°C. This occurrence can significantly impact the alloy’s characteristics, particularly its mechanical and thermal attributes. In the cases of the Al 6061-O and Al 6061-clad alloys, it is observed that the exothermic peak materializes within the temperature range of 350–400°C. On the other hand, the Al 6061-insp alloy specimens exhibit minimal phase alteration. It is characterized by an exothermic reaction at 400°C with minimal heat energy emission. As for the Al 6061-final alloy specimens, an exothermic peak is evident at a lower temperature spanning from 250 to 300°C, denoting the occurrence of the second phase transition.

AlMgSi alloys typically exhibit two distinct phases, namely α and β phases [49,51]. The binary system phase diagram of Al–Mg₂Si alloy, as illustrated in Figure 15, displays these phases. Upon reaching or surpassing the solvus line, the β (Mg₂Si) phase is known to dissolve entirely into the α (Al) phase, as depicted in Figure 15. This phenomenon arises due to the decomposition of the Mg₂Si compound during the heating process within the β phase, causing Mg and Si atoms to solidly dissolve into the α phase, thereby occupying substitutional or interstitial positions within the Al matrix crystal lattice (α).

Figure 15

Phase diagram of Al–Mg₂Si [51,52].

The solid solubility limit of the Mg₂Si metal compound in aluminum is 1.85% at 595°C. However, the solubility of the Mg₂Si compound tends to decrease as the temperature decreases, leading to a reduction in the soluble Mg₂Si content to 1.48% at 550°C, 1.08% at 500°C, and 0.51% at 400°C. This decline in solubility is coupled with the expulsion of the Mg₂Si intermetallic compound from the Al solid solution. According to Mondolfo, the formation of the Mg₂Si phase occurs when the Mg content exceeds 0.3% and the Si content surpasses 0.2%. The binary phase diagram further indicates that the eutectic temperature of this alloy is situated at 595°C.

In contrast to the Al–Mg₂Si phase diagram, the solid solubility limit of the Mg₂Si metal compound depicted in the ternary balance diagram is 1.84% at 559°C in aluminum, as illustrated in Figure 16. The ternary balance diagram further demonstrates that the Mg₂Si content experiences a decline as the temperature decreases. The solid soluble Mg₂Si content reduces to 1.70% at 550°C, 1.14% at 500°C, 0.52% at 400°C, and 0.09% at 200°C. This reduction in solubility is concomitant with the expulsion of the inter-metallic compound Mg₂Si from the aluminum solid solution.

Figure 16

Ternary phase diagram of the AlMgSi alloy [51,52].

3.4 AlFeNi alloy

AlFeNi alloy is an alternative aluminum alloy used for high-density fuel cladding. AlFeNi alloys were developed to replace aluminum alloys such as AlMg and Al 6061. The development of cladding structure materials by integrating AlFeNi alloys is focused on producing metal alloys with heightened strength features to counteract the hardness attributes of high-density fuels. Research indicates that the AlFeNi aluminum alloy exhibits improved mechanical properties and corrosion resistance [53]. Al–Fe–Ni alloys demonstrate good heat stability and are extensively employed in various industrial sectors as potential alternatives to Al alloys with high heat resistance and exceptional corrosion resistance at elevated temperatures [54–56].

A polytropic structure in the AlFeNi metal alloy can enhance the material’s mechanical properties, particularly its strength and toughness. On the other hand, the resistance of corrosion properties is influenced by the formation of phase structures. The presence of phases in AlFeNi alloys is notably affected by the alloy’s constituent elements and their concentrations. Common phases observed in Al–Fe–Ni alloys comprise Al₃Fe, Al₃Ni, and Al₉FeNi [57,58]. Both the phase structure and microstructure contribute to alterations in material properties, particularly thermal and mechanical characteristics. The anisotropic structure of AlFeNi alloys, characterized by monoclinic and orthorhombic structures, remains stable at 377°C [59]. The stability of the phase structure is significantly dependent on the quantity of solid-soluble elements in the alloy, with higher levels of these elements leading to a slowdown in the diffusion transformation process.

The formation of a phase in an alloy may arise when its composition consists of two or more elements with differing atomic radii, forming a solid solution as one of the phases. Furthermore, the resulting phases must exhibit distinct characteristics such as lattice dimensions, crystal structures, and melting points. Within the AlFeNi alloy, each constituent displays unique atomic dimensions, interatomic spacing, and crystal structure morphology. Al metal exhibits an FCC crystal structure characterized by lattice parameters of 4.0496 Å and an interatomic distance of 2.8635 Å. On the other hand, Fe metal displays a body-centered cubic (BCC) crystal structure with lattice parameters measuring 2.8664 Å and an interatomic distance of 2.4823 Å. Similarly, Ni metal shows an FCC unit with a lattice size of 3.52338 Å and an interatomic distance of 2.4919 Å [59].

The formation and transformation of phases are significantly influenced by the composition of the alloy and the prevailing temperature conditions. The interplay of phase reactions in the alloy substantially impacts the overall process. Figure 17 illustrates the development of phases and phase transitions in Al–Fe alloys according to the binary system phase diagram. At a temperature of 652°C and with a Fe content of 1.8%, the eutectic phase reaction between aluminum and iron commences, forming a solid phase known as α + θ, specifically Al + FeAl₃. The solid solubility limit of Fe in the α (Al) phase reaches a maximum of 0.04% Fe at 652°C. In the composition range of 0.04–37 wt% Fe below 652°C, the α + θ phase commences its formation. This phase transition from the fusion of Al and Fe into the α + θ phase is driven by the eutectic phase reaction L → α + θ.

Figure 17

Phase diagram of the Al–Fe binary system [60].

Figure 18 illustrates the phase equilibrium diagram of the Al–Fe–Ni ternary system. The initiation of phase τ (Al₉FeNi) formation occurs at a temperature of 640°C, as indicated in the phase equilibrium diagram of the ternary system. This situation is caused by the high contents of Ni and Fe present in the alloy. The X-ray diffraction pattern reveals the existence of two distinct phases, Al and Al₉FeNi phases, as depicted in Figure 19. The microstructure pattern demonstrates that the Al₉FeNi phase commences at 1% Fe and 1% Ni, exhibiting further growth at 2% Fe and 1% Ni. Additionally, the microstructure of the Al₉FeNi phase originates at the grain boundary, typically displaying an elongated morphology resembling a needle.

Figure 18

Phase diagram of the Al–Fe–Ni ternary system [60].

Figure 19

(a) XRD patterns of as-cast alloy samples, powders, and SLM samples; (b) XRD patterns with slow scanning speed (0.1°/min) of SLM (selective laser melting) samples [61].

Figure 20 shows Al–Fe–Ni alloys with varying composition levels, specifically Al–0.7Fe–0.5Ni, Al–1.4Fe–1.0Ni, Al–1.75Fe–1.25Ni, Al–2.1Fe–1.5Ni, Al–2.8Fe–2.0Ni, and Al–3.5Fe–2.5Ni [61]. The structural characteristics of Al–0.7Fe–0.5Ni and Al–1.4Fe–1.0Ni alloys are explained in Figure 20(a) and (b), emphasizing the α-Al primary phase and Al₉FeNi eutectic phase. The microstructure of the Al–1.75Fe–1.25Ni alloy, as presented in Figure 20(c), exhibits a fully eutectic composition. On the other hand, Figure 20 d–f shows that Al–2.1Fe–1.0Ni, Al–2.8Fe–2.0Ni, and Al–3.5Fe–2.5Ni alloys exhibit the Al₉FeNi phase in their microstructure. The presence of the Al₉FeNi phase is intensified with higher levels of Fe and Ni, resulting in an elongated needle-shaped alloy. Aluminum alloys containing the Al₉FeNi phase demonstrate enhanced heat resistance and strength at elevated temperatures [62,63].

Figure 20

OM images of (a) Al–0.7Fe–0.5Ni, (b) Al–1.4Fe–1.0Ni, (c) Al–1.75Fe–1.25Ni, (d) Al–2.1Fe–1.0Ni, (e) Al–2.8Fe–2.0Ni alloy, and (f) Al–3.5Fe–2.5Ni [61].

The Al₃Ni phase compounds are observed in the reaction phase involving Al and Ni when Ni contents are maintained at a low level. In the case of the Ni content surpassing the solid solubility threshold beyond 0.04% in the mixture, the formation of the κ (Al₃Ni) phase becomes achievable. The emergence of the κ phase initiates within the 0.04–42 wt% Ni composition range at temperatures below 640°C. This phase results from Al and Ni combining after the eutectic phase reaction, denoted as L → α + κ. The dimension of the κ phase is significantly influenced by the Ni content. A higher Ni content corresponds to a greater quantity of the κ phase within the alloy. Furthermore, the reaction phase involving Fe and Ni can lead to the development of Ni₃Fe at lower temperatures, commencing from 345°C.

Al–Fe–Ni alloy in the solid state will form several types of solid phases, including θ, κ, and τ phases. These phases exhibit distinct crystal structures [64]. The θ phase (Al₃Fe) possesses a monoclinic crystal structure characterized by lattice parameters a = 15.489 Å, b = 8.0831 Å, c = 12.476 Å, and an angle of 107.720. The FeAl₃ phase demonstrates exceptional high-temperature stability and remarkable wear resistance. Moreover, the composition’s influence plays a crucial role in enhancing the mechanical properties. The κ phase (Al₃Ni) displays an orthorhombic crystal structure with lattice parameters a = 6.1114 Å, b = 7.3662 Å, and c = 4,8112 Å. On the other hand, the τ phase (Al₉FeNi) showcases a monoclinic crystal structure with lattice parameters a = 8.598 Å, b = 6.271 Å, c = 6.207 Å, and an angle of 94.66°. These solid phases of the AlFeNi alloy each exhibit liquidus temperatures of ≤11,570°C for the θ phase, <854°C for the κ phase, and <1,133°C for the τ phase.

4.1 Mechanical and thermal properties of AlFeNi alloys

Figure 21 shows the changes in the mechanical properties of the AlFeNi alloy in response to variations in its composition. The initial strength is 87.78 MPa and shifts alongside the increase in Ni content. Conversely, the elongation of the alloy diminishes gradually as the Ni content increases. At 6% Ni, the maximum strength of 148.22 MPa is achieved before showing a gradual reduction. Introducing 6% Ni prompts an increase in the precipitation of a finer-sized second phase, consequently enhancing the alloy’s strength. Nevertheless, exceeding 6% Ni leads to a decrease in strength due to the saturation of second-phase precipitations, causing dispersion and cracking. This phenomenon also contributes to a reduction in elongation. Furthermore, the eutectic structure of Al₃Fe, combined with α-Al grains, creates pathways for microcrack propagation when the applied load surpasses a specific threshold [65]. Consequently, the elongation of Al–1Fe–xNi alloy tends to decrease.

Figure 21

Mechanical properties of Al–1Fe–xNi [66].

The change in Ni content within the alloy also affects the microhardness measurement of the alloy. Observations indicated an increase in the microhardness measurement from 161.1 to 198.5 HV as the Ni content increased from 4 to 6 wt%. Conversely, exceeding 6 wt% Ni led to a notable increase in the microhardness to approximately 509.4 ± 44 HV. Moreover, the introduction of Ni instigated a transformation from the Al₃Fe phase to the Al₉FeNi phase and facilitated the refinement of microstructural grains within the alloy.

Figure 22 illustrates the decline in electrical conductivity of Al–1Fe–xNi alloy attributed to the introduction of Ni. Moreover, lattice imperfections like vacancies, dislocations, and grain boundaries can impede the movement of electrons, reducing the alloy’s electrical conductivity [67]. The process of diminishing the electrical conductivity of the alloy is depicted schematically in Figure 23. The movement of free electrons encounters minimal resistance when traversing the crystal grain boundaries, resulting in a decrease in the number of free electrons that successfully pass through. A factor hindering electron mobility is the expansion of the second phase, which enlarges due to the crystal structure in the α-Al phase and fluctuations in the number of grain boundaries. The decrease in the number of free electrons will inevitably affect the decrease in electrical conductivity.

Figure 22

Electrical conductivity of the Al–1Fe–xNi alloy [66].

Figure 23

Illustration of the electron transport mechanism through grain boundaries and the existence of a secondary phase [66].

Besides changes in the mechanical and conductivity properties, alterations in thermal properties have also been recognized. The changes consist of density and heat capacity, which are associated with the increased Fe and Ni content in the aluminum alloy. The changes in thermal and electrical conductivity of the Al–x(1.75Fe–1.25Ni) alloy are shown in Figure 24. The graphical representation in Figure 24 illustrates a significant reduction in thermal and electrical conductivity from 206.5 to 65.23 W·m⁻¹·K⁻¹ and from 29.94 to 6.40 MS·m⁻¹, respectively [68,69]. One of the primary factors contributing to the decline in thermal properties is the emergence of phase precipitates. These phase precipitates act as barriers and impediments to the movement of electrons. The formation of phase precipitates is likely to occur in Al–Fe–Ni alloys with Fe and Ni concentrations exceeding 1%, leading to the formation of Al9FeNi precipitates.

Figure 24

Variations in the thermal and electrical conductivity of the Al–x(1.75Fe–1.25Ni) alloy [68].

4.2 Thermal stability (analysis of phase transitions)

Figure 25 illustrates the differential scanning calorimetry (DSC) data of the Al–1Fe–xNi alloy [66]. The optimization of microstructure holds significant importance in enhancing the mechanical characteristics of metal alloys. Among the factors influencing the solidification properties of metal alloys, the introduction of elements plays a crucial role. The solidification mechanism of metal alloys directly impacts the development of the alloy’s microstructure. The rapid cooling rate observed during the solidification process of the alloy leads to inadequate diffusion time between the liquid and solid phases, causing an imbalanced progression. Consequently, a substantial accumulation of Fe atoms occurs gradually at the interface region connecting the primary α-Al grains and the liquid phase, resulting in localized surpassing of the maximum Fe concentration. The soluble Fe in Al undergoes a phase transformation to form the Al₃Fe phase. This phase exerts an inhibitory influence on the primary α-Al, thereby inducing grain refinement.

Figure 25

DSC curve of the Al–1Fe–xNi alloy [66].

At the same time, the inclusion of Ni will induce alterations in the quantity and structure of Al₃Fe, consequently leading to the creation of a subsequent phase, commonly referred to as the second phase. Upon examination of the DSC data illustrated in Figure 25, it is evident that the temperature at which the peak heat absorption occurs is 929.5 K (656.5°C) once the Ni concentration reaches 6%. Moreover, a distinct bulge appears in the heat absorption peak (indicative of an endothermic reaction) as the Ni content is 6 and 10%, thereby instigating a phase transition within the alloy. It has been established that the resultant phase arising from this reaction is identified as the Al₉FeNi phase, as depicted in Figure 26 [66].

Figure 26

XRD pattern of the Al–1Fe–xNi alloy [66].

4.3 Radiation shielding properties

Aluminum alloys play a significant role in the nuclear industry, particularly in neutron radiation shielding and as part of radiation protection structures. Although aluminum is not as effective as high-density materials like lead or concrete in blocking gamma or X-rays, aluminum offers several advantages that make it highly valuable in nuclear applications. One of the main benefits of aluminum alloys is their ability to block neutron radiation, especially thermal neutrons, which are produced in the nuclear fission process. Aluminum has a relatively good neutron absorption cross-section, and the addition of certain elements, such as boron (B) and lithium (Li), can enhance its ability to absorb thermal neutrons. Boron, for example, is known to be very effective in neutron absorption and is often used in compounds like boron carbide (B₄C) to improve neutron radiation shielding in aluminum alloys, making it a popular choice for applications that require neutron radiation protection in nuclear reactors.

However, although aluminum is effective against neutron radiation, its ability to block gamma and X-ray radiation is more limited. Aluminum can reduce the intensity of gamma radiation through Compton scattering and photoelectric absorption, but this reduction is much smaller compared to denser materials. Therefore, aluminum is typically used in nuclear applications as part of a complementary shielding system, where it works alongside denser materials such as lead or concrete, which are more effective at blocking gamma radiation. For example, aluminum alloys are often used as outer layers or structural components in multi-layer radiation shields to add flexibility and reduce the overall weight of the radiation shielding system [70].

Additionally, aluminum alloys have the advantage of being more resistant to radiation-induced damage. In nuclear environments exposed to radiation for extended periods, certain materials may undergo structural degradation, such as changes in the crystal structure or mechanical properties due to ionizing radiation and neutrons. Aluminum alloys, although exposed to radiation, are generally more resistant to such degradation compared to many other materials, and with the right alloy selection, they can maintain their structural integrity. Some aluminum alloys containing elements such as magnesium (Mg) and copper (Cu) have been developed to improve resistance to radiation damage and strengthen their mechanical properties, making them more suitable for long-term use in nuclear reactors.

In addition to its radiation shielding properties, aluminum also offers practical benefits such as lightweight, ease of fabrication, and resistance to corrosion, all of which are essential in harsh nuclear environments. The lightweight nature of aluminum alloys allows for more efficient structural designs and reduces the load on other components of the nuclear system. The ease of fabrication also facilitates the production of complex shielding components, while its resistance to corrosion ensures the material’s longevity even in radiation-exposed and high-temperature environments.

Aluminum alloys are used in various components within nuclear reactors, including neutron shields, fuel elements, and other structural parts. In some types of nuclear reactors, such as light water reactors (LWRs), aluminum functions in fuel elements and thermal management systems. Aluminum alloys are also used in protective shields to protect other components from radiation exposure. While aluminum is not the primary choice for gamma radiation shielding in environments with high radiation exposure, it remains an important material in the design of complex radiation shielding systems, where efficiency and weight reduction are highly valued.

Overall, although aluminum alloys cannot replace materials with higher radiation shielding capabilities, their advantages in neutron radiation resistance, lightweight, and ease of fabrication make them essential materials in the nuclear industry. The right aluminum alloy can provide adequate radiation protection while maintaining its advantages in design and structural efficiency in nuclear applications that require radiation shielding, whether for reactor components or nuclear waste management [71].