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Materials Used in the Production of Small Weapons
Published in Jose Martin Herrera Ramirez, Luis Adrian Zuñiga Aviles, Designing Small Weapons, 2022
Jose Martin Herrera Ramirez, Luis Adrian Zuñiga Aviles
Phase diagrams provide valuable information on which phases are thermodynamically stable in an alloy and whether they could be present over a long time when the component is subjected to a particular temperature, for example, in a gun barrel made of steel. Other information obtained from phase diagrams is about melting, solidification, solubility limits, the presence of solid solutions, and phase transformations. These diagrams are useful to metallurgists, materials engineers, and materials scientists in different areas [4]: Development of new alloys for specific applicationsFabrication of alloys into useful configurationsDesign and control of heat treatment procedures for specific alloys to produce the required propertiesSolving performance problems with specific alloys in commercial applications, thus improving product predictability.
High Entropy Alloys
Published in T.S. Srivatsan, Manoj Gupta, High Entropy Alloys, 2020
Gaurav Kumar Bansal, Avanish Kumar Chandan, Gopi Kishor Mandal, Vikas Chandra Srivastava
From the fundamental of phase transformation, a negative change in Gibbs free energy (ΔG = GProduct − Gparent = driving force) of parent and product phases results in the formation of thermodynamically stable structures, which grow in the existing phase until the whole of the parent phase is transformed. This driving force, i.e., negative change in Gibbs free energy, is necessary for the spontaneous phase transformation. However, it is not a sufficient requirement as the progress of transformation is generally decided by the kinetic factors also. The Gibbs free energy change of parent and product phases includes the contribution from change in enthalpy as well as entropy. In multicomponent high entropy alloys, an accurate estimation of ΔG is extremely complicated due to the involvement of many variables, even at a fixed temperature and composition. For multicomponent systems, Gibbs free energy of mixing (ΔGmix) can be expressed as: ΔGmix=ΔHmix−TΔSmix
Virtual environments
Published in Paulo Jorge Bártolo, Artur Jorge Mateus, Fernando da Conceição Batista, Henrique Amorim Almeida, João Manuel Matias, Joel Correia Vasco, Jorge Brites Gaspar, Mário António Correia, Nuno Carpinteiro André, Nuno Fernandes Alves, Paulo Parente Novo, Pedro Gonçalves Martinho, Rui Adriano Carvalho, Virtual and Rapid Manufacturing, 2007
Paulo Jorge Bártolo, Artur Jorge Mateus, Fernando da Conceição Batista, Henrique Amorim Almeida, João Manuel Matias, Joel Correia Vasco, Jorge Brites Gaspar, Mário António Correia, Nuno Carpinteiro André, Nuno Fernandes Alves, Paulo Parente Novo, Pedro Gonçalves Martinho, Rui Adriano Carvalho
During LPD the deposited material is submitted to consecutive thermal cycles, as successive layers are deposited and heat flows from recently deposited material through the previously deposited layers into the substrate. These thermal cycles can induce solid-state phase transformations that affect considerably the properties of the material. The final microstructure is determined by the thermal history of the material, which varies from point to point and, as a consequence, the phase constitution and properties will present a complex distribution in the deposited part. To simulate the microstructure and properties distribution in the part, it is necessary to know both the influence of the deposition parameters on the thermal history at each point of the part and phase transformations that may occur as the temperature field changes. Although this can be a complex task to be undertaken experimentally, it is suited to a mathematical computational approach if a thermo-kinetic model is available coupling phase transformations kinetics with heat transfer calculations.
Influence of phase transformation on stress wave propagation in thin-walled tubes
Published in Waves in Random and Complex Media, 2023
When solid materials are subjected to strong impact loads—such as an explosion, thermal shock, or high-velocity impact—they may yield or even undergo phase transformation. Phase transformation can greatly affect the mechanical response of materials and structures because the transformed material can be regarded as another material. This is a common concern in the fields of materials science, solid mechanics, and industrial manufacturing [1]. Shape memory alloys (SMAs) are typical functional materials that have many excellent properties, such as the pseudo-elastic effect and shape memory effect. Owing to their unique thermal and mechanical properties, SMAs have received extensive engineering attention, especially in applications involving high precision and military purposes, including in the aerospace, automotive, biomedical, and oil-exploration industries [2]. With recent increases in their application in explosive and impact environments, the dynamic response and mechanical behavior of SMAs need to be further explored [3].
Exploring the origin of variant selection through martensite-austenite reconstruction
Published in Philosophical Magazine, 2019
Sushil K. Giri, A. Durgaprasad, K.V. Manikrishna, C.R. Anoop, S. Kundu, I. Samajdar
Austenite grains with a Σ3 boundary can produce six α-variants of similar orientation [28]. The number of α-variants is naturally restricted by the boundary OR. The nucleation process in any phase transformation, including invariant plane strain transformation, is influenced by changes in chemical free energy, strain energy, and interfacial energy [46, 51]. Both strain and interfacial energy increase the activation energy for nucleation. If nucleation happens from a common prior-austenite Σ3 boundary, the common α-variants would maintain OR with both austenite neighbours. This is expected [26, 28] to reduce the interfacial energy for nucleation. In addition, the energy of the interface between the common α-variants would also be very small: reducing the interfacial energy contribution even further. Interfacial energy, however, plays an important role in reconstructive phase transformation [47]. For displacive phase transformation, on the other hand, a stronger influence of the strain energy has often been stipulated [46]. More specifically, strain energy (∼600 J/m2) is much higher than the interfacial energy (∼0.03 J/m2). The former is thus expected to have a stronger influence on the nucleation of displacive phase transformation [46]. This explains, albeit qualitatively, the observation that all prior-austenite Σ3 boundaries were not associated with common α-variants. However, all common α-variants were related to the presence of such boundaries (Figure 9(d)) plus minimisation of the transformation strain (Figure 9(e)). Both Germain et al. [28] and Abbasi et al. [26] related presence of Σ3 boundary, in the prior austenite grain, to the selection of common α-variants. Abbasi et al. [26] explained this in terms of minimisation of the interfacial energy. The present reconstruction and estimation of transformation strains provided a different rationale. As shown in Figure 9(c) and Table 2, any two common α-variants inside one prior-austenite grain provided a transformation strain of 0.1412. However, if these α-variants existed between the two austenite neighbours the transformation strain reduced significantly to 0.0286–0.0276. Finally, when all four α-variants were considered between the two prior-austenite grains: the lowest transformation strain of only 0.0181 was obtained. The present manuscript thus provided clear evidence (see Figure 9(e)) of the selection of common α-variants through a combination of prior-austenite Σ3 boundary plus self-accommodation in the martensitic structure.
Predicting the cooperative effect of Mn–Si and Mn–Mo on the incomplete bainite formation in quaternary Fe–C alloys
Published in Philosophical Magazine Letters, 2018
Hussein Farahani, Wei Xu, Sybrand van der Zwaag
The GEB model as developed by Chen et al. [16] is based on two fundamental components. The first component is the chemical driving force of the phase transformation, which is taken to be the Gibbs energy change for the transformation. The chemical driving force due to compositional and thermodynamic differences between the austenite and bainitic ferrite phases can be generally calculated usingwhere is the chemical driving force per mole atom, i is the element in the alloy, n is the total number of elements in the alloy, is the composition of material transferred over the interface, and are the chemical potential of the element i ferrite and austenite, and is the mole fraction of element i in ferritic side of interface and is the mole fraction at austenitic sides of interface [17]. For substitutional alloying elements, and are chosen according to the Negligible Partitioning Local Equilibrium (NPLE) thermodynamic model, but for carbon, being a fast diffusing interstitial alloying element, is assumed to be equal to the equilibrium concentrations of carbon in ferrite. is calculated from the Zener–Hillert equation [18]:where is the average carbon concentration in austenite, v is the velocity of the migrating austenite/ferritic bainite interface, is the diffusion coefficient of carbon in austenite and T is the temperature. The equilibrium carbon concentration in bainitic ferrite compared to carbon content of austenite is assumed to be negligible. As the carbon diffusion in austenite is much faster than of the substitutional alloying elements, the carbon content in the remaining austenite can be calculated using the mean-field approximation leading towhere is the fraction of bainitic ferrite and is the initial carbon concentration in austenite. Combining Equations (2) and (3), can be calculated as a function and migration rate of austenite/bainitic ferrite interface.