## 1 Introduction  
Uranium metal is used in a variety of applications because of its high density and unique nuclear properties. In pure uranium metal, the body-centered-cubic (BCC) allotrope, \(\gamma\), is only stable at elevated temperatures (776 \({}^{\circ}\)C \(<\) T \(<\) 1135 \({}^{\circ}\)C) [1]. On cooling, uranium undergoes a solid-solid phase transformation to the tetragonal \(\beta\) phase (observed in a very small interval of temperatures and pressures), followed by a second transformation to the orthorhombic \(\alpha\) phase, which is stable at room temperature and below. Among these phases, the \(\gamma\) phase has the best technical properties for nuclear engineering purposes, due to its cubic symmetry and elastic properties. However, \(\gamma\)-uranium is known to be stabilized at high temperatures by phonon-phonon coupling that is strongly anharmonic [2]. Consequently, this phase is decidedly unstable at low temperatures (T \(<\) 776 \({}^{\circ}\)C) and quenching cannot suppress spontaneous transformation to non-cubic crystal structures [3; 4; 5]. Uranium metal at low temperatures, where machining and other material processing may take place, exists in the \(\alpha\)-uranium structure, which is less than ideal due to its brittle nature and tendency to form uranium-oxides. The poor corrosion resistance and undesirable combinations of strength and ductility observed in this structure have spurred numerous investigations into methods to destabilize this phase by introducing one or more alloying constituents [3]. Most known alloying elements dissolve to only a very limited degree in the \(\alpha\) and \(\beta\) forms of uranium, while, in \(\gamma\)-uranium, more extensive solid solubility can be achieved [6]. For most applications, mechanical isotropy is a desirable property and much work has been focused on broadening the range of \(\gamma\)-phase stability by alloying. To achieve stability at room temperature, attempts have been made to alloy uranium with other \(\gamma\) soluble elemental metals, such as Mo [7; 8], Nb [9], Zr [10], and Ti [11; 12]. These elements have minimal solubility in \(\beta\) or \(\alpha\) uranium [7]; however, they can form _metastable_ alloys by rapidly quenching from the high-temperature \(\gamma\) phase. Among this group of \(\gamma\) soluble metals, Nb is somewhat unique because it does not form an intermetallic compound with uranium, although the possibility of the \(\alpha^{\prime}\) metastable structure as an \"ordered\" phase has been addressed [13]. The addition of niobium to uranium is particularly beneficial, as it significantly increases oxidation resistance [14]. Figure 1 shows the equilibrium U-Nb phase diagram [15; 16; 17; 18; 19]. There is complete solubility across the equilibrium U-Nb phase diagram in the high-temperature BCC (\(\gamma\)) phase [18]. The uranium-niobium phase diagram displays a \(\gamma\)-phase miscibility gap at high temperatures with a critical temperature, \(T_{c}=950\pm 20\)\({}^{\circ}\)C, and a critical concentration, \(c_{c}=52.3\) at.% Nb. At 647 \({}^{\circ}\)C an equilibrium monotectoid reaction is exhibited according to the formula \(\gamma_{1}(\text{BCC})\rightarrow\gamma_{2}(\text{BCC})+\alpha(\text{orthorhombic})\). The \(\gamma_{1}\) and \(\gamma_{2}\) phases contain 13.3 and 70 \(\pm\) 2 at.% Nb, respectively, whereas the \(\alpha\) phase has less than 1 at.% Nb in solution. Upon slow cooling under equilibrium conditions, the \(\gamma_{1}\) phase undergoes diffusional decomposition at 647 \({}^{\circ}\)C in order to form a two-phase structure consisting of unalloyed \(\alpha\)-uranium and niobium-rich \(\gamma\)-phase (\(\gamma_{2}\)). At this composition, microstructures tend to consist of discrete niobium-rich \(\gamma\)-phase (\(\gamma_{2}\)) with a matrix of unalloyed \(\alpha\)-uranium. This structure offers poor corrosion resistance due to the lack of niobium in solid solution, and also the two-phase structure gives rise to microscopic anodic and cathodic regions [3].