- Email: [email protected]
The microstructure and hydriding characteristics of high temperature aged U–13 at.%Nb alloy
Journal of Nuclear Materials 464 (2015) 43–47
Contents lists available at ScienceDirect
Journal of Nuclear Materials journal homepage: www.elsevier.com/locate/jnucmat
The microstructure and hydriding characteristics of high temperature aged U–13 at.%Nb alloy Hefei Ji a, Peng Shi a, Ruiwen Li a, Chunli Jiang b, Jiangrong Yang b,⇑, Guichao Hu a a b
China Academy of Engineering Physics, Mianyang 621900, China Science and Technology on Surface Physics and Chemistry Laboratory, P.O. Box 718-35, Mianyang 621907, China
a r t i c l e
i n f o
Article history: Received 17 October 2014 Accepted 29 March 2015 Available online 3 April 2015
a b s t r a c t Niobium as alloying element significantly improves physical and chemical properties of metallic uranium, exhibiting great application potential in uranium alloy materials. The corrosion resistance performance as well as the internal alloy phase structure of uranium–niobium alloy is closely related to aging processes. Microstructure and hydriding characteristics of the 400 °C/9 h + 500 °C/2 h aged uranium–13 at.% niobium alloys (U–13 at.%Nb) were investigated from the point of view of relationship between the microstructure and growth of the hydriding areas. The microstructure, morphology and composition of the alloy phases before and after the hydriding were well characterized by the laser scanning confocal microscopy (LSCM), scanning electron microscopy (SEM) and X-ray diffraction (XRD), respectively. Experimental results indicated that the hydrogen preferentially reacted with the Nb-depleted phase alike-U to form monolithic b-UH3Nbx, and the alloy microstructure played an important role in hydride growth. Ó 2015 Elsevier B.V. All rights reserved.
1. Introduction Alloying additives including common metallic elements added to uranium can optimize the performance while maintaining the advantageous characteristics of this nuclear metal. Niobium is one kind of alloying metal added to strengthen the corrosion resistance of obtained alloy [1] which is often quenched to form metastable solid solutions to achieve desired properties and performances. It is worth noting that although the as-quenched U–Nb alloy possesses perfect oxidation and corrosion resistance, it does react with H2O and H2 under ambient conditions. Previous investigations demonstrated that the metal uranium and its alloys could readily react with hydrogen under particular pressure–temperature conditions [2–7]. The reaction product was extremely unstable in air and its pyrophoricity when exposed to atmosphere made it highly hazardous. So it is of great significance to study the corrosion resistance of uranium as well as its alloys. Numerous studies focused on the kinetics of hydriding reaction of uranium [4,8–11]. The nucleation rates of the corrosion pitting increased with ascending hydrogen pressure as well as reaction temperature [2]. Several reasons could be listed for the hydride formation, however, by now there still not exist a systemic nucleation ⇑ Corresponding author. E-mail address: [email protected] (J. Yang). http://dx.doi.org/10.1016/j.jnucmat.2015.03.051 0022-3115/Ó 2015 Elsevier B.V. All rights reserved.
mechanism explaining the location of the initial hydride nucleation and corresponding initial growth characteristics. In addition, some related researches were relatively less concentrated on the reaction of uranium alloys and hydrogen remained less studied than for pure uranium. Under equilibrium conditions, niobium has a complete solubility at high temperatures and can form bcc solid solution with uranium, while exhibiting a limited solubility at relatively lower temperature. The as-quenched uranium–niobium (U–Nb) alloy has a single phase a00 [12], of which the internal phase structure varies to form the niobium-rich phase and niobium-depleted phase after a long duration of natural aging or accelerating artificial aging, yielding the a + c2 phase eventually [13]. As U–Nb alloys are one kind of most widely used forms of uranium and exhibit different types of microstructures, we undertook U–13 at.%Nb alloy aged at 400 °C for 9 h followed by 500 °C for 2 h (marked as 400 °C/9 h + 500 °C/2 h), the bulk microstructure of the aged alloy were also well investigated.
2. Experimental details 2.1. Sample preparation All hydriding experiments investigating the nucleation and expansion of hydride were conducted in a hot stage microscope system (HSM) [14] equipped with sealed furnace chamber with a
44
H. Ji et al. / Journal of Nuclear Materials 464 (2015) 43–47
quartz glass window enabling in-situ monitoring of the reaction between U–Nb alloys and hydrogen using a Hirox KH 7700 digital microscope. A U–13 at.%Nb alloy sample in the form of 8 mm diameter rod was obtained using water quenching. Except for the niobium atoms, the total impurity content in the alloy is not more than 200 ppm, including carbon, the concentration of which turned out to be approximately 50 ppm. In order to investigate the relationship between the hydride and microstructure of the alloy, especially the ‘‘a + c12’’ microstructure, the alloy rod was then transferred to a pre-set condition, under which the alloy sample was processed by the aging treatment at temperature of 400 °C for 9 h followed by a successive aging treatment at 500 °C for 2 h. The rod was then cut into small discs with the diameter of 8 mm and thickness of 2 mm. X-ray diffraction (XRD), metallographic and scanning electron microscopy (SEM) analysis were performed on these samples for characterization of microstructure, composition and morphology. After a set of wet mechanical polishing treatment in the open laboratory, the samples were ultra-sonicated in acetone and ethanol successively to remove the oxidized surface and other impurities. Then the samples were loaded into the visual volumecalibrated stainless steel reaction chamber positioned within a tube furnace. Afterward the chamber was evacuated to the pressure of 105 bar using an oil-free scroll pump followed by keeping the system at constant pressure for 10 min to remove oxygen and other gas impurities to the greatest extent. 2.2. Hydriding As long as the pressure of 105 bar was achieved and held for 10 min, the chamber was heated to 160 °C to activate the sample under dynamic vacuum for 2 h. After the activation, the chamber was cooled down to the reaction temperature of 100 °C followed by the immediate injection of high purity H2. The hydriding experiments were operated under the initial hydrogen pressure of 0.1 bar. The surface appearance of the samples was monitored through the chamber window using a digital microscope. In this work, the hydrogen consumption quantified by the hydrogen pressure dropping to half the initial pressure value was thought to be sufficient. In the end, the reaction was terminated by evacuating hydrogen gas from the reaction chamber. After being cooled down to room temperature, the chamber was reaerated to oxidize the hydride followed by taking the samples out of the chamber. The hydride on the sample surface was then studied and investigated by a laser scanning confocal microscope (LSCM, Lext OLS-4000), a scanning electronic microscope (SEM, SIRION 200) with associated EDS and an X-ray diffractometer (XRD, Empyream). 3. Results 3.1. Microstructure Before the hydriding experiments, the microstructure and phase composition of the alloy sample was studied. Fig. 1 shows the metallographic morphology (Fig. 1(a)) and the higher magnification SEM BSE image for the 400 °C/9 h + 500 °C/2 h sample (Fig. 1(b)). Compared with the as-quenched and low-temperature aged sample [15], the surface morphology of the sample in this work is more coarse. In Fig. 1(b), the lamellar structures proved to be the parallel pearlite composed of the niobium-rich (Nb-rich) phase (the gray layers, eroded during the etching processing) and the Nb-depleted phase (the bright layers) [16] accordingly. The characteristic dimension of the fine lamellar pearlite structure is
approximately 50–100 nm, as seen from Fig. 1(b). The uranium concentration on the nanoscale cannot be quantified. Upon the heat treatment condition of 400 °C/9 h + 500 °C/2 h, the discontinuous coarsening happens along the prior-c boundaries [16] in the sample. Additional characterization using XRD (Cu Ka) demonstrated the composite character of the aged alloy, when compared to non-alloyed uranium. The results shown in Fig. 2 indicate that the phase composition should be c12 (Nb-rich) + a-uranium (depleted) for the aged alloy. However, there is a slight peak shift toward lower angles compared with pure uranium metal, implying a lattice parameter expansion. The lattice parameters of the analogous a-uranium in the aged alloy (a = 2.885 Å, b = 5.873 Å, c = 4.970 Å) are slightly larger than those observed in a-uranium (a = 2.853 Å, b = 5.872 Å, c = 4.952 Å). Isothermal aging at temperature between 300 °C and 650 °C causes phase separation generically depicted as the progressive evolution Nb-rich and Nb-depleted regions to the final a(0Nb) + c2(50Nb) phases obtained only under extreme conditions. Thus there still exists some residual niobium atoms dispersed in the Nb-depleted phase for the aged alloy in this work, expanding the lattice parameters of the Nb-depleted phase than that of a-uranium. So we define the Nb-depleted phase in the aged alloy sample as a-like-U instead of a-U. In that way, the composition of the aged alloy forementioned is optimized to be alike-U + c12.
3.2. Hydride growth During the hydriding process, the reaction time was defined as duration within which the hydrogen pressure decreased to 60% of its initial value, so as to acquire sufficient hydride areas. Fig. 3 shows the SEM characterization micrographs of the alloy surface after the hydriding reaction. A series of circular hydriding areas with different sizes formed on the surface of the alloy, indicating their different initial nucleation time. From the enlarged image, an expanded hydriding area can be clearly distinguished. However, we do not assume the hydrided areas to be hydride, as they do not exhibit the powder form and they have a composite character, which is different from previously reported results [17,18] indicating the reaction products of uranium and hydrogen to be b-UH3 powders. Then, do the circular areas signify a kind of new reaction product just like the similar reaction products (UH3Mo0.18) obtained from the reaction between the so-called UMo0.18 [19] and hydrogen? Or the product is not yielded by the reaction between uranium and hydrogen at all. The XRD results of the hydrided sample are shown in Fig. 4. The pattern reveals the peaks assigned to b-UH3, but they are shifted to lower diffraction angles, giving the lattice parameter a = 6.642 Å, indicating a small lattice expansion in comparison with b-UH3 (6.631 Å). The Nb atoms are smaller than uranium, niobium atoms dispersed in the reaction products, and in the hydride, there happened a charge transfer in hydrides, which may actually shrink the electropositive uranium atoms, while Nb is affected less, contributing to the lattice expansion. Compared with the similar result that a lattice expansion exist in a-like-U obtained before hydriding, the smaller atoms Nb leading lattice aberration is considered to be another factor. Here the reaction product, which we denote as bUH3Nbx, should be subjected to further studies for the reason that we cannot give the determined value of the niobium content. Here we can compare the situation with reaction between U–Ti alloys and hydrogen, illustrating that the hydride is nucleated preferentially in a-uranium and not in U2Ti [18]. So the reactants in this work are supposed to be a-like-U and hydrogen.
H. Ji et al. / Journal of Nuclear Materials 464 (2015) 43–47
45
Fig. 1. (a) LSCM bright-field micrograph of the 400 °C/9 h + 500 °C/2 h aged alloy sample and (b) corresponding SEM BSE micrograph.
Fig. 2. XRD results of the 400 °C/9 h + 500 °C/2 h aged alloy and uranium samples before hydriding experiment.
Fig. 3. SEM of the 400 °C/9 h + 500 °C/2 h aged alloy and with (inset) magnified SEM result for a hydride area on the surface.
Fig. 4. XRD result of the 400 °C/9 h + 500 °C/2 h aged alloy sample after hydriding experiment.
46
H. Ji et al. / Journal of Nuclear Materials 464 (2015) 43–47
4. Discussions In this work, we paid attention to the relation between the hydride growth and microstructure of the bulk U–13 at.%Nb alloy after the aging treatment. The 400 °C/9 h + 500 °C/2 h aged alloy samples were characterized and the results showed that the microstructure and phase composition of the aged alloy sample changed from fine twinned martensite [13,15] to bi-phase pearlite after the aging treatment under elevated temperature conditions. The bi-phase pearlite was composed of the Nb-rich (c12) phase and Nb-depleted phase (a-like-U), which alternately exhibited the parallel microstructure with the dimension ranging from 50 nm to 100 nm. The LSCM and SEM photographs show monolithic reaction products on the alloy surface, but some embossing reaction regions. The XRD results in Section 3.2 above demonstrated the formation of the hydride b-UH3Nbx after hydriding, and the density difference between the hydride (10.9 g/cm3) and the bulk alloy (17.5 g/cm3) led to a volume expansion and some subsequent outwards raised hydriding regions. The monolithic hydriding areas formed in the aged alloy are different from that in uranium and U–Ti alloy, but similar to those in U–Mo0.18 [19]. Our previous work also reported similar monolithic hydriding characteristics of the 400 °C/3 h aged U–13 at.%Nb alloy [15], which illustrated that the hydriding areas expanded along specific directions forming specific tree-like hydriding areas. The U–13 at.%Nb alloys with complex microstructure obtained by different aging treatment described in the present work have also a specific morphology, affecting the expansion dye to hydriding. It is well known that b-UH3 is sensitive to oxygen, still our hydrided product does not turn into powder. We can assume the Nb-rich areas, which are separated only by 50–100 nm of Nb-depleted phase, can maintain the integrity of material when the bUH3Nbx is formed in the Nb-depleted phase. Moreover, the lack of fragmentation prevents the b-UH3Nbx phase formed inside the bulk (see Fig. 5) to be exposed to atmosphere, and oxidation therefore does not proceed and is limited to the surface only. We assume that the phase boundary is a suitable diffusion path for hydrogen into the bulk.
The reaction products here cannot be defined as b-UH3, the XRD analysis results further confirmed it. During the hydriding reaction, hydrogen preferentially reacted with the a-like-U among the Nbrich phases. In addition, the quantity of the a-like-U is small for the reason that the distance of the two Nb-rich phases was just 50–100 nm. In this regard, the hydrogen only reacted with the little sized a-like-U to form a spot of b-UH3Nbx, which still existed among the Nb-rich phases with a little volume expansion. The bUH3Nbx quantity between every two Nb-rich phases was insufficient to make the material crack, and is too little to be observed at macro-scale. Hence the b-UH3Nbx was packed between the c12 phases, making it difficult to be observed and oxidized. The composition of the bulk alloy material contained more than the equable fine bi-phase, there also exist the block a-like-U or c12 parts. So the hydriding reaction arose in the block a-like-U parts contributed to the cracking of the bulk material. With the reaction in the bi-phase areas, the overall volume expansion accumulated from every individual little volume expansion, exhibiting the monolithic embossing hydriding areas on the sample surface. The reaction products b-UH3Nbx is concealed among the Nb-rich phases in the bulk material, so when the sample after hydriding was exposed to the ambient atmosphere, the b-UH3Nbx were not exposed to the atmosphere directly, then they still can be identified by the XRD diffractometer. Here the scheme of a hypothetical formation mechanism of the hydriding reaction of the alloy sample was illustrated in Fig. 5 to provide the detailed procedural model during the hydriding reaction. The reaction regions were composed of Nb-depleted a-likeU and the Nb-rich c12 as illustrated in step one, which is defined as the hydrogen steaming step, hydrogen was absorbed by the alloy surface and dissociated into hydrogen atoms [8,9]. In step two, the hydrogen atoms diffused into the bulk alloy along the bi-phase boundaries which played the role of diffusion path, and the a-like-U phases reacted with hydrogen to form the bUH3Nbx, the hydrides nucleated beneath the surface. As the reaction, the hydride grew up accompanied by volume expansion and consequent stress, leading the bulk material to emboss or even crack in the fierce reaction areas. Here the cause of the formation for b-UH3Nbx has not been acquired yet, which will be the key researching problems in further study.
Boundary Oxide
Boundary Oxide
Nb-enriched γ
Nb-enriched γ
Nb-depleted α
Nb-depleted α
Step 1
Step 2 Boundary
Oxide
Nb-enriched γ Nb-depleted α Fig. 5. Scheme of the possible formation mechanism of the hydriding reaction of the 400 °C/9 h + 500 °C/2 h aged alloy samples.
H. Ji et al. / Journal of Nuclear Materials 464 (2015) 43–47
5. Conclusion The microstructure investigation for the 400 °C/9 h + 500 °C/2 h aged U–13 at.%Nb alloy revealed a phase composition of a-likeU + c12 caused by relatively complete phase decomposition. The hydriding study results implied that the hydrogen preferentially reacted with the Nb-depleted phase a-like-U to form b-UH3Nbx. The hydrided areas did not turn into powder due to the special bi-phase microstructure. The hydriding growth is closely connected with the microstructure of the alloy, however, the evidence acquired here could not identify the initial nucleation sites of the hydrides, a series of subsequent study will be carried out in the next step for further information. Acknowledgements The authors gratefully acknowledge Jun Wu for alloy preparation, Xianglin Chen for optical metallography. This work was supported by the Development Fund for China Academy of Engineering Physics, No. 2014B0301049. References [1] K. Eckelmeyer, A. Romig, L. Weirick, Metall. Trans. A 15 (1984) 1319–1330.
47
[2] S.G. Bazley, J.R. Petherbridge, J. Glascott, Solid State Ionics 211 (2012) 1–4. [3] J. Glascott, Sci. Technol. J. AWE 6 (2003) 16–27. [4] C.P. Jones, T.B. Scott, J.R. Petherbridge, J. Glascott, Solid State Ionics 231 (2013) 81–86. [5] R.M. Harker, J. Alloys Compd. 426 (2006) 106–117. [6] R. Li, X. Wang, J. Nucl. Mater. 449 (2014) 49–53. [7] R. Lillard, C. Taylor, J. Wermer, N. Mara, J. Cooley, J. Nucl. Mater. 444 (2014) 49– 55. [8] C.D. Taylor, R. Scott Lillard, Acta Mater. 57 (2009) 4707–4715. [9] D. Christopher, T.L. Taylor, L. Scott, Acta Mater. 58 (2010) 1045–1055. [10] G. Powell, W. Harper, J. Kirkpatrick, J. Less Comm. Met. 172 (1991) 116–123. [11] J. Condon, J. Phys. Chem. 79 (1975) 392–397. [12] T.B. Massalski, H. Okamoto, P. Subramanian, L. Kacprzak, Binary Alloy Phase Diagrams, ASM International, 1990. [13] L. Hsiung, J. Zhou, Low-Temperature Aging Kinetics of a 15-Year Old WaterQuenched U-6 wt.% Nb alloy, UCRL-TR-235973Livermore, Lawrence Livermore National Laboratory, USA, 2007. [14] D. Moreno, R. Arkush, S. Zalkind, N. Shamir, J. Nucl. Mater. 230 (1996) 181– 186. [15] H. Ji, J. Yang, G. Hu, C. Jiang, J. Adv. Mater. Res. (2014) 81–92. [16] R.E. Hackenberg, K.D. Clarke, R. Forsyth, A.M. Kelly, T.J. Tucker, P.A. Papin, H.M. Volz, G.M. Hemphill, Discontinuous Reactions in U–Nb Alloys: Energy Sinks and Approach to Equilibrium, Los Alamos National Laboratory (LANL), 2011. [17] A. Loui, The Hydrogen Corrosion of Uranium: Identification of Underlying Causes and Proposed Mitigation Strategies, Livermore, 2012. [18] P. Shi, L. Shen, B. Bai, D. Lang, L. Lu, G. Li, X. Lai, P. Zhang, X. Wang, J. Nucl. Mater. 441 (2013) 1–5. [19] I. Tkach, S. Mašková, Z. Mateˇj, N.T.H. Kim-Ngan, A.V. Andreev, L. Havela, Phys. Rev. B 88 (2013).
Contents lists available at ScienceDirect
Journal of Nuclear Materials journal homepage: www.elsevier.com/locate/jnucmat
The microstructure and hydriding characteristics of high temperature aged U–13 at.%Nb alloy Hefei Ji a, Peng Shi a, Ruiwen Li a, Chunli Jiang b, Jiangrong Yang b,⇑, Guichao Hu a a b
China Academy of Engineering Physics, Mianyang 621900, China Science and Technology on Surface Physics and Chemistry Laboratory, P.O. Box 718-35, Mianyang 621907, China
a r t i c l e
i n f o
Article history: Received 17 October 2014 Accepted 29 March 2015 Available online 3 April 2015
a b s t r a c t Niobium as alloying element significantly improves physical and chemical properties of metallic uranium, exhibiting great application potential in uranium alloy materials. The corrosion resistance performance as well as the internal alloy phase structure of uranium–niobium alloy is closely related to aging processes. Microstructure and hydriding characteristics of the 400 °C/9 h + 500 °C/2 h aged uranium–13 at.% niobium alloys (U–13 at.%Nb) were investigated from the point of view of relationship between the microstructure and growth of the hydriding areas. The microstructure, morphology and composition of the alloy phases before and after the hydriding were well characterized by the laser scanning confocal microscopy (LSCM), scanning electron microscopy (SEM) and X-ray diffraction (XRD), respectively. Experimental results indicated that the hydrogen preferentially reacted with the Nb-depleted phase alike-U to form monolithic b-UH3Nbx, and the alloy microstructure played an important role in hydride growth. Ó 2015 Elsevier B.V. All rights reserved.
1. Introduction Alloying additives including common metallic elements added to uranium can optimize the performance while maintaining the advantageous characteristics of this nuclear metal. Niobium is one kind of alloying metal added to strengthen the corrosion resistance of obtained alloy [1] which is often quenched to form metastable solid solutions to achieve desired properties and performances. It is worth noting that although the as-quenched U–Nb alloy possesses perfect oxidation and corrosion resistance, it does react with H2O and H2 under ambient conditions. Previous investigations demonstrated that the metal uranium and its alloys could readily react with hydrogen under particular pressure–temperature conditions [2–7]. The reaction product was extremely unstable in air and its pyrophoricity when exposed to atmosphere made it highly hazardous. So it is of great significance to study the corrosion resistance of uranium as well as its alloys. Numerous studies focused on the kinetics of hydriding reaction of uranium [4,8–11]. The nucleation rates of the corrosion pitting increased with ascending hydrogen pressure as well as reaction temperature [2]. Several reasons could be listed for the hydride formation, however, by now there still not exist a systemic nucleation ⇑ Corresponding author. E-mail address: [email protected] (J. Yang). http://dx.doi.org/10.1016/j.jnucmat.2015.03.051 0022-3115/Ó 2015 Elsevier B.V. All rights reserved.
mechanism explaining the location of the initial hydride nucleation and corresponding initial growth characteristics. In addition, some related researches were relatively less concentrated on the reaction of uranium alloys and hydrogen remained less studied than for pure uranium. Under equilibrium conditions, niobium has a complete solubility at high temperatures and can form bcc solid solution with uranium, while exhibiting a limited solubility at relatively lower temperature. The as-quenched uranium–niobium (U–Nb) alloy has a single phase a00 [12], of which the internal phase structure varies to form the niobium-rich phase and niobium-depleted phase after a long duration of natural aging or accelerating artificial aging, yielding the a + c2 phase eventually [13]. As U–Nb alloys are one kind of most widely used forms of uranium and exhibit different types of microstructures, we undertook U–13 at.%Nb alloy aged at 400 °C for 9 h followed by 500 °C for 2 h (marked as 400 °C/9 h + 500 °C/2 h), the bulk microstructure of the aged alloy were also well investigated.
2. Experimental details 2.1. Sample preparation All hydriding experiments investigating the nucleation and expansion of hydride were conducted in a hot stage microscope system (HSM) [14] equipped with sealed furnace chamber with a
44
H. Ji et al. / Journal of Nuclear Materials 464 (2015) 43–47
quartz glass window enabling in-situ monitoring of the reaction between U–Nb alloys and hydrogen using a Hirox KH 7700 digital microscope. A U–13 at.%Nb alloy sample in the form of 8 mm diameter rod was obtained using water quenching. Except for the niobium atoms, the total impurity content in the alloy is not more than 200 ppm, including carbon, the concentration of which turned out to be approximately 50 ppm. In order to investigate the relationship between the hydride and microstructure of the alloy, especially the ‘‘a + c12’’ microstructure, the alloy rod was then transferred to a pre-set condition, under which the alloy sample was processed by the aging treatment at temperature of 400 °C for 9 h followed by a successive aging treatment at 500 °C for 2 h. The rod was then cut into small discs with the diameter of 8 mm and thickness of 2 mm. X-ray diffraction (XRD), metallographic and scanning electron microscopy (SEM) analysis were performed on these samples for characterization of microstructure, composition and morphology. After a set of wet mechanical polishing treatment in the open laboratory, the samples were ultra-sonicated in acetone and ethanol successively to remove the oxidized surface and other impurities. Then the samples were loaded into the visual volumecalibrated stainless steel reaction chamber positioned within a tube furnace. Afterward the chamber was evacuated to the pressure of 105 bar using an oil-free scroll pump followed by keeping the system at constant pressure for 10 min to remove oxygen and other gas impurities to the greatest extent. 2.2. Hydriding As long as the pressure of 105 bar was achieved and held for 10 min, the chamber was heated to 160 °C to activate the sample under dynamic vacuum for 2 h. After the activation, the chamber was cooled down to the reaction temperature of 100 °C followed by the immediate injection of high purity H2. The hydriding experiments were operated under the initial hydrogen pressure of 0.1 bar. The surface appearance of the samples was monitored through the chamber window using a digital microscope. In this work, the hydrogen consumption quantified by the hydrogen pressure dropping to half the initial pressure value was thought to be sufficient. In the end, the reaction was terminated by evacuating hydrogen gas from the reaction chamber. After being cooled down to room temperature, the chamber was reaerated to oxidize the hydride followed by taking the samples out of the chamber. The hydride on the sample surface was then studied and investigated by a laser scanning confocal microscope (LSCM, Lext OLS-4000), a scanning electronic microscope (SEM, SIRION 200) with associated EDS and an X-ray diffractometer (XRD, Empyream). 3. Results 3.1. Microstructure Before the hydriding experiments, the microstructure and phase composition of the alloy sample was studied. Fig. 1 shows the metallographic morphology (Fig. 1(a)) and the higher magnification SEM BSE image for the 400 °C/9 h + 500 °C/2 h sample (Fig. 1(b)). Compared with the as-quenched and low-temperature aged sample [15], the surface morphology of the sample in this work is more coarse. In Fig. 1(b), the lamellar structures proved to be the parallel pearlite composed of the niobium-rich (Nb-rich) phase (the gray layers, eroded during the etching processing) and the Nb-depleted phase (the bright layers) [16] accordingly. The characteristic dimension of the fine lamellar pearlite structure is
approximately 50–100 nm, as seen from Fig. 1(b). The uranium concentration on the nanoscale cannot be quantified. Upon the heat treatment condition of 400 °C/9 h + 500 °C/2 h, the discontinuous coarsening happens along the prior-c boundaries [16] in the sample. Additional characterization using XRD (Cu Ka) demonstrated the composite character of the aged alloy, when compared to non-alloyed uranium. The results shown in Fig. 2 indicate that the phase composition should be c12 (Nb-rich) + a-uranium (depleted) for the aged alloy. However, there is a slight peak shift toward lower angles compared with pure uranium metal, implying a lattice parameter expansion. The lattice parameters of the analogous a-uranium in the aged alloy (a = 2.885 Å, b = 5.873 Å, c = 4.970 Å) are slightly larger than those observed in a-uranium (a = 2.853 Å, b = 5.872 Å, c = 4.952 Å). Isothermal aging at temperature between 300 °C and 650 °C causes phase separation generically depicted as the progressive evolution Nb-rich and Nb-depleted regions to the final a(0Nb) + c2(50Nb) phases obtained only under extreme conditions. Thus there still exists some residual niobium atoms dispersed in the Nb-depleted phase for the aged alloy in this work, expanding the lattice parameters of the Nb-depleted phase than that of a-uranium. So we define the Nb-depleted phase in the aged alloy sample as a-like-U instead of a-U. In that way, the composition of the aged alloy forementioned is optimized to be alike-U + c12.
3.2. Hydride growth During the hydriding process, the reaction time was defined as duration within which the hydrogen pressure decreased to 60% of its initial value, so as to acquire sufficient hydride areas. Fig. 3 shows the SEM characterization micrographs of the alloy surface after the hydriding reaction. A series of circular hydriding areas with different sizes formed on the surface of the alloy, indicating their different initial nucleation time. From the enlarged image, an expanded hydriding area can be clearly distinguished. However, we do not assume the hydrided areas to be hydride, as they do not exhibit the powder form and they have a composite character, which is different from previously reported results [17,18] indicating the reaction products of uranium and hydrogen to be b-UH3 powders. Then, do the circular areas signify a kind of new reaction product just like the similar reaction products (UH3Mo0.18) obtained from the reaction between the so-called UMo0.18 [19] and hydrogen? Or the product is not yielded by the reaction between uranium and hydrogen at all. The XRD results of the hydrided sample are shown in Fig. 4. The pattern reveals the peaks assigned to b-UH3, but they are shifted to lower diffraction angles, giving the lattice parameter a = 6.642 Å, indicating a small lattice expansion in comparison with b-UH3 (6.631 Å). The Nb atoms are smaller than uranium, niobium atoms dispersed in the reaction products, and in the hydride, there happened a charge transfer in hydrides, which may actually shrink the electropositive uranium atoms, while Nb is affected less, contributing to the lattice expansion. Compared with the similar result that a lattice expansion exist in a-like-U obtained before hydriding, the smaller atoms Nb leading lattice aberration is considered to be another factor. Here the reaction product, which we denote as bUH3Nbx, should be subjected to further studies for the reason that we cannot give the determined value of the niobium content. Here we can compare the situation with reaction between U–Ti alloys and hydrogen, illustrating that the hydride is nucleated preferentially in a-uranium and not in U2Ti [18]. So the reactants in this work are supposed to be a-like-U and hydrogen.
H. Ji et al. / Journal of Nuclear Materials 464 (2015) 43–47
45
Fig. 1. (a) LSCM bright-field micrograph of the 400 °C/9 h + 500 °C/2 h aged alloy sample and (b) corresponding SEM BSE micrograph.
Fig. 2. XRD results of the 400 °C/9 h + 500 °C/2 h aged alloy and uranium samples before hydriding experiment.
Fig. 3. SEM of the 400 °C/9 h + 500 °C/2 h aged alloy and with (inset) magnified SEM result for a hydride area on the surface.
Fig. 4. XRD result of the 400 °C/9 h + 500 °C/2 h aged alloy sample after hydriding experiment.
46
H. Ji et al. / Journal of Nuclear Materials 464 (2015) 43–47
4. Discussions In this work, we paid attention to the relation between the hydride growth and microstructure of the bulk U–13 at.%Nb alloy after the aging treatment. The 400 °C/9 h + 500 °C/2 h aged alloy samples were characterized and the results showed that the microstructure and phase composition of the aged alloy sample changed from fine twinned martensite [13,15] to bi-phase pearlite after the aging treatment under elevated temperature conditions. The bi-phase pearlite was composed of the Nb-rich (c12) phase and Nb-depleted phase (a-like-U), which alternately exhibited the parallel microstructure with the dimension ranging from 50 nm to 100 nm. The LSCM and SEM photographs show monolithic reaction products on the alloy surface, but some embossing reaction regions. The XRD results in Section 3.2 above demonstrated the formation of the hydride b-UH3Nbx after hydriding, and the density difference between the hydride (10.9 g/cm3) and the bulk alloy (17.5 g/cm3) led to a volume expansion and some subsequent outwards raised hydriding regions. The monolithic hydriding areas formed in the aged alloy are different from that in uranium and U–Ti alloy, but similar to those in U–Mo0.18 [19]. Our previous work also reported similar monolithic hydriding characteristics of the 400 °C/3 h aged U–13 at.%Nb alloy [15], which illustrated that the hydriding areas expanded along specific directions forming specific tree-like hydriding areas. The U–13 at.%Nb alloys with complex microstructure obtained by different aging treatment described in the present work have also a specific morphology, affecting the expansion dye to hydriding. It is well known that b-UH3 is sensitive to oxygen, still our hydrided product does not turn into powder. We can assume the Nb-rich areas, which are separated only by 50–100 nm of Nb-depleted phase, can maintain the integrity of material when the bUH3Nbx is formed in the Nb-depleted phase. Moreover, the lack of fragmentation prevents the b-UH3Nbx phase formed inside the bulk (see Fig. 5) to be exposed to atmosphere, and oxidation therefore does not proceed and is limited to the surface only. We assume that the phase boundary is a suitable diffusion path for hydrogen into the bulk.
The reaction products here cannot be defined as b-UH3, the XRD analysis results further confirmed it. During the hydriding reaction, hydrogen preferentially reacted with the a-like-U among the Nbrich phases. In addition, the quantity of the a-like-U is small for the reason that the distance of the two Nb-rich phases was just 50–100 nm. In this regard, the hydrogen only reacted with the little sized a-like-U to form a spot of b-UH3Nbx, which still existed among the Nb-rich phases with a little volume expansion. The bUH3Nbx quantity between every two Nb-rich phases was insufficient to make the material crack, and is too little to be observed at macro-scale. Hence the b-UH3Nbx was packed between the c12 phases, making it difficult to be observed and oxidized. The composition of the bulk alloy material contained more than the equable fine bi-phase, there also exist the block a-like-U or c12 parts. So the hydriding reaction arose in the block a-like-U parts contributed to the cracking of the bulk material. With the reaction in the bi-phase areas, the overall volume expansion accumulated from every individual little volume expansion, exhibiting the monolithic embossing hydriding areas on the sample surface. The reaction products b-UH3Nbx is concealed among the Nb-rich phases in the bulk material, so when the sample after hydriding was exposed to the ambient atmosphere, the b-UH3Nbx were not exposed to the atmosphere directly, then they still can be identified by the XRD diffractometer. Here the scheme of a hypothetical formation mechanism of the hydriding reaction of the alloy sample was illustrated in Fig. 5 to provide the detailed procedural model during the hydriding reaction. The reaction regions were composed of Nb-depleted a-likeU and the Nb-rich c12 as illustrated in step one, which is defined as the hydrogen steaming step, hydrogen was absorbed by the alloy surface and dissociated into hydrogen atoms [8,9]. In step two, the hydrogen atoms diffused into the bulk alloy along the bi-phase boundaries which played the role of diffusion path, and the a-like-U phases reacted with hydrogen to form the bUH3Nbx, the hydrides nucleated beneath the surface. As the reaction, the hydride grew up accompanied by volume expansion and consequent stress, leading the bulk material to emboss or even crack in the fierce reaction areas. Here the cause of the formation for b-UH3Nbx has not been acquired yet, which will be the key researching problems in further study.
Boundary Oxide
Boundary Oxide
Nb-enriched γ
Nb-enriched γ
Nb-depleted α
Nb-depleted α
Step 1
Step 2 Boundary
Oxide
Nb-enriched γ Nb-depleted α Fig. 5. Scheme of the possible formation mechanism of the hydriding reaction of the 400 °C/9 h + 500 °C/2 h aged alloy samples.
H. Ji et al. / Journal of Nuclear Materials 464 (2015) 43–47
5. Conclusion The microstructure investigation for the 400 °C/9 h + 500 °C/2 h aged U–13 at.%Nb alloy revealed a phase composition of a-likeU + c12 caused by relatively complete phase decomposition. The hydriding study results implied that the hydrogen preferentially reacted with the Nb-depleted phase a-like-U to form b-UH3Nbx. The hydrided areas did not turn into powder due to the special bi-phase microstructure. The hydriding growth is closely connected with the microstructure of the alloy, however, the evidence acquired here could not identify the initial nucleation sites of the hydrides, a series of subsequent study will be carried out in the next step for further information. Acknowledgements The authors gratefully acknowledge Jun Wu for alloy preparation, Xianglin Chen for optical metallography. This work was supported by the Development Fund for China Academy of Engineering Physics, No. 2014B0301049. References [1] K. Eckelmeyer, A. Romig, L. Weirick, Metall. Trans. A 15 (1984) 1319–1330.
47
[2] S.G. Bazley, J.R. Petherbridge, J. Glascott, Solid State Ionics 211 (2012) 1–4. [3] J. Glascott, Sci. Technol. J. AWE 6 (2003) 16–27. [4] C.P. Jones, T.B. Scott, J.R. Petherbridge, J. Glascott, Solid State Ionics 231 (2013) 81–86. [5] R.M. Harker, J. Alloys Compd. 426 (2006) 106–117. [6] R. Li, X. Wang, J. Nucl. Mater. 449 (2014) 49–53. [7] R. Lillard, C. Taylor, J. Wermer, N. Mara, J. Cooley, J. Nucl. Mater. 444 (2014) 49– 55. [8] C.D. Taylor, R. Scott Lillard, Acta Mater. 57 (2009) 4707–4715. [9] D. Christopher, T.L. Taylor, L. Scott, Acta Mater. 58 (2010) 1045–1055. [10] G. Powell, W. Harper, J. Kirkpatrick, J. Less Comm. Met. 172 (1991) 116–123. [11] J. Condon, J. Phys. Chem. 79 (1975) 392–397. [12] T.B. Massalski, H. Okamoto, P. Subramanian, L. Kacprzak, Binary Alloy Phase Diagrams, ASM International, 1990. [13] L. Hsiung, J. Zhou, Low-Temperature Aging Kinetics of a 15-Year Old WaterQuenched U-6 wt.% Nb alloy, UCRL-TR-235973Livermore, Lawrence Livermore National Laboratory, USA, 2007. [14] D. Moreno, R. Arkush, S. Zalkind, N. Shamir, J. Nucl. Mater. 230 (1996) 181– 186. [15] H. Ji, J. Yang, G. Hu, C. Jiang, J. Adv. Mater. Res. (2014) 81–92. [16] R.E. Hackenberg, K.D. Clarke, R. Forsyth, A.M. Kelly, T.J. Tucker, P.A. Papin, H.M. Volz, G.M. Hemphill, Discontinuous Reactions in U–Nb Alloys: Energy Sinks and Approach to Equilibrium, Los Alamos National Laboratory (LANL), 2011. [17] A. Loui, The Hydrogen Corrosion of Uranium: Identification of Underlying Causes and Proposed Mitigation Strategies, Livermore, 2012. [18] P. Shi, L. Shen, B. Bai, D. Lang, L. Lu, G. Li, X. Lai, P. Zhang, X. Wang, J. Nucl. Mater. 441 (2013) 1–5. [19] I. Tkach, S. Mašková, Z. Mateˇj, N.T.H. Kim-Ngan, A.V. Andreev, L. Havela, Phys. Rev. B 88 (2013).