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Enhancing high-temperature strength of B4C–6061Al neutron absorber material by in-situ Mg(Al)B2
Journal of Nuclear Materials 526 (2019) 151788
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Journal of Nuclear Materials journal homepage: www.elsevier.com/locate/jnucmat
Enhancing high-temperature strength of B4Ce6061Al neutron absorber material by in-situ Mg(Al)B2 Y.N. Zan a, c, Q. Zhang a, c, Y.T. Zhou a, *, Q.Z. Wang b, B.L. Xiao a, Z.Y. Ma a, ** a
Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang, 110016, China Key Laboratory of Nuclear Materials and Safety Assessment, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang, 110016, China c School of Materials Science and Engineering, University of Science and Technology of China, 72 Wenhua Road, Shenyang, 110016, China b
h i g h l i g h t s B4Ce6061Al neutron absorber material was fabricated by powder metallurgy and then hot-rolling. The interfacial reactions were controlled by adjusting hot-pressing temperatures. The composites were subjected to microstructural examinations and mechanical property tests. Increasing amount of reaction products in the composites resulted in the enhanced strength at room and high temperatures. The reaction products, Mg(Al)B2 nanorods, can effectively impede the movements of grain boundaries and dislocations.
a r t i c l e i n f o
a b s t r a c t
Article history: Received 10 June 2019 Received in revised form 5 August 2019 Accepted 4 September 2019 Available online 5 September 2019
The industrial development of dry storage method of spent nuclear fuels strongly requires the low-cost B4CeAl neutron absorber materials (NAMs) with high strength and high stability at elevated temperatures. In the present work, the B4Ce6061Al composites were fabricated by powder metallurgy method with different hot-pressing temperatures. Microscopic analysis showed that as hot-pressing temperatures increased from 560 C to 630 C, the interfacial reactions became severe, resulting in the formation of Al3BC and Mg(Al)B2 compounds. The mechanical properties of the composites were enhanced by the interfacial reactions, specifically, the tensile strength of the sample hot-pressed at 630 C was 40e59% higher than that at 560 C at testing temperatures from room temperature up to 350 C. Moreover, the high-temperature-pressed sample maintained the invariant strength level after suffering annealing at 400 C for 8000 h. Characterization of the fractured composite indicated that the Mg(Al)B2 nanorods which effectively impeded the dislocation movement and grain boundary sliding provided a significant contribution to the composite strength. © 2019 Published by Elsevier B.V.
Keywords: B4CeAl composite In-situ reinforcement High-temperature strength Thermal stability TEM
1. Introduction B4CeAl composites are widely used as neutron absorber materials (NAMs) nowadays owning to the abundant 10B isotope (high thermal neutron absorption cross-section of 3837 b) in B4C particles [1e6]. For safety and economy considerations, there is a growing interest in developing lightweight B4CeAl NAMs for the novel dry storage [7]. For such application, the mechanical
* Corresponding author. ** Corresponding author. E-mail addresses: [email protected] (Y.T. Zhou), [email protected] (Z.Y. Ma). https://doi.org/10.1016/j.jnucmat.2019.151788 0022-3115/© 2019 Published by Elsevier B.V.
properties and thermal stability of the material at high temperatures (>300 C) are specifically required. For most of the commercially used matrix alloys, such as 2xxx and 6xxx series [8], their ambient-temperature strengths are dominated by the nano-sized precipitates. However, the deficiency of these alloys at higher temperatures is the rapid coarsening of the strengthening precipitates [9]. Even some commercial heatresistant precipitation-strengthened aluminium alloys, like AleSieCu or AleSieMg series cannot maintain good strength at the working temperatures [10e12]. In Lai et al.'s work, a high-cost AlSc-Zr alloy was used as the matrix of NAMs [13]. Because of the thermal stability of the precipitates Al3(Sc,Zr) in the matrix, the composite did not loss the strength after long-term annealing until
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it experienced thermal holding above 300 C. Alternatively, thermally stable ceramic nanoparticles can be used as reinforcement to overcome the high-temperature strength deficiency of the aluminum alloy matrixes [14e19]. However, when introducing the nanometric particles into Al matrices, extra difficulties and costs are inevitable. Most of the approaches therefore are not suitable for large-scale industrial production. In B4C/Al system, it is known that the B4C particles are very reactive with molten aluminum alloy [20]. The chemical reactions between B4C and the alloy matrix (interfacial reactions for short, although not always happened at the B4C/Al interfaces) produces complex phases, including AlB12, AlB2, Al3BC, MgB2, etc [20e22] under different reaction conditions. As usual, the reaction products possess high melting points and high stability at the working temperature for a neutron absorber. In our recent work [23], we fabricated B4Ce6061Al composite containing in-situ Mg(Al)B2 nanorods. The material presented highly stable microstructures during annealing at 400 C for as long as 8000 h. Moreover, the Mg(Al)B2 nanorods showed strong pinning effect against the migration of grain boundaries (GBs) which is often thought to be responsible for the high-temperature softening of alloys. It is expected that the reaction products would also improve the strength of B4CeAl composite at elevated temperatures (T > 0.5 Tm, Tm is the melting temperature of the alloy matrix). In the present study, we fabricated B4CeAl NAM using the commercial component of B4Ce6061Al to show the positive effect of the interfacial reactions to improve the high-temperature strength. The composites were produced by powder metallurgy method with different hot-pressing temperatures to control the interfacial reactions. It was shown that the composite containing Mg(Al)B2 dispersoids exhibited higher strength at both ambient and high temperatures. Microscopic analysis on the fractured samples indicated that the Mg(Al)B2 nanorods play a significant role in strengthening GBs and impeding dislocation activities. This study is believed to be helpful to clarify the effects of reaction products on the mechanical properties of B4CeAl composites and provide a new route to optimize the integrated properties of the composites for high temperature usage. 2. Experimental Commercial alloy powders (Angang Group Aluminum Powder Co., Ltd.) with a nominal composition of Al-1.0Mg-0.65Si-0.25Cu and B4C particles (Mudanjiang Jingangzuan Boron Carbide Co., Ltd.) were used as the raw materials. The mean sizes of Al powder and B4C powder were 13 and 7 mm, respectively. 26 wt% B4Ce6061Al composites were fabricated using powder metallurgy (PM) technique. The as-received B4C particles were mechanically mixed with the alloy powders in a bi-axis rotary mixer for 8 h with a ball to powder ratio of 2:1 and a rotation speed of 50 rpm. The asmixed powders were cold pressed into the green billets under a pressure of 50 MPa. Since interfacial reactions become severe with the appearance of liquid phase above 600 C [21], the hot-pressing temperatures of 560 C, 600 C, and 630 C were chosen to fabricate three composites (designated as HP 560, HP 600, and HP 630 hereafter, respectively). Hot pressing was carried out in a vacuum chamber of 102 Pa. The billets were heated up to the preset temperature with a heating rate of 5 C/min, and then held for 2 h, after which hot pressing was carried out under a pressure of 40 MPa. The billets were hot-rolled into plates at 450 C with a height reduction ratio of 10% for each pass. A total height reduction ratio of 90% was attained. The densities of the composite samples were measured by using the Archimedes principle. To test the long-term thermal stability,
HP 630 were annealed uninterruptedly at 400 C for up to 8000 h in a muffle furnace and then cooled inside the furnace. The microstructural studies were carried out by means of optical microscopy (OM, Zeiss Axiovert 200MAT), field-emission scanning electron microscopy (SEM, Leo Supra 55) and transmission electron microscopy (TEM, TECNAI G2 F20) equipped with high-angle annular-dark-field (HAADF) detectors. The chemical compositions of the secondary phases were analyzed by X-ray energy-dispersive spectroscopy (EDS) under SEM and HAADF-mode TEM. TEM specimens were prepared by mechanical polishing and ion milling using Leica RES101 system. Tensile specimens were machined from the as-rolled composite sheets along the rolling direction. For room temperature (RT) tension, Instron 8810 equipped with an extensometer was used to obtain accurate yield stress and elongation, and specimens with a gauge length of 15 mm, a width of 3 mm and a thickness of 2.5 mm were used. For high temperature tests, tensile specimens (5 mm in length, 1.5 mm in width and 1.2 mm in thickness) were tested on Instron 5582 equipped with a heating furnace. At least 5 tensile specimens were tested for each composite at a strain rate of 103 at both RT and high temperatures.
3. Results 3.1. Densities and microstructures Table 1 shows densities of the fabricated composites. The results indicate that all composites were almost completely densified and the densities increased slightly with increasing hot-pressing temperatures. Optical microscopic analysis of as-rolled composites is shown in Fig. 1. B4C particles were homogenously distributed in these composites and no micro-pores were observed. Compared to the ‘clear’ matrix in HP 560 and HP 600 (Fig. 1a and b, respectively), the alloy matrix in HP 630 (Fig. 1c) was decorated by high-density of fine dispersoids as marked by arrows in the inset, which were generated by the interfacial reactions. The characterization of the dispersed particles will be shown later. The XRD patterns of the samples hot-pressed at different temperatures are shown in Fig. 2. The patterns were normalized by adjusting the peaks of B4C to the same height. In HP 560 and HP 600, the peaks of Mg2Si were found besides the peaks of Al and B4C. In the latter sample, the reaction products, Mg(Al)B2 and Al3BC were also detected. When the hot-pressing temperature was increased to 630 C, the interfacial reactions became severer and the quantities of Mg(Al)B2 and Al3BC increased. Due to the consumption of element Mg by the interfacial reactions, the Mg2Si precipitate was too less to be detected. MgAl2O4 spinel which was frequently seen in Mg-containing aluminum matrix composites [24,25] was present in these three samples. The Mg(Al)B2 and Al3BC in different samples were quantified based on the relative diffraction intensity of Al (111), B4C (104), Al3BC (101) and Mg(Al)B2 (100) planes. The relative intensity fractions of the Mg(Al)B2 and Al3BC in HP 560, HP 600 and HP 630 were calculated by the following equation used in Ref. [21].
Table 1 Densities of B4Ce6061Al composites with various hot-pressing temperatures. Sample 3
Density (g/cm ) Relative density (%)
HP 560
HP 600
HP 630
2.637 99.5
2.648 99.9
2.658 100.3
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Fig. 1. Optical micrographs of the samples (a) HP 560, (b) HP 600 and (c) HP 630. The B4C particles (black) homogenously distribute in all of the three samples. The fine dispersoids in HP 630 were denoted by black arrows in the image inset.
Fig. 2. XRD patterns of the as-rolled B4Ce6061Al composites with different hotpressing temperatures.
DX ð%Þ ¼
IX ð%Þ IAl þ IX þ IB4 C
(1)
where DX is the relative intensity fraction of compound X (Mg(Al)B2 or Al3BC), Ix is the diffraction intensity of X, IAl and IB4C are the intensities of Al and B4C, respectively. Calculated results are listed in Table 2. It is revealed that the interfacial reactions were enhanced by increasing the hot-pressing temperatures, resulting in the formation of more Mg(Al)B2 and Al3BC. The secondary phases in these samples were further analyzed by SEM and TEM, as shown in Fig. 3. In HP 560, sharp and clean edges of B4C particles were observed (Fig. 3a). Some dark particles (as denoted by arrows) were found in the sample and also in HP 600 (shown in Fig. 3b). In HP 600, some of the particles have grown to larger scales, i.e. about 2e4 mm in diameter. The chemical composition analyzed by EDS mapping indicated that they contained element Mg and Si. TEM characterization and selected area electron diffraction (SAED) were carried out, identifying these particles (Fig. 3d) to be b-Mg2Si (lattice constant a ¼ 0.45 nm). In addition, Table 2 The relative intensity fractions of the reaction products calculated based on Al (111), B4C (104), Al3BC (101) and Mg(Al)B2 (100) in the composites. Sample
HP 560 HP 600 HP 630
Relative fraction (%) Al3BC
Mg(Al)B2
0.12 0.55 0.78
0.20 0.76 1.55
the serrated B4C-particle edges in HP 600 implied that the interface reactions have occurred. When the hot-pressing temperature was increased to 630 C, the reactions between B4C and the alloy matrix became more apparent and a high density of secondary phases were detected (Fig. 3c). EDS analysis suggested that the grey phases (as denoted by yellow arrow) were mainly composed of Mg and Al. Referring to the XRD result, they were speculated to be Mg(Al)B2. These phases were further confirmed by SAED and EDS analyses performed by TEM, as shown in Fig. 3e. The EDS signal was obtained from the compound located at the very edge of the perforation of the TEM foil to avoid interference from overlapped aluminum matrix. It should be pointed out that AlB2 and MgB2 crystals have the same crystallographic characteristics but a little variation in lattice parameters [26,27]. The compound in our samples probably contained different contents of metallic (Al or Mg) atoms. Another reaction product Al3BC was also frequently seen at the interface of B4C particles and Al matrix, as illustrated in the TEM image in Fig. 3f. To compare the reaction product Mg(Al)B2 in different samples, Fig. 4 shows the micrographs of Mg(Al)B2 in HP 600 and HP 630. In both the samples, the Mg(Al)B2 nanorods were dispersed uniformly in the alloy matrix. As marked by black arrows, the Mg(Al)B2 dispersoids in HP 600 were measured to be about 100 nm in diameter and 0.5e2 mm in length. It is clearly seen that the quantity of Mg(Al) B2 in HP 630 was much higher than that in HP 600. A part of the Mg(Al)B2 rods grew to larger sizes, so became visible under optical microscope (Fig. 1c). Fig. 5aec shows low-magnification TEM images of the three composites. A fine grained structure with grain sizes of 0.5e2 mm was observed for all the composites. The observations agreed with the EBSD analysis in our previous work [23]. The fine grains should be generated during the hot-rolling process through dynamic recrystallization. Thus, the grain sizes in these composites fabricated at different hot-pressing temperatures were in the same range. The dislocation lines could be visible under proper diffractioncontrast conditions. Fig. 5def shows the images of dislocations in the grains. The dislocation densities in these different composites were roughly estimated to be of the same order of magnitude (~1013 m2). Threading dislocation lines throughout the grain interior were frequently observed in HP 560 (Fig. 5a). In contrast, the dislocations in the composites hot-pressed at higher temperatures (Fig. 5b and c) showed more complex and tangled feature owing to the interactions with the dispersoids inside the grains (arrowed in Fig. 5c). The dislocation densities in HP 600 and HP 630 were higher than those in HP 560. It was also seen that dislocation accumulation easily developed into sub-grain boundaries during hot-rolling process, as indicated in Fig. 5a.
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Fig. 3. Microstructural analysis of the samples. (a) SEM image of HP 560. (b) SEM image of HP 600. The inset is the elemental mapping of a Mg2Si particle. (c) SEM image of HP 630 in which the Mg(Al)B2 and Si particles can be found in the matrix. (d) TEM image and corresponding SAED pattern of the Mg2Si particle in HP 560. (e) HAADF image, SAED pattern and the EDS result of Mg(Al)B2 nanorods as interfacial reaction products. (f) Al3BC compound can be frequently found at the periphery of B4C particles in both HP 600 and HP 630.
Fig. 4. HAADF TEM images showing the Mg(Al)B2 distributions in (a) HP 600 and (b) HP 630.
3.2. Mechanical properties at different temperatures The stress-strain curves of the composites are shown in Fig. 6. It can be seen that the samples fabricated at different hot-pressing temperatures showed significant differences in strength and ductility at RT (Fig. 6a). The yield strength (YS) and the ultimate stress (UTS) of the composites increased with increasing hotpressing temperatures. The YS and UTS values of HP 560 were only 120 MPa and 198 MPa, respectively. A simultaneous enhancement of strength and elongation (EL) was obtained for HP 600, with the YS, UTS and EL values respectively reaching 166 MPa, 254 MPa and 12%. Obviously, the appearance of interfacial reactions improved the mechanical properties of the composite. For HP 630, the strength was further increased to 188 MPa (YS) and 315 MPa (UTS), which were about 57% and 59% larger than that of HP 560. However, the elongation declined to 7%. Fig. 6b shows the stress-strain curves of different composites measured at temperatures from 150 C to 350 C. Since no extensometer was used for high-temperature testing, only UTS values of the composites were labelled next to the curves. At high temperatures, strength enhancement of the composites with increasing hot-pressing temperatures was still confirmed. The UTS values of HP 630 were about 40%e50% higher than those of HP 560 at the testing temperatures.
We further tested the mechanical stability of HP 630 suffered long-term thermal exposure at 400 C. The UTS values tested at 150 C and 350 C are shown in Fig. 7. It is seen that the tensile strength of HP 630 experienced as long as 8000 h annealing remained in the same level. For the commercially used precipitation-strengthened aluminum alloys, such extraordinary stability is hardly achieved. 3.3. Fractographs The fracture surface of the composite was characterized using SEM. Ductile fracture was observed at both 150 C and 350 C for HP 560 and HP 630, indicated by the dimple structure. No pull-out of the B4C particles from the matrix was found, implying that the B4C particle bonded well with the alloy matrix. For fractographs of both HP 560 (Fig. 8a) and HP 630 (Fig. 8b) deformed at 150 C, the morphological difference in between is quite noticeable. The larger and deeper dimples in HP 560 suggested the higher ductility of the composite. An enlarged area of the image showed that small particles could be found in the bottom of a dimple (inset in Fig. 8a). As revealed by the microscopic analysis, the Mg2Si particles in HP 560 acted as the stress concentration sites due to their brittle nature and easily resulted in local micro-void formation under external loading. For the case of HP 630 (Fig. 8b), the high density of in-situ
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Fig. 5. TEM images showing the microstructures of the aluminum alloy matrices. (a)e(c) Low-magnification TEM images of the fine grains in HP 560, HP 600 and HP 630, respectively. (d)e(f) TEM images of the dislocations inside the grains in HP 560, HP 600 and HP 630, respectively.
Fig. 6. Tensile stressestrain curves of the composites at (a) RT and (b) elevated temperatures.
formed particles via the interfacial reactions rendered higher strength of the sample, but also more sites for stress concentration and cracks nucleation. Therefore, the fracture surface showed more dimples with smaller sizes. When the testing temperature was increased to 350 C, the dimples for both composites exhibited larger size, which was in accordance with the higher EL values. The secondary phases were also found on the fracture surfaces, especially for HP 630. To identify these particles, EDS analysis was carried out (Fig. 8e). By comparing the elementary composition of matrix (e1) and particle (e2), in can be concluded that the particles were Mg2Si (Fig. 8c). In the case of HP 630 (Fig. 8d), EDS analysis results are shown in Fig. 8f. According to above microstructure analysis and the EDS profile of Al matrix (f1), the fine particles (f2) were identified as Mg(Al)B2. The fractograph at 350 C suggested that the nucleation of microvoids at the dimple bottoms was also closely related to the secondary phases in the composites.
4. Discussion It is well known that the chemical reactions between B4C and Al obviously happen when the liquid phase appear in the alloy matrix. For 6061Al, the liquid temperature was determined as about 600 C (Tm) [21]. Alloying element Mg in this composite is noteworthy. Below Tm, the interfacial reactions were quite mild. As determined by above results, Mg existed in the forms of Mg2Si in the matrix. When hot-pressing temperature was increased to 600 C, a small portion of liquid aluminum reacted with B4C, yielding Al3BC and Mg(Al)B2 nanorods. In HP 630, a large amount of liquid phase facilitated the interfacial reactions. Because of the larger chemical affinity between Mg and B, the reactions consumed most of element Mg. At the same time, element Si could not form compound with other elements, so they would precipitate as Si or AleSi eutectic phases in the cooling process. The strengthening effects of introducing particles are usually thought as direct strengthening by load-transfer (L-T) and
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Fig. 7. The values of ultimate tensile strength of HP 630 experienced various annealing periods. It shows the high stability of the composite at elevated temperature.
secondary strengthening relevant to dislocation activities, including Orowan strengthening, geometrically necessary dislocation strengthening, etc [23,28,29]. These mechanisms well explained the RT strength enhancement of HP 600 and HP 630 relative to that of HP 560. At elevated temperatures, however, the thermal energy is sufficient to activate diffusion, leading to the low efficiency of the above strengthening mechanisms [30,31]. To gain insight into the role played by Mg(Al)B2 nanorods during the high-temperature deformation process, TEM observations were performed on the sample taken from the fractured HP 630. Fig. 9a exhibits the bright-field TEM image in which several grains and dislocation tangles could be seen. Compared to that in the as-rolled sample displayed in Fig. 5c, a high-density of dislocations was found to accumulate at the periphery of GBs in the deformed sample, suggesting the impeding effect of GBs to the dislocation movement. More importantly, as the arrows denoted, several Mg(Al)B2 nanorods were also observed in front of GBs. It is known that GBs will lose their strength at elevated temperatures. GB sliding largely assists the high-temperature deformation of alloys [32]. The pinning effect provided by fine particles has been proved to be an effective way to hinder the GB motion [33,34]. Here the Mg(Al)B2 nanorods located at the GBs clearly indicated their roles
Fig. 8. SEM fractographs of the composites: (a) HP 560 deformed at 150 C and (b) HP 630 at 150 C. (c) HP 560 at 350 C and (d) HP 630 at 350 C. (e) and (f) are EDS results of the marked positions in (b) and (e), respectively.
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Fig. 9. TEM images of the microstructures in the fractured sample HP 630. (a) The grain boundaries are pinned by Mg(Al)B2 nanorods (denoted by white arrows) (b) The dislocations tangled around the Mg(Al)B2 suggests its impeding effect to dislocation activity.
of pinning the GBs. GB sliding requires accommodation by diffusion and the dislocation/GB interactions [35]. The dislocation accumulations in front of GBs will be released by GB sliding as a result. In such way, it is not easy to form stress concentration and the material is softened. In the presence of fine particles in the alloy matrix, the GB strength is apparently enhanced by the pinning of reinforcement. The dislocation pile-ups in Fig. 9a confirmed the stability of GBs in the present composite. Consequently, the high-temperature strength of the material can be largely improved. With increasing loading, however, the stress concentration at the GBs tends to result in the micro-void nucleation at the local sites and finally the fracture failure. The analysis is quite consistent with the fractographs shown in Fig. 8, i.e. the voids nucleated at the GBs where a massive of Mg(Al)B2 particles were exposed after fracture. On the other hand, the Mg(Al)B2 nanorods also served as obstacles for the dislocation motion. In general, the material strength at ambient temperature can be effectively improved by fine particles via Orowan bypassing mechanism. The increasing temperature can activate vacancy diffusion and the dislocation climbing weakens the strengthening effect by Orowan bypassing mechanism. Nevertheless, the high strain energy induced by the mismatch between reinforcements and Al matrix (modulus mismatch and lattice mismatch) also provides a certain barrier for dislocations to climb over [36,37]. According to Qin et al., the strength of the composite at elevated temperatures was controlled by the combination of dislocation climb and Orowan models [38]. One of the Mg(Al)B2 nanorods in Fig. 9a were extracted and displayed in Fig. 9b. The dislocations tangled around the nanorod implied the interactions in between. The hindrance to dislocation activities increases the yielding strength of the composite to some extent, although it is less effective than that at ambient temperature.
5. Conclusions 1. B4Ce6061Al composites were fabricated by the PM method with varying hot-pressing temperatures. Mg2Si is the main precipitate in the composite hot-pressed at lower temperatures of 560 C and 630 C. The composite produced at a higher hotpressing temperature of 630 C contained plenty of homogeneously dispersed Mg(Al)B2 nanorods.
2. The interfacial reactions are beneficial to improve the material strength at both room and elevated temperatures. The yield strength of HP 630 at RT is 57% higher than that of HP 560. Tensile strengths of HP 630 at high temperatures (150 C ~ 350 C) were increased by 40% ~ 50%. In addition, the mechanical properties of HP 630 were proven to be extremely stable after being annealed at 400 C for 8000 h. 3. The Mg(Al)B2 nanorods acted as obstacles to the GB sliding and the dislocation activities. As a consequence, the hightemperature strength of the composite was greatly enhanced.
Data availability The processed data required to reproduce these findings cannot be shared at this time as the data also forms part of an ongoing study. Acknowledgments The authors gratefully acknowledge the support of the National Natural Science Foundation of China (NSFC) under grant Nos. U1508216, 51501195 and 51771194, NSF of Liaoning Province [20180551101] and the Innovation Fund of IMR [2017-PY10]. Appendix A. Supplementary data Supplementary data to this article can be found online at https://doi.org/10.1016/j.jnucmat.2019.151788. References [1] X.-G. Chen, Application of Al-B4C metal matrix composites in the nuclear industry for neutron absorber materials, in: N. Gupta, W.H. Hunt (Eds.), Solidification Processing of Metal Matrix Composites, TMS, USA, 2006, pp. 343e350. [2] H.S. Chen, W.X. Wang, Y.L. Li, J. Zhou, H.H. Nie, Q.C. Wu, The design, microstructure and mechanical properties of B4C/6061Al neutron absorber composites fabricated by SPS, Mater. Des. 94 (2016) 360e367. [3] L. Zhou, C. Cui, Q.Z. Wang, C. Li, B.L. Xiao, Z.Y. Ma, Constitutive equation and model validation for a 31vol.% B4Cp/6061Al composite during hot compression, J. Mater. Sci. Technol. 34 (2018) 1730e1738. [4] Z.G. Xu, L.T. Jiang, Q. Zhang, J. Qiao, D. Gong, G.H. Wu, The design of a novel neutron shielding B4C/Al composite containing Gd, Mater. Des. 111 (2016) 375e381. [5] P. Zhang, Y. Li, W. Wang, Z. Gao, B. Wang, The design, fabrication and
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[22] Y.T. Zhou, Y.N. Zan, S.J. Zheng, Q.Z. Wang, B.L. Xiao, X.L. Ma, Z.Y. Ma, Distribution of the microalloying element Cu in B4C-reinforced 6061Al composites, J. Alloy. Comp. 728 (2017) 112e117. [23] Y. Zhou, Y. Zan, S. Zheng, X. Shao, Q. Jin, B. Zhang, Q. Wang, B. Xiao, X. Ma, Z. Ma, Thermally stable microstructures and mechanical properties of B4C-Al composite with in-situ formed Mg(Al)B2, J. Mater. Sci. Technol. 35 (2019) 1825e1830. [24] K.B. Lee, H.S. Sim, S.Y. Cho, H. Kwon, Reaction products of AleMg/B4C composite fabricated by pressureless infiltration technique, Mater. Sci. Eng. A 302 (2001) 227e234. [25] C. Nie, J. Gu, J. Liu, D. Zhang, Investigation on microstructures and interface character of B4C particles reinforced 2024Al matrix composites fabricated by mechanical alloying, J. Alloy. Comp. 454 (2008) 118e122. [26] J. Nagamatsu, N. Nakagawa, T. Muranaka, Y. Zenitani, J. Akimitsu, Superconductivity at 39 K in magnesium diboride, Nature 410 (2001) 63e64. [27] E.J. Felten, The preparation of aluminum diboride, AlB2, J. Am. Chem. Soc. 78 (1956) 5977e5978. [28] V.C. Nardone, K.M. Prewo, On the strength of discontinuous silicon carbide reinforced aluminum composites, Scr. Mater. 20 (1986) 43e48. [29] M. Tabandeh-Khorshid, J.B. Ferguson, B.F. Schultz, C.S. Kim, K. Cho, P.K. Rohatgi, Strengthening mechanisms of graphene- and Al2O3-reinforced aluminum nanocomposites synthesized by room temperature milling, Mater. Des. 92 (2016) 79e87. [30] J. Liu, G. Fan, Z. Tan, Q. Guo, Y. Su, Z. Li, D.-B. Xiong, Mechanical properties and failure mechanisms at high temperature in carbon nanotube reinforced copper matrix nanolaminated composite, Compos. Appl. Sci. Manuf. 116 (2019) 54e61. [31] Z.Y. Liu, B.L. Xiao, W.G. Wang, Z.Y. Ma, Elevated temperature tensile properties and thermal expansion of CNT/2009Al composites, Compos. Sci. Technol. 72 (2012) 1826e1833. [32] M.A. Meyers, K.K. Chawla, Mechanical Behavior of Materials, second ed., Combridge University Press, 2009. [33] C. Poletti, M. Balog, F. Simancik, H.P. Degischer, High-temperature strength of compacted sub-micrometer aluminium powder, Acta Mater. 58 (2010) 3781e3789. [34] M. Balog, P. Krizik, M. Nosko, Z. Hajovska, M. Victoria Castro Riglos, W. Rajner, D.-S. Liu, F. Simancik, Forged HITEMAL: Al-based MMCs strengthened with nanometric thick Al2O3 skeleton, Mater. Sci. Eng. A 613 (2014) 82e90. [35] Y. Qi, P.E. Krajewski, Molecular dynamics simulations of grain boundary sliding: the effect of stress and boundary misorientation, Acta Mater. 55 (2007) 1555e1563. [36] E.A. Marquis, D.C. Dunand, Model for creep threshold stress in precipitationstrengthened alloys with coherent particles, Scr. Mater. 47 (2002) 503e508. [37] A.J. Ardell, Precipitation hardening, Metall. Trans. A 16 (1985) 2131e2165. [38] J. Qin, Z. Zhang, X.G. Chen, Mechanical properties and strengthening mechanisms of Al-15 pct B4C composites with Sc and Zr at elevated temperatures, Metall. Mater. Trans. A 47A (2016) 4694e4708.
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Enhancing high-temperature strength of B4Ce6061Al neutron absorber material by in-situ Mg(Al)B2 Y.N. Zan a, c, Q. Zhang a, c, Y.T. Zhou a, *, Q.Z. Wang b, B.L. Xiao a, Z.Y. Ma a, ** a
Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang, 110016, China Key Laboratory of Nuclear Materials and Safety Assessment, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang, 110016, China c School of Materials Science and Engineering, University of Science and Technology of China, 72 Wenhua Road, Shenyang, 110016, China b
h i g h l i g h t s B4Ce6061Al neutron absorber material was fabricated by powder metallurgy and then hot-rolling. The interfacial reactions were controlled by adjusting hot-pressing temperatures. The composites were subjected to microstructural examinations and mechanical property tests. Increasing amount of reaction products in the composites resulted in the enhanced strength at room and high temperatures. The reaction products, Mg(Al)B2 nanorods, can effectively impede the movements of grain boundaries and dislocations.
a r t i c l e i n f o
a b s t r a c t
Article history: Received 10 June 2019 Received in revised form 5 August 2019 Accepted 4 September 2019 Available online 5 September 2019
The industrial development of dry storage method of spent nuclear fuels strongly requires the low-cost B4CeAl neutron absorber materials (NAMs) with high strength and high stability at elevated temperatures. In the present work, the B4Ce6061Al composites were fabricated by powder metallurgy method with different hot-pressing temperatures. Microscopic analysis showed that as hot-pressing temperatures increased from 560 C to 630 C, the interfacial reactions became severe, resulting in the formation of Al3BC and Mg(Al)B2 compounds. The mechanical properties of the composites were enhanced by the interfacial reactions, specifically, the tensile strength of the sample hot-pressed at 630 C was 40e59% higher than that at 560 C at testing temperatures from room temperature up to 350 C. Moreover, the high-temperature-pressed sample maintained the invariant strength level after suffering annealing at 400 C for 8000 h. Characterization of the fractured composite indicated that the Mg(Al)B2 nanorods which effectively impeded the dislocation movement and grain boundary sliding provided a significant contribution to the composite strength. © 2019 Published by Elsevier B.V.
Keywords: B4CeAl composite In-situ reinforcement High-temperature strength Thermal stability TEM
1. Introduction B4CeAl composites are widely used as neutron absorber materials (NAMs) nowadays owning to the abundant 10B isotope (high thermal neutron absorption cross-section of 3837 b) in B4C particles [1e6]. For safety and economy considerations, there is a growing interest in developing lightweight B4CeAl NAMs for the novel dry storage [7]. For such application, the mechanical
* Corresponding author. ** Corresponding author. E-mail addresses: [email protected] (Y.T. Zhou), [email protected] (Z.Y. Ma). https://doi.org/10.1016/j.jnucmat.2019.151788 0022-3115/© 2019 Published by Elsevier B.V.
properties and thermal stability of the material at high temperatures (>300 C) are specifically required. For most of the commercially used matrix alloys, such as 2xxx and 6xxx series [8], their ambient-temperature strengths are dominated by the nano-sized precipitates. However, the deficiency of these alloys at higher temperatures is the rapid coarsening of the strengthening precipitates [9]. Even some commercial heatresistant precipitation-strengthened aluminium alloys, like AleSieCu or AleSieMg series cannot maintain good strength at the working temperatures [10e12]. In Lai et al.'s work, a high-cost AlSc-Zr alloy was used as the matrix of NAMs [13]. Because of the thermal stability of the precipitates Al3(Sc,Zr) in the matrix, the composite did not loss the strength after long-term annealing until
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it experienced thermal holding above 300 C. Alternatively, thermally stable ceramic nanoparticles can be used as reinforcement to overcome the high-temperature strength deficiency of the aluminum alloy matrixes [14e19]. However, when introducing the nanometric particles into Al matrices, extra difficulties and costs are inevitable. Most of the approaches therefore are not suitable for large-scale industrial production. In B4C/Al system, it is known that the B4C particles are very reactive with molten aluminum alloy [20]. The chemical reactions between B4C and the alloy matrix (interfacial reactions for short, although not always happened at the B4C/Al interfaces) produces complex phases, including AlB12, AlB2, Al3BC, MgB2, etc [20e22] under different reaction conditions. As usual, the reaction products possess high melting points and high stability at the working temperature for a neutron absorber. In our recent work [23], we fabricated B4Ce6061Al composite containing in-situ Mg(Al)B2 nanorods. The material presented highly stable microstructures during annealing at 400 C for as long as 8000 h. Moreover, the Mg(Al)B2 nanorods showed strong pinning effect against the migration of grain boundaries (GBs) which is often thought to be responsible for the high-temperature softening of alloys. It is expected that the reaction products would also improve the strength of B4CeAl composite at elevated temperatures (T > 0.5 Tm, Tm is the melting temperature of the alloy matrix). In the present study, we fabricated B4CeAl NAM using the commercial component of B4Ce6061Al to show the positive effect of the interfacial reactions to improve the high-temperature strength. The composites were produced by powder metallurgy method with different hot-pressing temperatures to control the interfacial reactions. It was shown that the composite containing Mg(Al)B2 dispersoids exhibited higher strength at both ambient and high temperatures. Microscopic analysis on the fractured samples indicated that the Mg(Al)B2 nanorods play a significant role in strengthening GBs and impeding dislocation activities. This study is believed to be helpful to clarify the effects of reaction products on the mechanical properties of B4CeAl composites and provide a new route to optimize the integrated properties of the composites for high temperature usage. 2. Experimental Commercial alloy powders (Angang Group Aluminum Powder Co., Ltd.) with a nominal composition of Al-1.0Mg-0.65Si-0.25Cu and B4C particles (Mudanjiang Jingangzuan Boron Carbide Co., Ltd.) were used as the raw materials. The mean sizes of Al powder and B4C powder were 13 and 7 mm, respectively. 26 wt% B4Ce6061Al composites were fabricated using powder metallurgy (PM) technique. The as-received B4C particles were mechanically mixed with the alloy powders in a bi-axis rotary mixer for 8 h with a ball to powder ratio of 2:1 and a rotation speed of 50 rpm. The asmixed powders were cold pressed into the green billets under a pressure of 50 MPa. Since interfacial reactions become severe with the appearance of liquid phase above 600 C [21], the hot-pressing temperatures of 560 C, 600 C, and 630 C were chosen to fabricate three composites (designated as HP 560, HP 600, and HP 630 hereafter, respectively). Hot pressing was carried out in a vacuum chamber of 102 Pa. The billets were heated up to the preset temperature with a heating rate of 5 C/min, and then held for 2 h, after which hot pressing was carried out under a pressure of 40 MPa. The billets were hot-rolled into plates at 450 C with a height reduction ratio of 10% for each pass. A total height reduction ratio of 90% was attained. The densities of the composite samples were measured by using the Archimedes principle. To test the long-term thermal stability,
HP 630 were annealed uninterruptedly at 400 C for up to 8000 h in a muffle furnace and then cooled inside the furnace. The microstructural studies were carried out by means of optical microscopy (OM, Zeiss Axiovert 200MAT), field-emission scanning electron microscopy (SEM, Leo Supra 55) and transmission electron microscopy (TEM, TECNAI G2 F20) equipped with high-angle annular-dark-field (HAADF) detectors. The chemical compositions of the secondary phases were analyzed by X-ray energy-dispersive spectroscopy (EDS) under SEM and HAADF-mode TEM. TEM specimens were prepared by mechanical polishing and ion milling using Leica RES101 system. Tensile specimens were machined from the as-rolled composite sheets along the rolling direction. For room temperature (RT) tension, Instron 8810 equipped with an extensometer was used to obtain accurate yield stress and elongation, and specimens with a gauge length of 15 mm, a width of 3 mm and a thickness of 2.5 mm were used. For high temperature tests, tensile specimens (5 mm in length, 1.5 mm in width and 1.2 mm in thickness) were tested on Instron 5582 equipped with a heating furnace. At least 5 tensile specimens were tested for each composite at a strain rate of 103 at both RT and high temperatures.
3. Results 3.1. Densities and microstructures Table 1 shows densities of the fabricated composites. The results indicate that all composites were almost completely densified and the densities increased slightly with increasing hot-pressing temperatures. Optical microscopic analysis of as-rolled composites is shown in Fig. 1. B4C particles were homogenously distributed in these composites and no micro-pores were observed. Compared to the ‘clear’ matrix in HP 560 and HP 600 (Fig. 1a and b, respectively), the alloy matrix in HP 630 (Fig. 1c) was decorated by high-density of fine dispersoids as marked by arrows in the inset, which were generated by the interfacial reactions. The characterization of the dispersed particles will be shown later. The XRD patterns of the samples hot-pressed at different temperatures are shown in Fig. 2. The patterns were normalized by adjusting the peaks of B4C to the same height. In HP 560 and HP 600, the peaks of Mg2Si were found besides the peaks of Al and B4C. In the latter sample, the reaction products, Mg(Al)B2 and Al3BC were also detected. When the hot-pressing temperature was increased to 630 C, the interfacial reactions became severer and the quantities of Mg(Al)B2 and Al3BC increased. Due to the consumption of element Mg by the interfacial reactions, the Mg2Si precipitate was too less to be detected. MgAl2O4 spinel which was frequently seen in Mg-containing aluminum matrix composites [24,25] was present in these three samples. The Mg(Al)B2 and Al3BC in different samples were quantified based on the relative diffraction intensity of Al (111), B4C (104), Al3BC (101) and Mg(Al)B2 (100) planes. The relative intensity fractions of the Mg(Al)B2 and Al3BC in HP 560, HP 600 and HP 630 were calculated by the following equation used in Ref. [21].
Table 1 Densities of B4Ce6061Al composites with various hot-pressing temperatures. Sample 3
Density (g/cm ) Relative density (%)
HP 560
HP 600
HP 630
2.637 99.5
2.648 99.9
2.658 100.3
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Fig. 1. Optical micrographs of the samples (a) HP 560, (b) HP 600 and (c) HP 630. The B4C particles (black) homogenously distribute in all of the three samples. The fine dispersoids in HP 630 were denoted by black arrows in the image inset.
Fig. 2. XRD patterns of the as-rolled B4Ce6061Al composites with different hotpressing temperatures.
DX ð%Þ ¼
IX ð%Þ IAl þ IX þ IB4 C
(1)
where DX is the relative intensity fraction of compound X (Mg(Al)B2 or Al3BC), Ix is the diffraction intensity of X, IAl and IB4C are the intensities of Al and B4C, respectively. Calculated results are listed in Table 2. It is revealed that the interfacial reactions were enhanced by increasing the hot-pressing temperatures, resulting in the formation of more Mg(Al)B2 and Al3BC. The secondary phases in these samples were further analyzed by SEM and TEM, as shown in Fig. 3. In HP 560, sharp and clean edges of B4C particles were observed (Fig. 3a). Some dark particles (as denoted by arrows) were found in the sample and also in HP 600 (shown in Fig. 3b). In HP 600, some of the particles have grown to larger scales, i.e. about 2e4 mm in diameter. The chemical composition analyzed by EDS mapping indicated that they contained element Mg and Si. TEM characterization and selected area electron diffraction (SAED) were carried out, identifying these particles (Fig. 3d) to be b-Mg2Si (lattice constant a ¼ 0.45 nm). In addition, Table 2 The relative intensity fractions of the reaction products calculated based on Al (111), B4C (104), Al3BC (101) and Mg(Al)B2 (100) in the composites. Sample
HP 560 HP 600 HP 630
Relative fraction (%) Al3BC
Mg(Al)B2
0.12 0.55 0.78
0.20 0.76 1.55
the serrated B4C-particle edges in HP 600 implied that the interface reactions have occurred. When the hot-pressing temperature was increased to 630 C, the reactions between B4C and the alloy matrix became more apparent and a high density of secondary phases were detected (Fig. 3c). EDS analysis suggested that the grey phases (as denoted by yellow arrow) were mainly composed of Mg and Al. Referring to the XRD result, they were speculated to be Mg(Al)B2. These phases were further confirmed by SAED and EDS analyses performed by TEM, as shown in Fig. 3e. The EDS signal was obtained from the compound located at the very edge of the perforation of the TEM foil to avoid interference from overlapped aluminum matrix. It should be pointed out that AlB2 and MgB2 crystals have the same crystallographic characteristics but a little variation in lattice parameters [26,27]. The compound in our samples probably contained different contents of metallic (Al or Mg) atoms. Another reaction product Al3BC was also frequently seen at the interface of B4C particles and Al matrix, as illustrated in the TEM image in Fig. 3f. To compare the reaction product Mg(Al)B2 in different samples, Fig. 4 shows the micrographs of Mg(Al)B2 in HP 600 and HP 630. In both the samples, the Mg(Al)B2 nanorods were dispersed uniformly in the alloy matrix. As marked by black arrows, the Mg(Al)B2 dispersoids in HP 600 were measured to be about 100 nm in diameter and 0.5e2 mm in length. It is clearly seen that the quantity of Mg(Al) B2 in HP 630 was much higher than that in HP 600. A part of the Mg(Al)B2 rods grew to larger sizes, so became visible under optical microscope (Fig. 1c). Fig. 5aec shows low-magnification TEM images of the three composites. A fine grained structure with grain sizes of 0.5e2 mm was observed for all the composites. The observations agreed with the EBSD analysis in our previous work [23]. The fine grains should be generated during the hot-rolling process through dynamic recrystallization. Thus, the grain sizes in these composites fabricated at different hot-pressing temperatures were in the same range. The dislocation lines could be visible under proper diffractioncontrast conditions. Fig. 5def shows the images of dislocations in the grains. The dislocation densities in these different composites were roughly estimated to be of the same order of magnitude (~1013 m2). Threading dislocation lines throughout the grain interior were frequently observed in HP 560 (Fig. 5a). In contrast, the dislocations in the composites hot-pressed at higher temperatures (Fig. 5b and c) showed more complex and tangled feature owing to the interactions with the dispersoids inside the grains (arrowed in Fig. 5c). The dislocation densities in HP 600 and HP 630 were higher than those in HP 560. It was also seen that dislocation accumulation easily developed into sub-grain boundaries during hot-rolling process, as indicated in Fig. 5a.
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Fig. 3. Microstructural analysis of the samples. (a) SEM image of HP 560. (b) SEM image of HP 600. The inset is the elemental mapping of a Mg2Si particle. (c) SEM image of HP 630 in which the Mg(Al)B2 and Si particles can be found in the matrix. (d) TEM image and corresponding SAED pattern of the Mg2Si particle in HP 560. (e) HAADF image, SAED pattern and the EDS result of Mg(Al)B2 nanorods as interfacial reaction products. (f) Al3BC compound can be frequently found at the periphery of B4C particles in both HP 600 and HP 630.
Fig. 4. HAADF TEM images showing the Mg(Al)B2 distributions in (a) HP 600 and (b) HP 630.
3.2. Mechanical properties at different temperatures The stress-strain curves of the composites are shown in Fig. 6. It can be seen that the samples fabricated at different hot-pressing temperatures showed significant differences in strength and ductility at RT (Fig. 6a). The yield strength (YS) and the ultimate stress (UTS) of the composites increased with increasing hotpressing temperatures. The YS and UTS values of HP 560 were only 120 MPa and 198 MPa, respectively. A simultaneous enhancement of strength and elongation (EL) was obtained for HP 600, with the YS, UTS and EL values respectively reaching 166 MPa, 254 MPa and 12%. Obviously, the appearance of interfacial reactions improved the mechanical properties of the composite. For HP 630, the strength was further increased to 188 MPa (YS) and 315 MPa (UTS), which were about 57% and 59% larger than that of HP 560. However, the elongation declined to 7%. Fig. 6b shows the stress-strain curves of different composites measured at temperatures from 150 C to 350 C. Since no extensometer was used for high-temperature testing, only UTS values of the composites were labelled next to the curves. At high temperatures, strength enhancement of the composites with increasing hot-pressing temperatures was still confirmed. The UTS values of HP 630 were about 40%e50% higher than those of HP 560 at the testing temperatures.
We further tested the mechanical stability of HP 630 suffered long-term thermal exposure at 400 C. The UTS values tested at 150 C and 350 C are shown in Fig. 7. It is seen that the tensile strength of HP 630 experienced as long as 8000 h annealing remained in the same level. For the commercially used precipitation-strengthened aluminum alloys, such extraordinary stability is hardly achieved. 3.3. Fractographs The fracture surface of the composite was characterized using SEM. Ductile fracture was observed at both 150 C and 350 C for HP 560 and HP 630, indicated by the dimple structure. No pull-out of the B4C particles from the matrix was found, implying that the B4C particle bonded well with the alloy matrix. For fractographs of both HP 560 (Fig. 8a) and HP 630 (Fig. 8b) deformed at 150 C, the morphological difference in between is quite noticeable. The larger and deeper dimples in HP 560 suggested the higher ductility of the composite. An enlarged area of the image showed that small particles could be found in the bottom of a dimple (inset in Fig. 8a). As revealed by the microscopic analysis, the Mg2Si particles in HP 560 acted as the stress concentration sites due to their brittle nature and easily resulted in local micro-void formation under external loading. For the case of HP 630 (Fig. 8b), the high density of in-situ
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Fig. 5. TEM images showing the microstructures of the aluminum alloy matrices. (a)e(c) Low-magnification TEM images of the fine grains in HP 560, HP 600 and HP 630, respectively. (d)e(f) TEM images of the dislocations inside the grains in HP 560, HP 600 and HP 630, respectively.
Fig. 6. Tensile stressestrain curves of the composites at (a) RT and (b) elevated temperatures.
formed particles via the interfacial reactions rendered higher strength of the sample, but also more sites for stress concentration and cracks nucleation. Therefore, the fracture surface showed more dimples with smaller sizes. When the testing temperature was increased to 350 C, the dimples for both composites exhibited larger size, which was in accordance with the higher EL values. The secondary phases were also found on the fracture surfaces, especially for HP 630. To identify these particles, EDS analysis was carried out (Fig. 8e). By comparing the elementary composition of matrix (e1) and particle (e2), in can be concluded that the particles were Mg2Si (Fig. 8c). In the case of HP 630 (Fig. 8d), EDS analysis results are shown in Fig. 8f. According to above microstructure analysis and the EDS profile of Al matrix (f1), the fine particles (f2) were identified as Mg(Al)B2. The fractograph at 350 C suggested that the nucleation of microvoids at the dimple bottoms was also closely related to the secondary phases in the composites.
4. Discussion It is well known that the chemical reactions between B4C and Al obviously happen when the liquid phase appear in the alloy matrix. For 6061Al, the liquid temperature was determined as about 600 C (Tm) [21]. Alloying element Mg in this composite is noteworthy. Below Tm, the interfacial reactions were quite mild. As determined by above results, Mg existed in the forms of Mg2Si in the matrix. When hot-pressing temperature was increased to 600 C, a small portion of liquid aluminum reacted with B4C, yielding Al3BC and Mg(Al)B2 nanorods. In HP 630, a large amount of liquid phase facilitated the interfacial reactions. Because of the larger chemical affinity between Mg and B, the reactions consumed most of element Mg. At the same time, element Si could not form compound with other elements, so they would precipitate as Si or AleSi eutectic phases in the cooling process. The strengthening effects of introducing particles are usually thought as direct strengthening by load-transfer (L-T) and
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Fig. 7. The values of ultimate tensile strength of HP 630 experienced various annealing periods. It shows the high stability of the composite at elevated temperature.
secondary strengthening relevant to dislocation activities, including Orowan strengthening, geometrically necessary dislocation strengthening, etc [23,28,29]. These mechanisms well explained the RT strength enhancement of HP 600 and HP 630 relative to that of HP 560. At elevated temperatures, however, the thermal energy is sufficient to activate diffusion, leading to the low efficiency of the above strengthening mechanisms [30,31]. To gain insight into the role played by Mg(Al)B2 nanorods during the high-temperature deformation process, TEM observations were performed on the sample taken from the fractured HP 630. Fig. 9a exhibits the bright-field TEM image in which several grains and dislocation tangles could be seen. Compared to that in the as-rolled sample displayed in Fig. 5c, a high-density of dislocations was found to accumulate at the periphery of GBs in the deformed sample, suggesting the impeding effect of GBs to the dislocation movement. More importantly, as the arrows denoted, several Mg(Al)B2 nanorods were also observed in front of GBs. It is known that GBs will lose their strength at elevated temperatures. GB sliding largely assists the high-temperature deformation of alloys [32]. The pinning effect provided by fine particles has been proved to be an effective way to hinder the GB motion [33,34]. Here the Mg(Al)B2 nanorods located at the GBs clearly indicated their roles
Fig. 8. SEM fractographs of the composites: (a) HP 560 deformed at 150 C and (b) HP 630 at 150 C. (c) HP 560 at 350 C and (d) HP 630 at 350 C. (e) and (f) are EDS results of the marked positions in (b) and (e), respectively.
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Fig. 9. TEM images of the microstructures in the fractured sample HP 630. (a) The grain boundaries are pinned by Mg(Al)B2 nanorods (denoted by white arrows) (b) The dislocations tangled around the Mg(Al)B2 suggests its impeding effect to dislocation activity.
of pinning the GBs. GB sliding requires accommodation by diffusion and the dislocation/GB interactions [35]. The dislocation accumulations in front of GBs will be released by GB sliding as a result. In such way, it is not easy to form stress concentration and the material is softened. In the presence of fine particles in the alloy matrix, the GB strength is apparently enhanced by the pinning of reinforcement. The dislocation pile-ups in Fig. 9a confirmed the stability of GBs in the present composite. Consequently, the high-temperature strength of the material can be largely improved. With increasing loading, however, the stress concentration at the GBs tends to result in the micro-void nucleation at the local sites and finally the fracture failure. The analysis is quite consistent with the fractographs shown in Fig. 8, i.e. the voids nucleated at the GBs where a massive of Mg(Al)B2 particles were exposed after fracture. On the other hand, the Mg(Al)B2 nanorods also served as obstacles for the dislocation motion. In general, the material strength at ambient temperature can be effectively improved by fine particles via Orowan bypassing mechanism. The increasing temperature can activate vacancy diffusion and the dislocation climbing weakens the strengthening effect by Orowan bypassing mechanism. Nevertheless, the high strain energy induced by the mismatch between reinforcements and Al matrix (modulus mismatch and lattice mismatch) also provides a certain barrier for dislocations to climb over [36,37]. According to Qin et al., the strength of the composite at elevated temperatures was controlled by the combination of dislocation climb and Orowan models [38]. One of the Mg(Al)B2 nanorods in Fig. 9a were extracted and displayed in Fig. 9b. The dislocations tangled around the nanorod implied the interactions in between. The hindrance to dislocation activities increases the yielding strength of the composite to some extent, although it is less effective than that at ambient temperature.
5. Conclusions 1. B4Ce6061Al composites were fabricated by the PM method with varying hot-pressing temperatures. Mg2Si is the main precipitate in the composite hot-pressed at lower temperatures of 560 C and 630 C. The composite produced at a higher hotpressing temperature of 630 C contained plenty of homogeneously dispersed Mg(Al)B2 nanorods.
2. The interfacial reactions are beneficial to improve the material strength at both room and elevated temperatures. The yield strength of HP 630 at RT is 57% higher than that of HP 560. Tensile strengths of HP 630 at high temperatures (150 C ~ 350 C) were increased by 40% ~ 50%. In addition, the mechanical properties of HP 630 were proven to be extremely stable after being annealed at 400 C for 8000 h. 3. The Mg(Al)B2 nanorods acted as obstacles to the GB sliding and the dislocation activities. As a consequence, the hightemperature strength of the composite was greatly enhanced.
Data availability The processed data required to reproduce these findings cannot be shared at this time as the data also forms part of an ongoing study. Acknowledgments The authors gratefully acknowledge the support of the National Natural Science Foundation of China (NSFC) under grant Nos. U1508216, 51501195 and 51771194, NSF of Liaoning Province [20180551101] and the Innovation Fund of IMR [2017-PY10]. Appendix A. Supplementary data Supplementary data to this article can be found online at https://doi.org/10.1016/j.jnucmat.2019.151788. References [1] X.-G. Chen, Application of Al-B4C metal matrix composites in the nuclear industry for neutron absorber materials, in: N. Gupta, W.H. Hunt (Eds.), Solidification Processing of Metal Matrix Composites, TMS, USA, 2006, pp. 343e350. [2] H.S. Chen, W.X. Wang, Y.L. Li, J. Zhou, H.H. Nie, Q.C. Wu, The design, microstructure and mechanical properties of B4C/6061Al neutron absorber composites fabricated by SPS, Mater. Des. 94 (2016) 360e367. [3] L. Zhou, C. Cui, Q.Z. Wang, C. Li, B.L. Xiao, Z.Y. Ma, Constitutive equation and model validation for a 31vol.% B4Cp/6061Al composite during hot compression, J. Mater. Sci. Technol. 34 (2018) 1730e1738. [4] Z.G. Xu, L.T. Jiang, Q. Zhang, J. Qiao, D. Gong, G.H. Wu, The design of a novel neutron shielding B4C/Al composite containing Gd, Mater. Des. 111 (2016) 375e381. [5] P. Zhang, Y. Li, W. Wang, Z. Gao, B. Wang, The design, fabrication and
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