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Ductility of Mo–12Si–8.5B alloys doped with lanthanum oxide by the liquid–liquid doping method
Journal of Alloys and Compounds 642 (2015) 34–39
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
Ductility of Mo–12Si–8.5B alloys doped with lanthanum oxide by the liquid–liquid doping method Wenhu Li a,b, Guojun Zhang a,⇑, Shixiong Wang a, Bin Li c, Jun Sun c a
School of Materials Science & Engineering, Xi’an University of Technology, Xi’an 710048, China School of Materials Science & Engineering, Shaanxi University of Technology, Hanzhong 723000, China c State Key Laboratory for Mechanical Behavior of Materials, Xi’an Jiaotong University, Xi’an 710049, China b
a r t i c l e
i n f o
Article history: Received 24 December 2014 Received in revised form 4 April 2015 Accepted 6 April 2015 Available online 9 April 2015 Keywords: Mo–Si–B alloy Lanthanum oxide Liquid–liquid doping Fine-grain strengthening
a b s t r a c t Mo–12Si–8.5B (Mo–Si–B) alloys doped with different mass fractions (0.3 wt%, 0.6 wt%, and 0.9 wt%) of lanthanum oxide (La2O3) were prepared by liquid–liquid (L–L) doping, mechanical alloying and hot pressing sintering techniques. The observation of the microstructures of the Mo–Si–B alloys reveals that the grain sizes of the alloys were refined with the increase in La2O3 doping. The fracture toughness values of the alloys of over 10 MPa m1/2 reveal that the addition of La2O3 via the L–L doping method can obviously improve the alloy fracture toughness compared to the alloys doped with La2O3 via the solid–solid (S–S) doping method. In addition, compression tests indicate that the compression strength of the alloys was improved compared to Mo–12Si–8.5B alloys. Ó 2015 Elsevier B.V. All rights reserved.
1. Introduction Due to the high melting point, low thermal expansion coefficient, excellent strength and high temperature oxidation resistance, Mo–Si–B alloy has attracted considerable interest and has been regarded as a potential structural material in high temperature applications for decades [1–3]. Currently, research studies of Mo–Si–B alloys are mainly concentrated on the three phase alloys of Mo3Si + Mo5Si3 + Mo5SiB2 and a-Mo + Mo3Si + Mo5SiB2 (T2) [3]; the different compositions of Mo–Si–B alloy are found to have different properties. Mo3Si + Mo5Si3 + Mo5SiB2 alloy has good high temperature oxidation resistance because the Mo5Si3 phase can provide the Si atoms required to form a protective coating (SiO2); however, it has poor toughness at room temperature [4]. aMo + Mo3Si + Mo5SiB2 (T2) alloy has better toughness as a result of the existence of the a-Mo phase, which is favourable for room temperature toughness [5]. In contrast, relative to the Mo alloy, the fracture toughness of Mo–12Si–8.5B has much room to improve. Therefore, new methods to improve the integrated property of Mo–Si–B alloy are required. As is known, strength and fracture toughness are the two important mechanical properties of alloys, and one of the means to improve the properties of an alloys is to dope it with rare earth oxide. Numerous studies have been performed reporting that ⇑ Corresponding author. Tel.: +86 29 82312592. E-mail address: [email protected] (G. Zhang). http://dx.doi.org/10.1016/j.jallcom.2015.04.047 0925-8388/Ó 2015 Elsevier B.V. All rights reserved.
adding rare earth oxide in Mo alloy can hinder the growth of grain boundaries and refine the grain size. These effects of doping can improve the strength and fracture toughness of Mo alloys [6,7]. In Zhang’s research, doping La2O3 into Mo alloys was found improve the fracture toughness values of the Mo Alloys from 9.9 MPa m1/2 up to 24.76 MPa m1/2 [8]. The room temperature fracture toughness of Mo–12Si–8.5B alloy was on the order of 5– 7 MPa m1/2 [9]. When the Mo–Si–B alloy was doped with a certain amount of La2O3, the fracture toughness of the Mo–Si–B alloy increased up to 9 MPa m1/2 [10,11]. Moreover, Mo–Si–B alloys with fine-grained microstructure have excellent oxidation resistance, as reviewed by Rioult [12]. Thus, doping Mo alloys with rare earth oxide is an effective method to improve the properties of Mo alloys. The typical methods to dope rare earth oxides into Mo alloy include solid–solid doping, solid–liquid doping and liquid–liquid doping. One of the most commonly used methods is solid–solid doping, but the studies of Liu et al. indicated that a molybdenum alloy doped with La2O3 using the liquid–liquid doping method has a smaller grain size and a smaller rare earth oxide particle size than such alloys doped using the traditional solid–liquid doping and solid–solid doping methods. The high temperature stability of the rare earth oxides allows the alloy to maintain its grain size, exhibiting no growth under high temperature [13]. Although much research exists on Mo–Si–B alloys doped with lanthanum oxide, little is known about the Mo–12Si–8.5B alloys doped with lanthanum oxide via the liquid–liquid doping method. In this paper,
W. Li et al. / Journal of Alloys and Compounds 642 (2015) 34–39
Fig. 1. XRD patterns of Mo–12Si–8.5B alloys with different mass fractions of La2O3.
we selected Mo–12Si–8.5B and used liquid–liquid doping, mechanical alloying and hot pressing sintering techniques to prepare the alloys. In addition, we studied the influence of the use of the liquid–liquid doping method on the mechanical properties of the alloys by comparing the mechanical properties of the Mo–Si–B alloys doped with different mass fractions of La2O3.
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powder to ball weight ratio of 1:10 for 15 h using a Retsch PM 400 apparatus. After mechanical alloying, the powders were hot-pressed at 50 MPa and then sintered in vacuum at 1600 °C for 2 h. The crystalline phases of the alloys were determined using an X-ray diffractometer (XRD) equipped with a Cu Ka radiation source. Microstructural analyses of the alloys were performed via optical metallography (GX71 of OLYMPUS) and scanning electron microscopy (SEM) (JSM-6700F). The specimens were polished and etched with Murakami’s etch (an aqueous solution of potassium ferricyanide and sodium hydroxide). The microstructure of the alloy was also observed via transmission electron microscopy (TEM). The foils for TEM observations were fabricated using a twin-jet apparatus with a solution of 5 vol% H2SO4 in ethanol at 15 °C. A voltage of 10 V produced a current density of 25 mA. The hardness was tested on a TUKON 2100 micro-Vickers’s hardness tester, and the result of every sample was the average value of five points. The densities of the samples were measured using the Archimedes technique. The room temperature compression tests were performed on a HT-2402 computer-controlled machine with a strain rate of 5 103 s1. The high temperature compressive was tested at Gleeble 1500D numerical control thermodynamic testing machine. The sample size was 5 mm 5 mm 10 mm. The fracture toughness values were determined via three-point bending tests. The specimens, which were a cross of 3 mm 4 mm 26 mm with a notch tip depth of 2 mm, were tested in a universal material testing machine (Instron 1185) with a 20-mm span at a crosshead speed of 0.05 mm/min. The fracture toughness values were determined on the basis of an energy criterion by integrating the load–displacement curves. The fracture toughness was determined as G = W/A, where W is the energy absorbed during the fracture of the area A swept out by the crack. Assuming the material to be linear elastic, the fracture energy can also be expressed in terms of the stress intensity by Kq = (G E0 )1/2 where E0 = E/(1 m2) is the plane strain Young’s modulus, and m is Poisson’s ratio (E = 327 GPa and m = 0.29) [1]. In this paper all samples were electro-discharge machined, ground, and then polished before testing. All of the experiment results are average values of at least three samples.
2. Experimental procedures The Mo–12Si–8.5B alloys of different mass fractions of La2O3 were fabricated via mechanical alloying and hot pressing. Molybdenum alloy powders containing different mass fractions of La2O3 particles were first prepared using the L–L doping method. In the L–L doping [13], first, La2O3 is dissolved into an aqueous solution of nitrate (La(NO3)3), and then the solution is mixed with ammonium dimolybdate solution ((NH4)2Mo2O7). Mo–La2O3 powders are subjected to drying roasting and then restored to the mixed solution. Next, Mo–La2O3, Si and B powders are mixed in a planetary ball mill with a speed of 300 rpm and a powder to ball weight ratio of 1:1 for 6 h. The elemental powders of Si and B were 99.99 wt% and 99.99 wt% purity, respectively. Mechanical alloying (MA, a high-energy milling process) was performed under a protective atmosphere (argon) at a speed of 250 rpm and a
3. Results and discussion 3.1. XRD analysis and microstructure characterisation The results of XRD analysis are presented in Fig. 1. These alloys are all composed of three phases of a-Mo, Mo3Si and T2. No other phases were observed, indicating that the phase composition is not affected by either the content of La2O3 or the doping method.
Fig. 2. Optical micrographs of the alloys: (a) Mo–12Si–8.5B + 0.3 wt%La2O3; (b) Mo–12Si–8.5B + 0.6 wt%La2O3; (c) Mo–12Si–8.5B + 0.9 wt%La2O3.
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W. Li et al. / Journal of Alloys and Compounds 642 (2015) 34–39
Moreover, the intensity of the diffraction peaks of the a-Mo phase are higher than those of the Mo3Si and T2 phases. Representative micrographs of the alloys doped with different mass fractions of La2O3 are provided in Fig. 2. The bright phase is the a-Mo matrix, which has the highest volume fraction. The grey phase and dark phase, which are dispersed within the continuous a-Mo matrix, are Mo3Si and T2, respectively. In addition, from Fig. 2, the intermetallic (Mo3Si/T2) phases are also observed to be distributed within the a-Mo matrix homogeneously. The microstructure and grain size of the alloys were investigated using TEM, as shown in Fig. 3. The TEM images clearly show that the grain sizes of the alloys are submicron. The a-Mo, Mo3Si and T2 phases were determined by analyses of the corresponding
selected electron diffraction area. The La2O3 particles, which are of nanometre size and are homogeneously dispersed in the grain interior, are highlighted in Fig. 3(e) and (g) by the solid arrows. The nanoscale particles distributed inside of the a-Mo and the intermetallic regions will be beneficial for both the strength and the toughness of the alloys. The observed grain sizes and statistical results for the alloys are shown in Table 1. The grain size, which is approximately 0.6 lm in the undoped alloys, is reduced with the addition of La2O3. In Zhang’s study [13], more La2O3 particles become homogeneously distributed in the grain interiors than in the grain boundaries when using the L–L doping process compared to the other doping processes. Because the La2O3 particles can act as nucleation sites (which can increase the nucleation rate during
Fig. 3. TEM images of the microstructural morphologies of the alloys (a) Mo3Si and T2 phase; (b–d) are the corresponding selected area diffraction patterns in (a); (e and f) show the images of La2O3 particles in a-Mo; (g and h) show the images of La2O3 particles in T2; (i and j) are the images of the a-Mo matrix.
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W. Li et al. / Journal of Alloys and Compounds 642 (2015) 34–39
Fig. 3 (continued)
the sintering process), this homogeneous distribution is the reason that the grain size in the L–L doping method is smaller than that in the S–S doping method. Additionally, some La2O3 particles may be present at the molybdenum grain boundaries to hinder the grain growth via the pinning effect. Both the pinning and the nucleation effects in La2O3-doped alloys decrease the grain size. 3.2. Fracture toughness The fracture toughness values of the alloys are shown in Fig. 4. All of the Kq values of the alloys are higher than that of the Mo– 12Si–8.5B alloy (9.132 MPa m1/2). Compared with Zhang’s study [10], these alloys prepared via the L–L doping process exhibit higher room temperature fracture toughness than the alloys prepared via the S–S doping process. SEM micrographs of the room temperature fracture surfaces are shown in Fig. 5. The images reveal that the principal failure mechanism is the mixed-mode fracture. Some pullout particles were also observed in the fracture surface. The different elastic modulus and yield strength of the aMo, Mo3Si and Mo5SiB2 phases are responsible for the mixed– mode failure. The transgranular fracture included the cleavage fracture of the larger-size intermetallic particles and of the larger-size a-Mo particles. The toughness of a-Mo is known to be better than that of the intermetallic phase [14,15]. As shown in
Fig. 5(b), the a-Mo phase regions exhibit obvious plastic deformation during fracture. Thus, the transgranular fracture of the ductile a-Mo phases was beneficial for the improvement of the alloy toughness. The toughening mechanisms include fine grain size toughening and particle toughening. The toughening mechanisms of the alloys mainly include finegrain toughening and particle toughening. Material toughness can be effectively improved by refining the grains. On the one hand, cracks are caused by large plastic deformation and mostly occur in grain boundaries and phase boundaries. The grain refinement leads to the increase per unit volume of the grain boundary area. Because the grain boundary atoms are arranged in a disordered manner and the dislocation structure is complex, when the deformation moves from a grain to an adjacent grain by grain boundary, the resistance of dislocation movement and the difficulty of the crack generation will be increased, thus increasing the consumption of energy. On the other hand, the increase in the grain boundary area can improve its effect of blocking the expansion of cracks. When a crack propagates to the grain boundary, different grain orientations on both sides of the grain boundary force the crack extension to change direction or stop. As a result, the alloy toughness is improved. Particle toughening consists of two aspects. The La2O3 particles during liquid–liquid doping of the molybdenum alloy are
Table 1 Grain sizes of the alloys with different mass fractions of La2O3. Alloy composition
Mo–12Si–8.5B + 0.3 wt%La2O3
Mo–12Si–8.5B + 0.6 wt%La2O3
Mo–12Si–8.5B + 0.9 wt%La2O3
Grain sizes (lm)
0.65
0.64
0.59
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W. Li et al. / Journal of Alloys and Compounds 642 (2015) 34–39
Fig. 4. The fracture toughness of Mo–12Si–8.5B alloys of different mass fractions of La2O3.
Fig. 6. Compression stress–strain curves of Mo–12Si–8.5B alloys doped with different mass fractions of La2O3 at ambient temperature.
material increases. Moreover, La2O3 particles refine the grain size of the matrix. When the crack extends to the La2O3 particles, each hard La2O3 particle leads to cracking deflection. As a result, a longer extension path is produced. The resulting energy consumption leads to toughening of the alloy. In summary, the improvement of fracture toughness in Mo alloys is due to the fine grain toughening and particle toughening via La2O3 introduced using the L–L doping method.
3.3. Compression
Fig. 5. SEM images of the fracture surfaces of the alloys: (a) Mo–12Si– 8.5B + 0.6 wt%La2O3; (b) Mo–12Si–8.5B + 0.9 wt%La2O3.
distributed both in the grains and at the grain boundaries. The second phase particles, regardless of being located in the grains or at the grain boundaries, must produce micro-cracks when the dislocation accumulates during the deformation process. This formation of micro-cracks tends to separate the matrix and the particles. Once micro-cracks are formed in the matrix (as indicated by the arrow in Fig. 5(a)), the stress surrounding the micro-cracks is relaxed. At the same time, the new surface formed will absorb a large amount of energy; consequently, the toughness of the
The compression stress–strain curves of the alloy at room temperature are shown in Fig. 6. No obvious yield point is observed, and the stress is observed to continue to increase until the material failed. The sample failed without plastic deformation. The behaviour of the alloy corresponds to the typical room temperature compression curve for brittle materials. The results of the compression strength measurements are listed in Table 2. The room temperature compression strengths of the alloys are improved with the addition of La2O3, and the room temperature compression strengths are also increased with the increase in the La2O3 content. The Mo–12Si–8.5B + 0.9 wt%La2O3 alloy exhibits the highest compression strength (2.97 GPa) compared with the other alloys. In Zhang’s work [10], the compression strength of the Mo–12Si–8.5B + 0.9 wt%La2O3 alloy, in which La2O3 was doped using the solid–solid doping process, was only 2.7 GPa. This result indicates that the alloy that is doped using the L–L doping process has a strength higher than the same composition alloy doped using the solid–solid doping process. As presented in Table 2, the doped alloys have high compression strength at high temperature. At 1000 °C, the Alur alloy had a yield stress of only 320–790 MPa [16]. In the L–L doping case, the La2O3 particles are homogeneously dispersed in the alloys; as a result, the numerous nanoscale particles interact with the migrating grain boundaries,
Table 2 The compression strengths of La2O3 doped Mo–12Si–8.5B alloys at 25 °C, 1000 °C and 1300 °C. Alloy composition
Mo–12Si–8.5B + 0.3 wt%La2O3 Mo–12Si–8.5B + 0.6 wt%La2O3 Mo–12Si–8.5B + 0.9 wt%La2O3
Compression strength (MPa) 25 °C
1000 °C
1300 °C
2565 2788 2872
1604 1539 1414
526 429 376
W. Li et al. / Journal of Alloys and Compounds 642 (2015) 34–39
causing the alloy grains to remain at small sizes of only 0.5 lm [13]. This refined microstructural morphology contributes to the high strength of the alloy.
4. Conclusions (1) The La2O3 doped Mo–12Si–8.5B alloys are composed of the a-Mo, Mo3Si and Mo5SiB2 phases. The intermetallic Mo3Si and Mo5SiB2 phases are distributed within a continuous aMo matrix. With doping of La2O3 into the alloys, the grain size and the intermetallic particles size of the alloys were refined and distributed more homogeneously. (2) The compression strength of Mo–12Si–8.5B alloys doped with La2O3 using the L–L doping method is improved to 2.87 GPa. In addition, the compression strength is increased with the addition of La2O3; in this case, the major strengthening mechanisms are the strengthening of fine grains and increased particle dispersion. (3) The fracture mode of the alloy is brittle cleavage fracture. The fracture toughness of doped Mo–12Si–8.5B alloys is improved by using the L–L doping method. The primary toughening mechanisms are toughening of fine grains and toughening of particles.
Acknowledgements This subject was supported by the National Natural Science Foundation (Grant Nos. 51171149 and 51371141), the National Science Technology Supporting Program of China (Grant No. 2012BAE06B02).
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References [1] J.H. Schneibel, M.J. Kramer, O. Unal, R.N. Wright, Processing and mechanical properties of a molybdenum silicide with the composition Mo–12Si–8.5B, Intermetallics 9 (2001) 25–31. [2] H. Choe, D. Chen, J.H. Schneibel, R.O. Ritchie, Ambient to high temperature fracture toughness and fatigue-crack propagation behavior in a Mo–12Si–8.5B (at.%) intermetallic, Intermetallics 9 (2001) 319–329. [3] J.J. Huebsch, M.J. Kramer, H.L. Zhao, et al., Solubility of boron in Mo5+ySi3y, Intermetallics 8 (2) (2000) 143–150. [4] C.T. Liu, J.H. Schneibel, L. Heatherly, High-temperature ordered intermetallic alloys VII, in: MRS. MRS Symposium Proceedings, Warrendale, MRS, 1999, pp. 621–627. [5] J.H. Schneibel, M.J. Kramer, D.S. Easton, A Mo–Si–B intermetallic alloy with a continuous a-Mo matrix, Scripta Mater. 46 (2002) 217–221. [6] G.J. Zhang, Y.J. Sun, C. Zuo, et al., Microstructure and mechanical properties of multi-components rare earth oxide-doped molybdenum alloys, Mater. Sci. Eng. A 384 (15) (2008) 350–352. [7] B. Xu, D.Z. Wang, Z.Z. Wu, A.K. Sun, Q.J. Cheng, X.Q. Zan, Preparation and properties of sintered molybdenum doped with La2O3/MoSi2, Int J Refract Met Hard Mater 28 (2010) 150–154. [8] J.X. Zhang, L. Liu, M.L. Zhou, Y.C. Hu, T.Y. Zuo, Fracture toughness of sintered Mo–La2O3 alloy and the toughening mechanism, Int. J. Refract. Met. Hard Mater. 17 (1999) 405–409. [9] H. Choe, J.H. Schneibel, R.O. Ritchie, On the fracture and fatigue properties of Mo–Mo3Si–Mo5SiB2refractory intermetallic alloys at ambient to elevated temperature (25 °C to 1300 °C), Metall. Mater. Trans. A 34 (2003) 225–239. [10] G.J. Zhang, Y. Zha, B. Li, W. He, J. Sun, Effects of lanthanum oxide content on mechanical properties of mechanical alloying Mo–12Si–8.5B (at.%) alloys, J. Refract. Met. Hard Mater. 41 (2013) 585–589. [11] G.J. Zhang, Q. Dang, H. Kou, R.H. Wang, G. Liu, J. Sun, Microstructure and mechanical properties of lanthanum oxide-doped Mo–12Si–8.5B (at.%) alloys, J. Alloys Comp. 577 (2013) S493–S498. [12] F.A. Rioult, S.D. Imhoff, R. Sakidja, J.H. Perepezko, Transient oxidation of Mo– Si–B alloys: effect of the microstructure size scale, Acta Mater. 57 (2009) 4600–4613. [13] G. Liu, G.J. Zhang, F. Jiang, X.D. Ding, Y.J. Sun, J. Sun, E. Ma, Nanostructured high-strength molybdenum alloys with unprecedented tensile ductility, Nat. Mater. 12 (2013) 344–350. [14] L.A. Joseph, R.M. Michael, W. Tobias, G. Bernd, K.C. Joe, O.R. Robert, On the fracture toughness of ne-grained Mo–3Si–1B (wt.%) alloys at ambient to elevated (1300 °C) temperatures, Intermetallics 20 (2012) 141–154. [15] J.H. Schneibel, R.O. Ritchie, J.J. Kruzic, P.F. Tortorelli, Metall. Mater. Trans. 36 (2005) 525–531. [16] A.P. Alur, N. Chollacoop, K.S. Kumar, High-temperature compression behavior of Mo–Si–B alloys, Acta Mater. 52 (2004) 5571–5587.
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
Ductility of Mo–12Si–8.5B alloys doped with lanthanum oxide by the liquid–liquid doping method Wenhu Li a,b, Guojun Zhang a,⇑, Shixiong Wang a, Bin Li c, Jun Sun c a
School of Materials Science & Engineering, Xi’an University of Technology, Xi’an 710048, China School of Materials Science & Engineering, Shaanxi University of Technology, Hanzhong 723000, China c State Key Laboratory for Mechanical Behavior of Materials, Xi’an Jiaotong University, Xi’an 710049, China b
a r t i c l e
i n f o
Article history: Received 24 December 2014 Received in revised form 4 April 2015 Accepted 6 April 2015 Available online 9 April 2015 Keywords: Mo–Si–B alloy Lanthanum oxide Liquid–liquid doping Fine-grain strengthening
a b s t r a c t Mo–12Si–8.5B (Mo–Si–B) alloys doped with different mass fractions (0.3 wt%, 0.6 wt%, and 0.9 wt%) of lanthanum oxide (La2O3) were prepared by liquid–liquid (L–L) doping, mechanical alloying and hot pressing sintering techniques. The observation of the microstructures of the Mo–Si–B alloys reveals that the grain sizes of the alloys were refined with the increase in La2O3 doping. The fracture toughness values of the alloys of over 10 MPa m1/2 reveal that the addition of La2O3 via the L–L doping method can obviously improve the alloy fracture toughness compared to the alloys doped with La2O3 via the solid–solid (S–S) doping method. In addition, compression tests indicate that the compression strength of the alloys was improved compared to Mo–12Si–8.5B alloys. Ó 2015 Elsevier B.V. All rights reserved.
1. Introduction Due to the high melting point, low thermal expansion coefficient, excellent strength and high temperature oxidation resistance, Mo–Si–B alloy has attracted considerable interest and has been regarded as a potential structural material in high temperature applications for decades [1–3]. Currently, research studies of Mo–Si–B alloys are mainly concentrated on the three phase alloys of Mo3Si + Mo5Si3 + Mo5SiB2 and a-Mo + Mo3Si + Mo5SiB2 (T2) [3]; the different compositions of Mo–Si–B alloy are found to have different properties. Mo3Si + Mo5Si3 + Mo5SiB2 alloy has good high temperature oxidation resistance because the Mo5Si3 phase can provide the Si atoms required to form a protective coating (SiO2); however, it has poor toughness at room temperature [4]. aMo + Mo3Si + Mo5SiB2 (T2) alloy has better toughness as a result of the existence of the a-Mo phase, which is favourable for room temperature toughness [5]. In contrast, relative to the Mo alloy, the fracture toughness of Mo–12Si–8.5B has much room to improve. Therefore, new methods to improve the integrated property of Mo–Si–B alloy are required. As is known, strength and fracture toughness are the two important mechanical properties of alloys, and one of the means to improve the properties of an alloys is to dope it with rare earth oxide. Numerous studies have been performed reporting that ⇑ Corresponding author. Tel.: +86 29 82312592. E-mail address: [email protected] (G. Zhang). http://dx.doi.org/10.1016/j.jallcom.2015.04.047 0925-8388/Ó 2015 Elsevier B.V. All rights reserved.
adding rare earth oxide in Mo alloy can hinder the growth of grain boundaries and refine the grain size. These effects of doping can improve the strength and fracture toughness of Mo alloys [6,7]. In Zhang’s research, doping La2O3 into Mo alloys was found improve the fracture toughness values of the Mo Alloys from 9.9 MPa m1/2 up to 24.76 MPa m1/2 [8]. The room temperature fracture toughness of Mo–12Si–8.5B alloy was on the order of 5– 7 MPa m1/2 [9]. When the Mo–Si–B alloy was doped with a certain amount of La2O3, the fracture toughness of the Mo–Si–B alloy increased up to 9 MPa m1/2 [10,11]. Moreover, Mo–Si–B alloys with fine-grained microstructure have excellent oxidation resistance, as reviewed by Rioult [12]. Thus, doping Mo alloys with rare earth oxide is an effective method to improve the properties of Mo alloys. The typical methods to dope rare earth oxides into Mo alloy include solid–solid doping, solid–liquid doping and liquid–liquid doping. One of the most commonly used methods is solid–solid doping, but the studies of Liu et al. indicated that a molybdenum alloy doped with La2O3 using the liquid–liquid doping method has a smaller grain size and a smaller rare earth oxide particle size than such alloys doped using the traditional solid–liquid doping and solid–solid doping methods. The high temperature stability of the rare earth oxides allows the alloy to maintain its grain size, exhibiting no growth under high temperature [13]. Although much research exists on Mo–Si–B alloys doped with lanthanum oxide, little is known about the Mo–12Si–8.5B alloys doped with lanthanum oxide via the liquid–liquid doping method. In this paper,
W. Li et al. / Journal of Alloys and Compounds 642 (2015) 34–39
Fig. 1. XRD patterns of Mo–12Si–8.5B alloys with different mass fractions of La2O3.
we selected Mo–12Si–8.5B and used liquid–liquid doping, mechanical alloying and hot pressing sintering techniques to prepare the alloys. In addition, we studied the influence of the use of the liquid–liquid doping method on the mechanical properties of the alloys by comparing the mechanical properties of the Mo–Si–B alloys doped with different mass fractions of La2O3.
35
powder to ball weight ratio of 1:10 for 15 h using a Retsch PM 400 apparatus. After mechanical alloying, the powders were hot-pressed at 50 MPa and then sintered in vacuum at 1600 °C for 2 h. The crystalline phases of the alloys were determined using an X-ray diffractometer (XRD) equipped with a Cu Ka radiation source. Microstructural analyses of the alloys were performed via optical metallography (GX71 of OLYMPUS) and scanning electron microscopy (SEM) (JSM-6700F). The specimens were polished and etched with Murakami’s etch (an aqueous solution of potassium ferricyanide and sodium hydroxide). The microstructure of the alloy was also observed via transmission electron microscopy (TEM). The foils for TEM observations were fabricated using a twin-jet apparatus with a solution of 5 vol% H2SO4 in ethanol at 15 °C. A voltage of 10 V produced a current density of 25 mA. The hardness was tested on a TUKON 2100 micro-Vickers’s hardness tester, and the result of every sample was the average value of five points. The densities of the samples were measured using the Archimedes technique. The room temperature compression tests were performed on a HT-2402 computer-controlled machine with a strain rate of 5 103 s1. The high temperature compressive was tested at Gleeble 1500D numerical control thermodynamic testing machine. The sample size was 5 mm 5 mm 10 mm. The fracture toughness values were determined via three-point bending tests. The specimens, which were a cross of 3 mm 4 mm 26 mm with a notch tip depth of 2 mm, were tested in a universal material testing machine (Instron 1185) with a 20-mm span at a crosshead speed of 0.05 mm/min. The fracture toughness values were determined on the basis of an energy criterion by integrating the load–displacement curves. The fracture toughness was determined as G = W/A, where W is the energy absorbed during the fracture of the area A swept out by the crack. Assuming the material to be linear elastic, the fracture energy can also be expressed in terms of the stress intensity by Kq = (G E0 )1/2 where E0 = E/(1 m2) is the plane strain Young’s modulus, and m is Poisson’s ratio (E = 327 GPa and m = 0.29) [1]. In this paper all samples were electro-discharge machined, ground, and then polished before testing. All of the experiment results are average values of at least three samples.
2. Experimental procedures The Mo–12Si–8.5B alloys of different mass fractions of La2O3 were fabricated via mechanical alloying and hot pressing. Molybdenum alloy powders containing different mass fractions of La2O3 particles were first prepared using the L–L doping method. In the L–L doping [13], first, La2O3 is dissolved into an aqueous solution of nitrate (La(NO3)3), and then the solution is mixed with ammonium dimolybdate solution ((NH4)2Mo2O7). Mo–La2O3 powders are subjected to drying roasting and then restored to the mixed solution. Next, Mo–La2O3, Si and B powders are mixed in a planetary ball mill with a speed of 300 rpm and a powder to ball weight ratio of 1:1 for 6 h. The elemental powders of Si and B were 99.99 wt% and 99.99 wt% purity, respectively. Mechanical alloying (MA, a high-energy milling process) was performed under a protective atmosphere (argon) at a speed of 250 rpm and a
3. Results and discussion 3.1. XRD analysis and microstructure characterisation The results of XRD analysis are presented in Fig. 1. These alloys are all composed of three phases of a-Mo, Mo3Si and T2. No other phases were observed, indicating that the phase composition is not affected by either the content of La2O3 or the doping method.
Fig. 2. Optical micrographs of the alloys: (a) Mo–12Si–8.5B + 0.3 wt%La2O3; (b) Mo–12Si–8.5B + 0.6 wt%La2O3; (c) Mo–12Si–8.5B + 0.9 wt%La2O3.
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Moreover, the intensity of the diffraction peaks of the a-Mo phase are higher than those of the Mo3Si and T2 phases. Representative micrographs of the alloys doped with different mass fractions of La2O3 are provided in Fig. 2. The bright phase is the a-Mo matrix, which has the highest volume fraction. The grey phase and dark phase, which are dispersed within the continuous a-Mo matrix, are Mo3Si and T2, respectively. In addition, from Fig. 2, the intermetallic (Mo3Si/T2) phases are also observed to be distributed within the a-Mo matrix homogeneously. The microstructure and grain size of the alloys were investigated using TEM, as shown in Fig. 3. The TEM images clearly show that the grain sizes of the alloys are submicron. The a-Mo, Mo3Si and T2 phases were determined by analyses of the corresponding
selected electron diffraction area. The La2O3 particles, which are of nanometre size and are homogeneously dispersed in the grain interior, are highlighted in Fig. 3(e) and (g) by the solid arrows. The nanoscale particles distributed inside of the a-Mo and the intermetallic regions will be beneficial for both the strength and the toughness of the alloys. The observed grain sizes and statistical results for the alloys are shown in Table 1. The grain size, which is approximately 0.6 lm in the undoped alloys, is reduced with the addition of La2O3. In Zhang’s study [13], more La2O3 particles become homogeneously distributed in the grain interiors than in the grain boundaries when using the L–L doping process compared to the other doping processes. Because the La2O3 particles can act as nucleation sites (which can increase the nucleation rate during
Fig. 3. TEM images of the microstructural morphologies of the alloys (a) Mo3Si and T2 phase; (b–d) are the corresponding selected area diffraction patterns in (a); (e and f) show the images of La2O3 particles in a-Mo; (g and h) show the images of La2O3 particles in T2; (i and j) are the images of the a-Mo matrix.
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Fig. 3 (continued)
the sintering process), this homogeneous distribution is the reason that the grain size in the L–L doping method is smaller than that in the S–S doping method. Additionally, some La2O3 particles may be present at the molybdenum grain boundaries to hinder the grain growth via the pinning effect. Both the pinning and the nucleation effects in La2O3-doped alloys decrease the grain size. 3.2. Fracture toughness The fracture toughness values of the alloys are shown in Fig. 4. All of the Kq values of the alloys are higher than that of the Mo– 12Si–8.5B alloy (9.132 MPa m1/2). Compared with Zhang’s study [10], these alloys prepared via the L–L doping process exhibit higher room temperature fracture toughness than the alloys prepared via the S–S doping process. SEM micrographs of the room temperature fracture surfaces are shown in Fig. 5. The images reveal that the principal failure mechanism is the mixed-mode fracture. Some pullout particles were also observed in the fracture surface. The different elastic modulus and yield strength of the aMo, Mo3Si and Mo5SiB2 phases are responsible for the mixed– mode failure. The transgranular fracture included the cleavage fracture of the larger-size intermetallic particles and of the larger-size a-Mo particles. The toughness of a-Mo is known to be better than that of the intermetallic phase [14,15]. As shown in
Fig. 5(b), the a-Mo phase regions exhibit obvious plastic deformation during fracture. Thus, the transgranular fracture of the ductile a-Mo phases was beneficial for the improvement of the alloy toughness. The toughening mechanisms include fine grain size toughening and particle toughening. The toughening mechanisms of the alloys mainly include finegrain toughening and particle toughening. Material toughness can be effectively improved by refining the grains. On the one hand, cracks are caused by large plastic deformation and mostly occur in grain boundaries and phase boundaries. The grain refinement leads to the increase per unit volume of the grain boundary area. Because the grain boundary atoms are arranged in a disordered manner and the dislocation structure is complex, when the deformation moves from a grain to an adjacent grain by grain boundary, the resistance of dislocation movement and the difficulty of the crack generation will be increased, thus increasing the consumption of energy. On the other hand, the increase in the grain boundary area can improve its effect of blocking the expansion of cracks. When a crack propagates to the grain boundary, different grain orientations on both sides of the grain boundary force the crack extension to change direction or stop. As a result, the alloy toughness is improved. Particle toughening consists of two aspects. The La2O3 particles during liquid–liquid doping of the molybdenum alloy are
Table 1 Grain sizes of the alloys with different mass fractions of La2O3. Alloy composition
Mo–12Si–8.5B + 0.3 wt%La2O3
Mo–12Si–8.5B + 0.6 wt%La2O3
Mo–12Si–8.5B + 0.9 wt%La2O3
Grain sizes (lm)
0.65
0.64
0.59
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Fig. 4. The fracture toughness of Mo–12Si–8.5B alloys of different mass fractions of La2O3.
Fig. 6. Compression stress–strain curves of Mo–12Si–8.5B alloys doped with different mass fractions of La2O3 at ambient temperature.
material increases. Moreover, La2O3 particles refine the grain size of the matrix. When the crack extends to the La2O3 particles, each hard La2O3 particle leads to cracking deflection. As a result, a longer extension path is produced. The resulting energy consumption leads to toughening of the alloy. In summary, the improvement of fracture toughness in Mo alloys is due to the fine grain toughening and particle toughening via La2O3 introduced using the L–L doping method.
3.3. Compression
Fig. 5. SEM images of the fracture surfaces of the alloys: (a) Mo–12Si– 8.5B + 0.6 wt%La2O3; (b) Mo–12Si–8.5B + 0.9 wt%La2O3.
distributed both in the grains and at the grain boundaries. The second phase particles, regardless of being located in the grains or at the grain boundaries, must produce micro-cracks when the dislocation accumulates during the deformation process. This formation of micro-cracks tends to separate the matrix and the particles. Once micro-cracks are formed in the matrix (as indicated by the arrow in Fig. 5(a)), the stress surrounding the micro-cracks is relaxed. At the same time, the new surface formed will absorb a large amount of energy; consequently, the toughness of the
The compression stress–strain curves of the alloy at room temperature are shown in Fig. 6. No obvious yield point is observed, and the stress is observed to continue to increase until the material failed. The sample failed without plastic deformation. The behaviour of the alloy corresponds to the typical room temperature compression curve for brittle materials. The results of the compression strength measurements are listed in Table 2. The room temperature compression strengths of the alloys are improved with the addition of La2O3, and the room temperature compression strengths are also increased with the increase in the La2O3 content. The Mo–12Si–8.5B + 0.9 wt%La2O3 alloy exhibits the highest compression strength (2.97 GPa) compared with the other alloys. In Zhang’s work [10], the compression strength of the Mo–12Si–8.5B + 0.9 wt%La2O3 alloy, in which La2O3 was doped using the solid–solid doping process, was only 2.7 GPa. This result indicates that the alloy that is doped using the L–L doping process has a strength higher than the same composition alloy doped using the solid–solid doping process. As presented in Table 2, the doped alloys have high compression strength at high temperature. At 1000 °C, the Alur alloy had a yield stress of only 320–790 MPa [16]. In the L–L doping case, the La2O3 particles are homogeneously dispersed in the alloys; as a result, the numerous nanoscale particles interact with the migrating grain boundaries,
Table 2 The compression strengths of La2O3 doped Mo–12Si–8.5B alloys at 25 °C, 1000 °C and 1300 °C. Alloy composition
Mo–12Si–8.5B + 0.3 wt%La2O3 Mo–12Si–8.5B + 0.6 wt%La2O3 Mo–12Si–8.5B + 0.9 wt%La2O3
Compression strength (MPa) 25 °C
1000 °C
1300 °C
2565 2788 2872
1604 1539 1414
526 429 376
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causing the alloy grains to remain at small sizes of only 0.5 lm [13]. This refined microstructural morphology contributes to the high strength of the alloy.
4. Conclusions (1) The La2O3 doped Mo–12Si–8.5B alloys are composed of the a-Mo, Mo3Si and Mo5SiB2 phases. The intermetallic Mo3Si and Mo5SiB2 phases are distributed within a continuous aMo matrix. With doping of La2O3 into the alloys, the grain size and the intermetallic particles size of the alloys were refined and distributed more homogeneously. (2) The compression strength of Mo–12Si–8.5B alloys doped with La2O3 using the L–L doping method is improved to 2.87 GPa. In addition, the compression strength is increased with the addition of La2O3; in this case, the major strengthening mechanisms are the strengthening of fine grains and increased particle dispersion. (3) The fracture mode of the alloy is brittle cleavage fracture. The fracture toughness of doped Mo–12Si–8.5B alloys is improved by using the L–L doping method. The primary toughening mechanisms are toughening of fine grains and toughening of particles.
Acknowledgements This subject was supported by the National Natural Science Foundation (Grant Nos. 51171149 and 51371141), the National Science Technology Supporting Program of China (Grant No. 2012BAE06B02).
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