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Microstructure and mechanical properties of molybdenum silicides with Al additions
Materials Science and Engineering A 459 (2007) 132–136
Microstructure and mechanical properties of molybdenum silicides with Al additions I. Rosales a,∗ , D. Bahena a , J. Col´ın b a
Centro de Investigaci´on en Ingenier´ıa y Ciencias Aplicadas, UAEM, Av. Univ. 1001 Col. Chamilpa 62210, Cuernavaca, Morelos, Mexico b Instituto de Ciencias F´ısicas, Universidad Nacional Autonoma de M´ exico, Apdo. Postal 48-3, 62251 Cuernavaca, Morelos, Mexico Received 5 July 2006; received in revised form 19 December 2006; accepted 11 January 2007
Abstract Several molybdenum silicides alloys with different aluminum additions were produced by the arc-cast method. Microstructure observed in the alloys presented a variation of the precipitated second phase respect to the aluminum content. Evaluation of the compressive behavior at high temperature of the alloys shows an important improvement in its ductility, approximately of 20%. Fracture toughness was increased proportionally with Al content. In addition at room temperature the alloys show a better mechanical behavior in comparison with the sample unalloyed. In general, Al additions result to be a good alternative to improve the resistance of these intermetallic alloys. The results are interpreted on the base of the analysis of second phase strengthening. © 2007 Elsevier B.V. All rights reserved. Keywords: Mechanical properties; Molybdenum silicides; Hardness; Fracture toughness; Intermetallic alloys
1. Introduction Significant progress has been reached during the last century in applications for high temperature materials in the aerospace industry, energy, materials processing and many other fields. Essentials to the success of these projects are the new materials such as the intermetallic compounds, which can work at marginal temperatures and stresses under extreme environments. Many intermetallic alloys have been studied extensively, such as nickel aluminides, titanium aluminides, and transition metal disilicides as potential high temperature candidates [1–2]. Intermetallic compounds, specially the silicides, offer the desired properties for high temperature structural applications. The silicides are a new class of materials considered as promissory in structural applications at high temperatures as for aerospace engineering as energy industry due to the combination of high melting point, relative low density and good corrosion resistance [3]. The MoSi2 intermetallic compound and specially MoSi2 base composites are materials used as refractory, wear and
∗
Corresponding author. Tel.: +52 777 3297984; fax: +52 777 3297084. E-mail address: [email protected] (I. Rosales).
0921-5093/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2007.01.048
corrosion resistant materials, however the main disadvantage of the MoSi2 is its low fracture toughness at room temperature and pest corrosion effect. The before characteristics can be enhanced through the alloying with other elements or forming composites. Several elements has been alloyed with the goal of enhance the mechanical properties of silicides such as V, Cr, Nb, Al, W, Re [4] and B [5] with good results. It is reported that Mo-rich Silicides Boron alloyed posses high melting point (up to 2000 ◦ C), good toughness at room temperature and great creep resistance at elevated temperature, however oxidation problems between 650 and 750 ◦ C has been reported, therefore, the application of an Al pack cementation coating is needed [5]. In addition it also has been found that Al additions induces a crystal structure transformation from C11b to a C40 which enhances its high temperature properties [4]. Until now many studies has been performed on Molibdenum silicides [6–11] and specifically the molybdenum disilicides [3–4,12–13], remaining compounds such as Mo3 Si without high attention. In the present investigation, we study this kind of intermetallic compound with A15 structure which has very limited information in such a way that the purpose of the current work is to develop a systematic study of the characterization of molybdenum silicides alloys with Al additions, correlating the effect of
I. Rosales et al. / Materials Science and Engineering A 459 (2007) 132–136
the microstructure with their mechanical properties for possible structural applications.
2. Experimental procedure The alloys were produced by the arc-melting technique under an argon atmosphere (99.999% purity) with a constant 75 at.% Mo composition and Aluminum additions ranging from 8 to 16 at.%. Due to the higher grade of aluminum evaporation and based in weight losses calculations, an extra amount of this element was added (up to 1 at.%). The alloys were drop-cast into water-cooled copper molds with a diameter of 12.5 mm. After that, the specimens were annealed in a vacuum of 10−4 Pa for 24 h at 1400 ◦ C, and then furnace cooled. After metallographic polishing the specimens were etched with Murakami’s reagent during 1–2 s and observed in an optical microscope. The measurement of the area fraction was performed using the IPA (Image Processing Analysis) software of the equipment. Qualitative chemical analysis of the microstructures was performed in a scanning electron microscope equipped with an energy dispersive spectroscopy system (EDS) with internal standards for determining the Mo:Si:Al ratios, then X-ray diffraction analyses were carried out in PHILIPS XRG-3100 equipment. Hardness test were realized in a Leco 300 microhardness tester with a load test of 200 g and a holding time of 15 s. Compression specimens
133
with dimensions of 2 mm × 2 mm × 4 mm were also machined and compressed in an Instron 4501 testing machine at 1400 ◦ C in argon atmosphere at a compression rate of 10−3 s−1 , the machine have an internal strain gage to measure de displacement of the compression rods.. Fracture toughness tests were developed based on the ASTM E399 [14] standard using a notched sample (three point bending procedure).
3. Results and discussion 3.1. Microstructure Fig. 1a–c shows the different microstructures obtained in the alloys under study related to the Al content after annealed for 1 day at 1400 ◦ C in vacuum atmosphere. In this series it can be observed the formation and evolution of a second phase, both were identified first by EDAX (energy disperse X-ray analysis) and then by X-ray diffraction analysis as Mo3 Al8 into the matrix of Mo3 (Si, Al). According to the micrographs in Fig. 1a, the Mo3 Al8 phase precipitates in the alloy containing 8 at.% Al and such phase is softer than the matrix of Mo3 (Si, Al) being the last one the fragile phase located in the interdendritic zone as the microhardness tests shows. In addition, it is observed that the area fraction of the Mo3 Al8 phase (dendritic zone in Fig. 1a and b) grew up proportionally with the increment of the Al in
Fig. 1. (a) Microstructure of the sample with 8 at.% Al. (b) Microstructure of the sample with 12 at.% Al. (c) Microstructure of the sample with 16 at.% Al. (d) Surface of the compressed surface sample 8 at.% Al showing the crack trapping.
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Fig. 2. Area fraction analysis of the different alloys.
the alloys (see Fig. 1a–c). The grain size of the alloys are 12 m for 8 at.% Al, 25 m for 12 at.% Al and 44 m for 16 at.% Al approximately. The presence of the Mo3 (Si, Al) and Mo3 Al8 phases are in good agreement with the AlMoSi ternary phase diagram reported by Yanagihara et al. [15]. Area fraction analysis were performed in the samples, the results obtained are shown in Fig. 2. It can be seen that the softer phase Mo3 Al8 increases from 16.3% for the alloy with 8 at.% Al to 36.13% for the alloy with 12 at.% Al and reaching the highest proportion to 45.5% in the alloy with 16 at.% Al. An EDAX spectrum (Fig. 3) shows the chemical analysis of the two present phases in the sample with 12 at.% Al (A and B for dendritic and interdendritic phases, respectively), where the peaks for spectrum A and B present the proportions of the elements of the constitutive phases identified by X-ray analysis. In Fig. 4, the X-ray diffraction pattern show the analyses for the different alloy compositions appearing the peaks of mainly two predominant phases identified as Mo3 Si–Mo3 Al and Mo3 Al8 corresponding to the two phases observed in Fig. 1b. This observation induce to the assumption that the alloys will get a bigger resistance presumably due to the combination of the
Fig. 3. EDAX spectrum of the sample with 12 at.% Al.
Fig. 4. X-ray diffraction analysis of the powder from samples with 8, 12 and 16 at.% Al, from the bottom to the top of the figure, respectively.
mechanical properties of the hard (Si, Al)Mo3 and the ductile Mo3 Al8 phases as it will be showed in the next paragraphs. 4. Mechanical properties 4.1. Hardness The values obtained of the micro-hardness tests on the matrix sample are shown in Table 1, it can be observed a hardness decrement when Al additions increase, being the lowest value of 1094 HV corresponding to the alloy with 16 at.% Al; hardness values of the alloys before this one are higher. From the information showed in Table 1, it can be observed a tendency to decrease the microhardness values of the alloys in relation to the Al content, assuming that the hardness of the alloys is lower than those without Al additions (Mo3 Si single phase) due to the contiguous presence of the ductile phase Mo3 Al8 interacting with the hard phase Mo3 (Si, Al), which allows that the material can be deformed when the load is applied, producing this reduction in the hardness value. Yanagihara et al. [15], studied the cracking behavior of two-phase Mo3 Al–Mo3 Al8 , quenched at 1623 K, pointing out that indentation cracking is suppressed in samples with a coarse lamellar structure. The hardness value of the softer phase is on the order of 650 HVN in annealed condition (since Mo3 Al8 is a hardly deformable phase due to its complex crystal structure (mC22), in this case we attribute this relative low hardness to the heat treatment for such reason we
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Table 1 Resume of the nominal compositions and mechanical properties of the alloys with Al additions Alloy (at.%)
Microhardness in matrix (HVN)
Mo–24Si Mo–17Si–8Al Mo–13Si–12Al Mo–11Si–16Al
1318 1221 1134 1094
± ± ± ±
27 30 25 18
Fracture toughness (MPa m1/2 ) 2.5 4.0 3.3 4.5
± ± ± ±
0.4 0.2 0.3 0.5
Compressive yield strength at 1400 ◦ C (MPa) 540 384 401 420
called it as a relative softer phase). The difference of the indentations marks observed in Fig. 1c shows the crack development in the corners of the indentation on the brittle phase, while some grade of deformation is evident on the adjacent zones of the indentation mark of the softer phase. 5. Fracture toughness Table 1 shows the plot of the fracture toughness at room temperature of the samples as a function of the aluminum content. It is evident that the aluminum addition contributed to the increased resistance of fracture toughness behavior. This phenomenon is attributed to the fact that the ductile phase Mo3 Al8 provide a direct interaction with the crack propagation developed in the matrix (crack trapping) and also as is pointed by Yanagihara et al. [15], that dislocation motion around crack contribute to high toughness, in such way that the toughness improvement in comparison with the sample with 0 at.% Al is at least twice, which can be considered as an important approach to the strengthening of the alloys. If well is true that the obtained results shows a significant increment in the fracture toughness of the alloys, such increment is not enough to considering the alloys for structural applications where values greater than 10 MPa m1/2 are needed, however a possible application could be as a coating. Currently several studies at this respect are being performed as Al–Mo and Al–Mo–Si coatings [16,17]. 6. Compresi´on test In Fig. 5 is presented the plot of the curves σ–ε for the different alloys at 1400 ◦ C and a strain rate of 10−3 s−1 . The curves of the compressed samples with Al additions shown a similar pattern in the yielding zone, changing the behavior to reach a plateau, with the exception of the sample with 8 at.% Al that do not present this behavior. The first consideration is the fact that the yield strength value of all of the alloys increases as a function of the Al content being the lowest value of 384 MPa in the alloy with 8 at.% Al and 420 MPa for the alloy with 16 at.% Al. In spite of the before, all of the three alloys shows lower values of the yield strength in comparison with the sample without Al additions, being the yield strength diminished approximately one order of magnitude. A second consideration is related with the deformation produced in the alloys with higher aluminum content, which, due to the fact that the area fraction percent of the ductile phase is higher in the structure when the Al increases, this produces also higher deformation rates, specifically 2, 8 and 10% approximately for 8, 12 and 16 at.% Al, respectively, in comparison with the sample unalloyed, as is observed in Fig. 5.
Fig. 5. Compression test at 1400 ◦ C and a strain rate of 10−3 s−1 .
The above results shows that the optimal aluminum content permissible to reach the best mechanical performance of the alloys is around of 16 at.% Al, as expected from the alloy design and the obtained microstructure, the increment of the area fraction of the softer phase result to be an efficient way to arrest the crack propagation as is observed in Fig. 1d. 7. Conclusions Aluminum additions to the Mo3 (Si, Al) matrix, result to be a good alternative to improve the mechanical properties of this intermetallic alloy. Microstructures obtained for samples after annealing, shows the formation a second phase with the remarkable presence of an Aluminum phase precipitation (softer phase) which has been identified by EDS and X-ray analysis as Mo3 Al8 . Aluminum alloying produces a slightly decrement in the microhardness values (matrix) in comparison to the unalloyed material; in addition, fracture toughness results shown an increment in the value of approximately 20% respect to the Mo3 Si single phase. Finally, the compression tests results at 1400 ◦ C shown an important increment in deformation when the aluminum content is increased. Studies of oxidation at high temperature for these alloys are currently being developed. Acknowledgements The authors want to thank to ORNL in sample preparation. Thanks to R. Guardian, and O. Flores for technical assistance. This project was supported by the PROMEP-UAE-074 project.
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I. Rosales et al. / Materials Science and Engineering A 459 (2007) 132–136
References [1] [2] [3] [4] [5] [6] [7] [8] [9]
R. Tanaka, Mater. High Temp. 17 (4) (2000) 457–464. J.J. Petrovic, MRS Bull. XVIII (1993) 35. P.V. Krakhmalev, J. Bergstr¨on, Wear 260 (2006) 450–457. H. Inui, K. Ishikawa, M. Yamaguchi, Intermetallics 8 (2000) 1131–1145. R. Sakidja, F. Rioult, J. Werner y, J.H. Perepezko, Scripta Mater. 55 (2006) 903–906. A.K. Vasudevan, J.J. Petrovic, Mater. Sci. Eng. A 155 (1992) 1–17. M.K. Meyer, M. Akin, J. Am. Ceram. Soc. 79 (1996) 938. M. Akinc, M.K. Meyer, M.J. Kramer, J.A. Thom, J.J. Huebsch, B. Cook, Mater. Sci. Eng. A 261 (1999) 16–23. J.H. Schneibel, M.J. Kramer, D.S. Easton, Scripta Mater. 46 (2002) 217–221.
[10] C.T. Liu, J.H. Schneibel, L. Heatherly, High temperature order intermetallics alloys VIII, E.P., George, et al. (ed.), Mat. Res. Soc. Sym. Proc 1999, 552, KK6.2. [11] F. Peizhong, Q. Xuanhui, A. Farid, S.H. Islam, Rare Metals 25 (3) (2006) 225–230. [12] P.V. Krakhmalev, E. Str¨om, C. Li, Intermetallics 12 (2004) 225–233. [13] J.J. Petrovic, Intermetallics 8 (2000) 1175–1182. [14] “Annual Book of ASTM Standards” Metals Test Methods and Analytical Procedures, 1991, p. 497. [15] K. Yanagihara, T. Maruyama, K. Nagata, Mater. Trans. JIM 34 (1993) 1200, 499. [16] R. Nino, S. Miura, T. Mor´ı, Intermetallics 9 (2001) 113–118. [17] B. Wendler, M. Danielewski, J. Dabek, A. Rylski, R. Filipek, Thin Solid Films 459 (2004) 178–182.
Microstructure and mechanical properties of molybdenum silicides with Al additions I. Rosales a,∗ , D. Bahena a , J. Col´ın b a
Centro de Investigaci´on en Ingenier´ıa y Ciencias Aplicadas, UAEM, Av. Univ. 1001 Col. Chamilpa 62210, Cuernavaca, Morelos, Mexico b Instituto de Ciencias F´ısicas, Universidad Nacional Autonoma de M´ exico, Apdo. Postal 48-3, 62251 Cuernavaca, Morelos, Mexico Received 5 July 2006; received in revised form 19 December 2006; accepted 11 January 2007
Abstract Several molybdenum silicides alloys with different aluminum additions were produced by the arc-cast method. Microstructure observed in the alloys presented a variation of the precipitated second phase respect to the aluminum content. Evaluation of the compressive behavior at high temperature of the alloys shows an important improvement in its ductility, approximately of 20%. Fracture toughness was increased proportionally with Al content. In addition at room temperature the alloys show a better mechanical behavior in comparison with the sample unalloyed. In general, Al additions result to be a good alternative to improve the resistance of these intermetallic alloys. The results are interpreted on the base of the analysis of second phase strengthening. © 2007 Elsevier B.V. All rights reserved. Keywords: Mechanical properties; Molybdenum silicides; Hardness; Fracture toughness; Intermetallic alloys
1. Introduction Significant progress has been reached during the last century in applications for high temperature materials in the aerospace industry, energy, materials processing and many other fields. Essentials to the success of these projects are the new materials such as the intermetallic compounds, which can work at marginal temperatures and stresses under extreme environments. Many intermetallic alloys have been studied extensively, such as nickel aluminides, titanium aluminides, and transition metal disilicides as potential high temperature candidates [1–2]. Intermetallic compounds, specially the silicides, offer the desired properties for high temperature structural applications. The silicides are a new class of materials considered as promissory in structural applications at high temperatures as for aerospace engineering as energy industry due to the combination of high melting point, relative low density and good corrosion resistance [3]. The MoSi2 intermetallic compound and specially MoSi2 base composites are materials used as refractory, wear and
∗
Corresponding author. Tel.: +52 777 3297984; fax: +52 777 3297084. E-mail address: [email protected] (I. Rosales).
0921-5093/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2007.01.048
corrosion resistant materials, however the main disadvantage of the MoSi2 is its low fracture toughness at room temperature and pest corrosion effect. The before characteristics can be enhanced through the alloying with other elements or forming composites. Several elements has been alloyed with the goal of enhance the mechanical properties of silicides such as V, Cr, Nb, Al, W, Re [4] and B [5] with good results. It is reported that Mo-rich Silicides Boron alloyed posses high melting point (up to 2000 ◦ C), good toughness at room temperature and great creep resistance at elevated temperature, however oxidation problems between 650 and 750 ◦ C has been reported, therefore, the application of an Al pack cementation coating is needed [5]. In addition it also has been found that Al additions induces a crystal structure transformation from C11b to a C40 which enhances its high temperature properties [4]. Until now many studies has been performed on Molibdenum silicides [6–11] and specifically the molybdenum disilicides [3–4,12–13], remaining compounds such as Mo3 Si without high attention. In the present investigation, we study this kind of intermetallic compound with A15 structure which has very limited information in such a way that the purpose of the current work is to develop a systematic study of the characterization of molybdenum silicides alloys with Al additions, correlating the effect of
I. Rosales et al. / Materials Science and Engineering A 459 (2007) 132–136
the microstructure with their mechanical properties for possible structural applications.
2. Experimental procedure The alloys were produced by the arc-melting technique under an argon atmosphere (99.999% purity) with a constant 75 at.% Mo composition and Aluminum additions ranging from 8 to 16 at.%. Due to the higher grade of aluminum evaporation and based in weight losses calculations, an extra amount of this element was added (up to 1 at.%). The alloys were drop-cast into water-cooled copper molds with a diameter of 12.5 mm. After that, the specimens were annealed in a vacuum of 10−4 Pa for 24 h at 1400 ◦ C, and then furnace cooled. After metallographic polishing the specimens were etched with Murakami’s reagent during 1–2 s and observed in an optical microscope. The measurement of the area fraction was performed using the IPA (Image Processing Analysis) software of the equipment. Qualitative chemical analysis of the microstructures was performed in a scanning electron microscope equipped with an energy dispersive spectroscopy system (EDS) with internal standards for determining the Mo:Si:Al ratios, then X-ray diffraction analyses were carried out in PHILIPS XRG-3100 equipment. Hardness test were realized in a Leco 300 microhardness tester with a load test of 200 g and a holding time of 15 s. Compression specimens
133
with dimensions of 2 mm × 2 mm × 4 mm were also machined and compressed in an Instron 4501 testing machine at 1400 ◦ C in argon atmosphere at a compression rate of 10−3 s−1 , the machine have an internal strain gage to measure de displacement of the compression rods.. Fracture toughness tests were developed based on the ASTM E399 [14] standard using a notched sample (three point bending procedure).
3. Results and discussion 3.1. Microstructure Fig. 1a–c shows the different microstructures obtained in the alloys under study related to the Al content after annealed for 1 day at 1400 ◦ C in vacuum atmosphere. In this series it can be observed the formation and evolution of a second phase, both were identified first by EDAX (energy disperse X-ray analysis) and then by X-ray diffraction analysis as Mo3 Al8 into the matrix of Mo3 (Si, Al). According to the micrographs in Fig. 1a, the Mo3 Al8 phase precipitates in the alloy containing 8 at.% Al and such phase is softer than the matrix of Mo3 (Si, Al) being the last one the fragile phase located in the interdendritic zone as the microhardness tests shows. In addition, it is observed that the area fraction of the Mo3 Al8 phase (dendritic zone in Fig. 1a and b) grew up proportionally with the increment of the Al in
Fig. 1. (a) Microstructure of the sample with 8 at.% Al. (b) Microstructure of the sample with 12 at.% Al. (c) Microstructure of the sample with 16 at.% Al. (d) Surface of the compressed surface sample 8 at.% Al showing the crack trapping.
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Fig. 2. Area fraction analysis of the different alloys.
the alloys (see Fig. 1a–c). The grain size of the alloys are 12 m for 8 at.% Al, 25 m for 12 at.% Al and 44 m for 16 at.% Al approximately. The presence of the Mo3 (Si, Al) and Mo3 Al8 phases are in good agreement with the AlMoSi ternary phase diagram reported by Yanagihara et al. [15]. Area fraction analysis were performed in the samples, the results obtained are shown in Fig. 2. It can be seen that the softer phase Mo3 Al8 increases from 16.3% for the alloy with 8 at.% Al to 36.13% for the alloy with 12 at.% Al and reaching the highest proportion to 45.5% in the alloy with 16 at.% Al. An EDAX spectrum (Fig. 3) shows the chemical analysis of the two present phases in the sample with 12 at.% Al (A and B for dendritic and interdendritic phases, respectively), where the peaks for spectrum A and B present the proportions of the elements of the constitutive phases identified by X-ray analysis. In Fig. 4, the X-ray diffraction pattern show the analyses for the different alloy compositions appearing the peaks of mainly two predominant phases identified as Mo3 Si–Mo3 Al and Mo3 Al8 corresponding to the two phases observed in Fig. 1b. This observation induce to the assumption that the alloys will get a bigger resistance presumably due to the combination of the
Fig. 3. EDAX spectrum of the sample with 12 at.% Al.
Fig. 4. X-ray diffraction analysis of the powder from samples with 8, 12 and 16 at.% Al, from the bottom to the top of the figure, respectively.
mechanical properties of the hard (Si, Al)Mo3 and the ductile Mo3 Al8 phases as it will be showed in the next paragraphs. 4. Mechanical properties 4.1. Hardness The values obtained of the micro-hardness tests on the matrix sample are shown in Table 1, it can be observed a hardness decrement when Al additions increase, being the lowest value of 1094 HV corresponding to the alloy with 16 at.% Al; hardness values of the alloys before this one are higher. From the information showed in Table 1, it can be observed a tendency to decrease the microhardness values of the alloys in relation to the Al content, assuming that the hardness of the alloys is lower than those without Al additions (Mo3 Si single phase) due to the contiguous presence of the ductile phase Mo3 Al8 interacting with the hard phase Mo3 (Si, Al), which allows that the material can be deformed when the load is applied, producing this reduction in the hardness value. Yanagihara et al. [15], studied the cracking behavior of two-phase Mo3 Al–Mo3 Al8 , quenched at 1623 K, pointing out that indentation cracking is suppressed in samples with a coarse lamellar structure. The hardness value of the softer phase is on the order of 650 HVN in annealed condition (since Mo3 Al8 is a hardly deformable phase due to its complex crystal structure (mC22), in this case we attribute this relative low hardness to the heat treatment for such reason we
I. Rosales et al. / Materials Science and Engineering A 459 (2007) 132–136
135
Table 1 Resume of the nominal compositions and mechanical properties of the alloys with Al additions Alloy (at.%)
Microhardness in matrix (HVN)
Mo–24Si Mo–17Si–8Al Mo–13Si–12Al Mo–11Si–16Al
1318 1221 1134 1094
± ± ± ±
27 30 25 18
Fracture toughness (MPa m1/2 ) 2.5 4.0 3.3 4.5
± ± ± ±
0.4 0.2 0.3 0.5
Compressive yield strength at 1400 ◦ C (MPa) 540 384 401 420
called it as a relative softer phase). The difference of the indentations marks observed in Fig. 1c shows the crack development in the corners of the indentation on the brittle phase, while some grade of deformation is evident on the adjacent zones of the indentation mark of the softer phase. 5. Fracture toughness Table 1 shows the plot of the fracture toughness at room temperature of the samples as a function of the aluminum content. It is evident that the aluminum addition contributed to the increased resistance of fracture toughness behavior. This phenomenon is attributed to the fact that the ductile phase Mo3 Al8 provide a direct interaction with the crack propagation developed in the matrix (crack trapping) and also as is pointed by Yanagihara et al. [15], that dislocation motion around crack contribute to high toughness, in such way that the toughness improvement in comparison with the sample with 0 at.% Al is at least twice, which can be considered as an important approach to the strengthening of the alloys. If well is true that the obtained results shows a significant increment in the fracture toughness of the alloys, such increment is not enough to considering the alloys for structural applications where values greater than 10 MPa m1/2 are needed, however a possible application could be as a coating. Currently several studies at this respect are being performed as Al–Mo and Al–Mo–Si coatings [16,17]. 6. Compresi´on test In Fig. 5 is presented the plot of the curves σ–ε for the different alloys at 1400 ◦ C and a strain rate of 10−3 s−1 . The curves of the compressed samples with Al additions shown a similar pattern in the yielding zone, changing the behavior to reach a plateau, with the exception of the sample with 8 at.% Al that do not present this behavior. The first consideration is the fact that the yield strength value of all of the alloys increases as a function of the Al content being the lowest value of 384 MPa in the alloy with 8 at.% Al and 420 MPa for the alloy with 16 at.% Al. In spite of the before, all of the three alloys shows lower values of the yield strength in comparison with the sample without Al additions, being the yield strength diminished approximately one order of magnitude. A second consideration is related with the deformation produced in the alloys with higher aluminum content, which, due to the fact that the area fraction percent of the ductile phase is higher in the structure when the Al increases, this produces also higher deformation rates, specifically 2, 8 and 10% approximately for 8, 12 and 16 at.% Al, respectively, in comparison with the sample unalloyed, as is observed in Fig. 5.
Fig. 5. Compression test at 1400 ◦ C and a strain rate of 10−3 s−1 .
The above results shows that the optimal aluminum content permissible to reach the best mechanical performance of the alloys is around of 16 at.% Al, as expected from the alloy design and the obtained microstructure, the increment of the area fraction of the softer phase result to be an efficient way to arrest the crack propagation as is observed in Fig. 1d. 7. Conclusions Aluminum additions to the Mo3 (Si, Al) matrix, result to be a good alternative to improve the mechanical properties of this intermetallic alloy. Microstructures obtained for samples after annealing, shows the formation a second phase with the remarkable presence of an Aluminum phase precipitation (softer phase) which has been identified by EDS and X-ray analysis as Mo3 Al8 . Aluminum alloying produces a slightly decrement in the microhardness values (matrix) in comparison to the unalloyed material; in addition, fracture toughness results shown an increment in the value of approximately 20% respect to the Mo3 Si single phase. Finally, the compression tests results at 1400 ◦ C shown an important increment in deformation when the aluminum content is increased. Studies of oxidation at high temperature for these alloys are currently being developed. Acknowledgements The authors want to thank to ORNL in sample preparation. Thanks to R. Guardian, and O. Flores for technical assistance. This project was supported by the PROMEP-UAE-074 project.
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I. Rosales et al. / Materials Science and Engineering A 459 (2007) 132–136
References [1] [2] [3] [4] [5] [6] [7] [8] [9]
R. Tanaka, Mater. High Temp. 17 (4) (2000) 457–464. J.J. Petrovic, MRS Bull. XVIII (1993) 35. P.V. Krakhmalev, J. Bergstr¨on, Wear 260 (2006) 450–457. H. Inui, K. Ishikawa, M. Yamaguchi, Intermetallics 8 (2000) 1131–1145. R. Sakidja, F. Rioult, J. Werner y, J.H. Perepezko, Scripta Mater. 55 (2006) 903–906. A.K. Vasudevan, J.J. Petrovic, Mater. Sci. Eng. A 155 (1992) 1–17. M.K. Meyer, M. Akin, J. Am. Ceram. Soc. 79 (1996) 938. M. Akinc, M.K. Meyer, M.J. Kramer, J.A. Thom, J.J. Huebsch, B. Cook, Mater. Sci. Eng. A 261 (1999) 16–23. J.H. Schneibel, M.J. Kramer, D.S. Easton, Scripta Mater. 46 (2002) 217–221.
[10] C.T. Liu, J.H. Schneibel, L. Heatherly, High temperature order intermetallics alloys VIII, E.P., George, et al. (ed.), Mat. Res. Soc. Sym. Proc 1999, 552, KK6.2. [11] F. Peizhong, Q. Xuanhui, A. Farid, S.H. Islam, Rare Metals 25 (3) (2006) 225–230. [12] P.V. Krakhmalev, E. Str¨om, C. Li, Intermetallics 12 (2004) 225–233. [13] J.J. Petrovic, Intermetallics 8 (2000) 1175–1182. [14] “Annual Book of ASTM Standards” Metals Test Methods and Analytical Procedures, 1991, p. 497. [15] K. Yanagihara, T. Maruyama, K. Nagata, Mater. Trans. JIM 34 (1993) 1200, 499. [16] R. Nino, S. Miura, T. Mor´ı, Intermetallics 9 (2001) 113–118. [17] B. Wendler, M. Danielewski, J. Dabek, A. Rylski, R. Filipek, Thin Solid Films 459 (2004) 178–182.