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magnesium composites: Processing, microstructure and properties
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Journal of the European Ceramic Society 32 (2012) 3477–3483
Boron carbide/magnesium composites: Processing, microstructure and properties M. Cafri, H. Dilman, M.P. Dariel, N. Frage ∗ Department of Materials Engineering, Ben-Gurion University of the Negev, P.O. Box 653, Beer-Sheva 84105, Israel Received 14 December 2011; received in revised form 27 March 2012; accepted 1 April 2012 Available online 3 May 2012
Abstract Composites with a high fraction of the ceramic phase were fabricated by infiltration of 80% dense preforms with liquid Mg or Mg alloys, under an Mg vapor atmosphere. The infiltration was performed at 1123 K in a semi-hermetically closed container, from which gas had been evacuated. The Mg vapor atmosphere was achieved by heating subsequently to the relatively high temperature of the Mg infiltration process. The specific weight of the composites is about 2.44 g/cm3 . The microstructure of the composite consists of the newly formed MgB2 and ternary carbide MgB2 C2 phases that connect the initial B4 C particles and some residual metal. A thermodynamic analysis of the interaction between B4 C and liquid Mg was conducted and its results are in good agreement with the experimental observations. The mechanical properties of the composites were investigated and discussed. © 2012 Elsevier Ltd. All rights reserved. Keywords: Composites; Carbides; Microstructure; Mechanical properties
1. Introduction Ceramic metal composites are advanced materials of interest on account of the potential synergy provided by the combination of metallic and ceramic properties. This group includes metal matrix composites (MMCs) and ceramic matrix composites (CMCs). Boron carbide (B4 C) with its low density, elevated modulus and hardness values is an attractive ceramic for applications such as light armor and as a wear resistant component in structural materials for land borne and airborne transport vehicles. MMCs may be fabricated by a stir-casting process or by infiltration of porous ceramic preforms with a liquid metal. The fraction of ceramic particles that may be distributed homogenously in the melt in the course of the stir-casting process is of the order of 30 vol.%. This value may be significantly higher for a free infiltration approach. The metal-to-ceramic ratio in the composites determines their mechanical properties and their possible applications. Metal ceramic composites (MMCs) based on aluminum alloys strengthened with B4 C particles have been the subject of
∗
Corresponding author. Tel.: +972 8 6461468; fax: +972 8 6479350. E-mail address: [email protected] (N. Frage).
0955-2219/$ – see front matter © 2012 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.jeurceramsoc.2012.04.007
extensive studies.1–7 It was of obvious interest to extend these studies to Mg-based MMCs, considering the reduced specific weight of the latter. The fabrication conditions of the boron carbide–magnesium composites were discussed in Refs. 8–11. A maximal fraction of 50 vol% ceramic phase in these composites was attained by Kevorkijan and Scapin10 by free infiltration of porous boron carbide preforms. Higher B4 C content composites have been successfully manufactured by the so-called reaction bonding approach, which consists of infiltrating a green B4 C powder preform with molten silicon with or without the addition of free carbon.12–15 The final microstructure of the reaction bonded boron carbide composites (RBBC) composite consists of a continuous B4 C skeleton, silicon carbide (SiC), ternary B12 (Si,C,B)3 and residual Si. The high strength of the resulting material is due to the in situ reaction between carbon that has been added or that originates in the carbide, with molten silicon to form a SiC network. It was reported in our previous publication16 that the fraction of the residual Si may be significantly reduced (to 8–10 vol.%) by the infiltration of porous preforms with a relatively high green density (of about 80%), fabricated from a mixture of boron carbide powders with an appropriate particle size distribution. In the present study, we have used this approach in order to manufacture B4 C-based composites infiltrated with molten Mg and Mg alloys. A similar
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conceptual approach using Mg or Mg alloys infiltration instead of molten Si presents, in addition to the reduced specific weight, the advantage of lowering the processing temperature from close to 1773 K down to 1123 K. On the other hand it remains to be seen to what extent the reaction between B4 C and molten Mg or its alloys provides and adequate cohesive strength to the composite. Three conditions have to be fulfilled in order for the spontaneous infiltration to take place, namely: (a) adequate wetting of the porous preform by the liquid metal; (b) gas evacuation from the pores; (c) moderate chemical interaction to avoid the formation of new phases that may plug up open channels in the porous preform and impede, thereby, infiltration. The contact wetting angle between the ceramic substrate and the liquid infiltrant provides a measure of the wetting in the system. Its direct measurement in the Mg/B4 C system is a complex task due to the high partial vapor pressure and the reactivity of Mg. Zhang et al.16 have recently reported that a transition from nonwetting to wetting condition occurs in the 972–1173 K temperature range and have attributed it to the chemical interaction at the metal/carbide interface. The attempts to fabricate MMCs based on Mg or its alloys and B4 C particles8–11 have shown that under the appropriate atmosphere, the wetting angle is sufficient for free infiltration. For instance, in Ref. 10 the infiltration was conducted under a flow of oxygen-free nitrogen and with Ti addition as a wetting agent. Under these conditions B4 C preforms, prepared from a bimodal mixture of particles (0.8 and 50 m), with a fraction of fine powder less than 60 vol.% were successfully infiltrated with liquid Mg at 1123 K. Preforms with a higher fraction of fine powder cannot be infiltrated under similar conditions. In the latter case, such powder mixtures have a relatively high surface area and a higher oxygen content due to the presence of the thin boron oxide films on the particles surface. An increased content of oxygen leads to massive MgO formation that impedes the infiltration of the B4 C porous body. The presence of MgO at the Mg/B4 C interface was detected by Jiang et al.8 Due to high Mg affinity to oxygen it is impossible to prevent its oxidation under any gaseous atmosphere. In order to circumvent the issue we made use of Mg vapor atmosphere generated by Mg evaporation. The infiltration process was conducted in a semi-hermetically closed container, from which gas had been evacuated before heating and a Mg vapor atmosphere was generated at the relatively high temperature of the infiltration process. In the present work, the interface interaction in the Mg/B4 C system and the mechanical and physical properties of the infiltrated composites with about 80 vol.% of B4 C were investigated. 2. Experimental 2.1. Samples preparation Two sets of samples were fabricated from commercial B4 C powder mixtures that allowed reaching about 80% green density preforms. Fine B4 C powder (∼3 m, grade HS, 97.5% purity) was supplied by H.C. Starck and B4 C powders with average particle sizes of 13, 50 and 100 m (97% purity) were supplied by a
Chinese Company “Mudanjiang”. Samples with the appropriate particle size distribution (3 m – 19 vol.%, 13 m – 18 vol.%, 50 m – 10 vol.%, 106 m – 53 vol.%) were dry mixed for 16 h. Preforms, 20 mm diameter and 4 mm height, were compacted uniaxially (180 MPa) yielding compacts with about 20 vol.% porosity. Infiltration was done by placing Mg or Mg alloy AZ91 (Mg–9 wt.%Al–0.7 wt.%Zn–0.3 wt.%Mn) pieces on the top of preforms. The preforms with metal pieces on their top were positioned in a semi-hermetically closed graphite container (the cover was not completely sealed). The latter had been evacuated by pumping to ∼1 Pa vacuum for 2 h prior heating. The infiltration process was conducted at 1123 K for 20 min. At this temperature, the equilibrium Mg vapor pressure (∼6.6 kPa) is higher than the vacuum level in the furnace chamber and, although Mg escapes partly from the crucible, the infiltration process takes place under Mg vapor. The initial amount of the metal for infiltration was twice that required for filling the free volume in the preforms. The infiltrated composites were cooled down to room temperature at a 10 ◦ C/min cooling rate. The preforms, infiltrated with pure Mg and AZ91, are denoted as M1 and M2, respectively. 2.2. Wetting experiments Fully dense boron carbide substrates were used for the wetting experiments. Small pieces of Mg and AZ91 (2–3 g) were placed on the substrate and melted under a Mg vapor atmosphere in a semi-hermetically closed graphite container. The contact angles were measured using the optical image of the drops after cooling and solidification. 2.3. Characterization 2.3.1. Microstructure and composition The microstructure of the samples was studied by scanning electron microscopy (SEM, JEOL-5600) combined with an energy-dispersive spectrometer (EDS). The samples for the microstructural characterization were prepared using a standard metallographic procedure that included a last stage of polishing by 0.5 m diamond paste. The phase composition of the infiltrated composites was determined by X-ray diffraction (XRD) using Rigaku RINT 2100 Bragg-Brentano diffractometer with Cu K␣ radiation, the scanning angle 2θ was increased in 0.02◦ steps. 2.3.2. Mechanical and physical properties The micro-hardness was determined using a Buehler tester (Micromet 2100) with a Vickers indenter under 20 N load. The flexural strength was determined by three-point bending test with a LRX plus LLOYD machine (Lloyd Instruments, Fareham Hants, U.K.), on 1.5 mm × 2 mm × 10 mm samples. The elastic modulus of the composites was derived from ultrasonic sound velocity measurements using the “Pulse Echo” method. The density of the infiltrated composites was determined by the Archimedes method.
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Fig. 1. Optical images of cross-sectioned drops and B4 C substrate: (a) pure Mg and (b) AZ91.
Fig. 2. SEM images of the composites microstructure: (a) M1 and (b) M2.
3. Results and discussion 3.1. Wetting and infiltration
diffraction (Fig. 3) showed that both the M1 and M2 composites contain the Mg, B4 C, MgB2 and MgB2 C2 phases. The M2 composite also contains the Mg17 Al12 phase (␥phase), which is expected to precipitate from the supersaturated
The optical images of the drops after solidification and cooling are presented in Fig. 1. The contact angle of pure Mg is about 25–30◦ , while that angle for AZ91 Mg-alloy is significantly lower (∼12◦ ). The wetting in these systems is governed by an interface chemical interaction and the formation of new phases at the interface (see below). The important conclusion is that the values of the contact angle are low enough for free infiltration. The microstructure of the composites (Fig. 2) indeed supports the successful infiltration of the porous preforms with liquid Mg and AZ91 and the absence of any residual porosity. 3.2. Phase composition of the infiltrated composites In order to investigate the phase composition of the composites and to clarify the chemical interaction that takes place at the metal/ceramic interface, the composites were fabricated by infiltration of porous preforms that had been compacted exclusively of coarse B4 C powder particles (of about 50 m). The X-ray
Fig. 3. XRD patterns of Mg, AZ91, M1 and M2 composites.
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Table 1 The estimated phase composition of the composites. Composite
Mg (vol.%)
B4 C (vol.%)
MgB2 (vol.%)
MgB2 C2 (vol.%)
Mg17 Al12 (vol.%)
M1 M2
10 5
76 75
14 13
<1 3
0 4
Mg(Al,Zn,Mn) solid solution. In the AZ91 and M2 composites the Mg diffraction peaks are shifted to higher angels than in the pure Mg and M1 composites on account of the presence of alloying elements, namely Al and Zn, in the Mg solid solution. Alumina, which was identified by XRD analysis, originates from the powder preparation stage form the bulk composite specimens (milling in an alumina vessel using alumina balls). The diffraction data were used for the semi-quantitative analysis of the volume fraction of each phase within the composites (Table 1). Composite M1 contains 10 vol.% of residual Mg and the dominant newly formed phase is MgB2 . The M2 composite contains only 5 vol.% of residual metallic phase and a relatively high fractions of the ternary MgB2 C2 carbide and Mg17 Al12 intermetallic. We suggest that due to the high affinity of Al to carbon and boron, it enhances the formation of the ternary carbide, which according to the EDS analysis (Fig. 4d) contains some fraction of aluminum.
The results of the XRD analysis are supported by the microstructural characterization of the infiltrated composites (Fig. 4). The surface of each B4 C particle is surrounded by a continuous layer of newly formed phases that reflect a chemical interaction at the metal/ceramic interface. This interaction is more massive in the B4 C/AZ91 composite, probably, due to the presence of Al in the melt and its involvement in the interface reaction. EDS analysis indicates that Al is actually present (Fig. 4d) in the interfacial layer. This observation is a good agreement with results on Al and B4 C interaction reported in the literature.17–19 The important point is that a new interfacial layer connects the B4 C particles and gives rise to the formation of a so-called reaction bonded composite. The cohesive strength between the newly formed intermediate layer and the B4 C particles stands behind the relatively elevated mechanical properties of the composite. The thermodynamic analysis presented
Fig. 4. SEM images of the composites (a – M1, b – M2) fabricated using coarse B4 C. Graphs c and d show the results of line scans along the white lines in the inserts.
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below provides the underlying reasons for the interfacial layer formation. 3.3. Interaction between B4 C and liquid Mg: thermodynamic considerations According to the binary Mg–C phase diagram,20 the solubility of C in liquid Mg is relatively high (of about 20 at.% at 1123 K) and magnesium carbide is stable only at temperatures lower than 873 K. The binary Mg–B phase diagram was calculated and discussed by Xi et al.21 In the temperature interval of our interest, only MgB2 coexists with liquid and gaseous Mg. The boron solubility limit in liquid magnesium at 1123 K is about 0.01 at.%. The calculated22 isothermal section of the ternary Mg–C–B system at 1000 K and 133.3 Pa is presented in Fig. 5. This isothermal section corresponds to the phase equilibria between gaseous Mg (Mg vapor pressure at this temperature is about 6.6 kPa) and various nonmetallic phases, including the ternary MgB2 C2 carbide. The stability of the ternary carbide was also reported in Ref. 23. The isothermal section is very useful for evaluation of the processes, which occur during infiltration. Actually, the infiltration of the porous B4 C preforms was performed at 1123 K in a semi-hermetically closed container, from which the evacuation of Mg vapor is significantly reduced. Under these conditions, a non-stable liquid phase may coexist with its vapor and with B4 C for a relatively prolonged period, for at least 20 min, as required for full infiltration. In order to analyze the interaction of liquid Mg with B4 C under the given conditions, which correspond to spontaneous evaporation of Mg, a quasi-equilibrium isothermal section of the Mg–B–C system in the Mg-rich corner was constructed (Fig. 6). This isothermal section was used for analyzing the interaction that takes place during the infiltration process.
Fig. 5. Calculated isothermal section of the Mg–B–C ternary phase diagram at 1000 K and 133.3 Pa.22 The phases that are in equilibrium within the various regions are marked in the figure.
Fig. 6. A schematic view of the Mg rich corner of the isothermal section of the Mg–B–C phase diagram at 1123 K.
The points 1 and 4 were taken from the binary phase diagrams20,21 and correspond to the solubility of C and B in liquid Mg. The points 2 and 3 correspond to the carbon and boron solubility in the ternary Mg–B–C liquid solution. The dotted curve 2–3 divides the diagram into two regions: the L-single phase region that corresponds to a liquid Mg–C–B solution and the two-phase region L + MgB2 C2 , in which the liquid solution coexists with the ternary carbide. The three phase regions are also shown in the diagram. The straight line 6–7 connects the Mg corner with the point corresponding to B4 C and each point on this line denotes the ratio of B4 C and liquid Mg in the starting mixture. We assume that in our case only a thin surface layer of B4 C particles reacts with liquid Mg and the total composition of the system corresponds to point A, which is invariant, since it represents the initial composition, which remains constant. This point is quite close to Mg corner and located in the three phase L + MgB2 C2 + MgB2 region. The composition of the liquid solution, which is in equilibrium with two non-metallic phases, corresponds to the point 3. According to this diagram, B4 C cannot be in equilibrium with the Mg–C–B liquid solution and has to transform into the ternary MgB2 C2 carbide and MgB2 . We suggest that this transformation occurs through congruent B4 C dissolution, followed by precipitation of new phases on the surface of the initial B4 C particles. This mechanism was discussed in details in Ref. 15 in connection with the B4 C–silicon system. In the present case, the congruent B4 C dissolution dictates the release of B and C from the carbide particles in the 4:1 atomic ratio, which does not corresponds to equilibrium conditions, defined by the point 3 in Fig. 6. In order to achieve the equilibrium state the composition of liquid solution has to be shifted toward point 3 and the precipitation of the ternary carbide MgB2 C2 and MgB2 has to take place. The process will stop when the free surface of the B4 C particles is completely covered by the new phases. This suggested scenario of the interaction is in good agreement with the results of the XRD analysis and the microstructural characterization.
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Table 2 The properties of the raw materials and the Mg infiltrated composites (the processing temperature and the properties of B4 C–Si composites are presented for comparison). Composite
Process temperature (K)
Fraction of a metallic phase (vol.%)
Density (g/cm3 )
Microhardness (Hv)
Bending strength (MPa)
Young modulus (GPa)
Weibull modulus
Mg AZ91 B4 C [24] B4 C–Si [25] M1 M2
– – 2473 1773 1123 1123
– – 0 ∼10 ∼10 ∼5
1.74 1.81 2.52 2.54 2.40 ± 0.02 2.44 ± 0.02
43 70 2910–3770 2300 ± 250 857 ± 63 1148 ± 141
95 130 350–480 318 ± 20 341 ± 16 327 ± 19
45 45 390–440 400 ± 10 269 ± 4 295 ± 3
– – 5 18.9 11.3 11.7
3.4. Mechanical properties The properties of the composites are summarized in Table 2. The mechanical and physical properties reflect the phase composition of the composites. The low density is determined by the density of the starting materials, while the other properties depend on the ceramic to metal ratio in the final composites. A high fraction of the nonmetallic phases stands behind the relatively high Young modulus and hardness values, while the bending strength of the composites and the high value of the Weibull modulus reflect the homogeneous microstructure and the absence of coarse defects with a dimension larger than their critical size. The differences in the hardness and Young modulus values of M1 and M2 composites are attributed to the higher fraction of the nonmetallic phases in M2 composites, due to the presence of Al involved in the interface interaction. Although, the properties of the magnesium infiltrated composites are lower than that for B4 C–Si composites, the significant advantage is the lowering of the process temperature from 1773 K to 1123 K. The mechanical properties presented in Table 2 may be useful and provide guidelines for engineers dealing with materials selection for various applications. 4. Conclusions Fully dense Mg/B4 C and AZ91/B4 C composites, with more than 80 vol.% nonmetallic phases were fabricated by free infiltration under a Mg vapor atmosphere. The interaction between B4 C and the liquid metal leads to the formation of the MgB2 and MgB2 C2 phases as a new interfacial layer, which connects the initial B4 C particles and gives rise the formation of so-called reaction bonded composites. A thermodynamic analysis of the Mg–B–C system provides an explanation for the interfacial layer formation. The composites display low density and relatively high values of the Young modulus and hardness. References 1. Chapman TR, Niesz DE, Fox RT, Fawcett T. Wear-resistant aluminum–boron–carbide cermets for automotive brake applications. Wear 1999;236:81–7. 2. Shorowordi KM, Laoui T, Haseeb ASMA, Celis JP, Froyen L. Microstructure and interface characteristics of B4 C, SiC and Al2 O3 reinforced Al matrix composites: a comparative study. J Mater Process Technol 2003;142:738–43.
3. Tang F, Wu X, Ge S, Ye J, Zhu H, Hagiwara M, et al. Dry sliding friction and wear properties of B4 C particulate-reinforced Al-5083 matrix composites. Wear 2008;264:555–61. 4. Topcu I, Gulsoy H, Kadioglu N, Gulluoglu A. Processing and mechanical properties of B4 C reinforced Al matrix composites. J Alloys Compd 2009;482:516–21. 5. Canakci A. Microstructure and abrasive wear behaviour of B4 C particle reinforced 2014 Al matrix composites. J Mater Sci 2011;46: 2805–13. 6. Kalaiselvan K, Murugan N, Parameswaran S. Production and characterization of AA6061-B4 C stir cast composite. Mater Des 2011. 7. Mohanty R, Balasubramanian K, Seshadri S. Boron carbide-reinforced alumnium 1100 matrix composites: fabrication and properties. Mater Sci Eng A 2008;498:42–52. 8. Jiang Q, Wang H, Ma B, Wang Y, Zhao F. Fabrication of B4 C particulate reinforced magnesium matrix composite by powder metallurgy. J Alloys Compd 2005;386:177–81. 9. Badini C, Marino F, Montorsi M, Guo X. Precipitation phenomena in B4 C-reinforced magnesium-based composite. Mater Sci Eng A 1992;157: 53–61. 10. Kevorkijan V, Skapin S. Characterisation of Mg–B4 C composites with a high volume fraction of fine ceramic reinforcement fabricated by pressureless infiltration of porous ceramic preforms. J Adv Mater 2010;42: 35–47. 11. Aghajanian M, McCormick A, Marshall A, Waggoner W, Karandikar P. Static and dynamic properties of Mg/ceramic MMCS. Adv Ceram Armor VI 2010:79–86. 12. Taylor KM. Dense carbide composite for armor and abrasives; 1973. 13. Aghajanian M, Morgan B, Singh J, Mears J, Wolffe R. A new family of reaction bonded ceramics for armor applications. Ceram Trans 2002;134:527–39. 14. Hayun S, Dilman H, Dariel M, Frage N, Dub S. The Effect of Carbon Source on the Microstructure and the Mechanical Properties of Reaction Bonded Boron Carbide, Advances in Sintering Science and Technology. John Wiley & Sons Inc.; 2010. pp. 29–39. 15. Hayun S, Weizmann A, Dariel MP, Frage N. Microstructural evolution during the infiltration of boron carbide with molten silicon. J Eur Ceram Soc 2010;30:1007–14. 16. Zhang D, Shen P, Shi L, Jiang Q. Wetting of B4 C, TiC and graphite substrates by molten Mg. Mater Chem Phys 2011;130:665–71. 17. Frage N, Levin L, Frumin N, Gelbstein M, Dariel M. Manufacturing B4 C–(Al, Si) composite materials by metal alloy infiltration. J Mater Process Technol 2003;143:486–90. 18. Halverson DC, Pyzik AJ, Aksay IIA, Snowden WE. Processing of boron carbide–aluminum composites. J Am Ceram Soc 1989;72: 775–80. 19. Pyzik AJ, Beaman DR. Al–B–C phase development and effects on mechanical properties of B4 C/Al-derived composites. J Am Ceram Soc 1995;78:305–12. 20. Hu B, Du Y, Xu H, Sun W, Zhang W, Zhao D. Thermodynamic description of the C–Ge and C–Mg systems. J Min Metall B Metall 2010;46: 97–103.
M. Cafri et al. / Journal of the European Ceramic Society 32 (2012) 3477–3483 21. Xi X, Zeng X, Soukiassian A, Jones J, Hotchkiss J, Zhong Y, et al. Thermodynamics and thin film deposition of MgB2 superconductors. Supercond Sci Technol 2002;15:451. 22. Saengdeejing A, Saal JE, Wang Y, Liu ZK. Effects of carbon in MgB2 thin films: intrinsic or extrinsic. Appl Phys Lett 2007;90:151920–3. 23. Wörle M, Nesper R. MgB2 C2 , a new graphite-related refractory compound. J Alloys Compd 1994;216:75–83.
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24. Thévenot F. Boron carbide—a comprehensive review. J Eur Ceram Soc 1990;6:205–25. 25. Hayun S, Weizmann A, Dariel MP, Frage N. The effect of particle size distribution on the microstructure and the mechanical properties of boron carbide-based reaction-bonded composites. Int J Appl Ceram Technol 2009;6:492–500.
Journal of the European Ceramic Society 32 (2012) 3477–3483
Boron carbide/magnesium composites: Processing, microstructure and properties M. Cafri, H. Dilman, M.P. Dariel, N. Frage ∗ Department of Materials Engineering, Ben-Gurion University of the Negev, P.O. Box 653, Beer-Sheva 84105, Israel Received 14 December 2011; received in revised form 27 March 2012; accepted 1 April 2012 Available online 3 May 2012
Abstract Composites with a high fraction of the ceramic phase were fabricated by infiltration of 80% dense preforms with liquid Mg or Mg alloys, under an Mg vapor atmosphere. The infiltration was performed at 1123 K in a semi-hermetically closed container, from which gas had been evacuated. The Mg vapor atmosphere was achieved by heating subsequently to the relatively high temperature of the Mg infiltration process. The specific weight of the composites is about 2.44 g/cm3 . The microstructure of the composite consists of the newly formed MgB2 and ternary carbide MgB2 C2 phases that connect the initial B4 C particles and some residual metal. A thermodynamic analysis of the interaction between B4 C and liquid Mg was conducted and its results are in good agreement with the experimental observations. The mechanical properties of the composites were investigated and discussed. © 2012 Elsevier Ltd. All rights reserved. Keywords: Composites; Carbides; Microstructure; Mechanical properties
1. Introduction Ceramic metal composites are advanced materials of interest on account of the potential synergy provided by the combination of metallic and ceramic properties. This group includes metal matrix composites (MMCs) and ceramic matrix composites (CMCs). Boron carbide (B4 C) with its low density, elevated modulus and hardness values is an attractive ceramic for applications such as light armor and as a wear resistant component in structural materials for land borne and airborne transport vehicles. MMCs may be fabricated by a stir-casting process or by infiltration of porous ceramic preforms with a liquid metal. The fraction of ceramic particles that may be distributed homogenously in the melt in the course of the stir-casting process is of the order of 30 vol.%. This value may be significantly higher for a free infiltration approach. The metal-to-ceramic ratio in the composites determines their mechanical properties and their possible applications. Metal ceramic composites (MMCs) based on aluminum alloys strengthened with B4 C particles have been the subject of
∗
Corresponding author. Tel.: +972 8 6461468; fax: +972 8 6479350. E-mail address: [email protected] (N. Frage).
0955-2219/$ – see front matter © 2012 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.jeurceramsoc.2012.04.007
extensive studies.1–7 It was of obvious interest to extend these studies to Mg-based MMCs, considering the reduced specific weight of the latter. The fabrication conditions of the boron carbide–magnesium composites were discussed in Refs. 8–11. A maximal fraction of 50 vol% ceramic phase in these composites was attained by Kevorkijan and Scapin10 by free infiltration of porous boron carbide preforms. Higher B4 C content composites have been successfully manufactured by the so-called reaction bonding approach, which consists of infiltrating a green B4 C powder preform with molten silicon with or without the addition of free carbon.12–15 The final microstructure of the reaction bonded boron carbide composites (RBBC) composite consists of a continuous B4 C skeleton, silicon carbide (SiC), ternary B12 (Si,C,B)3 and residual Si. The high strength of the resulting material is due to the in situ reaction between carbon that has been added or that originates in the carbide, with molten silicon to form a SiC network. It was reported in our previous publication16 that the fraction of the residual Si may be significantly reduced (to 8–10 vol.%) by the infiltration of porous preforms with a relatively high green density (of about 80%), fabricated from a mixture of boron carbide powders with an appropriate particle size distribution. In the present study, we have used this approach in order to manufacture B4 C-based composites infiltrated with molten Mg and Mg alloys. A similar
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conceptual approach using Mg or Mg alloys infiltration instead of molten Si presents, in addition to the reduced specific weight, the advantage of lowering the processing temperature from close to 1773 K down to 1123 K. On the other hand it remains to be seen to what extent the reaction between B4 C and molten Mg or its alloys provides and adequate cohesive strength to the composite. Three conditions have to be fulfilled in order for the spontaneous infiltration to take place, namely: (a) adequate wetting of the porous preform by the liquid metal; (b) gas evacuation from the pores; (c) moderate chemical interaction to avoid the formation of new phases that may plug up open channels in the porous preform and impede, thereby, infiltration. The contact wetting angle between the ceramic substrate and the liquid infiltrant provides a measure of the wetting in the system. Its direct measurement in the Mg/B4 C system is a complex task due to the high partial vapor pressure and the reactivity of Mg. Zhang et al.16 have recently reported that a transition from nonwetting to wetting condition occurs in the 972–1173 K temperature range and have attributed it to the chemical interaction at the metal/carbide interface. The attempts to fabricate MMCs based on Mg or its alloys and B4 C particles8–11 have shown that under the appropriate atmosphere, the wetting angle is sufficient for free infiltration. For instance, in Ref. 10 the infiltration was conducted under a flow of oxygen-free nitrogen and with Ti addition as a wetting agent. Under these conditions B4 C preforms, prepared from a bimodal mixture of particles (0.8 and 50 m), with a fraction of fine powder less than 60 vol.% were successfully infiltrated with liquid Mg at 1123 K. Preforms with a higher fraction of fine powder cannot be infiltrated under similar conditions. In the latter case, such powder mixtures have a relatively high surface area and a higher oxygen content due to the presence of the thin boron oxide films on the particles surface. An increased content of oxygen leads to massive MgO formation that impedes the infiltration of the B4 C porous body. The presence of MgO at the Mg/B4 C interface was detected by Jiang et al.8 Due to high Mg affinity to oxygen it is impossible to prevent its oxidation under any gaseous atmosphere. In order to circumvent the issue we made use of Mg vapor atmosphere generated by Mg evaporation. The infiltration process was conducted in a semi-hermetically closed container, from which gas had been evacuated before heating and a Mg vapor atmosphere was generated at the relatively high temperature of the infiltration process. In the present work, the interface interaction in the Mg/B4 C system and the mechanical and physical properties of the infiltrated composites with about 80 vol.% of B4 C were investigated. 2. Experimental 2.1. Samples preparation Two sets of samples were fabricated from commercial B4 C powder mixtures that allowed reaching about 80% green density preforms. Fine B4 C powder (∼3 m, grade HS, 97.5% purity) was supplied by H.C. Starck and B4 C powders with average particle sizes of 13, 50 and 100 m (97% purity) were supplied by a
Chinese Company “Mudanjiang”. Samples with the appropriate particle size distribution (3 m – 19 vol.%, 13 m – 18 vol.%, 50 m – 10 vol.%, 106 m – 53 vol.%) were dry mixed for 16 h. Preforms, 20 mm diameter and 4 mm height, were compacted uniaxially (180 MPa) yielding compacts with about 20 vol.% porosity. Infiltration was done by placing Mg or Mg alloy AZ91 (Mg–9 wt.%Al–0.7 wt.%Zn–0.3 wt.%Mn) pieces on the top of preforms. The preforms with metal pieces on their top were positioned in a semi-hermetically closed graphite container (the cover was not completely sealed). The latter had been evacuated by pumping to ∼1 Pa vacuum for 2 h prior heating. The infiltration process was conducted at 1123 K for 20 min. At this temperature, the equilibrium Mg vapor pressure (∼6.6 kPa) is higher than the vacuum level in the furnace chamber and, although Mg escapes partly from the crucible, the infiltration process takes place under Mg vapor. The initial amount of the metal for infiltration was twice that required for filling the free volume in the preforms. The infiltrated composites were cooled down to room temperature at a 10 ◦ C/min cooling rate. The preforms, infiltrated with pure Mg and AZ91, are denoted as M1 and M2, respectively. 2.2. Wetting experiments Fully dense boron carbide substrates were used for the wetting experiments. Small pieces of Mg and AZ91 (2–3 g) were placed on the substrate and melted under a Mg vapor atmosphere in a semi-hermetically closed graphite container. The contact angles were measured using the optical image of the drops after cooling and solidification. 2.3. Characterization 2.3.1. Microstructure and composition The microstructure of the samples was studied by scanning electron microscopy (SEM, JEOL-5600) combined with an energy-dispersive spectrometer (EDS). The samples for the microstructural characterization were prepared using a standard metallographic procedure that included a last stage of polishing by 0.5 m diamond paste. The phase composition of the infiltrated composites was determined by X-ray diffraction (XRD) using Rigaku RINT 2100 Bragg-Brentano diffractometer with Cu K␣ radiation, the scanning angle 2θ was increased in 0.02◦ steps. 2.3.2. Mechanical and physical properties The micro-hardness was determined using a Buehler tester (Micromet 2100) with a Vickers indenter under 20 N load. The flexural strength was determined by three-point bending test with a LRX plus LLOYD machine (Lloyd Instruments, Fareham Hants, U.K.), on 1.5 mm × 2 mm × 10 mm samples. The elastic modulus of the composites was derived from ultrasonic sound velocity measurements using the “Pulse Echo” method. The density of the infiltrated composites was determined by the Archimedes method.
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Fig. 1. Optical images of cross-sectioned drops and B4 C substrate: (a) pure Mg and (b) AZ91.
Fig. 2. SEM images of the composites microstructure: (a) M1 and (b) M2.
3. Results and discussion 3.1. Wetting and infiltration
diffraction (Fig. 3) showed that both the M1 and M2 composites contain the Mg, B4 C, MgB2 and MgB2 C2 phases. The M2 composite also contains the Mg17 Al12 phase (␥phase), which is expected to precipitate from the supersaturated
The optical images of the drops after solidification and cooling are presented in Fig. 1. The contact angle of pure Mg is about 25–30◦ , while that angle for AZ91 Mg-alloy is significantly lower (∼12◦ ). The wetting in these systems is governed by an interface chemical interaction and the formation of new phases at the interface (see below). The important conclusion is that the values of the contact angle are low enough for free infiltration. The microstructure of the composites (Fig. 2) indeed supports the successful infiltration of the porous preforms with liquid Mg and AZ91 and the absence of any residual porosity. 3.2. Phase composition of the infiltrated composites In order to investigate the phase composition of the composites and to clarify the chemical interaction that takes place at the metal/ceramic interface, the composites were fabricated by infiltration of porous preforms that had been compacted exclusively of coarse B4 C powder particles (of about 50 m). The X-ray
Fig. 3. XRD patterns of Mg, AZ91, M1 and M2 composites.
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Table 1 The estimated phase composition of the composites. Composite
Mg (vol.%)
B4 C (vol.%)
MgB2 (vol.%)
MgB2 C2 (vol.%)
Mg17 Al12 (vol.%)
M1 M2
10 5
76 75
14 13
<1 3
0 4
Mg(Al,Zn,Mn) solid solution. In the AZ91 and M2 composites the Mg diffraction peaks are shifted to higher angels than in the pure Mg and M1 composites on account of the presence of alloying elements, namely Al and Zn, in the Mg solid solution. Alumina, which was identified by XRD analysis, originates from the powder preparation stage form the bulk composite specimens (milling in an alumina vessel using alumina balls). The diffraction data were used for the semi-quantitative analysis of the volume fraction of each phase within the composites (Table 1). Composite M1 contains 10 vol.% of residual Mg and the dominant newly formed phase is MgB2 . The M2 composite contains only 5 vol.% of residual metallic phase and a relatively high fractions of the ternary MgB2 C2 carbide and Mg17 Al12 intermetallic. We suggest that due to the high affinity of Al to carbon and boron, it enhances the formation of the ternary carbide, which according to the EDS analysis (Fig. 4d) contains some fraction of aluminum.
The results of the XRD analysis are supported by the microstructural characterization of the infiltrated composites (Fig. 4). The surface of each B4 C particle is surrounded by a continuous layer of newly formed phases that reflect a chemical interaction at the metal/ceramic interface. This interaction is more massive in the B4 C/AZ91 composite, probably, due to the presence of Al in the melt and its involvement in the interface reaction. EDS analysis indicates that Al is actually present (Fig. 4d) in the interfacial layer. This observation is a good agreement with results on Al and B4 C interaction reported in the literature.17–19 The important point is that a new interfacial layer connects the B4 C particles and gives rise to the formation of a so-called reaction bonded composite. The cohesive strength between the newly formed intermediate layer and the B4 C particles stands behind the relatively elevated mechanical properties of the composite. The thermodynamic analysis presented
Fig. 4. SEM images of the composites (a – M1, b – M2) fabricated using coarse B4 C. Graphs c and d show the results of line scans along the white lines in the inserts.
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below provides the underlying reasons for the interfacial layer formation. 3.3. Interaction between B4 C and liquid Mg: thermodynamic considerations According to the binary Mg–C phase diagram,20 the solubility of C in liquid Mg is relatively high (of about 20 at.% at 1123 K) and magnesium carbide is stable only at temperatures lower than 873 K. The binary Mg–B phase diagram was calculated and discussed by Xi et al.21 In the temperature interval of our interest, only MgB2 coexists with liquid and gaseous Mg. The boron solubility limit in liquid magnesium at 1123 K is about 0.01 at.%. The calculated22 isothermal section of the ternary Mg–C–B system at 1000 K and 133.3 Pa is presented in Fig. 5. This isothermal section corresponds to the phase equilibria between gaseous Mg (Mg vapor pressure at this temperature is about 6.6 kPa) and various nonmetallic phases, including the ternary MgB2 C2 carbide. The stability of the ternary carbide was also reported in Ref. 23. The isothermal section is very useful for evaluation of the processes, which occur during infiltration. Actually, the infiltration of the porous B4 C preforms was performed at 1123 K in a semi-hermetically closed container, from which the evacuation of Mg vapor is significantly reduced. Under these conditions, a non-stable liquid phase may coexist with its vapor and with B4 C for a relatively prolonged period, for at least 20 min, as required for full infiltration. In order to analyze the interaction of liquid Mg with B4 C under the given conditions, which correspond to spontaneous evaporation of Mg, a quasi-equilibrium isothermal section of the Mg–B–C system in the Mg-rich corner was constructed (Fig. 6). This isothermal section was used for analyzing the interaction that takes place during the infiltration process.
Fig. 5. Calculated isothermal section of the Mg–B–C ternary phase diagram at 1000 K and 133.3 Pa.22 The phases that are in equilibrium within the various regions are marked in the figure.
Fig. 6. A schematic view of the Mg rich corner of the isothermal section of the Mg–B–C phase diagram at 1123 K.
The points 1 and 4 were taken from the binary phase diagrams20,21 and correspond to the solubility of C and B in liquid Mg. The points 2 and 3 correspond to the carbon and boron solubility in the ternary Mg–B–C liquid solution. The dotted curve 2–3 divides the diagram into two regions: the L-single phase region that corresponds to a liquid Mg–C–B solution and the two-phase region L + MgB2 C2 , in which the liquid solution coexists with the ternary carbide. The three phase regions are also shown in the diagram. The straight line 6–7 connects the Mg corner with the point corresponding to B4 C and each point on this line denotes the ratio of B4 C and liquid Mg in the starting mixture. We assume that in our case only a thin surface layer of B4 C particles reacts with liquid Mg and the total composition of the system corresponds to point A, which is invariant, since it represents the initial composition, which remains constant. This point is quite close to Mg corner and located in the three phase L + MgB2 C2 + MgB2 region. The composition of the liquid solution, which is in equilibrium with two non-metallic phases, corresponds to the point 3. According to this diagram, B4 C cannot be in equilibrium with the Mg–C–B liquid solution and has to transform into the ternary MgB2 C2 carbide and MgB2 . We suggest that this transformation occurs through congruent B4 C dissolution, followed by precipitation of new phases on the surface of the initial B4 C particles. This mechanism was discussed in details in Ref. 15 in connection with the B4 C–silicon system. In the present case, the congruent B4 C dissolution dictates the release of B and C from the carbide particles in the 4:1 atomic ratio, which does not corresponds to equilibrium conditions, defined by the point 3 in Fig. 6. In order to achieve the equilibrium state the composition of liquid solution has to be shifted toward point 3 and the precipitation of the ternary carbide MgB2 C2 and MgB2 has to take place. The process will stop when the free surface of the B4 C particles is completely covered by the new phases. This suggested scenario of the interaction is in good agreement with the results of the XRD analysis and the microstructural characterization.
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Table 2 The properties of the raw materials and the Mg infiltrated composites (the processing temperature and the properties of B4 C–Si composites are presented for comparison). Composite
Process temperature (K)
Fraction of a metallic phase (vol.%)
Density (g/cm3 )
Microhardness (Hv)
Bending strength (MPa)
Young modulus (GPa)
Weibull modulus
Mg AZ91 B4 C [24] B4 C–Si [25] M1 M2
– – 2473 1773 1123 1123
– – 0 ∼10 ∼10 ∼5
1.74 1.81 2.52 2.54 2.40 ± 0.02 2.44 ± 0.02
43 70 2910–3770 2300 ± 250 857 ± 63 1148 ± 141
95 130 350–480 318 ± 20 341 ± 16 327 ± 19
45 45 390–440 400 ± 10 269 ± 4 295 ± 3
– – 5 18.9 11.3 11.7
3.4. Mechanical properties The properties of the composites are summarized in Table 2. The mechanical and physical properties reflect the phase composition of the composites. The low density is determined by the density of the starting materials, while the other properties depend on the ceramic to metal ratio in the final composites. A high fraction of the nonmetallic phases stands behind the relatively high Young modulus and hardness values, while the bending strength of the composites and the high value of the Weibull modulus reflect the homogeneous microstructure and the absence of coarse defects with a dimension larger than their critical size. The differences in the hardness and Young modulus values of M1 and M2 composites are attributed to the higher fraction of the nonmetallic phases in M2 composites, due to the presence of Al involved in the interface interaction. Although, the properties of the magnesium infiltrated composites are lower than that for B4 C–Si composites, the significant advantage is the lowering of the process temperature from 1773 K to 1123 K. The mechanical properties presented in Table 2 may be useful and provide guidelines for engineers dealing with materials selection for various applications. 4. Conclusions Fully dense Mg/B4 C and AZ91/B4 C composites, with more than 80 vol.% nonmetallic phases were fabricated by free infiltration under a Mg vapor atmosphere. The interaction between B4 C and the liquid metal leads to the formation of the MgB2 and MgB2 C2 phases as a new interfacial layer, which connects the initial B4 C particles and gives rise the formation of so-called reaction bonded composites. A thermodynamic analysis of the Mg–B–C system provides an explanation for the interfacial layer formation. The composites display low density and relatively high values of the Young modulus and hardness. References 1. Chapman TR, Niesz DE, Fox RT, Fawcett T. Wear-resistant aluminum–boron–carbide cermets for automotive brake applications. Wear 1999;236:81–7. 2. Shorowordi KM, Laoui T, Haseeb ASMA, Celis JP, Froyen L. Microstructure and interface characteristics of B4 C, SiC and Al2 O3 reinforced Al matrix composites: a comparative study. J Mater Process Technol 2003;142:738–43.
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