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MgAl2O4 composites by single-stage pressureless infiltration
Accepted Manuscript Tailoring microstructure and properties of bilayer-graded Al/B4C/MgAl2O4 composites by single-stage pressureless infiltration A. Bahrami, M.I. Pech-Canul, N. Soltani, C.A. Gutiérrez, P.H. Kamm, A. Gurlo PII:
S0925-8388(16)33047-X
DOI:
10.1016/j.jallcom.2016.09.284
Reference:
JALCOM 39118
To appear in:
Journal of Alloys and Compounds
Received Date: 18 April 2016 Revised Date:
8 September 2016
Accepted Date: 25 September 2016
Please cite this article as: A. Bahrami, M.I. Pech-Canul, N. Soltani, C.A. Gutiérrez, P.H. Kamm, A. Gurlo, Tailoring microstructure and properties of bilayer-graded Al/B4C/MgAl2O4 composites by single-stage pressureless infiltration, Journal of Alloys and Compounds (2016), doi: 10.1016/j.jallcom.2016.09.284. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
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Tailoring microstructure and properties of bilayer-graded Al/B4C/MgAl2O4 composites by single-stage pressureless infiltration A. Bahrami 1,2, M. I. Pech-Canul1*, N. Soltani 1,2, C.A. Gutiérrez 1, P. H. Kamm3, A. Gurlo4 1
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Centro de Investigación y de Estudios Avanzados del IPN Unidad Saltillo, Ave. Industria Metalúrgica No. 1062, Parque Industrial Saltillo-Ramos Arizpe, Ramos Arizpe, Coahuila, México, 25900. 2 Technische Universität Berlin, Fakultät III Prozesswissenschaften, Institut für Werkstoffwissenschaften und technologien, Fachgebiet Keramische Werkstoffe, Hardenbergstraße 40, 10623 Berlin, Germany. 3 Applied Materials, Helmholtz-Zentrum Berlin für Materialien und Energie, Germany. 4 Chair of Advanced Ceramic Materials, Technische Universität Berlin, Germany.
Abstract: The quantitative effect of the following parameters on the one single step pressureless
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infiltration characteristics of bilayer (B4C)p/rice-husk ash (RHA) porous preforms by aluminum alloys was investigated using the Taguchi method and analysis of variance (ANOVA): infiltration temperature and time, B4C particle size, RHA percentage, preform porosity and alloy magnesium content. Contribution determinations of each of the parameters on the bending
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strength of the resulting bilayer composites indicated that the parameter with the most significant effect is process temperature, with a contribution percentage of 79 %. Verification tests conducted using the established optimum parameters show good agreement with those of
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projected values. A detailed examination of microstructure and potential chemical reactions, aided by thermodynamic analysis, allowed establishing conditions for suppressing or
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diminishing the impact of the deleterious Al4C3 phase. Keywords: One-step pressureless infiltration; Graded materials; Rice-husk ash; Mechanical properties.
1 * Corresponding author. Tel.: +52 (844) 4-38-96-00; Ext.: 8678 Email address: [email protected] (Martin I. Pech-Canul); [email protected]
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1 Introduction In spite of the outstanding physical and mechanical properties, monolithic B4C seems to not live up to its true potential as an advanced structural ceramic due its extreme sensitivity to brittle
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fracture and the difficulties associated with fabricating fully dense microstructures[1]. The experiments have demonstrated that these issues can be significantly reduced by introducing a metal phase [2]. By far, the most popular metal used in conjunction with B4C for this purpose is
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aluminum [3-6]. In addition to availability and lightness of Al-Mg-Si alloys, they wet B4C well at elevated temperatures. Molten aluminum alloy has been shown to form a variety of binary and
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ternary phases when in contact with B4C such as Al3BC, AlB10 and Al4C3 [7]. The manufacturing process of a functionally graded materials (FGM) can usually be divided into building the spatially inhomogeneous structure “gradation” and transformation of this structure into a bulk material “consolidation”. The fabrication methods of functionally graded composite
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materials can be similar or totally different from the conventional methods used for fabrication of metal matrix composites [8-15]. These methods include but are not limited to: powder stacking, centrifugal method, electromagnetic separation, casting, PVD, CVD, etc.
[16-21]. Die
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compaction of layers (powder stacking), which was used in this study to make graded reinforcement preforms, is a simple and well established method in which a gradient is formed
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by the deposition of powder layers with changing the size and composition of powder in the compacting die [22].
Although the infiltration of ceramic preforms by liquid metals has been typically applied for the processing of composite materials with homogeneous reinforcement shape, size and composition [23, 24], as a constitutive route, it offers the potential for the production of graded materials by variation of shape, size, and volume fraction of the reinforcement in each layer.
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In this investigation, it is postulated that graded metal matrix composite materials may be produced by one-step infiltration process. In spite of various efforts devoted to the optimization of wettability and pressureless infiltration parameters [25, 26], it is expected that in pressureless
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infiltration-based processing, changes of key parameters (reinforcement composition, size, percentage porosity in the preform) from one section to another within the ceramic preform (layer by layer) will increase the inherent barriers and challenges. It was reported that some of
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the problems associated with the infiltration process can be solved by addition of special elements or phases [25, 27]. Furthermore, addition of reinforcement phases with special
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morphology and chemical composition can enhance the mechanical and physical properties of the composites. Rice husk ash (RHA) as a source of silicon and oxygen with truss like and spiral like structure can be used as a unique reinforcement phase to satisfy these expectations [28, 29]. The aim of this investigation is to gain further insight into the processing of functionally graded
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Al/B4Cp/MgAl2O4 composites by pressureless infiltration. In this article, using Taguchi method, the quantitative effect of process temperature and time, B4Cp particle size in the preform layer, crystallinity of SiO2 obtained from rice husk ash, ceramic preform porosity and Mg
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concentration in the aluminum alloys on the microstructure, physical and mechanical properties of composites prepared from RHA/B4Cp ceramic preform stacked layers, is investigated.
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Analysis of variance (ANOVA) is a method to determine the variability in the response parameters in an experiment, by systematically changing the levels in the independent variables, identifying sources of variation and the associated degrees of freedom [30, 31]. In this study ANOVA was performed on the data in order to determine the quantitative effect of each parameter and each interaction on the modulus of elasticity and bending strength, CTE value and electrical resistivity of the graded composites. ANOVA provides a means of estimating the
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percent contribution of each of the parameters tested to the variability in the measured response variable (i.e., bending strength, modulus of elasticity, etc.). In addition, ANOVA provides insight
variable produced when using the optimum process parameters. 2
Experimental procedure
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into the optimum process parameters, and allows the estimation of the magnitude of the response
Infiltration tests for Al alloys/B4C/RHA composites were conducted under the conditions shown
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in Table 1. The effect of various parameters on the physical and mechanical properties of composite were investigated using the Taguchi method for design of experiments. Predetermined
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amounts of ceramic powders with 5wt. % Dextrin from potato starch (Fluka®, Switzerland) as a binder were soaked and mixed with 5 wt. % of distilled water. The mixture of first layer was put in the mold and uniaxially pressed to a pressure of 20 MPa and up to 5mm height and then the mixture of second layer was added and pressed to the final height of preform (10mm). The
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amount of needed mixture to fabricate porous preform was calculated according to density formula (ρ=m/V) considering the porosity percentage in the volume of each layer. Using adequate amount of binder (Dextrin), prevents cracks formation and changes in the desirable
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porosity values in the preform. The preforms were heated up to 400 °C in a furnace open to the atmosphere to remove the organic components. In rows 1 and 2 of Table 1, the numbers on the
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left and the right, correspond to the values of the top and bottom preform layers, respectively. Accordingly, each of the selected parameters was tested at two levels. (Table 1)
Table 2 shows the parameters and levels that are investigated in this work arranged in a standard Taguchi L12 orthogonal array. (Table 2) 4
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The chemical compositions of the two alloys used in the experiment are presented in table 3. The structure and phase type of silica derived from rice husk depend on calcining temperature, atmosphere and chemical treatment of rice husk before and after burning. Two procedures were
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selected for extraction of silica from rice husk. The extraction procedure and physical properties of obtained silica corresponding to each method are reported elsewhere [25]. Infiltration trials were performed in a horizontal tube furnace with a 6.5 cm diameter alumina tube closed at both
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ends with end-cap fittings to control the process atmosphere. More details about the process are presented elsewhere [25]. The mold-preform-metal assembly was placed in the center of the
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alumina tube and heated with a heating rate of 20 °C/min in ultra-high purity argon up to test temperature. The three dimensional structure of the preform was analyzed by using a Hamamatsu microfocus X-ray source and flat panel sensor. The voxel size of the obtained reconstructed volume was 7 µm. The volume was visualized by VGStudio MAX 2.2. Specimens for
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microstructure analysis were mounted and polished using standard metallurgical procedures, and the analysis was done using X-ray diffraction (XRD), scanning electron microscopy (SEM) and energy dispersive X-rays (EDXs). Thermodynamic calculations were carried out using the
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FactSageTM software and FTlite databases. FactSageTM is the fusion of the FACTW in /F*A*C*T and ChemSage/mn SOLGASMIX thermochemical packages. The FTIR analysis was
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performed using Bruker FT-IR Spectrometer (Equinox 55). The bending strength of prepared monolayer and bilayer composites was determined using four-point bending tests according to the ASTM 1161 standard procedures. Furthermore, the modulus of elasticity of composites was determined using data from mechanical testing and from ultrasonic tests conducted according to the ASTM: E494-95 standard procedures. The hardness of the samples was determined by
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Vickers indentation tests (Tuckson-300) on the polished surfaces, using a load of 1 kg and a dwell time of 15 s.
3 3.1
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(Table 3) Result and Discussion Microstructure examination
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Cracks and non-uniformity in reinforcement particles distribution can be generated during preform fabrication and/or melt infiltration. Therefore, the preforms should be made sound and
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without any cracks to the extent possible. Figure 1 shows the three-dimensional image of B4C/RHA bilayer preforms before infiltration. As it can be observed from figure 1 (a), the layer with 60 vol. % porosity containing fine B4C particles is represented by the darker tone. As the atomic number of B4C is lower than that of the RHA, the RHA appears in lighter shades of gray
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(Figure 1a). The inspection shows that the inner structure is dense and particles and porosities (Figure 1b) are distributed uniformly. No defects in the scale of higher than 5 mm, such as
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microcracks are found in the inner structure of the preform. (Figure 1)
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The layered structures of composites fabricated according to configurations 7 and 4, with abrupt microstructural changes, are shown in figures 2 and 3, respectively. As it can be confirmed by XRD (figure 4), both layers of fabricated composites contain different phases that depend on the processing parameters. As it can be seen from figure 2, by choosing adequate processing parameters, the presence of defects in the fabricated composite, such as those shown in figure 3, can be eliminated. It can be observed that in the layer with small size of B4C, due to the high surface area of the fine B4C particles, the reaction of Al and B4C is intensified and more ceramic 6
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phases can be observed than in the layer which contains B4C of large size. On the other hand, transformation of RHA to MgAl2O4 strongly depends on the crystallinity and procedure conditions. The EDS results presented in figure 2 confirm the complete transformation of RHA
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to MgAl2O4. Also the good bonding between matrix alloy and B4C particles is illustrated in figure 2. In addition, neither debonding nor particle displacements can be observed at the
smoothly without displacing particles by the molten metal.
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interface of top and bottom layers, suggesting that the infiltration front passed through the layers
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Figure 3 shows the areas devoid of reinforcement particles which can be formed as a result of preform breakage during its fabrication or from particle displacement by the molten alloy during pressureless infiltration. The abundant porosities and unwanted reaction products distributed or agglomerated throughout the matrix can have a diverse effect on the final physical and
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mechanical properties of fabricated composites.
(Figure 2)
The boron carbide particles react strongly with Al, resulting in a variety of binary and ternary
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compounds, including Al3BC, AlB24C4 (commonly designated as AlB10), Al3BC3 (Al8B4C7), Al3B48C2 (often known as β-AlB12), AlB2, AlB12C2, AlB48C2, Al4C3 and α-AlB12 [32]. However,
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by shifting the boron carbide composition to the limit of the boron carbide’s solid solution range at the carbon-rich end, the kinetics of the boron carbide – Al reaction and types of reactions are drastically changed [33]. Thus, at a temperature range of 900 to 1225 °C, carbon-rich boron carbide reacts with Al, forming principally AlB12C2 and Al3BC, rather than a multitude of phases. Halverson et al. [34] stated that Al3BC is replaced by other Al–B–C compounds richer in boron (AlB24C2, AlB12C2) and Al–C compounds (Al4C3) when infiltrated at higher temperatures and/or applied a post-heat treatment. 7
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(Figure 3) MgAl2O4, Al4C3, Al3BC, Al8B4C7 and AlB2 are the main ceramic phases that were detected by XRD in the composites prepared. The phase composition in B4C–Al composites depends on the
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starting material [32] and processing variables such as (a) thermal or chemical treatment of boron carbide prior to metal infiltration, (b) the densification temperature and time, and (c) the postdensification heat treatment process [35]. As it can be seen from figure 4, by increasing
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processing temperature and time, AlB2 phase tends to be undetectable. Pyzik and Beaman [35] reported that the reaction between B4C and Al starts at about 450 °C with the formation of
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Al4BC. Above 600 °C, AlB2 forms and Al is rapidly depleted. Between 600 and 700 °C, AlB2 and B4C are the predominant phases. Above 700 °C, AlB2 and Al4BC are both present and the relative amount of Al4BC increases with increasing temperature. Viala et al. [32], claimed that the aluminum borocarbide which they have prepared and characterized as Al3BC and the
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compound Al4BC, actually correspond to the same phase. Between 900 and 1000 °C, the predominant reaction product is Al4BC. At about 1000 °C, AlB2 decomposes and generates free Al. Heating above 1000 °C produces mainly AlB24C4 and small amount of Al4C3. Phases formed
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below 1000 °C are Al-rich and their formation leads to the rapid depletion of metal whereas phases formed above 1000 °C are rich in boron and carbon, resulting in B4C depletion and
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composites with a large amount of free metal and a small amount of B4C. Provided that all the experimental temperatures are above 1000 °C, it is expected a decrease of B4C and in the formation of different Al-B-C binary or ternary ceramic phases. (Figure 4) Although Similar to Al/SiC/RHA-derived composite [25], the degradation can also happen in Al/B4C/MgAl2O4, it was observed that the level of degradation in the latter is much slower and
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milder than that of the former. However, it was found that among all the bilayer and mono layer composites, the bottom layer of L3-Al/B4C/MgAl2O4 exposed to air suffered degradation
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gradually after two years from the fabrication date. Figure 5 illustrates the isothermal Al-B-C ternary phase diagram at 1150 °C. The approximate starting molar positions of both layers of L3-bilayer composite in the Al-B-C compositional triangle, are indicated in figure 5. To examine the robustness of the conclusions, the position of
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the initial composition for both layers was re-calculated assuming that the stoichiometry of B4C
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to be B4C0.5, with all else remaining the same. When plotted, both composition points shifted by a negligible amount. A calculation was also carried out without incorporating the presence of silicon and magnesium in alloys and silica in the preform in the starting composition. The resulting location point for intact layer (hollow circle) was nearly indiscernible from the location of the red solid circle (intact layer) shown in figure 5 and the result for degraded layer (hollow
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triangle) was almost close and in the same region of the red solid triangle (degraded layer). Based on these findings, it can be concluded that neither the B4C stoichiometry nor the presence of other elements and silica in the initial mixture affects significantly the location of the initial
(Figure 5)
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composition in the Al-B-C ternary phase diagram.
As it can be seen, the position of the degraded layer is closer to the pointed area with red arrow which is the equilibrium area of Al3BC3 (Al8B4C7) + Al4C3 + liquid. Considering the presence of Al4C3 in the final composition of degraded composite, it can be assumed that over the process time, the chemical composition of degraded layer shifts to the equilibrium area of Al3BC3 (Al8B4C7) + Al4C3 + liquid, while that of intact layer does not reach. The postulated shifting of
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composition point can be due to increment in aluminum concentration as a result of magnesium evaporation during the process.
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Figure 6 shows the FTIR spectra of both intact and degraded layers of L3-Al/B4C/MgAl2O4 bilayer composite after two years from the tests as well as their photographs. It was reported that the formed Al4C3 reacts with water or moisture in the atmosphere and forms aluminum hydroxide and eventually degrades the composite according to following reactions [25, 36]: ∆G˚1100 = -293.9 kJ/mol
(2)
∆G˚1100 = -634.6 kJ/mol
(3)
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Al4C3 (s) + 12 H2O (g) → 4Al(OH)3 (s) + 3CH4 (g)
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Al4C3 (s) + 18 H2O (g) → 4Al(OH)3 (s) + 3CO2 (g) + 12 H2
FTIR spectra of degraded layer confirms the presence of C-H and Al-O-H bonds in the degradation products of composites. According to this observation it can be concluded that the degradation reaction is happening according to the reaction 2. (Figure 6)
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To sum up, although presence of Al4C3 deleterious phase was detected in the XRD patterns of most of the composites, it can be concluded that the amount of this phase was not sufficient to
composite.
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cause degradation in fabricated composite except for L3-Al/B4C/MgAl2O4 degraded layer
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3.2 Hardness measurements
Figure 7 shows the hardness profile of composites fabricated based on two configurations. It has been reported that hardness distribution in composites fabricated by pressureless infiltration might be undesirably non-uniform [37]. Displacement of reinforcement particles during the fabrication process, due to advancing of molten metal front, would be exacerbated when the alloy is placed on top of the preform. In this configuration, the amount of reinforcement particles generally increases as one moves from the top to the bottom of the samples and in the case of
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graded preforms, alters and annihilates the preform desirable gradient. In order to prevent the occurrence of this particle movement and as a result non-uniformity, in this experiment the preform was placed on top of the alloy. Nevertheless, drawing the hardness profile can assist to
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better understand the properties distribution, in particular hardness in both layers. The average of
supplementary materials section. (Figure 7)
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hardness value of each layer of all fabricated composites is presented in figure 1 in the
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The architecture of reinforcement phases and also the content of each phase have impact on hardness of the substance. Most likely, the large difference in hardness results from the stronger carbide network in the layer, due to finer particles [30] or to the higher volume fraction of B4C particles. The indentation process results in the fracture of the brittle B4C particles, but the
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fracture appears to arrest at the interface between the B4C and the aluminum metal matrix. Flow of the ductile aluminum does not significantly reduce the hardness of the composites, but does increase toughness [38]. As is reported by Halverson et al. [34], the magnitude relation of
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hardness in B4C/Al composites is as follows: B4C>Al3BC>AlB2>Al. Thus, the amount and continuity of formed ceramic phases can characterize the hardness value of each layer. As it can
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be understood from the obtained data (figures 1 of supplementary data), the high value of hardness belongs to the layers that contain small or high percentage of B4C particles. These two factors can provoke formation of a ceramic network in the matrix which inherently has a higher hardness value than aluminum matrix. 3.3 Bending strength and modulus of elasticity Typical stress–strain curves recorded during the four-point bending tests of bilayer and monolayer Al/B4C/MgAl2O4 composites are presented in supplementary materials section, figure
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2. The characteristic bending strength is based on the results of a minimum of 12 test bars. Bending strength values, modulus of elasticity obtained from both stress-strain curves and by Pulse-Echo ultrasonic method are presented in table 4. In those curves the plastic deformation
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zone is absent due to the brittle nature of ceramic particles and as a consequence MMCs with high percentage of ceramic reinforcements.
Among all fabricated Al/B4C/ MgAl2O4 composites, L9 bilayer composite has the highest
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bending strength (318 MPa) followed by L2 (248 MPa) and L4 (243 MPa). The bending strength
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value obtained for L9 composite is higher than that reported for similar systems in the literature [39]. In both L9 and L2 Al/B4C/MgAl2O4 composites the load bearing layer has a preform porosity of 40%. It means that besides reaction products in layer, the high percentage of reinforcement makes the layer stronger. It can be attributed to the formation of ceramic networks during the process. Reaction products of B4C (mainly Al3BC phase), can act as a bridge to
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connect B4C particles together.
(Table 4)
In order to verify the accuracy of the modulus of elasticity obtained from stress-strain curves of
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monolayer and bilayer composites, the Pulse-Echo ultrasonic determination method was used.
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Comparison between Young’s modulus data obtained by Pulse-Echo technique and from stressstrain curves shows that there is a negligible difference between these two methods, attributable to the difference between the tested areas in the samples. The probe in the Pulse-Echo method covers an area of 0.78 cm2 while the bending sample fabricated according to configuration A of ASTM 1161 standard covers an area of 0.4 cm2. Nonetheless, a difference of less than 6 % between the two methods, as measured, falls within an acceptable determination error.
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Results of the pooled ANOVA for the bending strength of bilayer Al/B4C/MgAl2O4 composites are shown in Table 5. This table indicates that, at the levels studied, the parameter that affects the bending strength of graded composites most significantly is infiltration temperature, with a
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relative contribution to the variance in the bending strength of 79 %. Other factors like preform porosity and silica content have almost similar but not significant contribution percentage on bending strength of the composites. It is expected that, due to strong dependency of Al-B-C
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reaction products to temperature, pressureless infiltration temperature has a significant contribution to bending strength values. As it was mentioned before, AlB2 formation can be
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diminished in relatively high temperatures while formation of some ternary phases like Al3BC are confined in low temperatures. On the other hand, chemical composition of alloy along with temperature can define level of transformation of SiO2 from rice husk to MgAl2O4 or MgO phases [25]. The percent contribution due to the error term in the ANOVA is 3%. It provides an
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estimate of the adequacy of the experiment. If the error term is low, i.e., 15% or less, it is assumed that no important factors were omitted from the experiment, and no measurement errors
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were significant.
(Table 5)
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Considering the parameters and levels of the designed experiment, the maximum bending strength can be obtained by using the process parameters shown in table 6. When using these process parameters, the projected bending strength is 341±9 MPa. As it can be inferred from this table,
(Table 6)
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Results of the pooled ANOVA for the modulus of elasticity obtained from stress-strain curves are shown in Table 7. It shows that, at the levels studied, the parameter that affects the modulus of elasticity of graded composites most significantly is chemical composition of Al alloy
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followed by porosity of preforms with contribution percentages of 32 and 22 %, respectively. Process temperature, B4C particle size and silica content have the same contribution percentage of 11% on the modulus of elasticity of fabricated alloys. As mentioned before, chemical
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composition of alloy has an important effect on formation of different phases during pressureless infiltration. Preform porosity, on the other hand, also affects the modulus of elasticity of the
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composites notably. It is expected that the original porosity designed in the preform has a relatively high effect on the composites modulus of elasticity, since this porosity becomes infiltrated with the less elastic component of the composite, i.e., the aluminum alloy. Consequently, the higher the preform porosity levels, the higher the aluminum alloy content of
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the composite, and the lower the modulus of elasticity of the composites. The percent contribution due to the error term in the ANOVA is 4%. This indicates that no important factors
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were omitted in the design of the experiment and measurement errors are trivial. (Table 7)
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Considering experiment parameters and levels designed for this study, the maximum modulus of elasticity can be obtained by using the process parameters shown in Table 8. When using these process parameters, the projected modulus of elasticity is 153±5 GPa. (Table 8) Results of the pooled ANOVA for the modulus of elasticity determined by Pulse-Echo ultrasonic method are shown in Table 9. Similar to the results obtained from stress-strain curves, the 14
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parameter that affects the modulus of elasticity of graded composites most significantly is chemical composition of Al alloy followed by porosity of preforms and B4C particle size. The percent contribution due to the error term in this ANOVA is 2.84% which indicates that no
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important factors were neglected in the design of the experiment. Beffort et al. [40] stated that in contrast to bending strength, modulus of elasticity primarily depends on the elastic properties and
independent on matrix alloy.
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(Table 9)
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the respective volume fractions of the composites constituents and is thus expected to be
Considering only the parameters included in Table 1 in the ranges tested, the maximum modulus of elasticity can be obtained by using the process parameters shown in table 10 which are in a good agreement with the projected factor obtained from the data of the first method of
GPa.
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measuring. When using these process parameters, the projected modulus of elasticity is 144±2
(Table 10)
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The results from pooled ANOVA table for modulus of elasticity obtained by both methods show
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that chemical composition of alloy has a higher influence than volume fraction of reinforcement (preform porosity). This can be attributed to role of magnesium in defining the level of secondary in-situ formed phases during pressureless infiltration process [25]. 3.3.1
Verification of projected optimum results
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As it can be observed from tables all the projected parameters for optimum value of bending strength, modulus of elasticity and retained porosity, propose the same processing condition for
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porosity of preform, chemical composition of alloy, amount and crystallinity of RHA. Since both sets of processing parameters proposed by ANOVA – both methods of Pulse-Echo ultrasonic and stress-strain curve - for maximum modulus of elasticity are the same, they were validated in one single verification test under the experimental conditions established in Tables 6
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and 8. The fact that both optimum processing parameter sets obtained in separate ANOVA tables
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are the same, provides an increased level of confidence in the obtained results. The measured modulus of elasticity determined by Pulse-Echo ultrasonic and stress-strain curve are 149 ± 3 and 156 ± 4 GPa, respectively. These results are in good agreement with the projected values (144 ± 2 and 153 ± 5 GPa, respectively).
The verification tests have been done according to the projected parameters proposed by
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ANOVA for maximum bending strength and minimum retained porosity. The measured bending strength and retained porosity are 345 ± 6 MPa and 2.36 ± 0.7 %, respectively. These results are
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in good agreement with the projected values (341 ± 9 MPa and 2.02 ± 0.4 %). 3.3.2 Fracture characteristics The fracture micrographs of monolayer and bilayer L9-Al/B4C/MgAl2O4 composite are
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presented in figure 8. The fracture surface shows both decohered and cracked B4C particles. It can be seen that almost all the B4C particles on the fracture surfaces belong to the cracked particles rather than decohered ones and most of the particles run parallel to the fracture surfaces of the composites, suggesting a brittle mode of failure. Thus, in both layers, fracture is dominated by transgranular fracture across the particles.
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According to the fracture behavior model proposed by Babout et al. [41], in metal matrix composites, regardless of volume fraction of particles, the strength of the interface is a parameter which obviously controls the selection of the dominant mechanism, and decohesion occurs when
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the average tensile stress in the particles surpasses a critical value known as the interface strength. In spite of the extensive plastic deformation and ductile fracture undergone by the metallic phase (fine dimples that can be observed in the fracture surface), due to semicoherent
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nature of both B4C–Al and reactant–Al interfaces and absence of any other phases and impurities [42], interface is strong enough and load is transferred to B4C particles. As a result, fracture
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occurs as soon as the threshold stress is reached. However, if the matrix is soft and the interface is weak, decohesion prevails. Even though crack propagation in-between the closely packed B4C particles proceeds adjacent to the interfaces, as exemplarily depicted in figure 8 with red arrow, effective interfacial debonding between the matrix alloy and the B4C particles has not been
B4C particles’ surface.
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observed in any case. Some remaining thin matrix material layer is always found to adhere to the
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(Figure 8)
Figure 9 shows fracture micrographs of monolayer and bilayer L8 - Al/B4C/MgAl2O4
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composites. The presence of defects such as large porosities and cracks is clear even in lower magnifications. The surfaces of the ceramic particles in the intergranular infiltration region were free of adhering alloy, (Figure 9, second layer). The metal ligaments between the particles were fissured and are marked with an arrow. This indicated ductile fracture of the metal alloy. Existence of matrix free regions as well as undesirable bonding between RHA and matrix alloy can be mentioned as the main reasons for low bending strength of both monolayer and bilayer L8-Al/B4C/MgAl2O4 composites. The observed matrix free regions could be explained by non17
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equilibrium between liquid aluminum movement speed through the boron carbides agglomerates and the reaction rate between molten aluminum and reinforcement. If the reaction rate is higher
consequence, the densification will be interrupted [43]. (Figure 9)
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than that of the liquid movement, the channels between particles will be closed and as a
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4 Conclusions In this study, RHA as an important oxygen source was incorporated into B4C porous preforms. The presence of RHA causes the in situ formation of MgAl2O4 phase during the pressureless
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infiltration of aluminum alloys into the preforms. Using the Taguchi method, the quantitative effect of different parameters on the infiltration characteristics of bilayer B4Cp-RHA porous preforms by aluminum alloys was investigated. The contributions of each of the parameters on the variability of the retained porosity, hardness, bending strength and modulus of elasticity of
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the monolayer and bilayer composites were determined. In addition, an optimized process for enhanced infiltration was established and validated. Degradation of composites due to the reaction of Al4B3 with the atmospheric moisture was observed just for the bottom layer of L3-
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Al/B4C/MgAl2O4 composite. The reaction results in formation of CH4 and aluminum hydroxide as degradation product, whose presence was confirmed by FTIR. While process temperature
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impacts most significantly the bending strength of the bilayer composites, alloy chemical composition affects both the modulus of elasticity (E) and the retained porosity of the composites. Verification tests using the optimum parameters showed a good agreement between the projected values of bending strength, modulus of elasticity and retained porosity, with those measured experimentally. Due to the high percentage of ceramic particles in the matrix and to brittle nature of the ceramic particles, in all stress-strain graphs, the plastic deformation zone is
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absent. It was observed that besides the abovementioned processing factors, architecture of load bearing layer has a significant influence on the bending strength of the bilayer composites.
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Acknowledgment Mr. Amin Bahrami and Ms. Niloofar Soltani gratefully acknowledge Conacyt (National Council of Science and Technology, in Mexico) for granting a doctoral scholarship. The authors are also
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thankful to Cinvestav IPN-Saltillo for support in the research activities in the field of advanced materials. Thanks also to Dr. Denis Shishin (The University of Queensland), Mr. Evgenii
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Nekhoroshev (Université De Montréal) and Dr. Héctor Manuel Hernández García (Corporación Mexicana de Investigación en Materiales) for technical assistance during the thermodynamic calculations. 5
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[31] N. Soltani, M.I. Pech-Canul, L.A. González, A. Bahrami, Mechanism and Parameters Controlling the Decomposition Kinetics of Na2SiF6 Powder to SiF4, International Journal of Chemical Kinetics, 48 (2016) 379-395. [32] J. Viala, J. Bouix, G. Gonzalez, C. Esnouf, Chemical reactivity of aluminium with boron carbide, Journal of materials science, 32 (1997) 4559-4573. [33] I.A. Aksay, A.J. Pyzik, Multipurpose boron carbide-aluminum composite and its manufacture via the control of the microstructure, US Patents, US4702770 A, 1987. [34] D.C. Halverson, A.J. Pyzik, I.A. Aksay, Boron-carbide-aluminum and boron-carbide-reactive metal cermets, US Patents, US4605440 A, 1986. [35] A.J. Pyzik, D.R. Beaman, Al‐B‐C Phase Development and Effects on Mechanical Properties of B4C/Al‐Derived Composites, Journal of the American Ceramic Society, 78 (1995) 305-312. [36] J. Park, J. Lucas, Moisture effect on SiC P/6061 Al MMC: dissolution of interfacial Al4C3, Scripta materialia, 37 (1997) 511-516. [37] A. Zulfia, R.J. Hand, The production of Al-Mg alloy/SiC metal matrix composites by pressureless infiltration, Journal of Materials Science, 37 (2002) 955-961. [38] H.J. Kim, K.P. Trumble, K.J. Bowman, Layered boron carbide-aluminum composites with large changes in microstructure, in: Materials Science Forum, Trans Tech Publ, 2005, pp. 673-678. [39] J. Abenojar, P. Díaz, M. Martinez, F. Velasco, Influence of Forming on the Mechanical Properties of the Al+ 50% B4C System, in: Materials science forum, Trans Tech Publ, 2006, pp. 304-309. [40] O. Beffort, S. Long, C. Cayron, J. Kuebler, P.-A. Buffat, Alloying effects on microstructure and mechanical properties of high volume fraction SiC-particle reinforced Al-MMCs made by squeeze casting infiltration, Composites Science and Technology, 67 (2007) 737-745. [41] L. Babout, Y. Brechet, E. Maire, R. Fougères, On the competition between particle fracture and particle decohesion in metal matrix composites, Acta Materialia, 52 (2004) 4517-4525. [42] Z. Luo, Y. Song, S. Zhang, D.J. Miller, Interfacial Microstructure in a B4C/Al Composite Fabricated by Pressureless Infiltration, Metall and Mat Trans A, 43 (2012) 281-293. [43] A.J. Pyzik, I.A. Aksay, Multipurpose boron carbide-aluminum composite and its manufacture via the control of the microstructure, US Patents, US4702770 A, 1987.
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Figure captions
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Figure 1. a) A 3D tomographic image of B4C/RHA preform, and b) porosity distribution through the layers obtained from the µCT (Top layer: large B4C particles with 40% porosity, Bottom layer: Small B4C particles with 60% porosity) Figure 2. SEM micrograph of both layers of fabricated composite according to configuration L7. Figure 3. SEM micrographs of both layers of prepared composite according to configuration L4. Figure 4. X-Ray diffraction patterns of Al/B4C/MgAl2O4 fabricated composites.
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Figure 5. The Al-B-C isothermal section at 1150 °C, constructed by FactSageTM software. Positions of chemical compositions of different layers of L3-Al/B4C/MgAl2O4 composite are indicated by hollow and solid triangle and circle.
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Figure 6. a) The FTIR spectra and b) photograph of both intact and degraded layers of L3Al/B4C/MgAl2O4 bilayer composite, two years after infiltration tests. Figure 7. 3D hardness profile for composites fabricated based on configurations (a) L2 and (b) L7. Figure 8. The fracture micrographs of monolayer and bilayer L9-Al/B4C/MgAl2O4 composite.
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Figure 9. The fracture micrographs of monolayer and bilayer L8-Al/B4C/MgAl2O4 composite.
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Tables Table 1. Parameters and levels tested in the experiment
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Constants
Ar flow rate: 50 cm3/min. Alloy amount: 50g.
Level 2 130 µm /9 µm 60 % / 40 % High Mg 5% 120 min 1250 °C Amorphous • Geometry and size of cuboid preform: for each layer: 20×10×50 mm.
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Level 1 9 µm /130 µm 40 % / 60 % Low Mg 10% 60 min 1150 °C Crystalline
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Parameter B4C particle size Porosity of preform Alloy SiO2 content Time Temperature SiO2 crystallinity
Porosity 1 1 1 2 2 2 1 1 1 2 2 2
Alloy 1 1 2 1 2 2 2 2 1 2 1 1
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Particle size 1 1 1 1 1 1 2 2 2 2 2 2
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Number 1 2 3 4 5 6 7 8 9 10 11 12
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Table 2. Standard L12 Taguchi table for pressureless infiltration of this study. SiO2% 1 1 2 2 1 2 2 1 2 1 2 1
Time 1 1 2 2 2 1 1 2 2 1 1 2
Temp. 1 2 1 1 2 2 1 2 2 1 2 1
Crystallinity 1 2 1 2 1 2 2 2 1 1 1 1
Table 3. Chemical composition of the alloys used in the experiment (wt. %). Si 13.1 12.7
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Element Alloy 1 Alloy 2
Mg 9.2 13.5
Fe 0.244 0.277
Mn 0.038 0.038
Zn 0.066 0.055
Al Balance Balance
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Table 4. Young’s modulus (calculated by pulse-echo ultrasonic method and stress-strain curve) and bending strength of Al/B4C/MgAl2O4 monolayer and bilayer composites
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L6
L7
L8
L9
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Bending strength (MPa) 268 180 253 175 152 270 244 193 246 86 67 77 222 203 173 203 188 180 138 184 153 203 243 320 206 127 180
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Young’s modulus (GPa) (Stress-strain curve) 220.83 120.45 130.16 134.62 109.35 128.49 195.20 175.45 125.43 111.66 175.51 110.71 190.03 156.15 123.67 205.05 188.72 135.43 136.54 140.36 123.79 151.78 127.89 138.61 154.82 158.36 134.51
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L2
1st Layer 2nd Layer Bilayer 1st Layer 2nd Layer Bilayer 1st Layer 2nd Layer Bilayer 1st Layer 2nd Layer Bilayer 1st Layer 2nd Layer Bilayer 1st Layer 2nd Layer Bilayer 1st Layer 2nd Layer Bilayer 1st Layer 2nd Layer Bilayer 1st Layer 2nd Layer Bilayer
Young’s modulus (GPa) (Pulse-Echo ultrasonic determination) 235.12 125.71 123.21 125.92 112.67 131.47 190.76 181.23 119.59 105.87 163.65 117.11 178.61 148.05 120.63 200.12 179.91 139.74 138.40 135.37 119.65 158.39 124.54 125.43 150.83 149.39 126.80
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B4C particle size Porosity of preform Alloy SiO2 content Time Temperature SiO2 crystallinity
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Column 1 2 3 4 5 6 7
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Table 5. Pooled ANOVA table for maximum bending strength of Al/B4C/MgAl2O4 composites
Error Total
Sum of squares 44.0 9738.8 20334.7 11536.0 619.8 169567.6 256.8 6383
212791.6
Variance Pooled 9738.8 20334.7 11536.0 Pooled 169567.6 Pooled
Contribution percentage Pooled 4 9 5 --79 --3 100
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Table 6. Optimum process parameters for maximum bending strength of Al/B4C/MgAl2O4 composites. Parameters
Proposed levels
B4C particle size (µm) Porosity of preform (%) Alloy SiO2 content (wt. %) Time (min) Temperature (°C) SiO2 crystallinity
9-130
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40-60 High Mg 5 120 1250 Amorphous
Factors
B4C particle size Porosity of preform Alloy SiO2 content Time Temperature SiO2 crystallinity
Variance 69.11762 132.3037 196.0632 68.04361 56.18 68.11741 Pooled
Contribution percentage 11 22 32 11 9 11 --4 100
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Error Total
Sum of squares 69.11762 132.3037 196.0632 68.04361 56.18 68.11741 0.32768 24 596.9248
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Column 1 2 3 4 5 6 7
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Table 7. Pooled ANOVA table for maximum modulus of elasticity of Al/B4C/MgAl2O4 composites
Table 8. Optimum process parameters for maximum modulus of elasticity of Al/B4C/MgAl2O4 composites Parameters
Proposed levels
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B4C particle size Porosity of preform (%) Alloy SiO2 content (wt. %) Time (min) Temperature (°C) SiO2 crystallinity
9-130 40-60 High Mg 5 60 1150 Amorphous
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Table 9. Pooled ANOVA table for modulus of elasticity determined by Pulse-Echo ultrasonic method of Al/B4C/MgAl2O4 composites Factors
B4C particle size Porosity of preform Alloy SiO2 content Time Temperature SiO2 crystallinity
Variance 129.31 157.68 180.18 77.52 11.02 46.02 79.30 1.26
Contribution percentage 18 22 26 11 2 7 11 3 100
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Error Total
Sum of squares 129.31 157.68 180.18 77.52 11.02 46.02 79.30 20.17 701.23
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Column 1 2 3 4 5 6 7
Table 10. Optimum process parameters for maximum modulus of elasticity of Al/B4C/MgAl2O4 composites Proposed levels 9-130 40-60 High Mg 5 60 1150 Amorphous
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Parameters
Figure captions:
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Supplementary material:
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B4C particle size (µm) Porosity of preform (%) Alloy SiO2 content (wt. %) Time (min) Temperature (°C) SiO2 crystallinity
Figure 1. The average of hardness value of each layer of all fabricated composites.
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Figure 2. Typical stress–strain curves recorded during the four-point bending tests of bilayer and monolayer Al/B4C/MgAl2O4 composites.
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Research highlights Graded Al/B4C/MgAl2O4 composites were fabricated via pressureless infiltration.
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L12 Taguchi method was applied to obtain optimum processing conditions.
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Optimal conditions to maximize the Young’s moduli were established.
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Optimum process condition for maximum bending strength were determined.
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S0925-8388(16)33047-X
DOI:
10.1016/j.jallcom.2016.09.284
Reference:
JALCOM 39118
To appear in:
Journal of Alloys and Compounds
Received Date: 18 April 2016 Revised Date:
8 September 2016
Accepted Date: 25 September 2016
Please cite this article as: A. Bahrami, M.I. Pech-Canul, N. Soltani, C.A. Gutiérrez, P.H. Kamm, A. Gurlo, Tailoring microstructure and properties of bilayer-graded Al/B4C/MgAl2O4 composites by single-stage pressureless infiltration, Journal of Alloys and Compounds (2016), doi: 10.1016/j.jallcom.2016.09.284. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
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Tailoring microstructure and properties of bilayer-graded Al/B4C/MgAl2O4 composites by single-stage pressureless infiltration A. Bahrami 1,2, M. I. Pech-Canul1*, N. Soltani 1,2, C.A. Gutiérrez 1, P. H. Kamm3, A. Gurlo4 1
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Centro de Investigación y de Estudios Avanzados del IPN Unidad Saltillo, Ave. Industria Metalúrgica No. 1062, Parque Industrial Saltillo-Ramos Arizpe, Ramos Arizpe, Coahuila, México, 25900. 2 Technische Universität Berlin, Fakultät III Prozesswissenschaften, Institut für Werkstoffwissenschaften und technologien, Fachgebiet Keramische Werkstoffe, Hardenbergstraße 40, 10623 Berlin, Germany. 3 Applied Materials, Helmholtz-Zentrum Berlin für Materialien und Energie, Germany. 4 Chair of Advanced Ceramic Materials, Technische Universität Berlin, Germany.
Abstract: The quantitative effect of the following parameters on the one single step pressureless
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infiltration characteristics of bilayer (B4C)p/rice-husk ash (RHA) porous preforms by aluminum alloys was investigated using the Taguchi method and analysis of variance (ANOVA): infiltration temperature and time, B4C particle size, RHA percentage, preform porosity and alloy magnesium content. Contribution determinations of each of the parameters on the bending
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strength of the resulting bilayer composites indicated that the parameter with the most significant effect is process temperature, with a contribution percentage of 79 %. Verification tests conducted using the established optimum parameters show good agreement with those of
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projected values. A detailed examination of microstructure and potential chemical reactions, aided by thermodynamic analysis, allowed establishing conditions for suppressing or
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diminishing the impact of the deleterious Al4C3 phase. Keywords: One-step pressureless infiltration; Graded materials; Rice-husk ash; Mechanical properties.
1 * Corresponding author. Tel.: +52 (844) 4-38-96-00; Ext.: 8678 Email address: [email protected] (Martin I. Pech-Canul); [email protected]
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1 Introduction In spite of the outstanding physical and mechanical properties, monolithic B4C seems to not live up to its true potential as an advanced structural ceramic due its extreme sensitivity to brittle
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fracture and the difficulties associated with fabricating fully dense microstructures[1]. The experiments have demonstrated that these issues can be significantly reduced by introducing a metal phase [2]. By far, the most popular metal used in conjunction with B4C for this purpose is
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aluminum [3-6]. In addition to availability and lightness of Al-Mg-Si alloys, they wet B4C well at elevated temperatures. Molten aluminum alloy has been shown to form a variety of binary and
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ternary phases when in contact with B4C such as Al3BC, AlB10 and Al4C3 [7]. The manufacturing process of a functionally graded materials (FGM) can usually be divided into building the spatially inhomogeneous structure “gradation” and transformation of this structure into a bulk material “consolidation”. The fabrication methods of functionally graded composite
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materials can be similar or totally different from the conventional methods used for fabrication of metal matrix composites [8-15]. These methods include but are not limited to: powder stacking, centrifugal method, electromagnetic separation, casting, PVD, CVD, etc.
[16-21]. Die
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compaction of layers (powder stacking), which was used in this study to make graded reinforcement preforms, is a simple and well established method in which a gradient is formed
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by the deposition of powder layers with changing the size and composition of powder in the compacting die [22].
Although the infiltration of ceramic preforms by liquid metals has been typically applied for the processing of composite materials with homogeneous reinforcement shape, size and composition [23, 24], as a constitutive route, it offers the potential for the production of graded materials by variation of shape, size, and volume fraction of the reinforcement in each layer.
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In this investigation, it is postulated that graded metal matrix composite materials may be produced by one-step infiltration process. In spite of various efforts devoted to the optimization of wettability and pressureless infiltration parameters [25, 26], it is expected that in pressureless
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infiltration-based processing, changes of key parameters (reinforcement composition, size, percentage porosity in the preform) from one section to another within the ceramic preform (layer by layer) will increase the inherent barriers and challenges. It was reported that some of
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the problems associated with the infiltration process can be solved by addition of special elements or phases [25, 27]. Furthermore, addition of reinforcement phases with special
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morphology and chemical composition can enhance the mechanical and physical properties of the composites. Rice husk ash (RHA) as a source of silicon and oxygen with truss like and spiral like structure can be used as a unique reinforcement phase to satisfy these expectations [28, 29]. The aim of this investigation is to gain further insight into the processing of functionally graded
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Al/B4Cp/MgAl2O4 composites by pressureless infiltration. In this article, using Taguchi method, the quantitative effect of process temperature and time, B4Cp particle size in the preform layer, crystallinity of SiO2 obtained from rice husk ash, ceramic preform porosity and Mg
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concentration in the aluminum alloys on the microstructure, physical and mechanical properties of composites prepared from RHA/B4Cp ceramic preform stacked layers, is investigated.
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Analysis of variance (ANOVA) is a method to determine the variability in the response parameters in an experiment, by systematically changing the levels in the independent variables, identifying sources of variation and the associated degrees of freedom [30, 31]. In this study ANOVA was performed on the data in order to determine the quantitative effect of each parameter and each interaction on the modulus of elasticity and bending strength, CTE value and electrical resistivity of the graded composites. ANOVA provides a means of estimating the
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percent contribution of each of the parameters tested to the variability in the measured response variable (i.e., bending strength, modulus of elasticity, etc.). In addition, ANOVA provides insight
variable produced when using the optimum process parameters. 2
Experimental procedure
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into the optimum process parameters, and allows the estimation of the magnitude of the response
Infiltration tests for Al alloys/B4C/RHA composites were conducted under the conditions shown
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in Table 1. The effect of various parameters on the physical and mechanical properties of composite were investigated using the Taguchi method for design of experiments. Predetermined
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amounts of ceramic powders with 5wt. % Dextrin from potato starch (Fluka®, Switzerland) as a binder were soaked and mixed with 5 wt. % of distilled water. The mixture of first layer was put in the mold and uniaxially pressed to a pressure of 20 MPa and up to 5mm height and then the mixture of second layer was added and pressed to the final height of preform (10mm). The
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amount of needed mixture to fabricate porous preform was calculated according to density formula (ρ=m/V) considering the porosity percentage in the volume of each layer. Using adequate amount of binder (Dextrin), prevents cracks formation and changes in the desirable
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porosity values in the preform. The preforms were heated up to 400 °C in a furnace open to the atmosphere to remove the organic components. In rows 1 and 2 of Table 1, the numbers on the
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left and the right, correspond to the values of the top and bottom preform layers, respectively. Accordingly, each of the selected parameters was tested at two levels. (Table 1)
Table 2 shows the parameters and levels that are investigated in this work arranged in a standard Taguchi L12 orthogonal array. (Table 2) 4
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The chemical compositions of the two alloys used in the experiment are presented in table 3. The structure and phase type of silica derived from rice husk depend on calcining temperature, atmosphere and chemical treatment of rice husk before and after burning. Two procedures were
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selected for extraction of silica from rice husk. The extraction procedure and physical properties of obtained silica corresponding to each method are reported elsewhere [25]. Infiltration trials were performed in a horizontal tube furnace with a 6.5 cm diameter alumina tube closed at both
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ends with end-cap fittings to control the process atmosphere. More details about the process are presented elsewhere [25]. The mold-preform-metal assembly was placed in the center of the
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alumina tube and heated with a heating rate of 20 °C/min in ultra-high purity argon up to test temperature. The three dimensional structure of the preform was analyzed by using a Hamamatsu microfocus X-ray source and flat panel sensor. The voxel size of the obtained reconstructed volume was 7 µm. The volume was visualized by VGStudio MAX 2.2. Specimens for
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microstructure analysis were mounted and polished using standard metallurgical procedures, and the analysis was done using X-ray diffraction (XRD), scanning electron microscopy (SEM) and energy dispersive X-rays (EDXs). Thermodynamic calculations were carried out using the
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FactSageTM software and FTlite databases. FactSageTM is the fusion of the FACTW in /F*A*C*T and ChemSage/mn SOLGASMIX thermochemical packages. The FTIR analysis was
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performed using Bruker FT-IR Spectrometer (Equinox 55). The bending strength of prepared monolayer and bilayer composites was determined using four-point bending tests according to the ASTM 1161 standard procedures. Furthermore, the modulus of elasticity of composites was determined using data from mechanical testing and from ultrasonic tests conducted according to the ASTM: E494-95 standard procedures. The hardness of the samples was determined by
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Vickers indentation tests (Tuckson-300) on the polished surfaces, using a load of 1 kg and a dwell time of 15 s.
3 3.1
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(Table 3) Result and Discussion Microstructure examination
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Cracks and non-uniformity in reinforcement particles distribution can be generated during preform fabrication and/or melt infiltration. Therefore, the preforms should be made sound and
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without any cracks to the extent possible. Figure 1 shows the three-dimensional image of B4C/RHA bilayer preforms before infiltration. As it can be observed from figure 1 (a), the layer with 60 vol. % porosity containing fine B4C particles is represented by the darker tone. As the atomic number of B4C is lower than that of the RHA, the RHA appears in lighter shades of gray
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(Figure 1a). The inspection shows that the inner structure is dense and particles and porosities (Figure 1b) are distributed uniformly. No defects in the scale of higher than 5 mm, such as
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microcracks are found in the inner structure of the preform. (Figure 1)
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The layered structures of composites fabricated according to configurations 7 and 4, with abrupt microstructural changes, are shown in figures 2 and 3, respectively. As it can be confirmed by XRD (figure 4), both layers of fabricated composites contain different phases that depend on the processing parameters. As it can be seen from figure 2, by choosing adequate processing parameters, the presence of defects in the fabricated composite, such as those shown in figure 3, can be eliminated. It can be observed that in the layer with small size of B4C, due to the high surface area of the fine B4C particles, the reaction of Al and B4C is intensified and more ceramic 6
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phases can be observed than in the layer which contains B4C of large size. On the other hand, transformation of RHA to MgAl2O4 strongly depends on the crystallinity and procedure conditions. The EDS results presented in figure 2 confirm the complete transformation of RHA
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to MgAl2O4. Also the good bonding between matrix alloy and B4C particles is illustrated in figure 2. In addition, neither debonding nor particle displacements can be observed at the
smoothly without displacing particles by the molten metal.
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interface of top and bottom layers, suggesting that the infiltration front passed through the layers
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Figure 3 shows the areas devoid of reinforcement particles which can be formed as a result of preform breakage during its fabrication or from particle displacement by the molten alloy during pressureless infiltration. The abundant porosities and unwanted reaction products distributed or agglomerated throughout the matrix can have a diverse effect on the final physical and
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mechanical properties of fabricated composites.
(Figure 2)
The boron carbide particles react strongly with Al, resulting in a variety of binary and ternary
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compounds, including Al3BC, AlB24C4 (commonly designated as AlB10), Al3BC3 (Al8B4C7), Al3B48C2 (often known as β-AlB12), AlB2, AlB12C2, AlB48C2, Al4C3 and α-AlB12 [32]. However,
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by shifting the boron carbide composition to the limit of the boron carbide’s solid solution range at the carbon-rich end, the kinetics of the boron carbide – Al reaction and types of reactions are drastically changed [33]. Thus, at a temperature range of 900 to 1225 °C, carbon-rich boron carbide reacts with Al, forming principally AlB12C2 and Al3BC, rather than a multitude of phases. Halverson et al. [34] stated that Al3BC is replaced by other Al–B–C compounds richer in boron (AlB24C2, AlB12C2) and Al–C compounds (Al4C3) when infiltrated at higher temperatures and/or applied a post-heat treatment. 7
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(Figure 3) MgAl2O4, Al4C3, Al3BC, Al8B4C7 and AlB2 are the main ceramic phases that were detected by XRD in the composites prepared. The phase composition in B4C–Al composites depends on the
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starting material [32] and processing variables such as (a) thermal or chemical treatment of boron carbide prior to metal infiltration, (b) the densification temperature and time, and (c) the postdensification heat treatment process [35]. As it can be seen from figure 4, by increasing
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processing temperature and time, AlB2 phase tends to be undetectable. Pyzik and Beaman [35] reported that the reaction between B4C and Al starts at about 450 °C with the formation of
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Al4BC. Above 600 °C, AlB2 forms and Al is rapidly depleted. Between 600 and 700 °C, AlB2 and B4C are the predominant phases. Above 700 °C, AlB2 and Al4BC are both present and the relative amount of Al4BC increases with increasing temperature. Viala et al. [32], claimed that the aluminum borocarbide which they have prepared and characterized as Al3BC and the
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compound Al4BC, actually correspond to the same phase. Between 900 and 1000 °C, the predominant reaction product is Al4BC. At about 1000 °C, AlB2 decomposes and generates free Al. Heating above 1000 °C produces mainly AlB24C4 and small amount of Al4C3. Phases formed
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below 1000 °C are Al-rich and their formation leads to the rapid depletion of metal whereas phases formed above 1000 °C are rich in boron and carbon, resulting in B4C depletion and
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composites with a large amount of free metal and a small amount of B4C. Provided that all the experimental temperatures are above 1000 °C, it is expected a decrease of B4C and in the formation of different Al-B-C binary or ternary ceramic phases. (Figure 4) Although Similar to Al/SiC/RHA-derived composite [25], the degradation can also happen in Al/B4C/MgAl2O4, it was observed that the level of degradation in the latter is much slower and
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milder than that of the former. However, it was found that among all the bilayer and mono layer composites, the bottom layer of L3-Al/B4C/MgAl2O4 exposed to air suffered degradation
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gradually after two years from the fabrication date. Figure 5 illustrates the isothermal Al-B-C ternary phase diagram at 1150 °C. The approximate starting molar positions of both layers of L3-bilayer composite in the Al-B-C compositional triangle, are indicated in figure 5. To examine the robustness of the conclusions, the position of
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the initial composition for both layers was re-calculated assuming that the stoichiometry of B4C
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to be B4C0.5, with all else remaining the same. When plotted, both composition points shifted by a negligible amount. A calculation was also carried out without incorporating the presence of silicon and magnesium in alloys and silica in the preform in the starting composition. The resulting location point for intact layer (hollow circle) was nearly indiscernible from the location of the red solid circle (intact layer) shown in figure 5 and the result for degraded layer (hollow
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triangle) was almost close and in the same region of the red solid triangle (degraded layer). Based on these findings, it can be concluded that neither the B4C stoichiometry nor the presence of other elements and silica in the initial mixture affects significantly the location of the initial
(Figure 5)
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composition in the Al-B-C ternary phase diagram.
As it can be seen, the position of the degraded layer is closer to the pointed area with red arrow which is the equilibrium area of Al3BC3 (Al8B4C7) + Al4C3 + liquid. Considering the presence of Al4C3 in the final composition of degraded composite, it can be assumed that over the process time, the chemical composition of degraded layer shifts to the equilibrium area of Al3BC3 (Al8B4C7) + Al4C3 + liquid, while that of intact layer does not reach. The postulated shifting of
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composition point can be due to increment in aluminum concentration as a result of magnesium evaporation during the process.
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Figure 6 shows the FTIR spectra of both intact and degraded layers of L3-Al/B4C/MgAl2O4 bilayer composite after two years from the tests as well as their photographs. It was reported that the formed Al4C3 reacts with water or moisture in the atmosphere and forms aluminum hydroxide and eventually degrades the composite according to following reactions [25, 36]: ∆G˚1100 = -293.9 kJ/mol
(2)
∆G˚1100 = -634.6 kJ/mol
(3)
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Al4C3 (s) + 12 H2O (g) → 4Al(OH)3 (s) + 3CH4 (g)
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Al4C3 (s) + 18 H2O (g) → 4Al(OH)3 (s) + 3CO2 (g) + 12 H2
FTIR spectra of degraded layer confirms the presence of C-H and Al-O-H bonds in the degradation products of composites. According to this observation it can be concluded that the degradation reaction is happening according to the reaction 2. (Figure 6)
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To sum up, although presence of Al4C3 deleterious phase was detected in the XRD patterns of most of the composites, it can be concluded that the amount of this phase was not sufficient to
composite.
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cause degradation in fabricated composite except for L3-Al/B4C/MgAl2O4 degraded layer
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3.2 Hardness measurements
Figure 7 shows the hardness profile of composites fabricated based on two configurations. It has been reported that hardness distribution in composites fabricated by pressureless infiltration might be undesirably non-uniform [37]. Displacement of reinforcement particles during the fabrication process, due to advancing of molten metal front, would be exacerbated when the alloy is placed on top of the preform. In this configuration, the amount of reinforcement particles generally increases as one moves from the top to the bottom of the samples and in the case of
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graded preforms, alters and annihilates the preform desirable gradient. In order to prevent the occurrence of this particle movement and as a result non-uniformity, in this experiment the preform was placed on top of the alloy. Nevertheless, drawing the hardness profile can assist to
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better understand the properties distribution, in particular hardness in both layers. The average of
supplementary materials section. (Figure 7)
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hardness value of each layer of all fabricated composites is presented in figure 1 in the
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The architecture of reinforcement phases and also the content of each phase have impact on hardness of the substance. Most likely, the large difference in hardness results from the stronger carbide network in the layer, due to finer particles [30] or to the higher volume fraction of B4C particles. The indentation process results in the fracture of the brittle B4C particles, but the
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fracture appears to arrest at the interface between the B4C and the aluminum metal matrix. Flow of the ductile aluminum does not significantly reduce the hardness of the composites, but does increase toughness [38]. As is reported by Halverson et al. [34], the magnitude relation of
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hardness in B4C/Al composites is as follows: B4C>Al3BC>AlB2>Al. Thus, the amount and continuity of formed ceramic phases can characterize the hardness value of each layer. As it can
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be understood from the obtained data (figures 1 of supplementary data), the high value of hardness belongs to the layers that contain small or high percentage of B4C particles. These two factors can provoke formation of a ceramic network in the matrix which inherently has a higher hardness value than aluminum matrix. 3.3 Bending strength and modulus of elasticity Typical stress–strain curves recorded during the four-point bending tests of bilayer and monolayer Al/B4C/MgAl2O4 composites are presented in supplementary materials section, figure
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2. The characteristic bending strength is based on the results of a minimum of 12 test bars. Bending strength values, modulus of elasticity obtained from both stress-strain curves and by Pulse-Echo ultrasonic method are presented in table 4. In those curves the plastic deformation
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zone is absent due to the brittle nature of ceramic particles and as a consequence MMCs with high percentage of ceramic reinforcements.
Among all fabricated Al/B4C/ MgAl2O4 composites, L9 bilayer composite has the highest
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bending strength (318 MPa) followed by L2 (248 MPa) and L4 (243 MPa). The bending strength
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value obtained for L9 composite is higher than that reported for similar systems in the literature [39]. In both L9 and L2 Al/B4C/MgAl2O4 composites the load bearing layer has a preform porosity of 40%. It means that besides reaction products in layer, the high percentage of reinforcement makes the layer stronger. It can be attributed to the formation of ceramic networks during the process. Reaction products of B4C (mainly Al3BC phase), can act as a bridge to
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connect B4C particles together.
(Table 4)
In order to verify the accuracy of the modulus of elasticity obtained from stress-strain curves of
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monolayer and bilayer composites, the Pulse-Echo ultrasonic determination method was used.
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Comparison between Young’s modulus data obtained by Pulse-Echo technique and from stressstrain curves shows that there is a negligible difference between these two methods, attributable to the difference between the tested areas in the samples. The probe in the Pulse-Echo method covers an area of 0.78 cm2 while the bending sample fabricated according to configuration A of ASTM 1161 standard covers an area of 0.4 cm2. Nonetheless, a difference of less than 6 % between the two methods, as measured, falls within an acceptable determination error.
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Results of the pooled ANOVA for the bending strength of bilayer Al/B4C/MgAl2O4 composites are shown in Table 5. This table indicates that, at the levels studied, the parameter that affects the bending strength of graded composites most significantly is infiltration temperature, with a
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relative contribution to the variance in the bending strength of 79 %. Other factors like preform porosity and silica content have almost similar but not significant contribution percentage on bending strength of the composites. It is expected that, due to strong dependency of Al-B-C
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reaction products to temperature, pressureless infiltration temperature has a significant contribution to bending strength values. As it was mentioned before, AlB2 formation can be
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diminished in relatively high temperatures while formation of some ternary phases like Al3BC are confined in low temperatures. On the other hand, chemical composition of alloy along with temperature can define level of transformation of SiO2 from rice husk to MgAl2O4 or MgO phases [25]. The percent contribution due to the error term in the ANOVA is 3%. It provides an
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estimate of the adequacy of the experiment. If the error term is low, i.e., 15% or less, it is assumed that no important factors were omitted from the experiment, and no measurement errors
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were significant.
(Table 5)
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Considering the parameters and levels of the designed experiment, the maximum bending strength can be obtained by using the process parameters shown in table 6. When using these process parameters, the projected bending strength is 341±9 MPa. As it can be inferred from this table,
(Table 6)
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Results of the pooled ANOVA for the modulus of elasticity obtained from stress-strain curves are shown in Table 7. It shows that, at the levels studied, the parameter that affects the modulus of elasticity of graded composites most significantly is chemical composition of Al alloy
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followed by porosity of preforms with contribution percentages of 32 and 22 %, respectively. Process temperature, B4C particle size and silica content have the same contribution percentage of 11% on the modulus of elasticity of fabricated alloys. As mentioned before, chemical
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composition of alloy has an important effect on formation of different phases during pressureless infiltration. Preform porosity, on the other hand, also affects the modulus of elasticity of the
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composites notably. It is expected that the original porosity designed in the preform has a relatively high effect on the composites modulus of elasticity, since this porosity becomes infiltrated with the less elastic component of the composite, i.e., the aluminum alloy. Consequently, the higher the preform porosity levels, the higher the aluminum alloy content of
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the composite, and the lower the modulus of elasticity of the composites. The percent contribution due to the error term in the ANOVA is 4%. This indicates that no important factors
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were omitted in the design of the experiment and measurement errors are trivial. (Table 7)
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Considering experiment parameters and levels designed for this study, the maximum modulus of elasticity can be obtained by using the process parameters shown in Table 8. When using these process parameters, the projected modulus of elasticity is 153±5 GPa. (Table 8) Results of the pooled ANOVA for the modulus of elasticity determined by Pulse-Echo ultrasonic method are shown in Table 9. Similar to the results obtained from stress-strain curves, the 14
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parameter that affects the modulus of elasticity of graded composites most significantly is chemical composition of Al alloy followed by porosity of preforms and B4C particle size. The percent contribution due to the error term in this ANOVA is 2.84% which indicates that no
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important factors were neglected in the design of the experiment. Beffort et al. [40] stated that in contrast to bending strength, modulus of elasticity primarily depends on the elastic properties and
independent on matrix alloy.
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(Table 9)
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the respective volume fractions of the composites constituents and is thus expected to be
Considering only the parameters included in Table 1 in the ranges tested, the maximum modulus of elasticity can be obtained by using the process parameters shown in table 10 which are in a good agreement with the projected factor obtained from the data of the first method of
GPa.
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measuring. When using these process parameters, the projected modulus of elasticity is 144±2
(Table 10)
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The results from pooled ANOVA table for modulus of elasticity obtained by both methods show
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that chemical composition of alloy has a higher influence than volume fraction of reinforcement (preform porosity). This can be attributed to role of magnesium in defining the level of secondary in-situ formed phases during pressureless infiltration process [25]. 3.3.1
Verification of projected optimum results
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As it can be observed from tables all the projected parameters for optimum value of bending strength, modulus of elasticity and retained porosity, propose the same processing condition for
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porosity of preform, chemical composition of alloy, amount and crystallinity of RHA. Since both sets of processing parameters proposed by ANOVA – both methods of Pulse-Echo ultrasonic and stress-strain curve - for maximum modulus of elasticity are the same, they were validated in one single verification test under the experimental conditions established in Tables 6
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and 8. The fact that both optimum processing parameter sets obtained in separate ANOVA tables
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are the same, provides an increased level of confidence in the obtained results. The measured modulus of elasticity determined by Pulse-Echo ultrasonic and stress-strain curve are 149 ± 3 and 156 ± 4 GPa, respectively. These results are in good agreement with the projected values (144 ± 2 and 153 ± 5 GPa, respectively).
The verification tests have been done according to the projected parameters proposed by
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ANOVA for maximum bending strength and minimum retained porosity. The measured bending strength and retained porosity are 345 ± 6 MPa and 2.36 ± 0.7 %, respectively. These results are
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in good agreement with the projected values (341 ± 9 MPa and 2.02 ± 0.4 %). 3.3.2 Fracture characteristics The fracture micrographs of monolayer and bilayer L9-Al/B4C/MgAl2O4 composite are
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presented in figure 8. The fracture surface shows both decohered and cracked B4C particles. It can be seen that almost all the B4C particles on the fracture surfaces belong to the cracked particles rather than decohered ones and most of the particles run parallel to the fracture surfaces of the composites, suggesting a brittle mode of failure. Thus, in both layers, fracture is dominated by transgranular fracture across the particles.
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According to the fracture behavior model proposed by Babout et al. [41], in metal matrix composites, regardless of volume fraction of particles, the strength of the interface is a parameter which obviously controls the selection of the dominant mechanism, and decohesion occurs when
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the average tensile stress in the particles surpasses a critical value known as the interface strength. In spite of the extensive plastic deformation and ductile fracture undergone by the metallic phase (fine dimples that can be observed in the fracture surface), due to semicoherent
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nature of both B4C–Al and reactant–Al interfaces and absence of any other phases and impurities [42], interface is strong enough and load is transferred to B4C particles. As a result, fracture
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occurs as soon as the threshold stress is reached. However, if the matrix is soft and the interface is weak, decohesion prevails. Even though crack propagation in-between the closely packed B4C particles proceeds adjacent to the interfaces, as exemplarily depicted in figure 8 with red arrow, effective interfacial debonding between the matrix alloy and the B4C particles has not been
B4C particles’ surface.
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observed in any case. Some remaining thin matrix material layer is always found to adhere to the
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(Figure 8)
Figure 9 shows fracture micrographs of monolayer and bilayer L8 - Al/B4C/MgAl2O4
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composites. The presence of defects such as large porosities and cracks is clear even in lower magnifications. The surfaces of the ceramic particles in the intergranular infiltration region were free of adhering alloy, (Figure 9, second layer). The metal ligaments between the particles were fissured and are marked with an arrow. This indicated ductile fracture of the metal alloy. Existence of matrix free regions as well as undesirable bonding between RHA and matrix alloy can be mentioned as the main reasons for low bending strength of both monolayer and bilayer L8-Al/B4C/MgAl2O4 composites. The observed matrix free regions could be explained by non17
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equilibrium between liquid aluminum movement speed through the boron carbides agglomerates and the reaction rate between molten aluminum and reinforcement. If the reaction rate is higher
consequence, the densification will be interrupted [43]. (Figure 9)
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than that of the liquid movement, the channels between particles will be closed and as a
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4 Conclusions In this study, RHA as an important oxygen source was incorporated into B4C porous preforms. The presence of RHA causes the in situ formation of MgAl2O4 phase during the pressureless
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infiltration of aluminum alloys into the preforms. Using the Taguchi method, the quantitative effect of different parameters on the infiltration characteristics of bilayer B4Cp-RHA porous preforms by aluminum alloys was investigated. The contributions of each of the parameters on the variability of the retained porosity, hardness, bending strength and modulus of elasticity of
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the monolayer and bilayer composites were determined. In addition, an optimized process for enhanced infiltration was established and validated. Degradation of composites due to the reaction of Al4B3 with the atmospheric moisture was observed just for the bottom layer of L3-
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Al/B4C/MgAl2O4 composite. The reaction results in formation of CH4 and aluminum hydroxide as degradation product, whose presence was confirmed by FTIR. While process temperature
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impacts most significantly the bending strength of the bilayer composites, alloy chemical composition affects both the modulus of elasticity (E) and the retained porosity of the composites. Verification tests using the optimum parameters showed a good agreement between the projected values of bending strength, modulus of elasticity and retained porosity, with those measured experimentally. Due to the high percentage of ceramic particles in the matrix and to brittle nature of the ceramic particles, in all stress-strain graphs, the plastic deformation zone is
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absent. It was observed that besides the abovementioned processing factors, architecture of load bearing layer has a significant influence on the bending strength of the bilayer composites.
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Acknowledgment Mr. Amin Bahrami and Ms. Niloofar Soltani gratefully acknowledge Conacyt (National Council of Science and Technology, in Mexico) for granting a doctoral scholarship. The authors are also
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thankful to Cinvestav IPN-Saltillo for support in the research activities in the field of advanced materials. Thanks also to Dr. Denis Shishin (The University of Queensland), Mr. Evgenii
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Nekhoroshev (Université De Montréal) and Dr. Héctor Manuel Hernández García (Corporación Mexicana de Investigación en Materiales) for technical assistance during the thermodynamic calculations. 5
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[31] N. Soltani, M.I. Pech-Canul, L.A. González, A. Bahrami, Mechanism and Parameters Controlling the Decomposition Kinetics of Na2SiF6 Powder to SiF4, International Journal of Chemical Kinetics, 48 (2016) 379-395. [32] J. Viala, J. Bouix, G. Gonzalez, C. Esnouf, Chemical reactivity of aluminium with boron carbide, Journal of materials science, 32 (1997) 4559-4573. [33] I.A. Aksay, A.J. Pyzik, Multipurpose boron carbide-aluminum composite and its manufacture via the control of the microstructure, US Patents, US4702770 A, 1987. [34] D.C. Halverson, A.J. Pyzik, I.A. Aksay, Boron-carbide-aluminum and boron-carbide-reactive metal cermets, US Patents, US4605440 A, 1986. [35] A.J. Pyzik, D.R. Beaman, Al‐B‐C Phase Development and Effects on Mechanical Properties of B4C/Al‐Derived Composites, Journal of the American Ceramic Society, 78 (1995) 305-312. [36] J. Park, J. Lucas, Moisture effect on SiC P/6061 Al MMC: dissolution of interfacial Al4C3, Scripta materialia, 37 (1997) 511-516. [37] A. Zulfia, R.J. Hand, The production of Al-Mg alloy/SiC metal matrix composites by pressureless infiltration, Journal of Materials Science, 37 (2002) 955-961. [38] H.J. Kim, K.P. Trumble, K.J. Bowman, Layered boron carbide-aluminum composites with large changes in microstructure, in: Materials Science Forum, Trans Tech Publ, 2005, pp. 673-678. [39] J. Abenojar, P. Díaz, M. Martinez, F. Velasco, Influence of Forming on the Mechanical Properties of the Al+ 50% B4C System, in: Materials science forum, Trans Tech Publ, 2006, pp. 304-309. [40] O. Beffort, S. Long, C. Cayron, J. Kuebler, P.-A. Buffat, Alloying effects on microstructure and mechanical properties of high volume fraction SiC-particle reinforced Al-MMCs made by squeeze casting infiltration, Composites Science and Technology, 67 (2007) 737-745. [41] L. Babout, Y. Brechet, E. Maire, R. Fougères, On the competition between particle fracture and particle decohesion in metal matrix composites, Acta Materialia, 52 (2004) 4517-4525. [42] Z. Luo, Y. Song, S. Zhang, D.J. Miller, Interfacial Microstructure in a B4C/Al Composite Fabricated by Pressureless Infiltration, Metall and Mat Trans A, 43 (2012) 281-293. [43] A.J. Pyzik, I.A. Aksay, Multipurpose boron carbide-aluminum composite and its manufacture via the control of the microstructure, US Patents, US4702770 A, 1987.
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Figure captions
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Figure 1. a) A 3D tomographic image of B4C/RHA preform, and b) porosity distribution through the layers obtained from the µCT (Top layer: large B4C particles with 40% porosity, Bottom layer: Small B4C particles with 60% porosity) Figure 2. SEM micrograph of both layers of fabricated composite according to configuration L7. Figure 3. SEM micrographs of both layers of prepared composite according to configuration L4. Figure 4. X-Ray diffraction patterns of Al/B4C/MgAl2O4 fabricated composites.
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Figure 5. The Al-B-C isothermal section at 1150 °C, constructed by FactSageTM software. Positions of chemical compositions of different layers of L3-Al/B4C/MgAl2O4 composite are indicated by hollow and solid triangle and circle.
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Figure 6. a) The FTIR spectra and b) photograph of both intact and degraded layers of L3Al/B4C/MgAl2O4 bilayer composite, two years after infiltration tests. Figure 7. 3D hardness profile for composites fabricated based on configurations (a) L2 and (b) L7. Figure 8. The fracture micrographs of monolayer and bilayer L9-Al/B4C/MgAl2O4 composite.
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Figure 9. The fracture micrographs of monolayer and bilayer L8-Al/B4C/MgAl2O4 composite.
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Tables Table 1. Parameters and levels tested in the experiment
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Constants
Ar flow rate: 50 cm3/min. Alloy amount: 50g.
Level 2 130 µm /9 µm 60 % / 40 % High Mg 5% 120 min 1250 °C Amorphous • Geometry and size of cuboid preform: for each layer: 20×10×50 mm.
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Level 1 9 µm /130 µm 40 % / 60 % Low Mg 10% 60 min 1150 °C Crystalline
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Parameter B4C particle size Porosity of preform Alloy SiO2 content Time Temperature SiO2 crystallinity
Porosity 1 1 1 2 2 2 1 1 1 2 2 2
Alloy 1 1 2 1 2 2 2 2 1 2 1 1
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Particle size 1 1 1 1 1 1 2 2 2 2 2 2
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Number 1 2 3 4 5 6 7 8 9 10 11 12
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Table 2. Standard L12 Taguchi table for pressureless infiltration of this study. SiO2% 1 1 2 2 1 2 2 1 2 1 2 1
Time 1 1 2 2 2 1 1 2 2 1 1 2
Temp. 1 2 1 1 2 2 1 2 2 1 2 1
Crystallinity 1 2 1 2 1 2 2 2 1 1 1 1
Table 3. Chemical composition of the alloys used in the experiment (wt. %). Si 13.1 12.7
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Element Alloy 1 Alloy 2
Mg 9.2 13.5
Fe 0.244 0.277
Mn 0.038 0.038
Zn 0.066 0.055
Al Balance Balance
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Table 4. Young’s modulus (calculated by pulse-echo ultrasonic method and stress-strain curve) and bending strength of Al/B4C/MgAl2O4 monolayer and bilayer composites
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L6
L7
L8
L9
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Bending strength (MPa) 268 180 253 175 152 270 244 193 246 86 67 77 222 203 173 203 188 180 138 184 153 203 243 320 206 127 180
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Young’s modulus (GPa) (Stress-strain curve) 220.83 120.45 130.16 134.62 109.35 128.49 195.20 175.45 125.43 111.66 175.51 110.71 190.03 156.15 123.67 205.05 188.72 135.43 136.54 140.36 123.79 151.78 127.89 138.61 154.82 158.36 134.51
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1st Layer 2nd Layer Bilayer 1st Layer 2nd Layer Bilayer 1st Layer 2nd Layer Bilayer 1st Layer 2nd Layer Bilayer 1st Layer 2nd Layer Bilayer 1st Layer 2nd Layer Bilayer 1st Layer 2nd Layer Bilayer 1st Layer 2nd Layer Bilayer 1st Layer 2nd Layer Bilayer
Young’s modulus (GPa) (Pulse-Echo ultrasonic determination) 235.12 125.71 123.21 125.92 112.67 131.47 190.76 181.23 119.59 105.87 163.65 117.11 178.61 148.05 120.63 200.12 179.91 139.74 138.40 135.37 119.65 158.39 124.54 125.43 150.83 149.39 126.80
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Factors
B4C particle size Porosity of preform Alloy SiO2 content Time Temperature SiO2 crystallinity
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Column 1 2 3 4 5 6 7
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Table 5. Pooled ANOVA table for maximum bending strength of Al/B4C/MgAl2O4 composites
Error Total
Sum of squares 44.0 9738.8 20334.7 11536.0 619.8 169567.6 256.8 6383
212791.6
Variance Pooled 9738.8 20334.7 11536.0 Pooled 169567.6 Pooled
Contribution percentage Pooled 4 9 5 --79 --3 100
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Table 6. Optimum process parameters for maximum bending strength of Al/B4C/MgAl2O4 composites. Parameters
Proposed levels
B4C particle size (µm) Porosity of preform (%) Alloy SiO2 content (wt. %) Time (min) Temperature (°C) SiO2 crystallinity
9-130
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B4C particle size Porosity of preform Alloy SiO2 content Time Temperature SiO2 crystallinity
Variance 69.11762 132.3037 196.0632 68.04361 56.18 68.11741 Pooled
Contribution percentage 11 22 32 11 9 11 --4 100
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Sum of squares 69.11762 132.3037 196.0632 68.04361 56.18 68.11741 0.32768 24 596.9248
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Column 1 2 3 4 5 6 7
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Table 7. Pooled ANOVA table for maximum modulus of elasticity of Al/B4C/MgAl2O4 composites
Table 8. Optimum process parameters for maximum modulus of elasticity of Al/B4C/MgAl2O4 composites Parameters
Proposed levels
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B4C particle size Porosity of preform (%) Alloy SiO2 content (wt. %) Time (min) Temperature (°C) SiO2 crystallinity
9-130 40-60 High Mg 5 60 1150 Amorphous
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Table 9. Pooled ANOVA table for modulus of elasticity determined by Pulse-Echo ultrasonic method of Al/B4C/MgAl2O4 composites Factors
B4C particle size Porosity of preform Alloy SiO2 content Time Temperature SiO2 crystallinity
Variance 129.31 157.68 180.18 77.52 11.02 46.02 79.30 1.26
Contribution percentage 18 22 26 11 2 7 11 3 100
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Sum of squares 129.31 157.68 180.18 77.52 11.02 46.02 79.30 20.17 701.23
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Column 1 2 3 4 5 6 7
Table 10. Optimum process parameters for maximum modulus of elasticity of Al/B4C/MgAl2O4 composites Proposed levels 9-130 40-60 High Mg 5 60 1150 Amorphous
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Figure captions:
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Supplementary material:
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B4C particle size (µm) Porosity of preform (%) Alloy SiO2 content (wt. %) Time (min) Temperature (°C) SiO2 crystallinity
Figure 1. The average of hardness value of each layer of all fabricated composites.
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Figure 2. Typical stress–strain curves recorded during the four-point bending tests of bilayer and monolayer Al/B4C/MgAl2O4 composites.
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Research highlights Graded Al/B4C/MgAl2O4 composites were fabricated via pressureless infiltration.
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L12 Taguchi method was applied to obtain optimum processing conditions.
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Optimal conditions to maximize the Young’s moduli were established.
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Optimum process condition for maximum bending strength were determined.
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