Synthesis, characterisation and mechanical properties of SiC ...

Synthesis, characterisation and mechanical properties of SiC reinforced Al based nanocomposites processed by MA and SPS N. Al-Aqeeli*1, K. Abdullahi1, A. S. Hakeem2, C. Suryanarayana3, T. Laoui1 and S. Nouari1 Al based alloys reinforced with different amounts (5, 12 and 20 wt-%) of nanosized SiC particulates were synthesised by mechanical alloying and consolidated by the spark plasma sintering (SPS) technique. The distribution of the reinforcement phase in the composite was evaluated as a function of the milling time and the amount of SiC. The processed materials were characterised by scanning electron microscopy and energy dispersive spectroscopy for the morphology and composition and X-ray diffraction. Continuous reduction in crystallite size was observed as milling progressed and after milling for 20 h the resulting powders reached a grain size of ,100 nm. These Al–SiC composites were successfully consolidated by the SPS method at different sintering temperatures of 400, 450 and 500uC. It is suggested that a higher hardness can be achieved even at 20 wt-%SiC when a higher sintering temperature, for example, above 500uC, is used. Keywords: Aluminium alloys, Silicon carbide, Mechanical alloying, Spark plasma sintering, Microstructure, Hardness

Introduction Metal matrix composites have continued to receive considerable attention from researchers due to their excellent physical and mechanical properties.1–3 Among these, aluminium alloys are widely used as matrixes for the production of these composites.4 Al–Si alloys possess high wear resistance, low coefficient of thermal expansion, good corrosion resistance and improved mechanical properties over a wide range of temperatures. These properties led to the application of Al–Si alloys in the automotive industry, especially for cylinder blocks, cylinder heads, pistons, and valve lifters.5 In certain applications, Al–Si–Mg alloys are considered in both cast and wrought forms. These alloys are age hardenable, and are routinely heat treated to the T6 condition to develop adequate strength.6 To further enhance the strength of Al based alloys and improve their thermal stability, it has been a common practice to reinforce different aluminium alloys with a variety of reinforcements. Some of the successful composites developed this way include Al-6061 containing SiC7 or graphite,8 Al-70759 and Al-202410 containing SiC, and pure Al containing SiC11,12 or Al2O3.13 Besides its density being slightly higher than the density of 1

Mechanical Engineering Department, King Fahd University of Petroleum and Minerals (KFUPM), Dhahran 31261, Saudi Arabia Center of Excellence in Nanotechnology (CENT), King Fahd University of Petroleum and Minerals (KFUPM), Dhahran 31261, Saudi Arabia 3 Department of Mechanical, Materials and Aerospace Engineering, University of Central Florida, Orlando, FL 32816-2450, USA 2

*Corresponding author, email [email protected]

ß 2013 Institute of Materials, Minerals and Mining Published by Maney on behalf of the Institute Received 18 May 2012; accepted 24 June 2012 DOI 10.1179/1743290112Y.0000000029

aluminium, SiC is primarily used because of its low cost and wide range of available sizes and grades. Addition of SiC increases the Young’s modulus and tensile strength of the composite materials. Another useful advantage of SiC reinforcement is the possible increase in the wear resistance.14 The improved mechanical properties of particulate SiC/Al are thought to be a result of the transfer of shear load at the matrix/reinforcement interface. Nevertheless, while the yield and ultimate strength of the matrix increases with increasing concentration of the ceramic micrometre sized particles, the ductility of the composites significantly deteriorates at higher concentrations. The use of nanosized ceramic particles has been shown to strengthen the matrix and at the same time retain some ductility. Yang and Lan15 reported that the yield strength of an Al–7Si–0?3Mg (all compositions are expressed in wt-% unless otherwise stated) alloy fabricated by an ultrasonic assisted casting method was improved by .50% with the addition of 2?0% of nanosized SiC particles. This result is significantly better than what the aluminium alloy with the same percentage of micrometre sized particle reinforcement can offer, with little change in elongation. Mechanical alloying (MA) is a powder processing method involving cold welding, fracturing, and rewelding of powder particles in a high energy ball mill16,17 leading to a uniform dispersion of second phase particles as obtained in oxide dispersion strengthened superalloys. It has also been shown that in addition to the improvement in mechanical properties, addition of SiC particles can also enhance the sinterability of the processed powders. In the work of Woo and Zhang,18

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SiC reinforced Al based nanocomposites

SiC reinforced Al–7Si–0?4Mg composite powders were synthesised by ball milling, and they reported an increase in the sintering rate of the composite powder due to the increased diffusion rate. The microstructure of the sintered specimen obtained using the milled composite powder was finer than that of the sintered specimen obtained from simple mixing. Moreover, the composite fabricated using the milled composite powder contained nanometre and submicrometre sized SiC particles and had higher hardness than the composite produced by sintering the mixed powder. This is indicative that MA can provide better properties rather than just achieving homogeneous structures. In another study,19 the effect of SiC particle size on the laser sintering behaviour of Al–7Si–0?3Mg alloy was investigated. It was shown that densification rates followed first order kinetics and the rate constant was found to be higher at low SiC fractions. The melt track became steadier and a more continuous sintered surface was attained in the presence of ceramic particles. Despite multiple studies in evaluating the improvement in properties due to the presence of SiC particles, not much work was done to study the effect of milling parameters and the concentration of SiC on the resulting structure. This will aid in optimising the synthesis method and coming up with a technological route for the production of improved Al–Si–Mg alloys. Consolidation of mechanically alloyed powder particles to full density is a complex phenomenon. The particles become very hard and strong due to the small size and also because of the incorporation of a high density of crystal defects. Consequently, the pressures required for compaction are expected to be higher. However, once compacted, the process of sintering is expected to take place quickly due to the small size of the particles, and therefore increased grain boundary area, in the mechanically alloyed powder particles. Occasionally, a combination of more than one technique is expected to be utilised to achieve full densification.20–22 During sintering, the density and average particle/grain size increases as a result of the thermal energy supplied to the powder compact. The sinterability and the sintered microstructure of a powder compact are dependent on two important variables, namely, material and process variables and the sintering equations are governed by these variables.23 The atmosphere maintained during sintering affects the sintering mechanism in addition to impurity/reinforcement content and composition. High gas transportation rates and pore structure, grain structure, impurity content, kinetics and surface structures are major changes caused by the sintering atmosphere. Therefore, determination of atmospheres that are most effective for control of impurity is necessary. However, consolidation of powders through isostatic technique is a feasible technique for composite materials which are intricate or costly to fabricate by other methods.24 Therefore, the goals of the present work are to study the effect of milling time and variable volume fraction of SiC on the distribution of the second phases and the final morphology and structure of the produced Al based composite using scanning electron microscopy (SEM)/energy dispersive spectroscopy (EDS). The effect of different sintering parameters on the mechanical properties and microstructure of the composite are also investigated. Prealloyed Al–7Si–0?3Mg powder was used

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to assess the effect of variable milling parameters on the final product. The technique of X-ray diffraction (XRD) was used to determine crystallite size and internal strain along with the evolution of phases, and the effect of sintering parameters in the spark plasma system (SPS) processed materials.

Experimental Materials Prealloyed Al–7Si–0?3Mg powder (procured from Aluminium Powder Co Ltd, Froge Lane, Sutton Coldfield, West Midlands, UK, with an average particle size of 40 mm) was used as the matrix material. The chemical composition of starting aluminium powder alloy is 92?06Al–6?5Si–0?30Mg–0?35Fe–0?040Mn–0?037Cu–0?015 Zn–0?015Pb–0?001Co–0?011Ni–0?004Cr–0?016Ti–0?001Zr– 0?010Ga (wt-%). SiC (b), in the size range of 20–40 nm (surface density: 90 m2 g21, and a purity of y99%) provided by Nanostructured and Amorphous Materials Inc (Houston, TX, USA).

Milling procedure The matrix Al based alloy powder was ball milled with SiC to produce a homogeneous mixture. For the milling experiments, the powder mix was loaded in a stainless steel vial and milled in a planetary ball mill (Fritsch Pulverisette 5) under argon gas (Ar) atmosphere to avoid contamination. A small amount of stearic acid (2%) was used as a process control agent to avoid excessive cold welding of aluminium powder to the walls of the container or the grinding medium. Stainless steel balls were used with each ball weighing 4?1 g and the ball to powder weight ratio was set at 10 : 1. Milling was performed at a speed of 200 rev min21 for different milling periods of 5, 12 and 20 h. The as received Al alloy powder was mixed with 5, 12, and 20% of SiC powder. The ball milling experiments were periodically halted after an interval of 3 h for 30 min to avoid temperature build-up in the milling vial. Furthermore, to eliminate any accumulation of unprocessed powders on the internal walls, the vial was opened at regular intervals of milling time and the deposited powders were scraped out from the vial walls. It is worth mentioning that the loading and unloading of powders was done inside a glove box to eliminate/minimise any chance of undesirable oxidation.

Characterisation of powder and sintered samples A small amount of the milled powder was taken out from the vials to follow the morphological changes. JEOL JSM-6460LV (Japan) and Tescan Lyra-3 (Czech Republic) scanning electron microscopes were used for morphological evaluations. Energy dispersive spectroscopy using the Oxford system along with the mapping capability was employed for compositional analysis. Optical microscopy (MEIJI-Techno microscope, Japan) was also used for further characterisation of the morphology of the sintered composite. The Bruker D8-X-ray diffractometer was employed to study phase evolution and to determine the crystallite size and lattice strain of the milled powder. Cu Ka radiation was used and the diffractometer was operating at a voltage of 40 kV and a current of 40 mA. The Williamson–Hall method was used for analysing the X-ray diffraction patterns. The Scherer relation and Williamson’s plot were

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1 Image (SEM) of as received Al–7Si–0?3Mg powder sample

used to determine the crystallite size and lattice strain from the broadening of XRD peaks.25 Powder samples were sintered in a SPS of FCT group (Systeme GmBH, Germany) at 400, 450 and 500uC, with a holding time of 20 min and a force of 11 kN (35 MPa) applied at a heating rate of 100uC min21. The microhardness of the sintered samples was measured using a Vickers hardness machine (MMT-3 digital Microhardness Tester, Buehler, USA) using 100 gf. The hardness values reported are an average of 10–12 readings. The densities of the sintered pure and reinforced Al alloy were measured using an electronic Densimeter (MD-300S, Alfa Mirage, SG resolution 0?001 g cm23, capacity 300 g).

Results and discussion As milled powder Figure 1 shows the SEM image of the prealloyed Al–7Si–0?3Mg alloy in the as received condition. The as received powder samples showed spherical shapes with a broad size distribution. The Al alloyznanoscaled SiC powder blends (at 5, 12 and 20% levels) were milled for different periods of time to study the effect of SiC on the microstructure evolution. Figure 2 shows the structure and morphology of the Al alloyz5%SiC powder milled for 5, 12 and 20 h. It may be noted that in the early stages of milling, the Al matrix powder was showing a high aspect ratio after deformation with marked irregularity, as reported elsewhere.26 As milling progressed the plastic deformation was more pronounced resulting in flake-like particles (Fig. 2). The extent of deformation was noticed to be higher as the concentration of nanosized SiC was increased from 5 to 20%, and continuous milling of the powder to long times has resulted in reduced average powder particle sizes (Fig 3a–c respectively). It may also be seen that at the higher SiC concentration of 20% (Fig. 3c) and extended milling time of 20 h, more equiaxed particles were formed. This can be attributed to the excessive and repeated grinding process which the powder particles have gone through. The increased milling time ensures that the SiC particles are more spherical and that they are increasingly embedded into the matrix of the nanocomposite. As confirmed by EDS analysis and X-ray mapping, no compositional fluctuation was observed in the resulting nanocomposites. Therefore, the MA technique is highly beneficial in achieving a

a 5 h; b 12 h; c 20 h 2 Images (SEM) of Al–7Si–0?3Mgz5 wt-%SiC powder blend milled for different periods of time

homogeneous distribution of SiC particles in the Al based matrix, and this was also evident elsewhere in other Al based alloy systems.7,13,16,17 X-ray diffraction patterns of SiC reinforced Al–7Si– 0?3Mg alloy are shown in Fig. 4 for different SiC contents and milling times. Owing to the nature of MA process, in which severe plastic deformation of the powder occurs, there is a continuous reduction in the crystallite size and accumulation of lattice plastic strain as milling experiments progressed. This is evident by the continuous broadening of the Al peaks (from the pure aluminium) as milling time is increased (Fig. 4). However, by comparing the results obtained for two different weight fractions of SiC (Figs. 2 and 3), it can be seen that more refinement of the microstructure is

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4 X-ray diffraction patterns of SiC/Al alloy composites as function of SiC content and milling time

a 5 h; b 12 h; c 20 h 3 Images (SEM) of Al–7Si–0?3Mgz20 wt-%SiC powder blend milled for different periods of time

observed and a continuous reduction of crystallite size occurred at a higher SiC concentration. This proves the role of the concentration of the reinforcement phase in refining the microstructure and achieving reduced crystallite sizes over the same milling periods. It can also be seen that there is a pronounced shift in the Al peaks towards higher angles in the milled powder containing 20%SiC for 12 and 20 h (Fig. 4). This is a clear indication of higher internal strains resulting from the accumulation of higher percentages of cold work. In addition to the broadening of the Al diffraction peaks with increasing deformation, another point to be noted is that the SiC peaks have a very low intensity in the 5%SiC composite. While the SiC peaks are hardly seen in the 5%SiC composite, they become very clear with increasing

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amount of SiC. While the SiC peaks can be observed in the 5%SiC composite, they are very clear in the 20%SiC composite (the SiC positions are labelled in the XRD patterns shown in Fig. 4). This can possibly be attributed to the detectability limit of the XRD machine used, in which 5%SiC is perhaps close to the minimum quantity required for its detection. No intermetallic phases were observed in any of the milled powder, irrespective of the SiC content in the powder blend; this observation is at variance with the result reported elsewhere.12 The crystallite size and lattice strain in the milled powder were evaluated from the broadening of the XRD peaks. Figure 5a compares the effect of different SiC concentrations and extended milling periods on the resulting crystallite size of the synthesised Al based alloy composite. It can be seen from the graph that milling was successful in bringing down the crystallite size of the nanocomposite to very small values (,100 nm) after 20 h of milling; measured from Al peaks. Another important point to be noted is that the minimum crystallite size achieved seems to be lower with increasing SiC content. Accordingly, the minimum crystallite size reached values close to 70–80 nm in the Al–7Si–3Mgz20%SiC composite. Further, it may be noted that the difference in the crystallite size is narrowing down in the composites with different SiC contents. Therefore, as the milling time increases, the minimum crystallite size obtained is the same, irrespective of the SiC content in the composite. This may be attributed to the fact that as milling time increases the possible reduction in crystallite size becomes more difficult as the structure becomes somewhat saturated with the accumulation of defects and the dislocation density reaches a saturation value. If the milling time was increased further, it will be expected that the powders will reach a steady state at different rates depending on the volume fraction of the reinforcement material. The induced internal lattice strain following milling is presented in Fig. 5b. The lattice strain increased with increasing milling time. Further, for any given milling time, the strain was always higher at higher SiC contents in the composite. For example, the strain in the composite with 20%SiC was almost twice that in the composite with 5%SiC. It should, however, be mentioned in this context that these results were obtained using indirect measurements from XRD and the inaccuracy in the values can be up to ¡5%.

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5 a reduction of crystallite size and b accumulation of internal strain as function of milling time it can be seen that addition of higher percentages of SiC reﬁned structure further and increased percentage of strain

As a verification of the results obtained using the indirect measurement of XRD, some direct measurements using TEM were taken for samples milled for 20 h and it was found that the final crystallite size was in the vicinity of 100 nm. Figure 6 shows a bright field TEM image of a powder sample with 5%SiC milled for 20 h and showing a crystallite with y100 nm in diameter. It is important to mention at this stage that the alloy with 12%SiC was following comparable trend with other alloys in both morphology and crystallite size measurements. However, some differences started to appear when consolidation was carried out and such the results of the alloy with 12%SiC are included.

Consolidated product 6 Bright Field TEM image of alloy with 5%SiC milled for 20 h

All the milled powders were consolidated using the spark plasma sintering method at three different temperatures, namely, 400, 450 and 500uC. The efficiency of consolidation was evaluated both by microstructural

7 Optical micrographs from a monolithic Al alloy sintered at 400uC and b Al–7Si–0?3Mg containing 20%SiC composite sintered at 400uC and 35 MPa pressure applied during SPS sintering and 20 min of holding time Table 1 Vickers hardness HV and relative density (%) of monolithic Al alloy and composite samples sintered at different temperatures and 35 MPa pressure Sample

Vickers hardness HV temperature/uC

Relative density (%) temperature/uC

SiC/wt-%

400

450

500

400

450

500

0 5 12 20

41 58 62 49

57 66 70 63

63 71 75 69

94.9 93.6 91.7 89.34

99.3 97.5 95.5 92.8

100 98 97.5 94.4

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8 Optical micrographs from a monolithic unmilled Al alloy sintered at 500uC and b Al–7Si–0?3Mg containing 20%SiC composite sintered at 500uC and 35 MPa pressure applied during SPS sintering and 20 min of holding time

observations and by measuring the density and hardness of the consolidated product. The monolithic Al alloy sample (without milling and without any SiC addition) was also consolidated under identical conditions so that a base line could be established.

Table 1 shows the hardness and density values of the SPS samples sintered at different temperatures at a pressure of 35 MPa. From Table 1, it may be observed that, as a general trend, the hardness and density increased with an increase in the sintering temperature for the

9 Images (SEM) of a pure alloy as received, b Alz5%SiC, c Alz12%SiC and d Alz20%SiC (all sintered at 400uC for 20 min holding time)

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10 Images (SEM) for a pure Al AR, b Alz5%SiC, c Alz12%SiC and d Alz20%SiC (all sintered at 500uC for 20 min holding time)

monolithic Al alloy and all the composites with different SiC levels. Another observation was that the hardness had also increased with increasing SiC content, up to 12%SiC, beyond which the hardness had decreased. For example, the hardness and density of the composite with 20%SiC (sample 4) is lower than that of the composite with 12%SiC. It should, however, be noted that the hardness and strength of the powder with 20%SiC are much higher than those at lower SiC contents. Even though the same set of temperatures were used in the present study for all alloy powders and composites, use of a much higher sintering temperatures will be required to achieve better densification. A higher hardness and density could, therefore, be achieved either by increasing the temperature to above 500uC or by employing higher pressures. Figures 7 and 8 show optical micrographs of the consolidated samples from the unmilled aluminium alloy and the Al alloy reinforced with 20%SiC and milled for 20 h. Figure 7 shows that some degree of porosity is observed when the samples were sintered at 400uC. This

is also verified from the density and hardness data (Table 1) for the sample sintered at 400uC. However, when the monolithic Al alloy samples were sintered at 500uC, no porosity was observed in the optical micrographs (Fig. 8a) suggesting that 100% densification was achieved; confirmed by the measured density value listed in Table 1. Metallographically, no porosity was observed in the Al alloy composite containing 20%SiC (Fig. 8b), even though the measured density was not 100% of the theoretical value. Figures 9 and 10 are SEM images of the samples sintered at 400 and 500uC for 20 min holding time at 35 MPa pressure and a heating rate of 100uC min21. The morphology of the phases observed here in the SEM images is similar to what was observed in the optical micrographs in Figs 7 and 8. It is noted that at the low temperature at 400uC there was some amount of porosity present in the sintered samples compared to the samples sintered at 500uC. The addition of SiC also affected the sintering and densification behaviour of the

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11 a densiﬁcation versus sintering temperature and b hardness versus sintering temperature; and all samples were milled for 20 h except sample containing 0%SiC and labelled as %

samples. These effects are more pronounced in the samples sintered at the lower temperature of 400uC. It is also noted from the SEM images that as the SiC content increased in the composite, the area of the dark grey regions also increased as it is presumed that these regions represent areas that were partially sintered compared to the light grey regions (fully sintered regions), as shown in Fig. 9 and more clearly in Fig. 10 at the higher temperature of sintering. Similar trends of (light grey regions representing the fully sintered condition decreasing as the SiC content increases) were also noted in the optical micrographs. This decreased sintering also affects the properties of the composites through a decrease in the density as well as a decrease in the hardness of the composite materials. Figure 11 shows the densification and hardness plots of samples with different SiC contents, as a function of the sintering temperatures for the samples milled for 20 h. Generally, at 500uC, samples show higher hardness and densification of .98%. Pure aluminium sample has the highest densification corresponding to 100% of the theoretical density. However, it was found that the composite with 12%SiC gives the highest values of hardness among all the samples (Fig. 10c). As the SiC increases to 20% both hardness and density start to decrease because of partially sintered clusters in the sintered regions. However, it was found that increasing the SiC content to .12% causes a decrease in densification at the evaluated temperature (500uC). Therefore, it may be appropriate to increase the sintering temperature at a higher concentration of SiC in the matrix or increase the pressure (from 35 MPa to higher valves) during sintering at the selected temperature of 500uC.

Conclusions Prealloyed Al matrix composites reinforced with nanosized SiC were synthesised using ball milling technique and sintered using the non-conventional method of SPS. The effect of different milling periods and the amount of the reinforcement phases on the resulting phase evolution and morphology was elucidated as well as the effect of

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sintering parameters on the microstructure and mechanical properties. The following conclusions could be drawn. 1. A homogeneous distribution of SiC reinforcement in Al–7Si–0?3Mg prealloyed powders was obtained by employing the ball milling technique. 2. As milling progressed there was continuous reduction in the crystallite size and accumulation of lattice plastic strain, and after 20 h of milling, the crystallite size reached a value of ,100 nm and the addition of 20%SiC induced more than double the amount of internal strain into the alloy matrix. 3. Solid state synthesis methods proved to be beneficial in retarding the formation of undesirable intermetallic phases which might cause embrittlement. 4. One hundred per cents densification can be obtained using the spark plasma sintering technique. 5. The hardness increases as the content of reinforcement and temperature are increased.

Acknowledgement The authors wish to express their gratitude to the King Fahd University of Petroleum and Minerals (KFUPM) for financial support provided under project no. NT-04-2008.

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