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Al composites fabricated by spark plasma sintering with powders prepared by mechanical alloying process
Materials Science and Engineering A 527 (2009) 218–224
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Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
The dynamic properties of SiCp /Al composites fabricated by spark plasma sintering with powders prepared by mechanical alloying process Jiangtao Zhang a,b,∗ , Huiji Shi a , Mingchun Cai a , Lisheng Liu b , Pengcheng Zhai b a b
Institute of Solid Mechanics, Department of Engineering Mechanics, Tsinghua University, Tsinghua East Road, Haidian Dist., Beijing 100084, PR China State Key Laboratory of Advanced Technology for Materials Synthesis and Processing, Wuhan University of Technology, Wuhan 430070, PR China
a r t i c l e
i n f o
Article history: Received 3 March 2009 Received in revised form 25 July 2009 Accepted 31 August 2009
Keywords: Dynamic compressive properties Spark plasma sintering Mechanical alloying Al matrix composites Strain rate Damage
a b s t r a c t The dynamic compressive properties of SiC particle reinforced pure Al matrix composites, fabricated by spark plasma sintering technique with mixture powders prepared by mechanical alloying process, were tested in this paper. Two different average SiC particle sizes of 12 m and 45 m were adopted, and the compressive tests of these composites at strain rates ranging from 800/s to 5200/s were conducted by split Hopkinson pressure bar. The damage mechanism of the SiCp /Al composites was analyzed through the microstructural observations and high-precision density measurements. Results show that the dynamic properties and damage accumulation of these composites are significantly affected by the particle distribution, size, particle cracking, particle/matrix interface debonding and adiabatic heat softening. The composites containing smaller SiC particles exhibit higher flow stress, lower strain rate sensitivity, and less damage at high strain rate deformation. © 2009 Elsevier B.V. All rights reserved.
1. Introduction Metal matrix composites (MMCs) reinforced with ceramic particles have the potential to provide desirable properties, such as high stiffness, high strength, good thermal stability and corrosion resistance. These superior properties make particle reinforced MMCs attractive in a wide range of applications, such as the automobile, aerospace and defense industries. The increasing applications of MMCs in high rate deformation conditions, such as structural impact, blast loading and metal processing, have made essential demand for a thorough understanding of constitutive and damage behaviors of MMCs at high rate deformation [1–4]. The dynamic properties of MMCs, produced by various traditional processes, have been tested over the last decades. The effects of matrix materials, fabricating processes and the type, size, distribution and shape of the particles on the strength and damage mechanism of MMCs have been considered [1–10]. It has been recognized that the uniform spatial distribution of the particles in the matrix is essential for improving the mechanical and physical properties of MMCs [11,12]. However, MMCs produced by the traditional processing techniques, such as casting, forging and powder metallurgy (PM) processes, suffer from some drawbacks, such as
∗ Corresponding author at: Institute of Solid Mechanics, Department of Engineering Mechanics, Tsinghua University, Tsinghua East Road, Haidian Dist., Beijing 100084, PR China. Tel.: +86 10 62771817; fax: +86 10 62781824. E-mail address: [email protected] (J. Zhang). 0921-5093/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2009.08.067
clustering of the particles, undesirable chemical reaction and poor adhesion between particles and matrix [9,13,14]. Recently, it has been reported that the high performance MMCs can be fabricated by using mechanical alloying (MA) process to mix the powders and then followed by rapid spark plasma sintering (SPS) [15,16]. MA process is originally developed by Benjamin [17] to produce a nickel based superalloy hardened by oxide dispersion by using high-energy ball mill. During MA process, two essential processes are involved, namely, cold-welding and fracturing of powder [17,18]. The former process makes the reinforcement particles wrapped up into the constituent powders, while the latter can separate the reinforcement particles from the constituent powders again. The constituent powders are repeatedly fractured and cold-welded throughout the milling, so that a homogenous distribution of the particles can be achieved [19]. Besides, the continuous collision between the reinforcement particles with balls, wall of container and other reinforcement particles can eliminate the possible defects contained in the reinforcement. Hence, the size of the particles is reduced and the strength is increased [12]. And so far, a large number of reports have dealt with the fabrication of particulates reinforced MMCs by MA process [12]. In order to produce bulk composites from powders prepared by MA process, various consolidation processes have been employed [12,20,21], such as sintering, hot-pressing (HP), hot-isostatic pressing (HIP) and hot extrusion of loose or prepressed powders. Recently, SPS technique has attracted extensive attention [22,23]. SPS is a novel consolidation technique for the manufacture of fully dense materials, such as metal alloys, ceramics, composites and
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functionally graded materials (FGMs). This technique combines plasma generation with resistive heating and pressure application. Self-heat generation by DC on-off pulse and spark plasma generated among the powder particles lead to the densification of the samples at low temperature within relatively short time [16,17]. The low sintering temperature and short consolidation time have been known to be very helpful in preventing chemical reaction at the interface of MMCs. The purpose of the present work is to investigate the dynamic properties of SiC (Vol. 20%)/Al matrix composites, which are fabricated from the mechanically alloyed (MAed) powders of pure Al and SiC particles and subsequently sintered by SPS. The dynamic properties of these composites are examined by using conventional split Hopkinson pressure bar (SHPB). The results of instrumented tests at high strain rates were compared with the microstructural observations. The strength, strain rate sensitivity, damage mechanism and heat softening of these composites were studied.
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A MTS-810 Material Testing System was used to perform quasistatic compression tests at constant crosshead speed with an initial strain rate of 0.0013/s, and strain gages were mounted on the samples to measure the strains until they detached from the samples, then larger strains were read from the machine output. A conventional SHPB equipment with 14.5 mm diameter bars was used for dynamic compression tests at strain rates ranging from 800/s to 5200/s. The detailed description about SHPB can be found in Ref. [24]. After SHPB tests, the density of the specimens were tested again, and the specimens were then sectioned parallel to the loading direction by using electro-discharge machining, and polished for identification of the microstructural damage. The load axis is parallel to the applied pressure direction during the SPS sintering. All tests were performed at room temperature with sufficient lubrication to prevent barrelling in the specimens. 3. Results and discussions
2. Experimental procedures 3.1. Microstructure 2.1. Materials and fabrication processes In the present study, a commercial pure Al powder with the average size of 6 m was used as the matrix material. SiC particles with the average size of 12 m and 45 m were used as the reinforcement respectively. The nominal volume fraction of SiC particles is 20%. The composite powders were first blended using MA process in a planetary high-energy ball mill (Fritsch, Pulverisette-4). In order to avoid excessive decrease of the SiC particle size, alumina balls were used as milling media due to their lower density compared with steel balls. The diameters of balls vary from 5 mm to 10 mm. The mixture of pure Al powder with SiC particles was put into a WC vial with alumina balls, and the weight ratio of balls to powder was fixed to 8:1. During the powder mixing, the rotational speed of the turn table of the ball mill was controlled at 240 rpm and the relative rotational speed ratio of the vial was 2. The milling time was 10 h. To prevent overheating, the system was constantly cooled using a fan, and ball milling was stopped periodically (every 1 h) and then resumed when the temperature of the WC vial decreased to room temperature. To prevent the oxidation of aluminum, the WC vial was full of argon atmosphere, and ethanol was used as a process control agent. Then the MAed powders were consolidated by SPS (Sumitomo Coal Mining, SPS-1050). Fourteen grams of MAed powder was placed into a graphite die with an inner diameter of 30 mm and heated under vacuum with an applied pressure of 30 MPa at 590 ◦ C for 10 min, and then followed by furnace cooling. The temperature was measured through a K-type thermocouple and the heating rate was equal to 50 ◦ C/min. The height of the SPS consolidated specimens was about 7 mm. No further heat-treatment was applied to the specimens.
After MA process, the MAed powders were firstly cold compacted at a pressure of 250 MPa in a steel die to prepare cylindrical compacts for the observation of the microstructure. The SEM micrographs of the compacts are shown in Fig. 1. The morphology of the MAed powders is rounded and composed of numerous layers, which are the typical features of the MA ending stage [25].
2.2. Mechanical tests The sintered composites were well polished and the densities were determined by Archimedes method prior to the test of mechanical properties. The microstructure of samples was analyzed by optical microscopy (OM) and scanning electron microscope (SEM). The chemical composition of SPS composites was verified by X-ray diffraction (XRD). The specimens which were machined in cylindrical shape with 6 mm in length and 8 mm in diameter by electro-discharge machining (EDM) were used for both dynamic and quasi-static compression tests. Both surfaces of the cylindrical specimens were finely ground and made as flat and as parallel as possible.
Fig. 1. SEM micrographs of compacts prepared by MAed powders with the pressure of 250 MPa. (a) 45 m SiC/Al and (b) 12 m SiC/Al.
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It can be clearly observed from Fig. 1 that the SiC particles are uniformly distributed in the MAed powders. The SiC particle size has great effect on the MAed powders. The size of MAed powder containing 45 m SiC particle is much smaller than that containing 12 m SiC particle. Besides, some large SiC particles are still not embedded into the composite powders in the former (as shown in Fig. 1(a)), while no free SiC particle can be observed in the latter (as shown in Fig. 1(b)). This result is consistent with that of Ruiz-Navas et al. [25]. Their experiments revealed that the bigger reinforcement leads to a smaller final size of composite powder when the same milling conditions were employed. The main reason is that small particles are easily stuck into, and then tightly wrapped in the composite powder. On the contrary, it is more difficult for the large particles to be trapped into the composite powders. Due to the low density of the alumina balls, the large SiC particles cannot be fragmented or comminuted under the ball-powder-ball collisions in a short milling time, so that the size of the SiC particles are not reduced too much. Only part of the large particle can be embedded in the surface of the composite powders under the ballpowder-ball collisions [26], and the constraint imposed by matrix on the SiC particles is quite weak. The embedding of large particles may cause large stress on the surface of the composite powders [5]. When mechanical alloying continues, the surface with SiC may fracture, or the SiC particles may separate from the composite powders again due to the constant impact of the balls and other composite powders. The processes occur repeatedly during milling process.
Fig. 2. The optical micrographs of SPS composites: (a) 45 m SiC/Al and (b) 12 m SiC/Al.
Table 1 Mechanical properties of SPS composites. Composites
Density (g/cm3 )
Modulus E (GPa)
Proof stress 0.2 (MPa)
45 m SiC/Al 12 m SiC/Al
2.774 2.781
104 113
228 273
Hence, the composite powders cannot grow as larger as those containing small SiC particle, and meanwhile, a portion of large SiC particles will be wrapped into the composite powders, as shown in Fig. 1(a). Fig. 2 shows the optical micrographs of SPS composites. No detectable microvoids can be observed, and the distribution of SiC particles is uniform in the composites. The size of SiC particles is apparently reduced. These excellent features are mainly attributed to the MA process. Due to the collision of the ceramic particles with balls, wall of container and other ceramic particles, the particles containing defects are more likely to be fractured during milling. So the particle size is reduced and there are large numbers of fine SiC particles in the composites, especially in those containing 45 m SiC particle. Since large SiC particles are difficult to be embedded in the matrix, they are subjected to the collision for a comparatively long time by comparison with small SiC particles. Therefore, the size reduction of larger SiC particles is more remarkable than that of smaller SiC particles, and thus more fine SiC particles are created in the MAed powder with 45 m SiC particles. Since the small SiC particles can be easily wrapped into the matrix, fracturing of SiC particles promotes the rate of embedding [5]. It can be seen in Fig. 2 that there are still many fissures and microvoids presented in the surface of the SiC particles, and larger SiC particles contain more defects. These defects may provoke the premature fracture of large SiC particle in the composites during the quasi-static and dynamic loadings. The densities of the SPS composites are shown in Table 1, and they are very close to the theoretical value. It indicates that the composites have been fully consolidated by SPS sintering. The composite of 45 m SiC/Al has a slightly lower density. It may be due to the more defects in the larger particles or the difference in the MA and sintering mechanisms caused by the difference of the particle size. In order to verify the composition of the SPS composites, X-ray diffraction (XRD) was applied and the XRD patterns of SPS composites were shown in Fig. 3. Except a little SiO2 contained in the raw material of SiC particles, no other impurity was detected in
Fig. 3. X-ray diffraction patterns of SPS composites.
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rate above 5% strain. At higher strain rates ε˙ ≥ 3400/s, however, the stress–strain curves present three distinct regions. Besides the two regions mentioned above, the last region (iii) of obvious strain softening can be observed. The flow stresses reach the maximum value near the strain of 23% for 45 m SiC/Al composite and 16% for 12 m SiC/Al composite, and then the flow stresses decrease gradually as the strain increases continuously. There are two main factors responsible for strain softening: (1) adiabatic heating, which is caused by the adiabatic character of the dynamic deformation process results in a significant temperature increase in the specimen at large strains, and (2) the successive microstructural damage, which takes place in the composites during the dynamic deformation process. Previous researches [2,27] have revealed that the strain rate has no obvious effect on the microstructural damage evolution in the composite, and the damage increases with the increasing strain. The strain hardening of the quasi-static stress–strain curves in the whole tested strain range indicates that it is the adiabatic heating that dominates the strain softening of the composites under dynamic loads. The hardening–softening transition strains decline with the decreasing particle size because of the dominating influence of the adiabatic heating softening, since the higher the flow stress, the more the deformation energy transferred to heat, and the higher the temperature of the specimen. At the strain rate 900/s < ε˙ < 3400/s, strain softening is not observed in the stress–strain curves. The main reason is that the total strains in these tests are less than the hardening–softening transition strains, where the strain softening caused by adiabatic heating and damage overwhelms the strain hardening of plastic deformation. 3.3. Strain rate sensitivity
Fig. 4. Stress–strain curves of SPS composites under quasi-static and dynamic loads.
SPS composites by XRD. This indicates that the interface reaction between the matrix and the particle is very weak or does not occur during SPS sintering. The alumina debris, which may be expected to be created from the alumina balls during milling process or derived from the oxide skin on the Al powder during the fabrication processes, is not found by XRD. It means that the fabrication processes are properly controlled. As a result, the content of alumina in the SPS composites is very low, although oxidation on the surface of the Al powders is considered inevitable [32]. So the effect of alumina debris on the mechanical properties of the SPS composites was not considered in the following discussions. 3.2. Stress–strain curves The stress–strain curves of the SPS composites at strain rates ranging from 0.0013/s to 5200/s are shown in Fig. 4. These curves exhibit the typical constitutive characteristics of the aluminum matrix composites as reported in the previous literatures: (1) strain rate sensitivity, namely the initial yield strength and the flow stress at a constant strain increase with the increasing strain rate, (2) particle size-dependence, namely the flow stress increases with the decrease of the particle size at certain strain rate, (3) similar strain hardening trends among different strain rate curves at strain less than the hardening–softening transition strain. Here the hardening–softening transition strain is defined as the strain at which the stress–strain curve is changed from strain hardening stage to strain softening stage. The quasi-static stress–strain curves can be divided into two distinct regions: (i) initial region of high work hardening rate until about 5% strain, and (ii) region of relatively low work hardening
In previous studies, various parameters have been defined to estimate the strain rate sensitivity of the rate dependent materials at high strain rate deformation. For example, Guden and Hall [6] and Tjong et al. [7] used the ratio of the increment of the dynamic flow stress to the quasi-static flow stress at certain strains as the strain rate sensitivity parameter: Rs =
d − q q
(1)
where Rs is the strain rate sensitivity parameter, d is the dynamic flow stress, q is quasi-static reference flow stress. The reference strain rate was chosen 0.0013/s, corresponding to the quasi-static strain rate adopted in this study. Fig. 5 shows the strain rate sensitivity parameter Rs as a function of strain for different strain rates. It should be noted that the low strain data (ε < 5%) of the stress–strain curves obtained from SHPB tests are unreliable, since the deformation in the tested specimens is inhomogeneous due to stress wave propagation [9], and usually the calculated stresses at a low strain decrease with the increasing strain rates from SHPB tests. The small values of Rs at low strains are not the true characteristic of the strain rate sensitivity of the composites. It can be seen from Fig. 5 that, in the strain ranging from 5% to the hardening–softening transition strain, namely region (ii) defined in Section 3.2, the values of Rs basically keep invariable for a given strain rate. This result confirms again the same strain hardening trend for the composites at different strain rate deformation. While in the region (iii), Rs decrease as the strain increases due to strain softening caused by the adiabatic heating and microstructural damage. As noted in Fig. 5, the values of Rs increase with increasing strain rate. This implies that Rs is not suitable to be used as rate sensitivity parameter for the present composites. On the other hand, it can be observed from Fig. 5 that the values of Rs almost increase linearly with the increasing strain rate in the strain range of region (ii).
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Fig. 6. Plot of flow stress at 6% strain vs. strain rate and linear interpolation.
high strain rate deformation [2,4,7]. Furthermore, San March et al. [2] deduced from the results of a designed reloading experiment that the strain rate sensitivity of particle reinforced aluminum composites is not linked with microstructural differences such as reinforcement size or damage accumulation, but associated with variations in the intrinsic mobility of dislocations in the aluminum matrix. The basic mechanical properties of the SPS composites are concluded in Table 1. It is obvious that composite reinforced by smaller particles has a higher flow stress, lower strain rate sensitivity. The higher modulus of 12 m SiC/Al may be attributed to its higher density and fewer defects in the SiC particles. 3.4. Damage Fig. 5. Rate sensitivity parameter Rs as a function of strain for different strain rates.
Therefore, the following equation can be used to describe the strain rate sensitivity of these composites reasonably: d = q + K ε˙
(2)
where ε˙ is the strain rate and K is the rate sensitivity parameter. Eq. (2) is constructed primarily based on the assumption that the dislocation-drag mechanism controls the deformation of metals at very high strain rate deformation (usually ε˙ > 1000/s) [8,9]. The flow stresses at 6% strain are plotted as function of strain rate as illustrated in Fig. 6. The choice of 6% strain is made so as to avoid the effect of the inhomogeneous deformation at lower strain during SHPB tests and to minimize the effect of damage accumulation and adiabatic heating at higher strain [2]. The rate sensitivity parameter K of the composites is determined from the slope of the linear regression line of the experimental results as shown in Fig. 6. With the decrease of average particle size, both the yield strengths and dynamic flow stresses of the composites increase significantly, while the strain rate sensitivity decreases. This result is consistent with that of San March et al. [2] and Chiem et al. [8]. They have summarized the strain rate sensitivity of the aluminum, aluminum alloys and their composites, including results from the published and their own works, and concluded that the strain rate sensitivity decreases as the flow stress increases. But so far the underlying microstructural deformation mechanism is still unclear. It is accepted for the present that the change of strain rate sensitivity of the alloys and composites is related to: (1) the change of the mobile dislocation density due to the mismatch of plastic strain between the matrix and the reinforcement; and (2) the resistance to the dislocation motion imposed by the reinforcement at
After SHPB tests, the specimens were compressed into thin disks, neither obvious barrelling deformation nor the adiabatic shear band was developed in the specimens. In order to identify the damage within the tested specimens under impact loading, the metallographic investigations and density measurements were carried out. The typical optical metallographies of tested specimens are shown in Fig. 7. The fractured SiC particles and cavities in the micrographs indicate that the microstructural damage is dominated by particle cracking and the interface debonding. Particles break primarily parallel to the loading direction, and this appears to be activated by stress transferred to particles across particle/particle contacts, since the fractured particles are mainly located in particle clustering regions. The presence of the cavities is caused by the debonding of the particle/matrix interface, because the debonded particles can be easily brushed away from the matrix during section polishing. The damage appears quite different between these two materials: the 12 m SiC/Al composite contains fewer broken particles and cavities comparing with the 45 m SiC/Al composite. There are two main factors responsible for the results in this study: (1) the larger particles contain more defects which can provoke the premature fracture under the dynamic loads as mentioned in Section 3.1; (2) the interfacial adhesion between the large SiC particle and the matrix is weak. It has been found in previous research that the large SiC particles are difficult to be wrapped into the composite powders and to be well cold-welded with the matrix during MA process. Since SiC particles are not electrically conductive, the spark plasma phenomenon is not generated between SiC particles and Al matrix during SPS. Besides, the sintering time is very short and the sintering temperature is beneath the melting point of the matrix
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material. The adsorptive gas and dust on the surface of SiC particles can also prevent the melted Al matrix adhering to SiC particles. As a consequence, the interfacial adhesion cannot be improved obviously through SPS sintering [28], and the bonding between large SiC particles and the matrix is very weak. Therefore, large particle debond easier than small one under impact loads. As a result, there are more debonded particles that clear away from the section and more cavities are left as shown in Fig. 7. Previous microstructural and fractographic analyses [2] on deformed specimens have revealed that the microstructural damage (including particle cracking, interface debonding and the microvoids nucleating and growing in the composites) can decrease the density of the particle reinforced MMCs. Therefore, damage accumulation can be quantified through a damage parameter D defined as a relative change in density of MMCs with strain [2,10]: D = 1 −
0
(3)
where 0 is the initial density and is the density of a tested specimen. The damage parameter D is plotted as a function of strain for dynamic compression as demonstrated in Fig. 8. In this plot, the higher the strain is, the higher the strain rate for the discrete points, namely the effect of the strain on the damage is coupled with that of the strain rate. Previous researches have revealed that the increasing strain rate has no obvious effect on the microstruc-
Fig. 8. Comparison of damage accumulation in composites due to dynamic compression.
tural damage in composites [2,27], so the tendency of the damage increasing with the strain is mainly with respect to the increasing strain as shown in Fig. 8. It is demonstrated that the value of damage parameter is higher for composite reinforced with larger particles, since the damage in the composite with larger particles is more evident as shown above. The basic trend of increasing damage with increasing particle size, as shown in Fig. 8, is also consistent with those found in previous work on particle reinforced MMCs [2,29–31]. 4. Conclusions In this paper, the 20% SiCp /pure Al composites were fabricated by SPS sintering with the mixture powders prepared by MA process, and the dynamic compressive properties of those composites were tested. The microstructure, strain rate sensitivity and damage of these composites are analyzed based on the results of the experiments, and the following conclusions can be drawn from this study:
Fig. 7. Optical metallography of compressed specimens: (a) 45 m SiC/Al and (b) 12 m SiC/Al. In both micrographs, the compressive load direction is nominally along the vertical. The strain rate and the final strain are 3000/s and 25%, respectively.
(1) The Al matrix composites with uniformly distributed SiC particles can be fabricated by using MA process to mix the powders followed by rapid SPS sintering. The MA process can effectively eliminate possible defects contained in the particles and reduce the average size of the particles. (2) These composites are strain rate sensitive and particle sizedependent. With the decrease of average particle size, both the yield strengths and dynamic flow stresses of the composites increase significantly, while their strain rate sensitivity decreases, and the hardening–softening transition strain is reduced. (3) The obvious strain softening in the dynamic stress–strain curves is mainly caused by the adiabatic heating and successive microstructural damage, and the adiabatic heating is the dominant factor for the strain softening under the dynamic loading. (4) Large particle debond easier than smaller one under the impact loads. The main reason is that the large SiC particles are difficult to be wrapped into the matrix during MA process and the interface bonding between large particles and matrix is quite weak. (5) The microstructural damage is dominated by particle cracking and interface debonding. The damage accumulation increases with the increase of the strain and the particle size in these composites.
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Acknowledgements This work was supported by the National Natural Science Foundation of China under Grant No. 10776019/A06, China Postdoctoral Science Foundation under Grant No. 20080430425 and by Open Foundation of State Key Laboratory of Explosion Science and Technology (Beijing Institute of Technology, China) under Grant No. KFJJ08-9. References [1] M. Guden, I.W. Hall, Mater. Sci. Eng. A 242 (1998) 141–152. [2] C. San Marchi, M. Fahe Cao, A. Kouleli, Mortensen, Mater. Sci. Eng. A 337 (2002) 202–211. [3] L.H. Dai, L.F. Liu, Y.L. Bai, Mater. Lett. 58 (2004) 1773–1776. [4] S. Yadav, D.R. Chichili, K.T. Ramesh, Acta Metall. Mater. 43 (12) (1995) 4453–4464. [5] N.Q. Zhao, P. Nash, X.J. Yang, J. Mater. Process. Technol. 170 (2005) 586–592. [6] M. Guden, I.W. Hall, Mater. Sci. Eng. A 232 (1997) 1–10. [7] S.C. Tjong, Z.Y. Ma, R.K.Y. Li, Mater. Lett. 38 (1999) 39–44. [8] C.Y. Chiem, X.W. Zhou, W.S. Lee, J. Phys. (1987), Colloque C3-577–586. [9] I˙ smail Tirtom, M. Güden, H. Yıldız, Comput. Mater. Sci. 42 (2008) 570–578. [10] M. Kouzeli, L. Weber, C.San. Marchi, A. Mortensen, Acta Mater. 49 (2001) 497–505.
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Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
The dynamic properties of SiCp /Al composites fabricated by spark plasma sintering with powders prepared by mechanical alloying process Jiangtao Zhang a,b,∗ , Huiji Shi a , Mingchun Cai a , Lisheng Liu b , Pengcheng Zhai b a b
Institute of Solid Mechanics, Department of Engineering Mechanics, Tsinghua University, Tsinghua East Road, Haidian Dist., Beijing 100084, PR China State Key Laboratory of Advanced Technology for Materials Synthesis and Processing, Wuhan University of Technology, Wuhan 430070, PR China
a r t i c l e
i n f o
Article history: Received 3 March 2009 Received in revised form 25 July 2009 Accepted 31 August 2009
Keywords: Dynamic compressive properties Spark plasma sintering Mechanical alloying Al matrix composites Strain rate Damage
a b s t r a c t The dynamic compressive properties of SiC particle reinforced pure Al matrix composites, fabricated by spark plasma sintering technique with mixture powders prepared by mechanical alloying process, were tested in this paper. Two different average SiC particle sizes of 12 m and 45 m were adopted, and the compressive tests of these composites at strain rates ranging from 800/s to 5200/s were conducted by split Hopkinson pressure bar. The damage mechanism of the SiCp /Al composites was analyzed through the microstructural observations and high-precision density measurements. Results show that the dynamic properties and damage accumulation of these composites are significantly affected by the particle distribution, size, particle cracking, particle/matrix interface debonding and adiabatic heat softening. The composites containing smaller SiC particles exhibit higher flow stress, lower strain rate sensitivity, and less damage at high strain rate deformation. © 2009 Elsevier B.V. All rights reserved.
1. Introduction Metal matrix composites (MMCs) reinforced with ceramic particles have the potential to provide desirable properties, such as high stiffness, high strength, good thermal stability and corrosion resistance. These superior properties make particle reinforced MMCs attractive in a wide range of applications, such as the automobile, aerospace and defense industries. The increasing applications of MMCs in high rate deformation conditions, such as structural impact, blast loading and metal processing, have made essential demand for a thorough understanding of constitutive and damage behaviors of MMCs at high rate deformation [1–4]. The dynamic properties of MMCs, produced by various traditional processes, have been tested over the last decades. The effects of matrix materials, fabricating processes and the type, size, distribution and shape of the particles on the strength and damage mechanism of MMCs have been considered [1–10]. It has been recognized that the uniform spatial distribution of the particles in the matrix is essential for improving the mechanical and physical properties of MMCs [11,12]. However, MMCs produced by the traditional processing techniques, such as casting, forging and powder metallurgy (PM) processes, suffer from some drawbacks, such as
∗ Corresponding author at: Institute of Solid Mechanics, Department of Engineering Mechanics, Tsinghua University, Tsinghua East Road, Haidian Dist., Beijing 100084, PR China. Tel.: +86 10 62771817; fax: +86 10 62781824. E-mail address: [email protected] (J. Zhang). 0921-5093/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2009.08.067
clustering of the particles, undesirable chemical reaction and poor adhesion between particles and matrix [9,13,14]. Recently, it has been reported that the high performance MMCs can be fabricated by using mechanical alloying (MA) process to mix the powders and then followed by rapid spark plasma sintering (SPS) [15,16]. MA process is originally developed by Benjamin [17] to produce a nickel based superalloy hardened by oxide dispersion by using high-energy ball mill. During MA process, two essential processes are involved, namely, cold-welding and fracturing of powder [17,18]. The former process makes the reinforcement particles wrapped up into the constituent powders, while the latter can separate the reinforcement particles from the constituent powders again. The constituent powders are repeatedly fractured and cold-welded throughout the milling, so that a homogenous distribution of the particles can be achieved [19]. Besides, the continuous collision between the reinforcement particles with balls, wall of container and other reinforcement particles can eliminate the possible defects contained in the reinforcement. Hence, the size of the particles is reduced and the strength is increased [12]. And so far, a large number of reports have dealt with the fabrication of particulates reinforced MMCs by MA process [12]. In order to produce bulk composites from powders prepared by MA process, various consolidation processes have been employed [12,20,21], such as sintering, hot-pressing (HP), hot-isostatic pressing (HIP) and hot extrusion of loose or prepressed powders. Recently, SPS technique has attracted extensive attention [22,23]. SPS is a novel consolidation technique for the manufacture of fully dense materials, such as metal alloys, ceramics, composites and
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functionally graded materials (FGMs). This technique combines plasma generation with resistive heating and pressure application. Self-heat generation by DC on-off pulse and spark plasma generated among the powder particles lead to the densification of the samples at low temperature within relatively short time [16,17]. The low sintering temperature and short consolidation time have been known to be very helpful in preventing chemical reaction at the interface of MMCs. The purpose of the present work is to investigate the dynamic properties of SiC (Vol. 20%)/Al matrix composites, which are fabricated from the mechanically alloyed (MAed) powders of pure Al and SiC particles and subsequently sintered by SPS. The dynamic properties of these composites are examined by using conventional split Hopkinson pressure bar (SHPB). The results of instrumented tests at high strain rates were compared with the microstructural observations. The strength, strain rate sensitivity, damage mechanism and heat softening of these composites were studied.
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A MTS-810 Material Testing System was used to perform quasistatic compression tests at constant crosshead speed with an initial strain rate of 0.0013/s, and strain gages were mounted on the samples to measure the strains until they detached from the samples, then larger strains were read from the machine output. A conventional SHPB equipment with 14.5 mm diameter bars was used for dynamic compression tests at strain rates ranging from 800/s to 5200/s. The detailed description about SHPB can be found in Ref. [24]. After SHPB tests, the density of the specimens were tested again, and the specimens were then sectioned parallel to the loading direction by using electro-discharge machining, and polished for identification of the microstructural damage. The load axis is parallel to the applied pressure direction during the SPS sintering. All tests were performed at room temperature with sufficient lubrication to prevent barrelling in the specimens. 3. Results and discussions
2. Experimental procedures 3.1. Microstructure 2.1. Materials and fabrication processes In the present study, a commercial pure Al powder with the average size of 6 m was used as the matrix material. SiC particles with the average size of 12 m and 45 m were used as the reinforcement respectively. The nominal volume fraction of SiC particles is 20%. The composite powders were first blended using MA process in a planetary high-energy ball mill (Fritsch, Pulverisette-4). In order to avoid excessive decrease of the SiC particle size, alumina balls were used as milling media due to their lower density compared with steel balls. The diameters of balls vary from 5 mm to 10 mm. The mixture of pure Al powder with SiC particles was put into a WC vial with alumina balls, and the weight ratio of balls to powder was fixed to 8:1. During the powder mixing, the rotational speed of the turn table of the ball mill was controlled at 240 rpm and the relative rotational speed ratio of the vial was 2. The milling time was 10 h. To prevent overheating, the system was constantly cooled using a fan, and ball milling was stopped periodically (every 1 h) and then resumed when the temperature of the WC vial decreased to room temperature. To prevent the oxidation of aluminum, the WC vial was full of argon atmosphere, and ethanol was used as a process control agent. Then the MAed powders were consolidated by SPS (Sumitomo Coal Mining, SPS-1050). Fourteen grams of MAed powder was placed into a graphite die with an inner diameter of 30 mm and heated under vacuum with an applied pressure of 30 MPa at 590 ◦ C for 10 min, and then followed by furnace cooling. The temperature was measured through a K-type thermocouple and the heating rate was equal to 50 ◦ C/min. The height of the SPS consolidated specimens was about 7 mm. No further heat-treatment was applied to the specimens.
After MA process, the MAed powders were firstly cold compacted at a pressure of 250 MPa in a steel die to prepare cylindrical compacts for the observation of the microstructure. The SEM micrographs of the compacts are shown in Fig. 1. The morphology of the MAed powders is rounded and composed of numerous layers, which are the typical features of the MA ending stage [25].
2.2. Mechanical tests The sintered composites were well polished and the densities were determined by Archimedes method prior to the test of mechanical properties. The microstructure of samples was analyzed by optical microscopy (OM) and scanning electron microscope (SEM). The chemical composition of SPS composites was verified by X-ray diffraction (XRD). The specimens which were machined in cylindrical shape with 6 mm in length and 8 mm in diameter by electro-discharge machining (EDM) were used for both dynamic and quasi-static compression tests. Both surfaces of the cylindrical specimens were finely ground and made as flat and as parallel as possible.
Fig. 1. SEM micrographs of compacts prepared by MAed powders with the pressure of 250 MPa. (a) 45 m SiC/Al and (b) 12 m SiC/Al.
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It can be clearly observed from Fig. 1 that the SiC particles are uniformly distributed in the MAed powders. The SiC particle size has great effect on the MAed powders. The size of MAed powder containing 45 m SiC particle is much smaller than that containing 12 m SiC particle. Besides, some large SiC particles are still not embedded into the composite powders in the former (as shown in Fig. 1(a)), while no free SiC particle can be observed in the latter (as shown in Fig. 1(b)). This result is consistent with that of Ruiz-Navas et al. [25]. Their experiments revealed that the bigger reinforcement leads to a smaller final size of composite powder when the same milling conditions were employed. The main reason is that small particles are easily stuck into, and then tightly wrapped in the composite powder. On the contrary, it is more difficult for the large particles to be trapped into the composite powders. Due to the low density of the alumina balls, the large SiC particles cannot be fragmented or comminuted under the ball-powder-ball collisions in a short milling time, so that the size of the SiC particles are not reduced too much. Only part of the large particle can be embedded in the surface of the composite powders under the ballpowder-ball collisions [26], and the constraint imposed by matrix on the SiC particles is quite weak. The embedding of large particles may cause large stress on the surface of the composite powders [5]. When mechanical alloying continues, the surface with SiC may fracture, or the SiC particles may separate from the composite powders again due to the constant impact of the balls and other composite powders. The processes occur repeatedly during milling process.
Fig. 2. The optical micrographs of SPS composites: (a) 45 m SiC/Al and (b) 12 m SiC/Al.
Table 1 Mechanical properties of SPS composites. Composites
Density (g/cm3 )
Modulus E (GPa)
Proof stress 0.2 (MPa)
45 m SiC/Al 12 m SiC/Al
2.774 2.781
104 113
228 273
Hence, the composite powders cannot grow as larger as those containing small SiC particle, and meanwhile, a portion of large SiC particles will be wrapped into the composite powders, as shown in Fig. 1(a). Fig. 2 shows the optical micrographs of SPS composites. No detectable microvoids can be observed, and the distribution of SiC particles is uniform in the composites. The size of SiC particles is apparently reduced. These excellent features are mainly attributed to the MA process. Due to the collision of the ceramic particles with balls, wall of container and other ceramic particles, the particles containing defects are more likely to be fractured during milling. So the particle size is reduced and there are large numbers of fine SiC particles in the composites, especially in those containing 45 m SiC particle. Since large SiC particles are difficult to be embedded in the matrix, they are subjected to the collision for a comparatively long time by comparison with small SiC particles. Therefore, the size reduction of larger SiC particles is more remarkable than that of smaller SiC particles, and thus more fine SiC particles are created in the MAed powder with 45 m SiC particles. Since the small SiC particles can be easily wrapped into the matrix, fracturing of SiC particles promotes the rate of embedding [5]. It can be seen in Fig. 2 that there are still many fissures and microvoids presented in the surface of the SiC particles, and larger SiC particles contain more defects. These defects may provoke the premature fracture of large SiC particle in the composites during the quasi-static and dynamic loadings. The densities of the SPS composites are shown in Table 1, and they are very close to the theoretical value. It indicates that the composites have been fully consolidated by SPS sintering. The composite of 45 m SiC/Al has a slightly lower density. It may be due to the more defects in the larger particles or the difference in the MA and sintering mechanisms caused by the difference of the particle size. In order to verify the composition of the SPS composites, X-ray diffraction (XRD) was applied and the XRD patterns of SPS composites were shown in Fig. 3. Except a little SiO2 contained in the raw material of SiC particles, no other impurity was detected in
Fig. 3. X-ray diffraction patterns of SPS composites.
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rate above 5% strain. At higher strain rates ε˙ ≥ 3400/s, however, the stress–strain curves present three distinct regions. Besides the two regions mentioned above, the last region (iii) of obvious strain softening can be observed. The flow stresses reach the maximum value near the strain of 23% for 45 m SiC/Al composite and 16% for 12 m SiC/Al composite, and then the flow stresses decrease gradually as the strain increases continuously. There are two main factors responsible for strain softening: (1) adiabatic heating, which is caused by the adiabatic character of the dynamic deformation process results in a significant temperature increase in the specimen at large strains, and (2) the successive microstructural damage, which takes place in the composites during the dynamic deformation process. Previous researches [2,27] have revealed that the strain rate has no obvious effect on the microstructural damage evolution in the composite, and the damage increases with the increasing strain. The strain hardening of the quasi-static stress–strain curves in the whole tested strain range indicates that it is the adiabatic heating that dominates the strain softening of the composites under dynamic loads. The hardening–softening transition strains decline with the decreasing particle size because of the dominating influence of the adiabatic heating softening, since the higher the flow stress, the more the deformation energy transferred to heat, and the higher the temperature of the specimen. At the strain rate 900/s < ε˙ < 3400/s, strain softening is not observed in the stress–strain curves. The main reason is that the total strains in these tests are less than the hardening–softening transition strains, where the strain softening caused by adiabatic heating and damage overwhelms the strain hardening of plastic deformation. 3.3. Strain rate sensitivity
Fig. 4. Stress–strain curves of SPS composites under quasi-static and dynamic loads.
SPS composites by XRD. This indicates that the interface reaction between the matrix and the particle is very weak or does not occur during SPS sintering. The alumina debris, which may be expected to be created from the alumina balls during milling process or derived from the oxide skin on the Al powder during the fabrication processes, is not found by XRD. It means that the fabrication processes are properly controlled. As a result, the content of alumina in the SPS composites is very low, although oxidation on the surface of the Al powders is considered inevitable [32]. So the effect of alumina debris on the mechanical properties of the SPS composites was not considered in the following discussions. 3.2. Stress–strain curves The stress–strain curves of the SPS composites at strain rates ranging from 0.0013/s to 5200/s are shown in Fig. 4. These curves exhibit the typical constitutive characteristics of the aluminum matrix composites as reported in the previous literatures: (1) strain rate sensitivity, namely the initial yield strength and the flow stress at a constant strain increase with the increasing strain rate, (2) particle size-dependence, namely the flow stress increases with the decrease of the particle size at certain strain rate, (3) similar strain hardening trends among different strain rate curves at strain less than the hardening–softening transition strain. Here the hardening–softening transition strain is defined as the strain at which the stress–strain curve is changed from strain hardening stage to strain softening stage. The quasi-static stress–strain curves can be divided into two distinct regions: (i) initial region of high work hardening rate until about 5% strain, and (ii) region of relatively low work hardening
In previous studies, various parameters have been defined to estimate the strain rate sensitivity of the rate dependent materials at high strain rate deformation. For example, Guden and Hall [6] and Tjong et al. [7] used the ratio of the increment of the dynamic flow stress to the quasi-static flow stress at certain strains as the strain rate sensitivity parameter: Rs =
d − q q
(1)
where Rs is the strain rate sensitivity parameter, d is the dynamic flow stress, q is quasi-static reference flow stress. The reference strain rate was chosen 0.0013/s, corresponding to the quasi-static strain rate adopted in this study. Fig. 5 shows the strain rate sensitivity parameter Rs as a function of strain for different strain rates. It should be noted that the low strain data (ε < 5%) of the stress–strain curves obtained from SHPB tests are unreliable, since the deformation in the tested specimens is inhomogeneous due to stress wave propagation [9], and usually the calculated stresses at a low strain decrease with the increasing strain rates from SHPB tests. The small values of Rs at low strains are not the true characteristic of the strain rate sensitivity of the composites. It can be seen from Fig. 5 that, in the strain ranging from 5% to the hardening–softening transition strain, namely region (ii) defined in Section 3.2, the values of Rs basically keep invariable for a given strain rate. This result confirms again the same strain hardening trend for the composites at different strain rate deformation. While in the region (iii), Rs decrease as the strain increases due to strain softening caused by the adiabatic heating and microstructural damage. As noted in Fig. 5, the values of Rs increase with increasing strain rate. This implies that Rs is not suitable to be used as rate sensitivity parameter for the present composites. On the other hand, it can be observed from Fig. 5 that the values of Rs almost increase linearly with the increasing strain rate in the strain range of region (ii).
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Fig. 6. Plot of flow stress at 6% strain vs. strain rate and linear interpolation.
high strain rate deformation [2,4,7]. Furthermore, San March et al. [2] deduced from the results of a designed reloading experiment that the strain rate sensitivity of particle reinforced aluminum composites is not linked with microstructural differences such as reinforcement size or damage accumulation, but associated with variations in the intrinsic mobility of dislocations in the aluminum matrix. The basic mechanical properties of the SPS composites are concluded in Table 1. It is obvious that composite reinforced by smaller particles has a higher flow stress, lower strain rate sensitivity. The higher modulus of 12 m SiC/Al may be attributed to its higher density and fewer defects in the SiC particles. 3.4. Damage Fig. 5. Rate sensitivity parameter Rs as a function of strain for different strain rates.
Therefore, the following equation can be used to describe the strain rate sensitivity of these composites reasonably: d = q + K ε˙
(2)
where ε˙ is the strain rate and K is the rate sensitivity parameter. Eq. (2) is constructed primarily based on the assumption that the dislocation-drag mechanism controls the deformation of metals at very high strain rate deformation (usually ε˙ > 1000/s) [8,9]. The flow stresses at 6% strain are plotted as function of strain rate as illustrated in Fig. 6. The choice of 6% strain is made so as to avoid the effect of the inhomogeneous deformation at lower strain during SHPB tests and to minimize the effect of damage accumulation and adiabatic heating at higher strain [2]. The rate sensitivity parameter K of the composites is determined from the slope of the linear regression line of the experimental results as shown in Fig. 6. With the decrease of average particle size, both the yield strengths and dynamic flow stresses of the composites increase significantly, while the strain rate sensitivity decreases. This result is consistent with that of San March et al. [2] and Chiem et al. [8]. They have summarized the strain rate sensitivity of the aluminum, aluminum alloys and their composites, including results from the published and their own works, and concluded that the strain rate sensitivity decreases as the flow stress increases. But so far the underlying microstructural deformation mechanism is still unclear. It is accepted for the present that the change of strain rate sensitivity of the alloys and composites is related to: (1) the change of the mobile dislocation density due to the mismatch of plastic strain between the matrix and the reinforcement; and (2) the resistance to the dislocation motion imposed by the reinforcement at
After SHPB tests, the specimens were compressed into thin disks, neither obvious barrelling deformation nor the adiabatic shear band was developed in the specimens. In order to identify the damage within the tested specimens under impact loading, the metallographic investigations and density measurements were carried out. The typical optical metallographies of tested specimens are shown in Fig. 7. The fractured SiC particles and cavities in the micrographs indicate that the microstructural damage is dominated by particle cracking and the interface debonding. Particles break primarily parallel to the loading direction, and this appears to be activated by stress transferred to particles across particle/particle contacts, since the fractured particles are mainly located in particle clustering regions. The presence of the cavities is caused by the debonding of the particle/matrix interface, because the debonded particles can be easily brushed away from the matrix during section polishing. The damage appears quite different between these two materials: the 12 m SiC/Al composite contains fewer broken particles and cavities comparing with the 45 m SiC/Al composite. There are two main factors responsible for the results in this study: (1) the larger particles contain more defects which can provoke the premature fracture under the dynamic loads as mentioned in Section 3.1; (2) the interfacial adhesion between the large SiC particle and the matrix is weak. It has been found in previous research that the large SiC particles are difficult to be wrapped into the composite powders and to be well cold-welded with the matrix during MA process. Since SiC particles are not electrically conductive, the spark plasma phenomenon is not generated between SiC particles and Al matrix during SPS. Besides, the sintering time is very short and the sintering temperature is beneath the melting point of the matrix
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material. The adsorptive gas and dust on the surface of SiC particles can also prevent the melted Al matrix adhering to SiC particles. As a consequence, the interfacial adhesion cannot be improved obviously through SPS sintering [28], and the bonding between large SiC particles and the matrix is very weak. Therefore, large particle debond easier than small one under impact loads. As a result, there are more debonded particles that clear away from the section and more cavities are left as shown in Fig. 7. Previous microstructural and fractographic analyses [2] on deformed specimens have revealed that the microstructural damage (including particle cracking, interface debonding and the microvoids nucleating and growing in the composites) can decrease the density of the particle reinforced MMCs. Therefore, damage accumulation can be quantified through a damage parameter D defined as a relative change in density of MMCs with strain [2,10]: D = 1 −
0
(3)
where 0 is the initial density and is the density of a tested specimen. The damage parameter D is plotted as a function of strain for dynamic compression as demonstrated in Fig. 8. In this plot, the higher the strain is, the higher the strain rate for the discrete points, namely the effect of the strain on the damage is coupled with that of the strain rate. Previous researches have revealed that the increasing strain rate has no obvious effect on the microstruc-
Fig. 8. Comparison of damage accumulation in composites due to dynamic compression.
tural damage in composites [2,27], so the tendency of the damage increasing with the strain is mainly with respect to the increasing strain as shown in Fig. 8. It is demonstrated that the value of damage parameter is higher for composite reinforced with larger particles, since the damage in the composite with larger particles is more evident as shown above. The basic trend of increasing damage with increasing particle size, as shown in Fig. 8, is also consistent with those found in previous work on particle reinforced MMCs [2,29–31]. 4. Conclusions In this paper, the 20% SiCp /pure Al composites were fabricated by SPS sintering with the mixture powders prepared by MA process, and the dynamic compressive properties of those composites were tested. The microstructure, strain rate sensitivity and damage of these composites are analyzed based on the results of the experiments, and the following conclusions can be drawn from this study:
Fig. 7. Optical metallography of compressed specimens: (a) 45 m SiC/Al and (b) 12 m SiC/Al. In both micrographs, the compressive load direction is nominally along the vertical. The strain rate and the final strain are 3000/s and 25%, respectively.
(1) The Al matrix composites with uniformly distributed SiC particles can be fabricated by using MA process to mix the powders followed by rapid SPS sintering. The MA process can effectively eliminate possible defects contained in the particles and reduce the average size of the particles. (2) These composites are strain rate sensitive and particle sizedependent. With the decrease of average particle size, both the yield strengths and dynamic flow stresses of the composites increase significantly, while their strain rate sensitivity decreases, and the hardening–softening transition strain is reduced. (3) The obvious strain softening in the dynamic stress–strain curves is mainly caused by the adiabatic heating and successive microstructural damage, and the adiabatic heating is the dominant factor for the strain softening under the dynamic loading. (4) Large particle debond easier than smaller one under the impact loads. The main reason is that the large SiC particles are difficult to be wrapped into the matrix during MA process and the interface bonding between large particles and matrix is quite weak. (5) The microstructural damage is dominated by particle cracking and interface debonding. The damage accumulation increases with the increase of the strain and the particle size in these composites.
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Acknowledgements This work was supported by the National Natural Science Foundation of China under Grant No. 10776019/A06, China Postdoctoral Science Foundation under Grant No. 20080430425 and by Open Foundation of State Key Laboratory of Explosion Science and Technology (Beijing Institute of Technology, China) under Grant No. KFJJ08-9. References [1] M. Guden, I.W. Hall, Mater. Sci. Eng. A 242 (1998) 141–152. [2] C. San Marchi, M. Fahe Cao, A. Kouleli, Mortensen, Mater. Sci. Eng. A 337 (2002) 202–211. [3] L.H. Dai, L.F. Liu, Y.L. Bai, Mater. Lett. 58 (2004) 1773–1776. [4] S. Yadav, D.R. Chichili, K.T. Ramesh, Acta Metall. Mater. 43 (12) (1995) 4453–4464. [5] N.Q. Zhao, P. Nash, X.J. Yang, J. Mater. Process. Technol. 170 (2005) 586–592. [6] M. Guden, I.W. Hall, Mater. Sci. Eng. A 232 (1997) 1–10. [7] S.C. Tjong, Z.Y. Ma, R.K.Y. Li, Mater. Lett. 38 (1999) 39–44. [8] C.Y. Chiem, X.W. Zhou, W.S. Lee, J. Phys. (1987), Colloque C3-577–586. [9] I˙ smail Tirtom, M. Güden, H. Yıldız, Comput. Mater. Sci. 42 (2008) 570–578. [10] M. Kouzeli, L. Weber, C.San. Marchi, A. Mortensen, Acta Mater. 49 (2001) 497–505.
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