Effect of Overaging and Particle Size on Tensile Deformation and Fracture of Particle-Reinforced Aluminum Matrix Composites J.J. WILLIAMS, G. PIOTROWSKI, R. SAHA, and N. CHAWLA The effect of reinforcement particle size and overaging treatment on the tensile behavior and fracture morphology of a 2080/SiC/20p composite was investigated. Tensile behavior was profoundly influenced by particle size and matrix strength. The composite strength increased with a decrease in particle size, while overaging greatly reduced the strength of the composite, independent of particle size. Almost all particles on the fracture plane were fractured, and the amount of particle fracture in the composites was insensitive to overaging and particle size, due to the excellent bonding between SiC particles and the Al matrix. Fractography showed that void nucleation in the matrix of peak-aged composites took place primarily at very fine SiC particles, which were much smaller than the average SiC particle size. Subsequent failure took place by the tearing topography surface (TTS) mechanism. In the overaged composite, composites failed by a more conventional void nucleation and growth process, where void nucleation took place at coarsened S precipitate particles, resulting in smaller and more elongated voids.
I. INTRODUCTION
DISCONTINUOUSLY reinforced aluminum (DRA) is emerging as an attractive lightweight alternative to conventional monolithic materials. It is well known that, in general, the incorporation of high stiffness ceramic particles into aluminum results in a composite with enhanced stiffness and strength over the unreinforced alloy.[1–7] Strengthening in the composite can be partitioned into direct and indirect strengthening. Direct strengthening is achieved by classical shear lag mechanisms, i.e., where load transfer takes place from the matrix to the stiffer reinforcement.[8,9,10] Indirect strengthening arises from changes to the matrix microstructure in the composite by the addition of the reinforcement.[11,12,13] An example of indirect strengthening is the dislocation punching that takes place due to the coefficient of thermal expansion mismatch between the ceramic particle and the aluminum matrix. The thermally induced dislocations significantly strengthen the matrix after processing. With a decrease in particle size or an increase in volume fraction, the level of indirect strengthening will increase because of an increase in the particle/matrix interfacial area. Other examples of indirect strengthening include grain size refinement due to pinning of grain boundaries by the ceramic particles and, in agehardenable matrix alloys, heterogeneous nucleation of precipitates at dislocations. The tensile behavior is profoundly influenced by reinforcement volume fraction and size. While there is general agreement that an increase in volume fraction and decrease in reinforcement particle size are conducive to strengthenJ.J. WILLIAMS, Postdoctoral Fellow, G. PIOTROWSKI, Undergraduate Research Assistant, R. SAHA, Graduate Research Assistant, and N. CHAWLA, Assistant Professor and Graduate Chair, are with the Department of Chemical and Materials Engineering, Arizona State University, Tempe, AZ 85287-6006. Contact e-mail:
[email protected] Manuscript submitted December 28, 2001. METALLURGICAL AND MATERIALS TRANSACTIONS A
ing in DRA, several mechanisms have been proposed for the modes and sequence of damage in DRA.[2,5,14–17] Lewandowski and co-workers[14–17] have shown that particle fracture takes place first, and that a triaxial state of stress in the matrix, due to plastic constraint imposed by the surrounding reinforcement particles, causes void nucleation and growth in the Al matrix. Whitehouse et al.[18,19] observed that during tensile loading, voids are nucleated at sharp corners or at the poles of the reinforcement, as means of stress relaxation in the matrix of the composite. They propose that void growth and coalescence precede particle fracture. It would appear that the earlier onset of voids, prior to particle fracture, is directly related to the highly ductile, pure aluminum matrix in the composites in that study. A variety of studies have examined the role and extent of particle fracture under tensile loading.[4–7,14–17] A relatively small number of these studies, however, have reported on the effects of overaging on the tensile behavior of the composite.[14,15] In this study, we have studied the combined effects of overaging and reinforcement particle size on the tensile behavior of an Al-Cu-Mg age-hardenable alloy reinforced with SiC particles. In particular, we have quantified and correlated the extent of particle fracture with tensile behavior for various aging conditions at two different particle sizes. The fracture behavior of the matrix was also examined as a function of overaging treatment. Following the work of Krajewski et al.[20] and Chawla et al.,[2] we have maintained a relatively uniform matrix microstructure in all composites, by applying a T8 treatment to all composites (solution treating, rolling, and aging). The rolling treatment produces a uniform distribution of dislocations in the matrix, which serve as heterogeneous nucleation sites for precipitation in the composite. In this manner, a homogeneous distribution of the precipitates, in the composites and the unreinforced alloy, was achieved, and the contribution from indirect strengthening kept relatively constant. VOLUME 33A, DECEMBER 2002—3861
Fig. 1—Powder metallurgy and extrusion process for fabricating 2080/SiCp composites.
II. MATERIALS AND EXPERIMENTAL PROCEDURE The composites in this study consisted of a 2080 Al alloy (Al-3.6Cu-1.9Mg-0.2Zr) reinforced with 20 pct SiC particles. The effect of particle size was examined on composites with an average particle size of 25 m (F280) and 9 m (F600), respectively. The composites were fabricated by powder metallurgy processing and extrusion (Aluminum Company of America, Alcoa Center, PA). Figure 1 shows the processing sequence for the composites. The powders were blended, vacuum degassed, and cold pressed, followed by hot pressing and extrusion. Particle size and aspect ratio distributions after extrusion were conducted using a conventional digital image analysis system (NIH Image). To measure aspect ratios, each particle was fitted as an ellipse, and the aspect ratio was taken as the ratio of major to minor axis. A minimum of 500 particles were analyzed for each composite to obtain the average aspect ratio. Particle size and aspect ratio measurements were conducted after extrusion because of changes in particle size and aspect ratio due to particle fracture. The matrix microstructure was controlled by a T8 treatment (solutionizing at 493 ⬚C for 2 h, water quenching, cold rolling to 5 pct reduction in thickness, and aging at 175 ⬚C for 24 h). Subsequent overaging was conducted for 24 hours, at 200 ⬚C, 225 ⬚C, and 250 ⬚C, respectively. The thermomechanical processing and heat treatment conditions were similar to those described by Chawla et al.[21] Tensile specimens were machined by low stress grinding. 3862—VOLUME 33A, DECEMBER 2002
Tensile testing was conducted at ambient temperature in strain control at a constant strain rate of 10⫺3/s. Details of the testing techniques and specimen geometry are described elsewhere.[2] Quantification of particle fracture was conducted by examining mirror fracture surfaces and matching particle halves. Due to the angular shape of the particles, it is necessary to use this approach to determine whether a given particle has indeed fractured. A statistically significant sample of about 200 particles per condition was taken in this analysis. Particle fracture quantification was also conducted on polished cross sections of fractured specimens. In the finer particle size composite, approximately 140 particles were counted at a given distance from the fracture surface. The sample size was slightly lower for the coarser particle size composite (the larger average particle size results in a smaller number of total particles, for a given volume fraction). III. RESULTS AND DISCUSSION A. Microstructure Figure 2 shows the microstructure of the composites along the extrusion direction. Note that the larger dimension of the particles is aligned parallel to the extrusion direction. Some Al-Cu dispersoids (finer particles in the matrix of the composite), which are a product of the powder metallurgy process, remain in the matrix. Table I lists the particle size distribution, of both composites, parallel to the extrusion direction. Characterization of the SiC particles indicated a METALLURGICAL AND MATERIALS TRANSACTIONS A
Fig. 2—Microstructure of 2080/SiC/20p composites: (a) F280 particle size and (b) F600 particle size. Note the alignment of particles along the extrusion direction.
Table I. Particle Size Distribution for 2080/SiC/20p-T8 Composites after Extrusion Material
D5 (m)
D50 (m)
D95 (m)
2080/SiC/20p-T8-F280 2080/SiC/20p-T8-F600
5.3 1.9
25 8.6
54.8 16.3
Fig. 3—Aspect ratio distributions for F280 and F600 composites.
broad distribution of particle sizes, particularly for the F280 particle size. While the average particle sizes of F600 and F280 were 9 and 25 m, respectively, the largest particles were on the order of 16 and 55 m, respectively. Thus, a significant fraction of the particles may have had lower particle strengths than those indicated by the average particle size. Aspect ratio analysis of the particles indicated that the particles had a mean aspect ratio of close to 2, although the F280 particles had a slightly lower aspect ratio than the F600 particles. Some degree of particle fracture due to rolling was observed in the F280 composites. As discussed in the introduction, the quenching step prior to aging promotes the formation of dislocations at the METALLURGICAL AND MATERIALS TRANSACTIONS A
particle/matrix interface to accommodate the thermal mismatch between matrix and reinforcement. As a result of the rolling step in the T8 process, a uniform dislocation distribution is obtained in the matrix, which results in a homogeneous distribution of precipitates after aging. Chawla et al.[21] studied the matrix microstructure of the 2080/SiC/ 20p-T8, in peak-aged and overaged conditions, by transmission electron microscopy. The relative size and spacing of the S⬘ (A12CuMg) precipitates in the unreinforced alloy and the matrix of the composite are quite similar. Size and spacing of the precipitates after the various aging treatments, from the work of Chawla et al.,[21] is shown in (Figure 4). All micrographs were taken along the 具100典 zone axis orientation, so that a valid comparison of precipitate sizes and spacing could be made. The three 具100典 directions of the S⬘ precipitates are shown as perpendicular needles in the plane, and as points coming out of the plane (Figure 4). The [100] growth directions of the S⬘ orthorhombic unit cell correspond to the [100] direction of the fcc Al matrix.[22] With increasing severity of overaging the precipitates became coarser, and the interprecipitate spacing increased. Chawla et al.[21] showed that the precipitate spacing increased from 28 nm for T8, to 73 nm after overaging for 24 hours at 250 ⬚C. From a strengthening point of view, the degree of strengthening contribution from the precipitates will significantly decrease with overaging treatment, as the precipitates become incoherent with the matrix, and Orowan bypass of dislocations becomes the dominant mechanism for dislocation motion. B. Tensile Behavior The tensile strength of the two composites for the different aging treatments is shown in (Table II). A comparison to tensile strength of the unreinforced alloy is also shown.[2] The composite reinforced with the finer particles exhibited significant strengthening over the unreinforced alloy, but the F280 composite did not. It is interesting to note, however, that particle size itself did not completely control the ultimate tensile strength of the composites. This is readily apparent by the significant decrease in composite strength with severity of overaging, for a given particle size. The unreinforced VOLUME 33A, DECEMBER 2002—3863
Fig. 4—Precipitate structure in the matrix of 2080/SiC/20p-T8 composites as a function of heat treatment and overaging—具100典 zone axis.[21] With increasing overaging temperature, the precipitates increase in size and spacing.
Table II. Ultimate Tensile Strength of 2080-T8 and 2080/ SiC/20p-T8 composites with F280 and F600 Particle Size Aging Treatment
2080-T8[2]
2080/SiC/ 20p-F280
2080/SiC/ 20p-F600
T8 T8 ⫹ 24 h at 200 ⬚C T8 ⫹ 24 h at 225 ⬚C T8 ⫹ 24 h at 250 ⬚C
522 473 411 333
484 414 348 290
576 489 410 —
alloy also showed a similar rate of decrease in strength with overaging. To understand the role of particle size and matrix microstructure on tensile deformation, it is, instructive to discuss the evolution of loading in particle-reinforced metal matrix composites (MMCs). Following elastic loading, where a significant fraction of the stress is initially borne by the reinforcement, the composite undergoes microplastic yielding. Microplasticity in the composites takes place at stress concentrations in the matrix at the poles of the reinforcement or at sharp corners of the reinforcing particles.[23–25] With an increase in applied strain, the particles are loaded to a higher stress. Due to much lower strain-to-failure of the SiC 3864—VOLUME 33A, DECEMBER 2002
particles, compared to the matrix, the particles will generally begin to fracture prior to ultimate tensile strength of the composite. In some cases, the onset of particle fracture has been observed at stresses slightly below the macroscopic yield strength of the composite,[26] and increases with size and volume fraction of particles.[27] The F280 composites had a lower strength than even the unreinforced alloy due to a fraction of the coarser particles in F280 being cracked during rolling. Cracked particles do not carry any load and can be effectively thought of as voids, so the strength of this composite was lower than that of the unreinforced material. The F600 composites had higher strengths because of the higher strength of the finer particles, which can be attributed to a lower probability of a strength-limiting flaw being present in the smaller material volume. Thus, the larger the particle size, the lower the particle strength, and the lower the ultimate tensile strength of the composite. It has also been proposed that the interparticle spacing (directly related to the particle size) influences the nature of dislocation structures in the matrix. The formation of dislocation tangles around the particles, due to plastic incompatibility between the reinforcement and matrix, and the formation of a dislocation cell structure with a cell size inversely proportional to METALLURGICAL AND MATERIALS TRANSACTIONS A
(a)
(b) Fig. 5—Scanning electron micrographs of tensile fracture surface for peak-aged composites: (a) 2080/SiC/20p (F280) and (b) 2080/SiC/20p (F600). The matched halves of the fracture surface are shown for each case.
the interparticle spacing could also contribute to the increase in strength with decreasing reinforcement particle size.[28] Quantification of the degree of particle fracture provided insight into the dependence of damage on particle size and overaging treatment. Scanning electron micrographs of the matched halves of the fracture surface for the peak-aged composites (F280 and F600) are shown in Figure 5. The fraction of fractured particles for all composites and aging treatments, obtained from the fracture surfaces, is shown in Table III. The extent of particle fracture was slightly higher in F280 composites because of the higher particle size. More importantly, however, the amount of particle fracture did not change significantly with aging treatment. Our results can be interpreted by examining the relationship between particle/matrix interface strength vs particle strength. If the particle strength is lower than the interface strength, then particle fracture will take place first and particle pullout will not be predominant. With an increase in particle strength (a METALLURGICAL AND MATERIALS TRANSACTIONS A
Table III. Statistics of Particle Fracture Observed on the Tensile Fracture Surface
Material 2080/SiC/20p 2080/SiC/20p 2080/SiC/20p 2080/SiC/20p 2080/SiC/20p 2080/SiC/20p 2080/SiC/20p
(F600) (F600) (F600) (F280) (F280) (F280) (F280)
Aging Treatment T8-peak-aged peak ⫹ 200 ⬚C peak ⫹ 225 ⬚C T8-peak-aged peak ⫹ 200 ⬚C peak ⫹ 225 ⬚C peak ⫹ 250 ⬚C
24 h 24 h 24 h 24 h 24 h
Fraction of Fractured Particles (Pct) 92.2 91.5 90.5 94.3 94.9 94.4 94.5
decrease in particle size), a larger fraction of the particles will have a strength higher than the interfacial strength, so a larger fraction of particles will be pulled out. It follows VOLUME 33A, DECEMBER 2002—3865
Fig. 6—Fraction of fractured particles as a function of distance from the fracture surface for F280 and F600 particle size. The higher fraction of fracture particles at the surface may be attributed to higher localization of strain.
that the fraction of pulled-out particles should also increase with a decrease in matrix strength, such as that caused by overaging, since the interfacial strength should also decrease. It is expected that the fraction of pulled-out particles would be more pronounced with a decrease in particle size, since this would increase the ratio of particle strength to interface strength. In our work, the relative insensitivity to particle fracture to overaging, even after severe overaging treatments, points to the excellent bonding and clean nature of the particle/matrix interface in the composites studied. This finding is further substantiated by fractography results presented later in this section, which confirm the absence of particle/matrix decohesion. While the fraction of particles fractured on the fracture surface provided a reasonable estimate of the degree of particle fracture, the number of particles fractured may be enhanced by the fact that a propagating crack will take the path of least resistance, and “artificially” increase the number of fractured particles. Thus, specimens were also sectioned and the number of fractured particles was also measured as a function of distance from the fracture surface (Figure 6). While there was still no significant difference in the extent of particle fracture, particularly with respect to aging treatment, the magnitude of particle fracture was significantly reduced, even 100 m away from the fracture plane, and steadily decreased with distance from the fracture surface. The reason for the drastic increase in particle fracture at the fracture plane lies in the intensification of strain at the fracture surface. This phenomenon has been observed in Al-SiMg/Sip composites[17] and Al/SiCp systems,[14,15] where a similar trend was observed with respect to the effect of particle size on the extent of particle fracture. While the nature and extent of particle fracture is very important in understanding the mechanisms of damage in the composites, the deformation in the matrix is equally important. Fractography of peak-aged and overaged composites, at both particle sizes, indicated very different fracture mechanisms. Figure 7 shows the fracture surface of peakaged F280 composite. Significant particle fracture was observed. Once the particle is cracked, it can be effectively 3866—VOLUME 33A, DECEMBER 2002
Fig. 7—Fracture surface of peak-aged 2080/SiC/20p (F280) composites in (a) lower magnification and (b) matrix fracture morphology. Evidence of TTS fracture mode is observed.
considered as a void in the matrix. Larger particles fail prior to smaller particles, nucleating voids in the matrix.[29] Because of the presence of rigid particles, the matrix is under a tensile triaxial state of stress and is unable to relax strain by plastic deformation. Rather, strain relaxation in the matrix takes place by void nucleation and propagation, which take place at a lower far-field applied strain than that observed in the unreinforced material. Stress triaxiality is further enhanced in particle clusters.[14,30,31] In the peak-aged composites, microvoids in the matrix seem to nucleate primarily at very fine SiC particles, presumably from the lower end of the reinforcement particle size distribution. The matrix fracture morphology seemed to be a result of microvoid nucleation and ductile tearing of the matrix. A similar microplastic mode of ductile tearing has been documented in steels and aluminum alloys, and is termed tearing topography surface (TTS).[32] In TTS, it has been proposed that microvoid nucleation occurs at very closely spaced nuclei, and that strain localization prevents any significant amount of subsequent void growth, which prevents the formation of well-developed voids. This is METALLURGICAL AND MATERIALS TRANSACTIONS A
Fig. 8—Fracture surface of overaged 2080/SiC/20p (F280) composites: (a) lower magnification and (b) matrix fracture morphology. Classical void nucleation and growth mechanisms are predominant.
Fig. 9—Fracture surface of peak-aged 2080/SiC/20p (F600) composites: (a) lower magnification and (b) matrix fracture morphology. Evidence of TTS fracture mode is observed.
indeed the case in the composites studied here, where the matrix between the finely dispersed particles was under a high triaxial state of stress. The T8 matrix in the peak-aged condition also had a very high strength, which could lend more credence to the hypothesis of fracture by a TTS type of ductile tearing mechanism. In the overaged composites, the fracture surface exhibited evidence of more classical microvoid growth and coalescence behavior observed in conventional engineering alloys[33] (Figure 8). The nucleation sites seemed to take place at non-reinforcement S-phase precipitate particles in the matrix, which coarsened significantly during overaging, rather than fine SiC particles. Thus, the number of voids in the overaged composites was larger, but the voids were much smaller in size and more elongated than in the peak-aged condition. Cavitation at SiC particle/matrix interfaces was also more predominant in the overaged condition. Qualitatively similar fracture characteristics were observed in peakaged and overaged F600 composites (Figures 9 and 10). The overaged fracture surface showed much more predominant void growth. The voids also appeared to be smaller and
more elongated. Figure 9(b) shows a coarsened precipitate particle at the base of a void in the matrix, which served as a nucleation site for a void. From a more microscopic point of view, void nucleation and matrix tearing could also be attributed to dislocation bypass of the precipitates in the overaged condition. It is interesting to note that the matrix fracture morphology is different at the particle/matrix interface and in the matrix of the composite (Figure 10). At the particle/matrix interface, there was less evidence of void formation and the fracture was relatively featureless, indicating severe ductile tearing that took place after particle fracture. Once the crack was nucleated into the matrix (from the fracture of the particle), conventional void growth mechanisms seemed to operate. Some evidence of particle/matrix decohesion, particularly at sharp corners of the particles, was also observed. Evidence of excellent bonding between the particle and matrix can be seen in Figure 11, which shows the particle impression (after pullout) in the matrix, which had much finer and larger number of voids than the rest of the matrix. Processinginduced defects within the SiC particle were also visible on
METALLURGICAL AND MATERIALS TRANSACTIONS A
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Fig. 10—Fracture surface of overaged 2080/SiC/20p (F600) composites: (a) lower magnification and (b) matrix fracture morphology. Classical void nucleation and growth mechanisms are predominant.
the fractured particle. Thus, the processes for void appeared smaller and seemed to grow in a somewhat more elongated fashion. The excellent bond strength between the particle and matrix in the extruded 2080/SiCp system would seem to explain the absence of large-scale particle/matrix interfacial debonding. IV. CONCLUSIONS The following conclusions can be made on this study of the effect of reinforcement particle size and overaging on the tensile behavior and fracture morphology of particlereinforced MMCs. 1. Tensile strength of the composites increased with a decrease in particle size. Smaller particles had a smaller volume and, thus, a lower probability of containing a strength-limiting flaw. Overaging greatly reduced the strength of the composite, regardless of particle size, because of the decrease in matrix strength. 2. Almost 95 pct of particles fractured on the fracture plane. 3868—VOLUME 33A, DECEMBER 2002
Fig. 11—Fracture morphology of overaged 2080/SiC/20p (F600) composites: (a) particle surface showing smaller equiaxed dimples in the Al matrix and (b) change in morphology from featureless fracture in the particle/ matrix interface to conventional void growth characteristics away from the particle/matrix interface. Note the processing-induced defects in the SiC particle.
The amount of particle fracture in the composites was also insensitive to overaging and particle size. This can be attributed to the excellent bonding between SiC particles and the Al matrix. Strain intensification at the fracture plane greatly increased the degree of particle fracture on the fracture surface, with respect to particles away from the fracture plane. 3. Fractography indicated that the high-strength T8-heattreated matrix failed by a TTS mechanism. Void nucleation took place primarily at very fine SiC particles, which were much smaller than the average SiC particle size. Void growth was quite limited, which was hypothesized to be caused by strain localization, and fracture by a tearing mode. Overaged composites failed by a more conventional void nucleation and growth mechanism, with voids being nucleated primarily at coarsened precipitate particles that resulted from overaging. 4. Fracture morphology also changed with distance from the particle/matrix interface, particularly in the overaged METALLURGICAL AND MATERIALS TRANSACTIONS A
composites. At the interface, fracture was relatively featureless, resulting from cracking of the matrix resulting from particle fracture. Fracture morphology in the matrix away from the interface exhibited characteristics of classical void nucleation and coalescence mechanisms.
ACKNOWLEDGMENTS The authors are grateful for the financial support from the United States Automotive Materials Partnership (USAMP), through a grant from the Department of Energy, and the Office of Naval Research (Dr. A.K. Vasudevan, Program Manager, Contract No. N000140110694). The authors also acknowledge Dr. Warren Hunt, Jr. (Aluminum Consultants Inc.), for supplying the materials used in this study. REFERENCES 1. N. Chawla and Y.-L. Shen: Adv. Mater. Eng., 2001, vol. 3, pp. 357-70. 2. N. Chawla, C. Andres, J.W. Jones, and J.E. Allison: Metall. Mater. Trans. A, 1998, vol. 29A, pp. 2843-54. 3. L.C. Davis, C. Andres, and J.E. Allison: Mater. Sci. Eng., 1998, vol. A249, pp. 40-45. 4. J. Llorca, S. Suresh, and A. Needleman: Metall. Trans. A, 1992, vol. 23A, pp. 919-34 5. P.M. Mummery, B. Derby, D.J. Buttle, and C.B. Scruby: Proc. Euromat 91, T.W. Clyne and P.J. Withers, eds., European Materials Society, Cambridge, United Kingdom, vol. 2, pp. 441-47. 6. T.S. Srivatsan: J. Mater. Sci., 1996, vol. 31, pp. 1375-88. 7. T.S. Srivatsan and J. Mattingly: J. Mater. Sci., 1993, vol. 28, pp. 611-20. 8. H.L. Cox: Br. J. Appl. Phys., 1952, vol. 3, p. 122. 9. A. Kelly: Strong Solids, Clarendon Press, Oxford, United Kingdom, 1973, p. 157. 10. V.C. Nardone and K.M. Prewo: Scripta Metall., 1989, United Kingdom, vol. 23, p. 291. 11. K.K. Chawla and M. Metzger: J. Mater. Sci., 1972, vol. 7, p. 34. 12. M. Vogelsang, R.J. Arsenault, and R.M. Fisher: Metall. Trans A, 1986,
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