- Email: [email protected]
Al composite
Scripta Materialia 49 (2003) 59–63 www.actamat-journals.com
Near-interface microstructure in a SiC/Al composite R.D. Evans *, J.D. Boyd Department of Materials and Metallurgical Engineering, QueenÕs University, Kingston, Canada ON K7L 3N6 Received 29 January 2003; received in revised form 27 March 2003; accepted 31 March 2003
Abstract The microstructure of a SiC particulate reinforced Al 2080 alloy has been characterized by focussed ion beam microscopy and transmission electron microscopy. A 40 nm thick amorphous interface layer develops during processing of the composite material by diffusion of Al and Mg into the pre-existing SiO2 layer on the 3 lm SiC particles. During tensile deformation, damage occurs in the near-interface zone of the matrix, at the particle–matrix interface and within the SiC particles. Ó 2003 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. Keywords: Composites; Aluminium; Powder processing; Microstructure; Damage
1. Introduction There has been considerable research on the effects of particle-induced damage in metal matrix composites (MMCs) [1–3]. Most of these studies have been on MMCs with particulate size >10 lm, where particle cracking is the dominant damage mechanism and the particle–matrix interface appears to have little effect on the overall damagefracture behaviour. In the current study, an MMC was produced in which interface decohesion was the dominant damage-initiation mechanism. The experimental design incorporated (a) small particulate size (3 lm) to suppress particle cracking, and (b)
* Corresponding author. Present address: Kingston Research and Development Centre, Alcan International Limited, 945 Princess Street, Kingston, Canada ON K7L 5L9. Tel: +1-613541-2137; fax: +1-613-541-2134. E-mail address: [email protected] (R.D. Evans).
powder processing to minimise interface reactions and produce a well-adhered, clean interface. This paper reports on microstructural evolution of the near-interface zone during deformation, and identification of an interface decohesion mechanism. Quantitative measurements of damage evolution and comparisons with existing damage models are reported elsewhere [4].
2. Experimental materials and methods The MMC material was 20 vol% SiC in Al 2080 (Al–4Cu–2Mg). The mean diameter of the SiC was 3.3 0.5 lm. The material was produced by a powder-processing route that included mixing, cold compaction, and hot extrusion at 500 °C to 16 mm diameter rod. Complete processing details are given in Ref. [4]. Tensile samples having an 18 mm 3 mm diam. gauge section were machined from the extruded rod, and solution treated for 2 h
1359-6462/03/$ - see front matter Ó 2003 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. doi:10.1016/S1359-6462(03)00180-5
60
R.D. Evans, J.D. Boyd / Scripta Materialia 49 (2003) 59–63
at 510 °C followed by water quenching (STQ condition). To study the evolution of damage and the nearinterface microstructure during deformation, tensile tests were carried out on an Instron 8502 servo-hydraulic testing machine at a strain rate of 104 s1 . Samples were deformed to various strains up to the fracture point, which occurred at about 2% plastic strain. Following deformation, all samples were aged for 24 h at 177 °C (DA condition). This treatment was employed to decorate the matrix dislocations with precipitates, and to represent a typical heat-treated condition for such an MMC material. The microstructures of the MMC prior to and after deformation were characterized by focussed ion beam (FIB) microscopy and transmission electron microscopy (TEM). Sections for FIB examination were cut parallel to the tensile direction. Final surface preparation by Gaþ ion milling and secondary electron imaging was carried out using a Micrion 2500 FIB system operating at 50 keV accelerating voltage. More details on this technique are given in Ref. [4]. TEM samples were prepared using conventional sample preparation techniques. Thin discs were produced and dimpled to 10 lm centre thickness on a Fischione Model 2000 dimpler. Final thinning to produce electron transparent samples was per-
formed on a Gatan Model 600 dual ion mill. Specimens were examined on a Philips CM20 TEM at an accelerating voltage of 200 kV, and X-ray microanalysis was carried out using a Noran Voyager EDS system with light-element (>B) capability. X-ray mapping of the particulate– matrix interface was carried out using a Philips CM20-FEG TEM equipped with an Oxford Instruments EDS system.
3. Results The microstructure of the MMC in the STQ condition is illustrated by the FIB micrograph in Fig. 1(a). Ion channelling produces orientation contrast of the matrix grains, and there is also atomic-number contrast of the SiC particulate, as well as other phases. Microstructural features indicated in Fig. 1(a) are SiC particles (S), dispersoid particles (D) and interface layer (I). TEM of the STQ material revealed a relatively low density of quenched-in dislocation loops and helices in the matrix grains, as well as large (100– 500 nm) particles within matrix grains and at the interfaces. Most of the particles were determined to be Cu-, Zr-, Fe-, or Mn-rich and were assumed to be dispersoid phases which were not dissolved during the solution treatment at 510 °C. The
Fig. 1. (a) STQ condition, FIB micrograph: S––SiC, D––dispersoid, I––interface layer; (b) STQ condition, TEM showing interface layer.
R.D. Evans, J.D. Boyd / Scripta Materialia 49 (2003) 59–63
remaining particles were identified as Al2 O3 , which likely resulted from the oxide layer on the initial Al 2080 powder. The SiC particles had hexagonal crystal structure and contained stacking faults characteristic of the 6H SiC polytype. There was also evidence of a thin layer at the interface between SiC particles and matrix grains, which appeared as a white line in FIB images (Fig. 1(a) at I). In TEM, this layer was clearly visible as a phase 20–40 nm in thickness, having a charac-
61
teristic mottled appearance (Fig. 1(b)). Selected area diffraction (SAD), and convergent beam diffraction (CBED) from the interface showed diffuse rings characteristic of an amorphous material. Nanoprobe EDS and X-ray mapping indicated that the interfacial layer was rich in Si, Mg, Al, and O. TEM of deformed material showed an increase in the overall dislocation density with increasing plastic strain. However, there was a zone of high
Fig. 2. (a) DA condition, TEM showing high dislocation density in grain adjacent to SiC; (b) DA condition, TEM showing crack initiating in the interfacial layer; (c) DA condition, TEM showing dislocations across an uncracked SiC particle; (d) DA condition, TEM showing a microcrack through a SiC particle.
62
R.D. Evans, J.D. Boyd / Scripta Materialia 49 (2003) 59–63
residual dislocation density extending 1–2 grain diameters from the particle interface (Fig. 2(a)). In addition to this increased dislocation density in matrix grains, the main damage mechanism was particle–matrix decohesion. In all cases, decohesion occurred by fracture through the thin interfacial layer. Fig. 2(b) shows an early stage of the ‘‘unzipping’’ mechanism where decohesion has initiated in the interfacial layer, and TEM images of decohered particles showed remnants of the interfacial layer on both surfaces. At low strains (<0.7%), there was no change in the internal structure of the SiC particles. At higher plastic strains, isolated rows of dislocations were observed in some SiC particles (Fig. 2(c)). The Burgers vector of these dislocations was determined to be b ¼ ½ 12 1 1 by a standard g b analysis. Microcracks were also observed in some SiC particles at plastic strains greater than 0.7% (Fig. 2(d)). Often, dislocations were observed at the ends of these microcracks.
4. Discussion The TEM observation that the maximum dislocation density occurs within one or two grains of the interface can be attributed to inhomogeneous plastic strain as the matrix material deforms around the rigid SiC particles. It is also likely that there is a higher than average dislocation density adjacent to the SiC particles in the STQ condition prior to deformation. The difference in the thermal expansion coefficient (CTE) of the two phases (approximately 7:1 for Al/SiC) causes dislocations to be ‘‘punched out’’ from the particle–matrix interface during cooling from the solution-treatment temperature [5,6]. The higher dislocation density in the near-interface region is likely a major contributing factor to the particle–matrix decohesion mechanism. Particle–matrix decohesion in the MMC studied appears to be associated with an amorphous Si– Mg–Al–O layer at the interface. This interface decohesion mechanism has not been reported previously for Al/SiC MMC systems. The natural layer of SiO2 ( 5 nm thick) on SiC particles is usually consumed by reaction with liquid Al when
Al/SiC MMCs are made by molten metal mixing or infiltration techniques. During the solid-state powder processing of the MMC in the present work, the SiO2 layer would likely survive. Furthermore, it is postulated that its thickness increases during the hot extrusion processing cycle by interdiffusion with the Al2 O3 surface layer on the Al 2080 powder. Ratnaparkhi and Howe [7] produced an Al/SiC composite by diffusion bonding surfaces and observed an amorphous interface layer having similar thickness and composition as in the present work. Similar, but thicker ( 200 nm) amorphous surface films are produced during dry sliding wear of Al/SiC [8]. Ratnaparkhi and Howe observed two distinct amorphous layers for thicknesses greater than 10 nm. The existence of two amorphous layers, each bonded strongly to either Al or SiC, would favour failure by tearing between them. While two different amorphous layers were not observed in the present work, interface decohesion occurs by fracture through the amorphous layer. Using a method based on an interfacial energy criterion [9]; the local critical stress for fracture of the interfacial layer is calculated to be 200–300 MPa. Prior to deformation, the only defects seen in the SiC particles are stacking faults. At a plastic strain of 0.7%, perfect dislocations are observed in the SiC particles and the stacking faults disappear. It appears that the partial dislocations forming the stacking faults in the SiC move together to form perfect dislocations at some critical applied stress. These dislocations are thought to be due to a combination of the applied stress and the residual microstresses present in the SiC particle after cooling. The difference in CTE of the two phases in the MMC has been shown to induce significant stresses in the SiC particles, and this combined with the applied stresses is likely sufficient to induce the formation of pure dislocations in the SiC. After the formation of the perfect dislocations, microcracks can form at the angular, stressconcentrating, sections of particles. After a critical strain level (approximately 0.7%), it is evident that some of the particles experience microcracking. In general, beyond 0.7% plastic strain, approximately 10% of the particles smaller than 5 lm have cracked.
R.D. Evans, J.D. Boyd / Scripta Materialia 49 (2003) 59–63
5. Conclusions Microstructural evolution in the near-interface region and within SiC particles has been observed during loading of a small-particle metal matrix composite. Particle–matrix decohesion was observed to occur through a 40 nm thick amorphous interfacial layer. The formation of this layer is considered to be due to diffusion of Al and Mg into the native SiO2 layer present on the surface of the SiC particles. The accumulation of perfect dislocations in the SiC leading to particle cracking was also observed, although particle cracking does not appear to be an important damage mechanism in this MMC.
Acknowledgements The authors thank Dr. Warren Hunt Jr., formerly of Alcoa, for supplying material, Drs. George Ruddle and Gianluigi Boton of
63
CANMET, for assistance with fabrication and characterization, respectively, and Mr. Mike Phaneuf of Fibics for provision of FIB facilities. The research was supported by the Natural Sciences and Engineering Research Council of Canada.
References [1] Lloyd DJ. Acta Metall Mater 1991;39:59. [2] Whitehouse AF, Clyne TW. Acta Metall Mater 1993;41:1701. [3] Ribes H, Da Silva R, Suery M, Bretheau T. Mater Sci Technol 1990;6:621. [4] Evans RD, Phaneuf MW, Boyd JD. J Microsc 1999;196:146. [5] Vogelsang M, Arsenault RJ, Fisher RM. Met Trans A 1986;17:379. [6] Barlow CY, Hansen N. Acta Metall Mater 1995;43:3633. [7] Ratnaparkhi PL, Howe JM. Acta Metall Mater 1994;42:81. [8] Li XY, Tandon KN. Acta Mater 1996;44:3611. [9] Sun J. Int J Frac 1990;44:R51.
Near-interface microstructure in a SiC/Al composite R.D. Evans *, J.D. Boyd Department of Materials and Metallurgical Engineering, QueenÕs University, Kingston, Canada ON K7L 3N6 Received 29 January 2003; received in revised form 27 March 2003; accepted 31 March 2003
Abstract The microstructure of a SiC particulate reinforced Al 2080 alloy has been characterized by focussed ion beam microscopy and transmission electron microscopy. A 40 nm thick amorphous interface layer develops during processing of the composite material by diffusion of Al and Mg into the pre-existing SiO2 layer on the 3 lm SiC particles. During tensile deformation, damage occurs in the near-interface zone of the matrix, at the particle–matrix interface and within the SiC particles. Ó 2003 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. Keywords: Composites; Aluminium; Powder processing; Microstructure; Damage
1. Introduction There has been considerable research on the effects of particle-induced damage in metal matrix composites (MMCs) [1–3]. Most of these studies have been on MMCs with particulate size >10 lm, where particle cracking is the dominant damage mechanism and the particle–matrix interface appears to have little effect on the overall damagefracture behaviour. In the current study, an MMC was produced in which interface decohesion was the dominant damage-initiation mechanism. The experimental design incorporated (a) small particulate size (3 lm) to suppress particle cracking, and (b)
* Corresponding author. Present address: Kingston Research and Development Centre, Alcan International Limited, 945 Princess Street, Kingston, Canada ON K7L 5L9. Tel: +1-613541-2137; fax: +1-613-541-2134. E-mail address: [email protected] (R.D. Evans).
powder processing to minimise interface reactions and produce a well-adhered, clean interface. This paper reports on microstructural evolution of the near-interface zone during deformation, and identification of an interface decohesion mechanism. Quantitative measurements of damage evolution and comparisons with existing damage models are reported elsewhere [4].
2. Experimental materials and methods The MMC material was 20 vol% SiC in Al 2080 (Al–4Cu–2Mg). The mean diameter of the SiC was 3.3 0.5 lm. The material was produced by a powder-processing route that included mixing, cold compaction, and hot extrusion at 500 °C to 16 mm diameter rod. Complete processing details are given in Ref. [4]. Tensile samples having an 18 mm 3 mm diam. gauge section were machined from the extruded rod, and solution treated for 2 h
1359-6462/03/$ - see front matter Ó 2003 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. doi:10.1016/S1359-6462(03)00180-5
60
R.D. Evans, J.D. Boyd / Scripta Materialia 49 (2003) 59–63
at 510 °C followed by water quenching (STQ condition). To study the evolution of damage and the nearinterface microstructure during deformation, tensile tests were carried out on an Instron 8502 servo-hydraulic testing machine at a strain rate of 104 s1 . Samples were deformed to various strains up to the fracture point, which occurred at about 2% plastic strain. Following deformation, all samples were aged for 24 h at 177 °C (DA condition). This treatment was employed to decorate the matrix dislocations with precipitates, and to represent a typical heat-treated condition for such an MMC material. The microstructures of the MMC prior to and after deformation were characterized by focussed ion beam (FIB) microscopy and transmission electron microscopy (TEM). Sections for FIB examination were cut parallel to the tensile direction. Final surface preparation by Gaþ ion milling and secondary electron imaging was carried out using a Micrion 2500 FIB system operating at 50 keV accelerating voltage. More details on this technique are given in Ref. [4]. TEM samples were prepared using conventional sample preparation techniques. Thin discs were produced and dimpled to 10 lm centre thickness on a Fischione Model 2000 dimpler. Final thinning to produce electron transparent samples was per-
formed on a Gatan Model 600 dual ion mill. Specimens were examined on a Philips CM20 TEM at an accelerating voltage of 200 kV, and X-ray microanalysis was carried out using a Noran Voyager EDS system with light-element (>B) capability. X-ray mapping of the particulate– matrix interface was carried out using a Philips CM20-FEG TEM equipped with an Oxford Instruments EDS system.
3. Results The microstructure of the MMC in the STQ condition is illustrated by the FIB micrograph in Fig. 1(a). Ion channelling produces orientation contrast of the matrix grains, and there is also atomic-number contrast of the SiC particulate, as well as other phases. Microstructural features indicated in Fig. 1(a) are SiC particles (S), dispersoid particles (D) and interface layer (I). TEM of the STQ material revealed a relatively low density of quenched-in dislocation loops and helices in the matrix grains, as well as large (100– 500 nm) particles within matrix grains and at the interfaces. Most of the particles were determined to be Cu-, Zr-, Fe-, or Mn-rich and were assumed to be dispersoid phases which were not dissolved during the solution treatment at 510 °C. The
Fig. 1. (a) STQ condition, FIB micrograph: S––SiC, D––dispersoid, I––interface layer; (b) STQ condition, TEM showing interface layer.
R.D. Evans, J.D. Boyd / Scripta Materialia 49 (2003) 59–63
remaining particles were identified as Al2 O3 , which likely resulted from the oxide layer on the initial Al 2080 powder. The SiC particles had hexagonal crystal structure and contained stacking faults characteristic of the 6H SiC polytype. There was also evidence of a thin layer at the interface between SiC particles and matrix grains, which appeared as a white line in FIB images (Fig. 1(a) at I). In TEM, this layer was clearly visible as a phase 20–40 nm in thickness, having a charac-
61
teristic mottled appearance (Fig. 1(b)). Selected area diffraction (SAD), and convergent beam diffraction (CBED) from the interface showed diffuse rings characteristic of an amorphous material. Nanoprobe EDS and X-ray mapping indicated that the interfacial layer was rich in Si, Mg, Al, and O. TEM of deformed material showed an increase in the overall dislocation density with increasing plastic strain. However, there was a zone of high
Fig. 2. (a) DA condition, TEM showing high dislocation density in grain adjacent to SiC; (b) DA condition, TEM showing crack initiating in the interfacial layer; (c) DA condition, TEM showing dislocations across an uncracked SiC particle; (d) DA condition, TEM showing a microcrack through a SiC particle.
62
R.D. Evans, J.D. Boyd / Scripta Materialia 49 (2003) 59–63
residual dislocation density extending 1–2 grain diameters from the particle interface (Fig. 2(a)). In addition to this increased dislocation density in matrix grains, the main damage mechanism was particle–matrix decohesion. In all cases, decohesion occurred by fracture through the thin interfacial layer. Fig. 2(b) shows an early stage of the ‘‘unzipping’’ mechanism where decohesion has initiated in the interfacial layer, and TEM images of decohered particles showed remnants of the interfacial layer on both surfaces. At low strains (<0.7%), there was no change in the internal structure of the SiC particles. At higher plastic strains, isolated rows of dislocations were observed in some SiC particles (Fig. 2(c)). The Burgers vector of these dislocations was determined to be b ¼ ½ 12 1 1 by a standard g b analysis. Microcracks were also observed in some SiC particles at plastic strains greater than 0.7% (Fig. 2(d)). Often, dislocations were observed at the ends of these microcracks.
4. Discussion The TEM observation that the maximum dislocation density occurs within one or two grains of the interface can be attributed to inhomogeneous plastic strain as the matrix material deforms around the rigid SiC particles. It is also likely that there is a higher than average dislocation density adjacent to the SiC particles in the STQ condition prior to deformation. The difference in the thermal expansion coefficient (CTE) of the two phases (approximately 7:1 for Al/SiC) causes dislocations to be ‘‘punched out’’ from the particle–matrix interface during cooling from the solution-treatment temperature [5,6]. The higher dislocation density in the near-interface region is likely a major contributing factor to the particle–matrix decohesion mechanism. Particle–matrix decohesion in the MMC studied appears to be associated with an amorphous Si– Mg–Al–O layer at the interface. This interface decohesion mechanism has not been reported previously for Al/SiC MMC systems. The natural layer of SiO2 ( 5 nm thick) on SiC particles is usually consumed by reaction with liquid Al when
Al/SiC MMCs are made by molten metal mixing or infiltration techniques. During the solid-state powder processing of the MMC in the present work, the SiO2 layer would likely survive. Furthermore, it is postulated that its thickness increases during the hot extrusion processing cycle by interdiffusion with the Al2 O3 surface layer on the Al 2080 powder. Ratnaparkhi and Howe [7] produced an Al/SiC composite by diffusion bonding surfaces and observed an amorphous interface layer having similar thickness and composition as in the present work. Similar, but thicker ( 200 nm) amorphous surface films are produced during dry sliding wear of Al/SiC [8]. Ratnaparkhi and Howe observed two distinct amorphous layers for thicknesses greater than 10 nm. The existence of two amorphous layers, each bonded strongly to either Al or SiC, would favour failure by tearing between them. While two different amorphous layers were not observed in the present work, interface decohesion occurs by fracture through the amorphous layer. Using a method based on an interfacial energy criterion [9]; the local critical stress for fracture of the interfacial layer is calculated to be 200–300 MPa. Prior to deformation, the only defects seen in the SiC particles are stacking faults. At a plastic strain of 0.7%, perfect dislocations are observed in the SiC particles and the stacking faults disappear. It appears that the partial dislocations forming the stacking faults in the SiC move together to form perfect dislocations at some critical applied stress. These dislocations are thought to be due to a combination of the applied stress and the residual microstresses present in the SiC particle after cooling. The difference in CTE of the two phases in the MMC has been shown to induce significant stresses in the SiC particles, and this combined with the applied stresses is likely sufficient to induce the formation of pure dislocations in the SiC. After the formation of the perfect dislocations, microcracks can form at the angular, stressconcentrating, sections of particles. After a critical strain level (approximately 0.7%), it is evident that some of the particles experience microcracking. In general, beyond 0.7% plastic strain, approximately 10% of the particles smaller than 5 lm have cracked.
R.D. Evans, J.D. Boyd / Scripta Materialia 49 (2003) 59–63
5. Conclusions Microstructural evolution in the near-interface region and within SiC particles has been observed during loading of a small-particle metal matrix composite. Particle–matrix decohesion was observed to occur through a 40 nm thick amorphous interfacial layer. The formation of this layer is considered to be due to diffusion of Al and Mg into the native SiO2 layer present on the surface of the SiC particles. The accumulation of perfect dislocations in the SiC leading to particle cracking was also observed, although particle cracking does not appear to be an important damage mechanism in this MMC.
Acknowledgements The authors thank Dr. Warren Hunt Jr., formerly of Alcoa, for supplying material, Drs. George Ruddle and Gianluigi Boton of
63
CANMET, for assistance with fabrication and characterization, respectively, and Mr. Mike Phaneuf of Fibics for provision of FIB facilities. The research was supported by the Natural Sciences and Engineering Research Council of Canada.
References [1] Lloyd DJ. Acta Metall Mater 1991;39:59. [2] Whitehouse AF, Clyne TW. Acta Metall Mater 1993;41:1701. [3] Ribes H, Da Silva R, Suery M, Bretheau T. Mater Sci Technol 1990;6:621. [4] Evans RD, Phaneuf MW, Boyd JD. J Microsc 1999;196:146. [5] Vogelsang M, Arsenault RJ, Fisher RM. Met Trans A 1986;17:379. [6] Barlow CY, Hansen N. Acta Metall Mater 1995;43:3633. [7] Ratnaparkhi PL, Howe JM. Acta Metall Mater 1994;42:81. [8] Li XY, Tandon KN. Acta Mater 1996;44:3611. [9] Sun J. Int J Frac 1990;44:R51.