- Email: [email protected]
Al composite materials
Materials Science and Engineering A282 (2000) 67 – 73 www.elsevier.com/locate/msea
Temperature induced worksoftening of Be/Al composite materials S.G. Song *, T.J. Garosshen, V.C. Nardone United Technologies Research Center, East Hartford, CT 06108, USA Received 1 July 1999; received in revised form 17 November 1999
Abstract Temperature introduced worksoftening has been observed in Be/Al composites with 70% volume fraction of Be reinforcement. The composites exhibited a normal workhardening behavior at room temperature but flow softening occurred at 232°C and above. Fractographic examinations showed that extensive matrix plastic flow and interface sliding occurred at elevated temperatures, which led to local rearrangement of beryllium particles. The rearrangement resulted in a reduction in flow stress of the composites. © 2000 Elsevier Science S.A. All rights reserved. Keywords: Temperature induced worksoftening; Be/Al composites; Fractographic examinations
1. Introduction Worksoftening has been observed in numerous metallic materials of both monolithic and composite form. Mechanisms of softening in the former are wellunderstood [1], with dynamic recovery and grain boundary sliding being the most common causes of worksoftening. For composite materials in general terms, the mechanisms of work softening vary widely from phase redistribution to spheroidization of second phase due to their multi-facet structural properties [1– 3]. Limited observations of worksoftening in discontinuously reinforced metal matrix composites (DMMCs) have been reported and all of them were found in low reinforcement-volume fraction materials (B 20%), where the plastic deformation is more or less dictated by the matrix [4,5] The present investigation reports worksoftening observed in Be/Al composites. Since beryllium is completely immiscible in aluminum, the alloy is virtually a composite material with beryllium phase (in particulate form) embedded in the aluminum matrix. The volume fraction of the beryllium phase is high, : 70%, so that neither the matrix nor the reinforcement dominates the plastic behavior of the composites. The beryllium rein* Corresponding author. Tel.: +1-860-610-7075; fax: + 1-860-6101697. E-mail address: [email protected] (S.G. Song)
forcement phase has some ductility and may participate in the plastic deformation of the composites. The interaction of the two phases in response to an external stress plays a key role in controlling the plastic properties of the composites.
2. Experimental Two Albemet alloys, i.e. 162 and 562, were received from Brush and Wellman for the present investigation. These alloys have nominally similar Be (62 wt.%) content, as indicated by the last two digits of the alloy designation, but with different matrix compositions. Alloy 162 has a pure aluminum matrix (1XXX series) and alloy 562 is of an Al–Mg alloy matrix (5XXX series). The Mg addition was made to strengthen the relatively weak aluminum matrix. The nominal compositions of the two alloys are listed in Table 1. Alloy 162 was received in the form of f13.5 cm extruded billets and alloy 562 was of 1.27×11.4-cm cross-section extruded plates. The reinforcement (Be) particle size of the two alloys differs due to different extrusion conditions and powder precursors, which are shown in Fig. 1(a) and (b), respectively. Tensile testing was conducted, pursuant to ASTM E-8 and E-21, at room temperature, 232, and 371°C in longitudinal (extrusion) direction. Because of the difference in the form of the as-received materials, different
0921-5093/00/$ - see front matter © 2000 Elsevier Science S.A. All rights reserved. PII: S 0 9 2 1 - 5 0 9 3 ( 9 9 ) 0 0 7 7 1 - 6
S.G. Song et al. / Materials Science and Engineering A282 (2000) 67–73
68
specimen sizes were adopted for tensile test. The tensile specimens had a nominal 0.64 cm gage diameter by 3.8 cm gage length for alloy 562 and 0.45 cm by 2 cm for alloy 162, respectively, both having threaded end attachment regions. The cross-head displacement rate is Table 1 Nominal composition of AlBemet 162 and 562 (wt.%) Alloy
Al
Mg
Be
Reinforcement (Be) volume fraction
162 562
Balance Balance
3
62 62
70 70
Fig. 2. Stress/strain curves of alloys 562 (a) and 162 (b) at varying temperatures. Worksoftening is evident in alloy 562 at elevated temperatures.
also slightly different at 0.13 and 0.1 cm min − 1, respectively. The fracture surfaces of the broken specimens was examined by scanning electron microscope (SEM).
3. Results
Fig. 1. Microstructures of alloy 162 (a) and 562 (b). The average beryllium particle size is :5 mm for alloy 162 and : 2 mm for alloy 562.
Fig. 2 displays the stress–strain curves of the two alloys tested in tension at different temperatures. Alloy 562 exhibited strain hardening at room temperature, see Fig. 2(a). Flow softening occurs at elevated temperatures (232 and 371°C). A transition from strain hardening to softening is expected to occur somewhere between ambient temperature and 232°C. Like alloy 562, alloy 162 strain hardens at ambient temperature with, however, a higher work hardening rate, see Fig. 2(b). At 232°C, alloy 162 exhibited initially a small workhardening rate and leveled off in a later stage, apparently resulting from the competing hardening and softening mechanisms that were very much in balance at that temperature.
S.G. Song et al. / Materials Science and Engineering A282 (2000) 67–73
69
SEM examinations of the tensile specimens of alloy 562 are shown in Figs. 3 – 5. The fracture surface of the room temperature samples appears to be flat in general, see Fig. 3(a). Fig. 3(b) shows substantial evidence of fractured beryllium particles on the fracture surface. The matrix phase exhibits dimples of dimensions commensurate with the beryllium particles. The planar fracture surface indicates that the crack front propagates in accordance to the maximum tensile stress and is less dependent on phases encountered in its path. At elevated temperatures, the fracture produced a rugged surface that did not follow the maximum stress direction, see Fig. 4(a) and (b). Sharp terrace and deep cavities are common features of the surface. A close examination of the fracture surface showed that extensive interfacial debonding had occurred at 232°C, see Fig. 5(a). Fractured particles are mostly of large sizes. At a higher temperature (371°C), the area fraction of the matrix on the fracture surface was considerably larger than its volume fraction (30%), see Fig. 5(b),
Fig. 4. Fractography of alloy 562 at 232°C (a) and 371°C (b). The rugged fracture surface is filled with bulges and cavities.
Fig. 3. Fractography of alloy 562 at room temperature, (a) flat fracture surface at a low magnification, and (b) particle fracture being the predominant fracture mode.
indicative of a preferred fracture path through the matrix. The dimple size of the matrix is smaller than that of the particle spacing, suggesting that an extensive plastic deformation had occurred in the matrix. Very few beryllium particles are seen fractured on the fractographic surface, see Fig. 5(b). Figs. 6 and 7 show the fracture surface of the alloy 162 at ambient temperature and 232°C. Like alloy 562, the room temperature fracture surface appears to be flat and it becomes rather uneven at elevated temperatures. The variation of dimple size with temperature is consistent with the observation in alloy 562, except that it is proportionally larger in accord with the greater beryllium particle size. Dislocation structure of a tensile specimen tested at room temperature is shown in Fig. 8. A high density of dislocations is seen in the matrix, see Fig. 8(a). The dislocations form cellular subgrain structures as in a monolithic matrix alloy. Few grain boundaries, however, were seen in the matrix, indicating that the grain size of the matrix is equal to or larger than the particle
70
S.G. Song et al. / Materials Science and Engineering A282 (2000) 67–73
spacing. A moderate dislocation density is also seen in the Be phase, where straight subgrain boundaries (dislocation walls) are evident, see Fig. 8(b). There are indications that local plastic deformation penetrated the interface, as indicated by the dislocation configuration in the Be phase, where dislocation pileups appeared to be initiated at the interface. Planar slip appears to be the predominant deformation mode in the reinforcement, as suggested by the straight sub-structure and dislocation arrays, since pyramidal slip in Be is rare and cross slip between basal and prism planes are unlikely [6] (unpublished research).
4. Discussion Owing to the high volume fraction of reinforcement in Be/Al composites, the plastic behavior of the composite is not fully controlled by the matrix. The observed worksoftening is, therefore, expected to result from the interaction of the two phases in response to
Fig. 6. Fractography of alloy 162 at room temperature, (a) flat fracture surfaces at a low magnification, and (b) particle fracture being the predominant fracture mode.
external loads. Fractographic examinations of the room temperature specimens showed that the surface portion of the fractured beryllium particles is proportional to the volume fraction of the phase. There is no apparent preferential path for crack propagation in the composite. Because of the difference in yield stress between the two phases, local yielding always begins in the matrix. The local stress buildup at the matrix/reinforcement interface by dislocation pileups in the matrix is sufficient to cause yielding of the beryllium phase at ambient temperature, as evidenced by the observed dislocation structure in the be phase (Fig. 8(b)). Unlike reinforcement phases in ceramic particulate reinforced MMCs, where the particles fracture in a brittle fashion, the beryllium phase is relatively ductile1 and can deform plastically. Dislocations will penetrate the interface when the matrix stress exceeds the yield stress of Fig. 5. A mixture of particle fracture and matrix tearing is seen on the fracture surface of alloy 562 at 232 °C (a), whereas the matrix plastic tearing is the dominant fracture mode at 371°C (b).
1 Extruded pure beryllium can deform plastically with up to 3% elongation [7].
S.G. Song et al. / Materials Science and Engineering A282 (2000) 67–73
the reinforcement, leading to plastic deformation of the beryllium phase. The workhardening behavior of the composite is therefore dictated by the beryllium reinforcement. In addition, the plastic deformation of the reinforcement also contributes to the overall plastic deformation of the composite. This result is consistent with predictions by a model proposed for high-reinforcement-volume fraction DMMCs (unpublished research). At elevated temperatures, the yield strength of the matrix declines faster than that of the reinforcement and the difference in yield strength increases with increasing temperature. As a result, the regular stress buildup by a dislocation pileup in the matrix may never reach a level comparable to the yield stress of the beryllium. The beryllium phase therefore does not fully participate in the plastic deformation of the composite, and the global yielding of the composite results from the matrix plastic flow over the scale of reinforcement particle dimension. This heterogeneous deformation
Fig. 7. Fractography of alloy 162 at 232°C, (a) rugged surface at a low magnification, and (b) matrix plastic tearing being the predominant fracture mode.
71
can lead to a very high plastic strain of the matrix and extensive interfacial debonding. The composite fails by matrix fracture plus interfacial debonding at 232°C for alloy 562. At 371°C, matrix plastic deformation appeared to be the predominant cause of fracture, as seen in Fig. 5(b). Alloy 162 showed similar matrix fracture at 232°C, since its matrix is not strengthened. The large plastic strain in the matrix may not transfer to significant plastic deformation of the composites due to its low volume fraction and thin-walled geometry. On the other hand, the extensive plastic flow in the matrix can bring about rearrangements of the particles (or particle clusters) in the composites. The particle rearrangement will lead to reduced resistance to the plastic deformation of the composites, resulting in a reduced flow stress. The reinforcement particles undergo minimum, if any, plastic deformation at elevated temperatures. The observed softening is expected to be inversely proportional to the particle (or grain) size, which is consistent with the present results. Dynamic recrystallization has been observed in highly strained Al–Mg alloys [8], TEM examination, however, did not indicate that the present softening behavior is due to dynamic recrystallization. The grain size of the plastically strained matrix is larger than its free path ( :1 mm) between reinforcement particles and recrystallization due to grain boundary migration is surpressed. Besides the matrix plastic flow, the matrix/particle interface sliding and debonding can also play a part in the observed softening. The sliding did not occur at room temperature as suggested by the fractographic examinations, of which the majority of the beryllium particles seen on the fracture surface were fractured. On the other hand, significant interfacial debonding was observed in alloy 562 tested at 232°C, indicating interface sliding prior to debonding. The bonding strength decreased with increasing temperature, partly because of a reduction in matrix squeezing force on the particles due to the disparity of coefficient of thermal expansion, and was comparable to the tearing strength of the matrix at 232°C. As temperature further increases, the difference in strength between the two phases becomes greater such that the matrix plastic flow becomes the primary driving force for particle rearrangement. This is consistent with the observation that little debonding is seen on the fracture surface of alloy 562 at 371°C, see Fig. 5(b). The fractographic observations are in agreement with the analysis that at elevated temperatures particle rearrangement had occurred as indicated by the bulges and cavities seen on the fracture surfaces. It should be noted that the rearrangement of particles is limited in Be/Al composites compared to grain switching seen in superplastic deformation because of the distinction between matrix flow, interfacial sliding and grain boundary slid-
72
S.G. Song et al. / Materials Science and Engineering A282 (2000) 67–73
Fig. 8. Bright filed image of room temperature deformation structure of AlBemet 162, (a) well-developed sub-structure in the matrix and planar slip in the reinforcement phase, (b) dislocation arrays in the reinforcement phase.
ing. The worksoftening of the composite is therefore expected to be less pronounced than that seen in superplasticity. The particle size of the alloys affects not only the tensile strength, consequent to Hall – Petch effect, but also the strain–softening behavior of the composites. The particle size of alloy 562 is significantly smaller than that of alloy 162. Accordingly, the softening behavior of the former is more conspicuous than the latter. Work hardening and softening are nearly balanced in alloy 162 at 232°C. It is of interest to note that the ductility of alloy 562 did not increase with increasing temperature as the matrix alloy became softer because the beryllium phase did not contribute to the plastic deformation at elevated temperatures. The stress buildup in the matrix was insufficient to cause yielding of the beryllium due to increased difference in strength between the two phases. Both composites were extruded at high temperatures and heat-treated afterward, matrix recrystallization is therefore expected to have occurred. The matrix grain size of alloy 162 should be greater than alloy 562 after recrystallization due to its larger beryllium particle size. The dislocation pileup in the matrix of alloy 162 is, therefore, expected to be more pronounced than that in alloy 562, which will lead to a lower yield stress of the former. Consequently, alloy 162 experienced lower interface shear stress at elevated temperatures than alloy 562. The higher interface stress in alloy 562 increases the probability of interface sliding so that both matrix plastic flow and interface sliding contribute to worksoftening. The difference in worksoftening between the two materials is reflected in fractographic characteristics, as seen earlier, that alloy 162 shows extensive matrix plastic strain with little sign of interfacial debonding whereas
alloy 562 exhibits both matrix plastic deformation and interfacial debonding. Although the two composites have distinct matrixes, the difference in the observed softening behavior is believed to be primarily a result of particle size effect, not the matrix property. Similarly, the large difference (\ 200 MPa) in yield strength between the two alloys cannot be accounted for by the strength variation of the matrix. However, the particle size strengthening effect cannot be calculated directly using the Hall–Petch relationship as for monolithic materials, because the dislocation initiation in the reinforcement phase by the matrix dislocation pileup at the interface is not only a function of the matrix dimension but also of the reinforcement particle size.
5. Conclusions The following conclusions can be drawn from the present investigation. Worksoftening behavior has been observed in Be/Al composites at elevated temperatures. Alloy 562 exhibited a pronounced softening behavior after yielding while alloy 162 workhardened slightly in the early stage of deformation and its stress/strain curve leveled off later. Worksoftening is always associated with uneven fracture surface exhibiting bulges and cavities. Extensive plastic deformation and interfacial debonding were commonly observed at elevated temperatures on the fracture surfaces. At room temperature however, the fracture surface is flat and the majority of the particles on the surface are fractured.
S.G. Song et al. / Materials Science and Engineering A282 (2000) 67–73
Reinforcement particle size plays an important part in the observed worksoftening. The smaller the particle size is, the more pronounced the softening behavior is seen. Particle rearrangement and interfacial sliding are the primary causes for the observed worksoftening of the Be/Al composites. Increasing temperature and decreasing particle size intensify worksoftening behavior.
Acknowledgements This work was supported by the Office of Naval Research (ONR/NW) of Department of Defense (DOD) under contract number N00014-94-C-0098/P00007 with the United Technologies Research Center (UTRC). The author would like to acknowledge J.T. Beals of Pratt and Whitney for his work on the mechanical testing of Be/Al.
.
73
References [1] J.J. Jonas, M.J. Luton, Advances in Deformation Processing, vol. 21, Plenum, New York, 1978, pp. 215 – 243. [2] E.A.McG. Chojnowski, W.J. Tegart, J. Mater. Sci. 2 (1968) 14 – 18. [3] N. Kanetake, N. Nakamura, J. Jpn. Soc. Technol. Plast. 31 (1990) 536 – 541. [4] A. Inoue, H.M. Kimura, M. Watanabe, A. Kawabata, Mater. Trans. JIM 38 (1997) 756 – 760. [5] J. Zhao, X. Xia, H.J. McQueen, Hot Workability of Steels and Light Alloys-Composites, Canadian Institute of Mining, Metallurgy and Petroleum, 1996, pp. 151 – 157. [6] J.F. Breedis, A. Lawley, J.A. Zeiger, in: L.M. Schetky, H.A. Johnson (Eds.), Beryllium Technology, vol. 1, Gordon and Breach, New York, 1966, pp. 197 – 226. [7] Designing with Beryllium, Product Brochure, Brush & Wellman Inc., Cleveland OH. [8] W. Blum, Q. Zhu, R. Merkel, H.J. McQueen, Z. Metallkd. 87 (1996) 14 – 23.
Temperature induced worksoftening of Be/Al composite materials S.G. Song *, T.J. Garosshen, V.C. Nardone United Technologies Research Center, East Hartford, CT 06108, USA Received 1 July 1999; received in revised form 17 November 1999
Abstract Temperature introduced worksoftening has been observed in Be/Al composites with 70% volume fraction of Be reinforcement. The composites exhibited a normal workhardening behavior at room temperature but flow softening occurred at 232°C and above. Fractographic examinations showed that extensive matrix plastic flow and interface sliding occurred at elevated temperatures, which led to local rearrangement of beryllium particles. The rearrangement resulted in a reduction in flow stress of the composites. © 2000 Elsevier Science S.A. All rights reserved. Keywords: Temperature induced worksoftening; Be/Al composites; Fractographic examinations
1. Introduction Worksoftening has been observed in numerous metallic materials of both monolithic and composite form. Mechanisms of softening in the former are wellunderstood [1], with dynamic recovery and grain boundary sliding being the most common causes of worksoftening. For composite materials in general terms, the mechanisms of work softening vary widely from phase redistribution to spheroidization of second phase due to their multi-facet structural properties [1– 3]. Limited observations of worksoftening in discontinuously reinforced metal matrix composites (DMMCs) have been reported and all of them were found in low reinforcement-volume fraction materials (B 20%), where the plastic deformation is more or less dictated by the matrix [4,5] The present investigation reports worksoftening observed in Be/Al composites. Since beryllium is completely immiscible in aluminum, the alloy is virtually a composite material with beryllium phase (in particulate form) embedded in the aluminum matrix. The volume fraction of the beryllium phase is high, : 70%, so that neither the matrix nor the reinforcement dominates the plastic behavior of the composites. The beryllium rein* Corresponding author. Tel.: +1-860-610-7075; fax: + 1-860-6101697. E-mail address: [email protected] (S.G. Song)
forcement phase has some ductility and may participate in the plastic deformation of the composites. The interaction of the two phases in response to an external stress plays a key role in controlling the plastic properties of the composites.
2. Experimental Two Albemet alloys, i.e. 162 and 562, were received from Brush and Wellman for the present investigation. These alloys have nominally similar Be (62 wt.%) content, as indicated by the last two digits of the alloy designation, but with different matrix compositions. Alloy 162 has a pure aluminum matrix (1XXX series) and alloy 562 is of an Al–Mg alloy matrix (5XXX series). The Mg addition was made to strengthen the relatively weak aluminum matrix. The nominal compositions of the two alloys are listed in Table 1. Alloy 162 was received in the form of f13.5 cm extruded billets and alloy 562 was of 1.27×11.4-cm cross-section extruded plates. The reinforcement (Be) particle size of the two alloys differs due to different extrusion conditions and powder precursors, which are shown in Fig. 1(a) and (b), respectively. Tensile testing was conducted, pursuant to ASTM E-8 and E-21, at room temperature, 232, and 371°C in longitudinal (extrusion) direction. Because of the difference in the form of the as-received materials, different
0921-5093/00/$ - see front matter © 2000 Elsevier Science S.A. All rights reserved. PII: S 0 9 2 1 - 5 0 9 3 ( 9 9 ) 0 0 7 7 1 - 6
S.G. Song et al. / Materials Science and Engineering A282 (2000) 67–73
68
specimen sizes were adopted for tensile test. The tensile specimens had a nominal 0.64 cm gage diameter by 3.8 cm gage length for alloy 562 and 0.45 cm by 2 cm for alloy 162, respectively, both having threaded end attachment regions. The cross-head displacement rate is Table 1 Nominal composition of AlBemet 162 and 562 (wt.%) Alloy
Al
Mg
Be
Reinforcement (Be) volume fraction
162 562
Balance Balance
3
62 62
70 70
Fig. 2. Stress/strain curves of alloys 562 (a) and 162 (b) at varying temperatures. Worksoftening is evident in alloy 562 at elevated temperatures.
also slightly different at 0.13 and 0.1 cm min − 1, respectively. The fracture surfaces of the broken specimens was examined by scanning electron microscope (SEM).
3. Results
Fig. 1. Microstructures of alloy 162 (a) and 562 (b). The average beryllium particle size is :5 mm for alloy 162 and : 2 mm for alloy 562.
Fig. 2 displays the stress–strain curves of the two alloys tested in tension at different temperatures. Alloy 562 exhibited strain hardening at room temperature, see Fig. 2(a). Flow softening occurs at elevated temperatures (232 and 371°C). A transition from strain hardening to softening is expected to occur somewhere between ambient temperature and 232°C. Like alloy 562, alloy 162 strain hardens at ambient temperature with, however, a higher work hardening rate, see Fig. 2(b). At 232°C, alloy 162 exhibited initially a small workhardening rate and leveled off in a later stage, apparently resulting from the competing hardening and softening mechanisms that were very much in balance at that temperature.
S.G. Song et al. / Materials Science and Engineering A282 (2000) 67–73
69
SEM examinations of the tensile specimens of alloy 562 are shown in Figs. 3 – 5. The fracture surface of the room temperature samples appears to be flat in general, see Fig. 3(a). Fig. 3(b) shows substantial evidence of fractured beryllium particles on the fracture surface. The matrix phase exhibits dimples of dimensions commensurate with the beryllium particles. The planar fracture surface indicates that the crack front propagates in accordance to the maximum tensile stress and is less dependent on phases encountered in its path. At elevated temperatures, the fracture produced a rugged surface that did not follow the maximum stress direction, see Fig. 4(a) and (b). Sharp terrace and deep cavities are common features of the surface. A close examination of the fracture surface showed that extensive interfacial debonding had occurred at 232°C, see Fig. 5(a). Fractured particles are mostly of large sizes. At a higher temperature (371°C), the area fraction of the matrix on the fracture surface was considerably larger than its volume fraction (30%), see Fig. 5(b),
Fig. 4. Fractography of alloy 562 at 232°C (a) and 371°C (b). The rugged fracture surface is filled with bulges and cavities.
Fig. 3. Fractography of alloy 562 at room temperature, (a) flat fracture surface at a low magnification, and (b) particle fracture being the predominant fracture mode.
indicative of a preferred fracture path through the matrix. The dimple size of the matrix is smaller than that of the particle spacing, suggesting that an extensive plastic deformation had occurred in the matrix. Very few beryllium particles are seen fractured on the fractographic surface, see Fig. 5(b). Figs. 6 and 7 show the fracture surface of the alloy 162 at ambient temperature and 232°C. Like alloy 562, the room temperature fracture surface appears to be flat and it becomes rather uneven at elevated temperatures. The variation of dimple size with temperature is consistent with the observation in alloy 562, except that it is proportionally larger in accord with the greater beryllium particle size. Dislocation structure of a tensile specimen tested at room temperature is shown in Fig. 8. A high density of dislocations is seen in the matrix, see Fig. 8(a). The dislocations form cellular subgrain structures as in a monolithic matrix alloy. Few grain boundaries, however, were seen in the matrix, indicating that the grain size of the matrix is equal to or larger than the particle
70
S.G. Song et al. / Materials Science and Engineering A282 (2000) 67–73
spacing. A moderate dislocation density is also seen in the Be phase, where straight subgrain boundaries (dislocation walls) are evident, see Fig. 8(b). There are indications that local plastic deformation penetrated the interface, as indicated by the dislocation configuration in the Be phase, where dislocation pileups appeared to be initiated at the interface. Planar slip appears to be the predominant deformation mode in the reinforcement, as suggested by the straight sub-structure and dislocation arrays, since pyramidal slip in Be is rare and cross slip between basal and prism planes are unlikely [6] (unpublished research).
4. Discussion Owing to the high volume fraction of reinforcement in Be/Al composites, the plastic behavior of the composite is not fully controlled by the matrix. The observed worksoftening is, therefore, expected to result from the interaction of the two phases in response to
Fig. 6. Fractography of alloy 162 at room temperature, (a) flat fracture surfaces at a low magnification, and (b) particle fracture being the predominant fracture mode.
external loads. Fractographic examinations of the room temperature specimens showed that the surface portion of the fractured beryllium particles is proportional to the volume fraction of the phase. There is no apparent preferential path for crack propagation in the composite. Because of the difference in yield stress between the two phases, local yielding always begins in the matrix. The local stress buildup at the matrix/reinforcement interface by dislocation pileups in the matrix is sufficient to cause yielding of the beryllium phase at ambient temperature, as evidenced by the observed dislocation structure in the be phase (Fig. 8(b)). Unlike reinforcement phases in ceramic particulate reinforced MMCs, where the particles fracture in a brittle fashion, the beryllium phase is relatively ductile1 and can deform plastically. Dislocations will penetrate the interface when the matrix stress exceeds the yield stress of Fig. 5. A mixture of particle fracture and matrix tearing is seen on the fracture surface of alloy 562 at 232 °C (a), whereas the matrix plastic tearing is the dominant fracture mode at 371°C (b).
1 Extruded pure beryllium can deform plastically with up to 3% elongation [7].
S.G. Song et al. / Materials Science and Engineering A282 (2000) 67–73
the reinforcement, leading to plastic deformation of the beryllium phase. The workhardening behavior of the composite is therefore dictated by the beryllium reinforcement. In addition, the plastic deformation of the reinforcement also contributes to the overall plastic deformation of the composite. This result is consistent with predictions by a model proposed for high-reinforcement-volume fraction DMMCs (unpublished research). At elevated temperatures, the yield strength of the matrix declines faster than that of the reinforcement and the difference in yield strength increases with increasing temperature. As a result, the regular stress buildup by a dislocation pileup in the matrix may never reach a level comparable to the yield stress of the beryllium. The beryllium phase therefore does not fully participate in the plastic deformation of the composite, and the global yielding of the composite results from the matrix plastic flow over the scale of reinforcement particle dimension. This heterogeneous deformation
Fig. 7. Fractography of alloy 162 at 232°C, (a) rugged surface at a low magnification, and (b) matrix plastic tearing being the predominant fracture mode.
71
can lead to a very high plastic strain of the matrix and extensive interfacial debonding. The composite fails by matrix fracture plus interfacial debonding at 232°C for alloy 562. At 371°C, matrix plastic deformation appeared to be the predominant cause of fracture, as seen in Fig. 5(b). Alloy 162 showed similar matrix fracture at 232°C, since its matrix is not strengthened. The large plastic strain in the matrix may not transfer to significant plastic deformation of the composites due to its low volume fraction and thin-walled geometry. On the other hand, the extensive plastic flow in the matrix can bring about rearrangements of the particles (or particle clusters) in the composites. The particle rearrangement will lead to reduced resistance to the plastic deformation of the composites, resulting in a reduced flow stress. The reinforcement particles undergo minimum, if any, plastic deformation at elevated temperatures. The observed softening is expected to be inversely proportional to the particle (or grain) size, which is consistent with the present results. Dynamic recrystallization has been observed in highly strained Al–Mg alloys [8], TEM examination, however, did not indicate that the present softening behavior is due to dynamic recrystallization. The grain size of the plastically strained matrix is larger than its free path ( :1 mm) between reinforcement particles and recrystallization due to grain boundary migration is surpressed. Besides the matrix plastic flow, the matrix/particle interface sliding and debonding can also play a part in the observed softening. The sliding did not occur at room temperature as suggested by the fractographic examinations, of which the majority of the beryllium particles seen on the fracture surface were fractured. On the other hand, significant interfacial debonding was observed in alloy 562 tested at 232°C, indicating interface sliding prior to debonding. The bonding strength decreased with increasing temperature, partly because of a reduction in matrix squeezing force on the particles due to the disparity of coefficient of thermal expansion, and was comparable to the tearing strength of the matrix at 232°C. As temperature further increases, the difference in strength between the two phases becomes greater such that the matrix plastic flow becomes the primary driving force for particle rearrangement. This is consistent with the observation that little debonding is seen on the fracture surface of alloy 562 at 371°C, see Fig. 5(b). The fractographic observations are in agreement with the analysis that at elevated temperatures particle rearrangement had occurred as indicated by the bulges and cavities seen on the fracture surfaces. It should be noted that the rearrangement of particles is limited in Be/Al composites compared to grain switching seen in superplastic deformation because of the distinction between matrix flow, interfacial sliding and grain boundary slid-
72
S.G. Song et al. / Materials Science and Engineering A282 (2000) 67–73
Fig. 8. Bright filed image of room temperature deformation structure of AlBemet 162, (a) well-developed sub-structure in the matrix and planar slip in the reinforcement phase, (b) dislocation arrays in the reinforcement phase.
ing. The worksoftening of the composite is therefore expected to be less pronounced than that seen in superplasticity. The particle size of the alloys affects not only the tensile strength, consequent to Hall – Petch effect, but also the strain–softening behavior of the composites. The particle size of alloy 562 is significantly smaller than that of alloy 162. Accordingly, the softening behavior of the former is more conspicuous than the latter. Work hardening and softening are nearly balanced in alloy 162 at 232°C. It is of interest to note that the ductility of alloy 562 did not increase with increasing temperature as the matrix alloy became softer because the beryllium phase did not contribute to the plastic deformation at elevated temperatures. The stress buildup in the matrix was insufficient to cause yielding of the beryllium due to increased difference in strength between the two phases. Both composites were extruded at high temperatures and heat-treated afterward, matrix recrystallization is therefore expected to have occurred. The matrix grain size of alloy 162 should be greater than alloy 562 after recrystallization due to its larger beryllium particle size. The dislocation pileup in the matrix of alloy 162 is, therefore, expected to be more pronounced than that in alloy 562, which will lead to a lower yield stress of the former. Consequently, alloy 162 experienced lower interface shear stress at elevated temperatures than alloy 562. The higher interface stress in alloy 562 increases the probability of interface sliding so that both matrix plastic flow and interface sliding contribute to worksoftening. The difference in worksoftening between the two materials is reflected in fractographic characteristics, as seen earlier, that alloy 162 shows extensive matrix plastic strain with little sign of interfacial debonding whereas
alloy 562 exhibits both matrix plastic deformation and interfacial debonding. Although the two composites have distinct matrixes, the difference in the observed softening behavior is believed to be primarily a result of particle size effect, not the matrix property. Similarly, the large difference (\ 200 MPa) in yield strength between the two alloys cannot be accounted for by the strength variation of the matrix. However, the particle size strengthening effect cannot be calculated directly using the Hall–Petch relationship as for monolithic materials, because the dislocation initiation in the reinforcement phase by the matrix dislocation pileup at the interface is not only a function of the matrix dimension but also of the reinforcement particle size.
5. Conclusions The following conclusions can be drawn from the present investigation. Worksoftening behavior has been observed in Be/Al composites at elevated temperatures. Alloy 562 exhibited a pronounced softening behavior after yielding while alloy 162 workhardened slightly in the early stage of deformation and its stress/strain curve leveled off later. Worksoftening is always associated with uneven fracture surface exhibiting bulges and cavities. Extensive plastic deformation and interfacial debonding were commonly observed at elevated temperatures on the fracture surfaces. At room temperature however, the fracture surface is flat and the majority of the particles on the surface are fractured.
S.G. Song et al. / Materials Science and Engineering A282 (2000) 67–73
Reinforcement particle size plays an important part in the observed worksoftening. The smaller the particle size is, the more pronounced the softening behavior is seen. Particle rearrangement and interfacial sliding are the primary causes for the observed worksoftening of the Be/Al composites. Increasing temperature and decreasing particle size intensify worksoftening behavior.
Acknowledgements This work was supported by the Office of Naval Research (ONR/NW) of Department of Defense (DOD) under contract number N00014-94-C-0098/P00007 with the United Technologies Research Center (UTRC). The author would like to acknowledge J.T. Beals of Pratt and Whitney for his work on the mechanical testing of Be/Al.
.
73
References [1] J.J. Jonas, M.J. Luton, Advances in Deformation Processing, vol. 21, Plenum, New York, 1978, pp. 215 – 243. [2] E.A.McG. Chojnowski, W.J. Tegart, J. Mater. Sci. 2 (1968) 14 – 18. [3] N. Kanetake, N. Nakamura, J. Jpn. Soc. Technol. Plast. 31 (1990) 536 – 541. [4] A. Inoue, H.M. Kimura, M. Watanabe, A. Kawabata, Mater. Trans. JIM 38 (1997) 756 – 760. [5] J. Zhao, X. Xia, H.J. McQueen, Hot Workability of Steels and Light Alloys-Composites, Canadian Institute of Mining, Metallurgy and Petroleum, 1996, pp. 151 – 157. [6] J.F. Breedis, A. Lawley, J.A. Zeiger, in: L.M. Schetky, H.A. Johnson (Eds.), Beryllium Technology, vol. 1, Gordon and Breach, New York, 1966, pp. 197 – 226. [7] Designing with Beryllium, Product Brochure, Brush & Wellman Inc., Cleveland OH. [8] W. Blum, Q. Zhu, R. Merkel, H.J. McQueen, Z. Metallkd. 87 (1996) 14 – 23.