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Strain-compensated nano-clusters in Al–Si–Ge alloys
Scripta Materialia 54 (2006) 1973–1978 www.actamat-journals.com
Strain-compensated nano-clusters in Al–Si–Ge alloys Velimir Radmilovic a
a,*
, Michael K. Miller b, David Mitlin c, Ulrich Dahmen
a
National Center for Electron Microscopy, Lawrence Berkeley Laboratory, MS-72, University of California, Berkeley, 1 Cyclotron Rd., CA 94720, USA b Oak Ridge National Laboratory, Metals and Ceramics Division, MS-6136, Oak Ridge, TN 37831-6136, USA c Department of Chemical and Materials Engineering, University of Alberta, Edmonton, Alberta, Canada T6G 2G6 Received 23 December 2005; received in revised form 27 January 2006; accepted 29 January 2006 Available online 9 March 2006
Abstract Atom probe tomography and high resolution transmission electron microscopy have been employed to reveal clustering of Si and Ge atoms in ternary Al–Si–Ge. No such clusters were observed in binary Al–Si. The clusters were on the order of five nanometers in diameter and contained Si, Ge and Al. This confirms a previous hypothesis that postulates the existence of such clusters due to atomic mismatch strain compensation between the Si and Ge atoms in an Al solid solution. 2006 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Al–Si–Ge; Clustering; Precipitation; Atom probe tomography; HREM
1. Introduction Because nano-clusters can significantly alter the properties of engineering materials, they have been the subject of numerous studies [1–4]. In alloy design it is accepted that clusters are indispensable in the formation of stable and metastable precipitates (which in turn determine the mechanical, electrical and corrosion properties of alloys) [5–8]. It has long been thought [6] that during nucleation there is a spontaneous transition from clusters, where the solute atoms occupy lattice sites of the host lattice, to precipitates where the atoms adopt their own crystal structure. A recent alternative view is that clusters represent another form of a heterogeneous nucleation site [7]. In this scenario, precipitates form directly from, on or near stable clusters, while unstable clusters do not survive long enough to influence the precipitation process. Some of the systems where pre-precipitate clusters have been detected include Al–Ag [9,10], Cu–Co [11], Al–Zn [12], Al–Cu [6], Cu–Be [13] and Au–Ni [14].
*
Corresponding author. E-mail address: [email protected] (V. Radmilovic).
Aluminum–germanium and aluminum–silicon alloys are metallurgically simple binary eutectics, and have historically been used as model systems to study mechanisms of precipitation. Both have limited and strongly temperature dependent Al-rich terminal solid solubilities, giving rise to classical age-hardening behavior. In both systems an appropriate quench/age treatment results in direct precipitation of the equilibrium Ge or Si phase without intermediate phase formation. Early work established the indispensable role of vacancies in the nucleation and growth processes of the diamond cubic precipitate phase from Al solid solution [15]. Later studies focused on the remarkably rich variety of morphologies in these two simple systems [16–18] and discovered that the key to most of the observed morphologies lies in the internal twin structure of the precipitates [19,20]. Both factors, the need for vacancies and the role of twinning, have been used to control the microstructure of these alloys through various heat treatment schedules [21]. Recent reports on precipitation in ternary Al–Si–Ge alloys showed a greatly enhanced density of nucleation and lower quench sensitivity than either of the binary alloys [22]. Subsequently, this effect was used to enhance dispersion strengthening in Al–Cu-based alloys [23]. Aluminum– copper alloys with minor additions of both Si and Ge are
1359-6462/$ - see front matter 2006 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.scriptamat.2006.01.046
1974
V. Radmilovic et al. / Scripta Materialia 54 (2006) 1973–1978
of scientific and commercial interest. Addition of both elements in amounts as small as 0.5 at.% each has resulted in the formation of certain intermetallic secondary phases (metastable Al2Cu known as h 0 ) at the expense of others (metastable Al3Cu known as h00 ) [24,25]. This effect is achieved by creating a dense distribution of ultra-fine Si–Ge precipitates as a template for h 0 nucleation. The extremely high number density of these precipitates is only present when both elements, Si and Ge are added. Addition of only a single element, either Si or Ge, results in precipitates that are too coarse to be useful [26]. However, a fundamental understanding of the enhanced density has been lacking. It has been proposed that misfit compensation during pre-precipitation clustering is the cause of the observed dramatic increase in the precipitate density in this ternary alloy [26]. Since Si and Ge atoms have opposite mismatch with the Al host lattice (Si is 3% smaller, Ge is 2% larger), clusters containing both elements are expected to be much more stable compared to clusters of either one. The main reason why this hypothesis has not been confirmed directly is that the Si–Ge clusters are extremely difficult to detect by conventional analytical electron microscopy or X-ray diffraction methods. In this study, atom probe tomography (APT) [27] has been combined with high resolution trans-
mission electron microscopy (HRTEM) to confirm the existence of such clusters by atomic-resolution observation. 2. Experimental procedure A bulk alloy of Al–0.5 at.%Si–0.5 at.%Ge was made by arc melting, 99.999 wt.% pure Si, 99.9999 Ge, and 99.99 Al. The sample was encapsulated in sealed quartz tubes backfilled with argon, annealed for 24 h at 500 C, and ice–water quenched. The final shape of the bulk alloy was a cylinder approximately 20 mm in length and 10 mm in diameter. Analysis by vacuum emission spectroscopy of a section of this alloy cylinder confirmed its composition with no other elements present in quantities greater than 0.005%. The samples were aged at room temperature and at 160 ± 1 C. To fabricate atom probe specimens, the sample was drawn to 0.1 mm diameter wire through a series of alternating drawing and annealing operations. These wires were again annealed for 24 h at 500 C and ice–water quenched, followed by a two stage electropolishing technique [27]. Atom probe tomography was performed at Oak Ridge National Laboratory with an energy-compensated threedimensional atom probe (ECOPOSAP) and a local electrode atom probe (LEAP). The experimental conditions
Fig. 1. Bright field (A) and dark field (B) images of Al–Si–Ge alloy quenched and aged 30 min at 160 C, recorded under 200 two beam condition close to 1 1 0 zone axis (inset); selected area diffraction patterns taken in 1 0 0Al (C) and 1 1 0Al (D) zone axis do not show any evidence for precipitation.
V. Radmilovic et al. / Scripta Materialia 54 (2006) 1973–1978
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Fig. 2. HRTEM images of (A) Si–Ge cluster (arrowed) viewed along 1 0 0 direction in a sample aged for 15 min at 160 C; (B) multiply twinned Si–Ge precipitate viewed along 1 1 0 direction in a sample aged for 5 h at 250 C (twin boundaries indicated by arrows).
used for these characterizations were a specimen temperature of 20–30 K, pulse repetition rates of 1.5 kHz (ECOPOSAP) and 200 kHz (LEAP), a pulse fraction of 20% and a detection rate of 2% in the LEAP. The maximum separation method that was used to detect clusters is based on the principle that the solute atoms in solute-enriched regions such as clusters or precipitates are closer together than the solute atoms in a random solid solution. Therefore a maximum separation distance between solute atoms can be used to distinguish between the solute atoms in clusters and in the matrix. Full details of this approach are given in reference [27]. Solute enriched regions were also investigated with isoconcentration surfaces and two-dimensional concentration maps that were constructed with the tracer method. In the tracer method, a three-dimensional array of 1 · 1 · 1 nm composition cells was constructed from the data. The maximum concentrations of the columns of cells perpendicular to one of the primary axes may be exhibited on a two-dimensional intensity map. The number density of the features was estimated from the number of features in the analyzed volume. The volume of the analyzed region was estimated from the number of atoms, detection efficiency of the mass spectrometer and the atomic density of the material.
typical of cube–cube related precipitates in these alloys indicates that at this stage these particles have not adopted the equilibrium diamond cubic structure and retain the face-centered cubic structure of the Al matrix. The clusters have diffuse boundaries, maintain the same crystal structure as the host lattice and are roughly spherical in shape. This is in contrast to Si–Ge precipitates, which are easily identified by their crystal structure, sharp outlines and internal twinning [28]. In addition, while Si–Ge precipitates are very stable during TEM examination, Si–Ge clusters dissolve when exposed to the electron beam for extended periods. For this reason, the size, distribution and composition of clusters have not been well characterized and there has been no unambiguous test of the strain compensation
3. Results and discussion Precipitates were not detected in samples aged up to 30 min. This is evident from Fig. 1, which shows transmission electron microscopy (TEM) images near a h1 1 0iAl zone axis of the Al–1Si–1Ge alloy in the as-quenched and quenched/aged condition (30 min at 160 C) along with diffraction patterns in 1 0 0 and 1 1 0 zone axis orientation. An HRTEM image of a solute cluster in a sample aged for 15 min at 160 C is shown in Fig. 2(A). The cluster is 3 nm in diameter and contains roughly 850 atoms (assuming a spherical geometry). The appearance of clusters in HRTEM is quite distinct from that of precipitates (see Fig. 2(B)). The absence of the Moire´ contrast that is
Fig. 3. A representative APT elemental map of the Si atoms in an Al–Si alloy that was quenched and aged for 15 min at 160 C prior to examination. Note the absence of Si clusters. The volume analyzed was 50 · 30 · 10 nm in size and contained 458,000 Al + Si atoms.
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Fig. 4. Representative APT elemental maps of Si and Ge atoms in an Al–Si–Ge alloy that was quenched and aged for 15 min at 160 C prior to examination. The correlation between Si and Ge clusters (arrowed) clearly indicates that alloying occurs at this stage. The size of the volume analyzed was identical to that in Fig. 3.
hypothesis. However, as shown below, atom probe tomography is particularly well suited to this question because it enables collection of atomic-level statistics from small sample volumes and allows a direct comparison of preprecipitation clustering in binary Al–1%Si and ternary Al–0.5%Si–0.5%Ge alloys. An atom map of a 10 nm-thick region in the binary Al– Si alloy aged for 15 min at 160 C is shown in Fig. 3. Atom maps of identically sized 10-nm-thick regions in the ternary Al–Si–Ge alloy after identical heat treatment are shown in Fig. 4. In this type of representation, the positions of the solute atoms are shown whereas the aluminum atoms have
been omitted for clarity. A random distribution of silicon atoms is apparent in the binary alloy (Fig. 3). Standard statistical tests [27] and comparisons of the experimental distribution of concentrations with computer simulated distributions of a random solid solution with the same silicon concentration using the maximum separation method at different maximum separation distances also indicated that the silicon was randomly distributed. In contrast, two silicon- and germanium-enriched regions are clearly evident in the ternary alloy (Fig. 4). Fig. 5(A) shows the 8%Si + Ge isoconcentration surface of the ternary alloy that was aged for 15 min at 160 C.
Fig. 5. (A) An 8%Si + Ge isoconcentration surface of the ternary alloy aged for 15 min at 160 C. Four spherical Si- and Ge-enriched regions are clearly evident in this volume and are fully encompassed in the volume. The corresponding two-dimensional concentration maps of the same volume for Si and Ge are shown in (B) and (C); size of the volume analyzed: 100 · 50 · 10 nm.
V. Radmilovic et al. / Scripta Materialia 54 (2006) 1973–1978
Four spherical silicon- and germanium-enriched regions are clearly visible in this volume. The corresponding twodimensional concentration maps of the same volume for Si and Ge are shown in Fig. 5(B) and (C), respectively. As these two-dimensional concentration maps are constructed from the maximum concentrations in each column, a high threshold of 10% solute was chosen in order to suppress the random fluctuations of solute in the matrix. A two-dimensional concentration map from a similarthickness region in the binary alloy, Fig. 6, does not reveal any significant solute-enriched regions, agreeing with the original hypothesis [22,26]. This observation clearly confirms the presence of Si–Ge clusters in the ternary Al–Si– Ge and the absence of Si clusters in the binary Al–Si alloy. An atom map of a well developed Si–Ge precipitate that was detected after aging for 9 h at 160 C is shown in Fig. 7. At this condition the microstructure contains a relatively dense distribution of multiply twinned Si–Ge
Fig. 6. Concentration map for Si in the binary Al–Si alloy aged for 15 min at 160 C. Note the absence of any Si enriched region above the random fluctuations of Si in the matrix; size of the volume analyzed: 100 · 50 · 10 nm.
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particles. The precipitate is spherically shaped with a diameter of 7 nm and a composition 67 at.%Si–33 at.%Ge. Thus, it can be concluded that, apart from their crystal structure, another critical difference between precipitates and clusters is that Si–Ge precipitates do not contain aluminum. These observations support the original hypothesis that in the ternary alloy the strain-compensated clusters transform directly to the Si–Ge precipitates [22,26], requiring fewer vacancies for the process than either binary alloy. The particle’s spherical shape is due to multiple twinning, which is ubiquitous in small precipitates (see for example Fig. 2(B)) [28–30]. The measured Si-rich composition agrees well with previous results obtained by analytical TEM for numerous Si–Ge precipitates [28] and with Calphad calculations of the ternary phase diagram [31]. 4. Conclusions The observations presented here give the first conclusive experimental evidence for pre-precipitation nano-clusters of Si–Ge in a ternary Al–Si–Ge alloy. Clusters were not observed in an identically aged binary Al–Si alloy. The clusters’ diameter was on the order of 3 to 5 nm, which is slightly less than the smallest reported sizes of the diamond cubic Si–Ge precipitates. The difference between the Si–Ge clusters and the Si–Ge precipitates was their crystal structure and composition. Clusters were face-centered cubic and contained up to 80 at.% Al, whereas the precipitates were diamond cubic and contained no Al. While the existence of strain-compensated Si–Ge clusters had been postulated previously, the direct observation presented here clarifies their role in the microstructural evolution and opens the door to a more systematic use of this phenomenon. Similar combinations of solute elements with opposite misfit could be used to form strain-compensated clusters in systems where elastic stresses prevent phase separation. For example, application to different host alloys may be useful in achieving novel microstructures that use these clusters as templates for heterogeneous nucleation.
Fig. 7. APT elemental distributions of spherical Si–Ge precipitate after ageing for 9 h at 160 C: (A) atom map; (B) corresponding concentration profile along z direction.
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V. Radmilovic et al. / Scripta Materialia 54 (2006) 1973–1978
Acknowledgments The authors would like to thank Kaye F. Russell, James Wu and John Jacobson for technical assistance. This work was funded by the Director, Office of Science, Office of Basic Energy Sciences, Materials Sciences and Engineering Division of the US Department of Energy under Contract No. DE-AC02-05CH1123. Research at the SHaRE User Facility was sponsored by the Division of Materials Sciences and Engineering, US Department of Energy, under Contract DE-AC05-00OR22725 with UT-Battelle, LLC.
[11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22]
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[1] Murayama M, Hono K. Acta Mater 1999;47:1537. [2] Bonnell DA, Liang Y, Wagner M, Carroll D, Ru¨hle M. Acta Mater 1998;46:2263. [3] Dahmen U, Xiao SQ, Paciornik S, Johnson E, Johansen A. Phys Rev Lett 1997;78:471. [4] Muller S, Wolverton C, Wang LW, Zunger A. Acta Mater 2000;48: 4007. [5] Rosenbaum HS, Turnbull D. Acta Metall 1959;7:664. [6] Cohen JB. Metall Trans A 1992;23:2685. [7] Nie JF, Muddle BC, Aaronson HI, Ringer SP, Hirth JP. Metall Mater Trans A 2002;33:1649. [8] Ringer SP, Hono K. Mater Charact 2000;44:101. [9] Alexander KB, LeGoues FK, Aaronson HI, Laughlin DE. Acta Metall 1984;32:2241. [10] Gragg JE, Cohen JB. Acta Metall 1971;19:507–12.
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Jiang XD, Wagner W, Wollenberger H. Z Metallkd 1991;82:192. Haeffner DR, Cohen JB. Acta Metall 1989;37:2185. Rioja RJ, Laughlin DE. Acta Metall 1980;23:1301. Wu TB, Cohen JB. Acta Metall 1983;31:1929. Saulnier A. Mem Sci Rev Metall 1961;58:615. Ko¨ster U. Mater Sci Eng 1969;5:174. Douin J, Dahmen U, Westmacott KH. Philos Mag B 1991;63:867. Hugo GR, Muddle BC. Acta Metall Mater 1990;38:365. Xiao SQ, Hinderberger S, Westmacott KH, Dahmen U. Philos Mag A 1996;73:1267. Hinderberger S, Xiao SQ, Westmacott KH, Dahmen U. Z Metallkd 1996;87:161. Hinderberger S, Dahmen U, Xiao SQ, Westmacott KH. Z Metallkd 2000;91:215. Hornbogen E, Mukhopadhyay AK, Starke EA. Z Metallkd 1992; 83:577. Mitlin D, Radmilovic V, Dahmen U, Morris JW. Metall Mater Trans A 2000;32:197. Mitlin D, Radmilovic V, Morris JW, Dahmen U. Metall Mater Trans A 2003;34:735. Mitlin D, Radmilovic V, Dahmen U, Morris JW. Metall Mater Trans A 2001;32:197. Hornbogen E, Mukhopadhyay AK, Starke EA. Z Metallkd 1992;92: 577. Miller MK. Atom probe tomography: analysis at the atomic level. New York (NY): Kluwer Academic/Plenum Press; 2000. Radmilovic V, Mitlin D, Dracup B, Miller MK, Dahmen U, Morris JW. Mater Sci Forum 2002;396–402:905. Radmilovic V, Mitlin D, Dahmen U, Morris JW. Metall Mater Trans A 2003;34:543. Mitlin D, Radmilovic V, Morris JW. Mater Sci Eng A 2001;301:231. Dracup B, Turchi PEA, Radmilovic V, Dahmen U, Morris JW. Metall Mater Trans 2004;35:2305.
Strain-compensated nano-clusters in Al–Si–Ge alloys Velimir Radmilovic a
a,*
, Michael K. Miller b, David Mitlin c, Ulrich Dahmen
a
National Center for Electron Microscopy, Lawrence Berkeley Laboratory, MS-72, University of California, Berkeley, 1 Cyclotron Rd., CA 94720, USA b Oak Ridge National Laboratory, Metals and Ceramics Division, MS-6136, Oak Ridge, TN 37831-6136, USA c Department of Chemical and Materials Engineering, University of Alberta, Edmonton, Alberta, Canada T6G 2G6 Received 23 December 2005; received in revised form 27 January 2006; accepted 29 January 2006 Available online 9 March 2006
Abstract Atom probe tomography and high resolution transmission electron microscopy have been employed to reveal clustering of Si and Ge atoms in ternary Al–Si–Ge. No such clusters were observed in binary Al–Si. The clusters were on the order of five nanometers in diameter and contained Si, Ge and Al. This confirms a previous hypothesis that postulates the existence of such clusters due to atomic mismatch strain compensation between the Si and Ge atoms in an Al solid solution. 2006 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Al–Si–Ge; Clustering; Precipitation; Atom probe tomography; HREM
1. Introduction Because nano-clusters can significantly alter the properties of engineering materials, they have been the subject of numerous studies [1–4]. In alloy design it is accepted that clusters are indispensable in the formation of stable and metastable precipitates (which in turn determine the mechanical, electrical and corrosion properties of alloys) [5–8]. It has long been thought [6] that during nucleation there is a spontaneous transition from clusters, where the solute atoms occupy lattice sites of the host lattice, to precipitates where the atoms adopt their own crystal structure. A recent alternative view is that clusters represent another form of a heterogeneous nucleation site [7]. In this scenario, precipitates form directly from, on or near stable clusters, while unstable clusters do not survive long enough to influence the precipitation process. Some of the systems where pre-precipitate clusters have been detected include Al–Ag [9,10], Cu–Co [11], Al–Zn [12], Al–Cu [6], Cu–Be [13] and Au–Ni [14].
*
Corresponding author. E-mail address: [email protected] (V. Radmilovic).
Aluminum–germanium and aluminum–silicon alloys are metallurgically simple binary eutectics, and have historically been used as model systems to study mechanisms of precipitation. Both have limited and strongly temperature dependent Al-rich terminal solid solubilities, giving rise to classical age-hardening behavior. In both systems an appropriate quench/age treatment results in direct precipitation of the equilibrium Ge or Si phase without intermediate phase formation. Early work established the indispensable role of vacancies in the nucleation and growth processes of the diamond cubic precipitate phase from Al solid solution [15]. Later studies focused on the remarkably rich variety of morphologies in these two simple systems [16–18] and discovered that the key to most of the observed morphologies lies in the internal twin structure of the precipitates [19,20]. Both factors, the need for vacancies and the role of twinning, have been used to control the microstructure of these alloys through various heat treatment schedules [21]. Recent reports on precipitation in ternary Al–Si–Ge alloys showed a greatly enhanced density of nucleation and lower quench sensitivity than either of the binary alloys [22]. Subsequently, this effect was used to enhance dispersion strengthening in Al–Cu-based alloys [23]. Aluminum– copper alloys with minor additions of both Si and Ge are
1359-6462/$ - see front matter 2006 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.scriptamat.2006.01.046
1974
V. Radmilovic et al. / Scripta Materialia 54 (2006) 1973–1978
of scientific and commercial interest. Addition of both elements in amounts as small as 0.5 at.% each has resulted in the formation of certain intermetallic secondary phases (metastable Al2Cu known as h 0 ) at the expense of others (metastable Al3Cu known as h00 ) [24,25]. This effect is achieved by creating a dense distribution of ultra-fine Si–Ge precipitates as a template for h 0 nucleation. The extremely high number density of these precipitates is only present when both elements, Si and Ge are added. Addition of only a single element, either Si or Ge, results in precipitates that are too coarse to be useful [26]. However, a fundamental understanding of the enhanced density has been lacking. It has been proposed that misfit compensation during pre-precipitation clustering is the cause of the observed dramatic increase in the precipitate density in this ternary alloy [26]. Since Si and Ge atoms have opposite mismatch with the Al host lattice (Si is 3% smaller, Ge is 2% larger), clusters containing both elements are expected to be much more stable compared to clusters of either one. The main reason why this hypothesis has not been confirmed directly is that the Si–Ge clusters are extremely difficult to detect by conventional analytical electron microscopy or X-ray diffraction methods. In this study, atom probe tomography (APT) [27] has been combined with high resolution trans-
mission electron microscopy (HRTEM) to confirm the existence of such clusters by atomic-resolution observation. 2. Experimental procedure A bulk alloy of Al–0.5 at.%Si–0.5 at.%Ge was made by arc melting, 99.999 wt.% pure Si, 99.9999 Ge, and 99.99 Al. The sample was encapsulated in sealed quartz tubes backfilled with argon, annealed for 24 h at 500 C, and ice–water quenched. The final shape of the bulk alloy was a cylinder approximately 20 mm in length and 10 mm in diameter. Analysis by vacuum emission spectroscopy of a section of this alloy cylinder confirmed its composition with no other elements present in quantities greater than 0.005%. The samples were aged at room temperature and at 160 ± 1 C. To fabricate atom probe specimens, the sample was drawn to 0.1 mm diameter wire through a series of alternating drawing and annealing operations. These wires were again annealed for 24 h at 500 C and ice–water quenched, followed by a two stage electropolishing technique [27]. Atom probe tomography was performed at Oak Ridge National Laboratory with an energy-compensated threedimensional atom probe (ECOPOSAP) and a local electrode atom probe (LEAP). The experimental conditions
Fig. 1. Bright field (A) and dark field (B) images of Al–Si–Ge alloy quenched and aged 30 min at 160 C, recorded under 200 two beam condition close to 1 1 0 zone axis (inset); selected area diffraction patterns taken in 1 0 0Al (C) and 1 1 0Al (D) zone axis do not show any evidence for precipitation.
V. Radmilovic et al. / Scripta Materialia 54 (2006) 1973–1978
1975
Fig. 2. HRTEM images of (A) Si–Ge cluster (arrowed) viewed along 1 0 0 direction in a sample aged for 15 min at 160 C; (B) multiply twinned Si–Ge precipitate viewed along 1 1 0 direction in a sample aged for 5 h at 250 C (twin boundaries indicated by arrows).
used for these characterizations were a specimen temperature of 20–30 K, pulse repetition rates of 1.5 kHz (ECOPOSAP) and 200 kHz (LEAP), a pulse fraction of 20% and a detection rate of 2% in the LEAP. The maximum separation method that was used to detect clusters is based on the principle that the solute atoms in solute-enriched regions such as clusters or precipitates are closer together than the solute atoms in a random solid solution. Therefore a maximum separation distance between solute atoms can be used to distinguish between the solute atoms in clusters and in the matrix. Full details of this approach are given in reference [27]. Solute enriched regions were also investigated with isoconcentration surfaces and two-dimensional concentration maps that were constructed with the tracer method. In the tracer method, a three-dimensional array of 1 · 1 · 1 nm composition cells was constructed from the data. The maximum concentrations of the columns of cells perpendicular to one of the primary axes may be exhibited on a two-dimensional intensity map. The number density of the features was estimated from the number of features in the analyzed volume. The volume of the analyzed region was estimated from the number of atoms, detection efficiency of the mass spectrometer and the atomic density of the material.
typical of cube–cube related precipitates in these alloys indicates that at this stage these particles have not adopted the equilibrium diamond cubic structure and retain the face-centered cubic structure of the Al matrix. The clusters have diffuse boundaries, maintain the same crystal structure as the host lattice and are roughly spherical in shape. This is in contrast to Si–Ge precipitates, which are easily identified by their crystal structure, sharp outlines and internal twinning [28]. In addition, while Si–Ge precipitates are very stable during TEM examination, Si–Ge clusters dissolve when exposed to the electron beam for extended periods. For this reason, the size, distribution and composition of clusters have not been well characterized and there has been no unambiguous test of the strain compensation
3. Results and discussion Precipitates were not detected in samples aged up to 30 min. This is evident from Fig. 1, which shows transmission electron microscopy (TEM) images near a h1 1 0iAl zone axis of the Al–1Si–1Ge alloy in the as-quenched and quenched/aged condition (30 min at 160 C) along with diffraction patterns in 1 0 0 and 1 1 0 zone axis orientation. An HRTEM image of a solute cluster in a sample aged for 15 min at 160 C is shown in Fig. 2(A). The cluster is 3 nm in diameter and contains roughly 850 atoms (assuming a spherical geometry). The appearance of clusters in HRTEM is quite distinct from that of precipitates (see Fig. 2(B)). The absence of the Moire´ contrast that is
Fig. 3. A representative APT elemental map of the Si atoms in an Al–Si alloy that was quenched and aged for 15 min at 160 C prior to examination. Note the absence of Si clusters. The volume analyzed was 50 · 30 · 10 nm in size and contained 458,000 Al + Si atoms.
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V. Radmilovic et al. / Scripta Materialia 54 (2006) 1973–1978
Fig. 4. Representative APT elemental maps of Si and Ge atoms in an Al–Si–Ge alloy that was quenched and aged for 15 min at 160 C prior to examination. The correlation between Si and Ge clusters (arrowed) clearly indicates that alloying occurs at this stage. The size of the volume analyzed was identical to that in Fig. 3.
hypothesis. However, as shown below, atom probe tomography is particularly well suited to this question because it enables collection of atomic-level statistics from small sample volumes and allows a direct comparison of preprecipitation clustering in binary Al–1%Si and ternary Al–0.5%Si–0.5%Ge alloys. An atom map of a 10 nm-thick region in the binary Al– Si alloy aged for 15 min at 160 C is shown in Fig. 3. Atom maps of identically sized 10-nm-thick regions in the ternary Al–Si–Ge alloy after identical heat treatment are shown in Fig. 4. In this type of representation, the positions of the solute atoms are shown whereas the aluminum atoms have
been omitted for clarity. A random distribution of silicon atoms is apparent in the binary alloy (Fig. 3). Standard statistical tests [27] and comparisons of the experimental distribution of concentrations with computer simulated distributions of a random solid solution with the same silicon concentration using the maximum separation method at different maximum separation distances also indicated that the silicon was randomly distributed. In contrast, two silicon- and germanium-enriched regions are clearly evident in the ternary alloy (Fig. 4). Fig. 5(A) shows the 8%Si + Ge isoconcentration surface of the ternary alloy that was aged for 15 min at 160 C.
Fig. 5. (A) An 8%Si + Ge isoconcentration surface of the ternary alloy aged for 15 min at 160 C. Four spherical Si- and Ge-enriched regions are clearly evident in this volume and are fully encompassed in the volume. The corresponding two-dimensional concentration maps of the same volume for Si and Ge are shown in (B) and (C); size of the volume analyzed: 100 · 50 · 10 nm.
V. Radmilovic et al. / Scripta Materialia 54 (2006) 1973–1978
Four spherical silicon- and germanium-enriched regions are clearly visible in this volume. The corresponding twodimensional concentration maps of the same volume for Si and Ge are shown in Fig. 5(B) and (C), respectively. As these two-dimensional concentration maps are constructed from the maximum concentrations in each column, a high threshold of 10% solute was chosen in order to suppress the random fluctuations of solute in the matrix. A two-dimensional concentration map from a similarthickness region in the binary alloy, Fig. 6, does not reveal any significant solute-enriched regions, agreeing with the original hypothesis [22,26]. This observation clearly confirms the presence of Si–Ge clusters in the ternary Al–Si– Ge and the absence of Si clusters in the binary Al–Si alloy. An atom map of a well developed Si–Ge precipitate that was detected after aging for 9 h at 160 C is shown in Fig. 7. At this condition the microstructure contains a relatively dense distribution of multiply twinned Si–Ge
Fig. 6. Concentration map for Si in the binary Al–Si alloy aged for 15 min at 160 C. Note the absence of any Si enriched region above the random fluctuations of Si in the matrix; size of the volume analyzed: 100 · 50 · 10 nm.
1977
particles. The precipitate is spherically shaped with a diameter of 7 nm and a composition 67 at.%Si–33 at.%Ge. Thus, it can be concluded that, apart from their crystal structure, another critical difference between precipitates and clusters is that Si–Ge precipitates do not contain aluminum. These observations support the original hypothesis that in the ternary alloy the strain-compensated clusters transform directly to the Si–Ge precipitates [22,26], requiring fewer vacancies for the process than either binary alloy. The particle’s spherical shape is due to multiple twinning, which is ubiquitous in small precipitates (see for example Fig. 2(B)) [28–30]. The measured Si-rich composition agrees well with previous results obtained by analytical TEM for numerous Si–Ge precipitates [28] and with Calphad calculations of the ternary phase diagram [31]. 4. Conclusions The observations presented here give the first conclusive experimental evidence for pre-precipitation nano-clusters of Si–Ge in a ternary Al–Si–Ge alloy. Clusters were not observed in an identically aged binary Al–Si alloy. The clusters’ diameter was on the order of 3 to 5 nm, which is slightly less than the smallest reported sizes of the diamond cubic Si–Ge precipitates. The difference between the Si–Ge clusters and the Si–Ge precipitates was their crystal structure and composition. Clusters were face-centered cubic and contained up to 80 at.% Al, whereas the precipitates were diamond cubic and contained no Al. While the existence of strain-compensated Si–Ge clusters had been postulated previously, the direct observation presented here clarifies their role in the microstructural evolution and opens the door to a more systematic use of this phenomenon. Similar combinations of solute elements with opposite misfit could be used to form strain-compensated clusters in systems where elastic stresses prevent phase separation. For example, application to different host alloys may be useful in achieving novel microstructures that use these clusters as templates for heterogeneous nucleation.
Fig. 7. APT elemental distributions of spherical Si–Ge precipitate after ageing for 9 h at 160 C: (A) atom map; (B) corresponding concentration profile along z direction.
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Acknowledgments The authors would like to thank Kaye F. Russell, James Wu and John Jacobson for technical assistance. This work was funded by the Director, Office of Science, Office of Basic Energy Sciences, Materials Sciences and Engineering Division of the US Department of Energy under Contract No. DE-AC02-05CH1123. Research at the SHaRE User Facility was sponsored by the Division of Materials Sciences and Engineering, US Department of Energy, under Contract DE-AC05-00OR22725 with UT-Battelle, LLC.
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