- Email: [email protected]
Metastable phases synthesized by inert-gas-condensation
NanoSm~ctmed Materials, Vol. 9, pp. I09-I 12,1997 Elsevier Scie~e Lid @ 1997 Acta Metallurgica Inc. Printed ~I the USb. All li~ats rmerved 0965-9TI3~7 $17.00 + .00
Pergamon
pII S09~-9773(97)00030.5
METASTABLE PHASES SYNTHESIZED BY INERT-GASCONDENSATION W. Krauss and R. Birringer Institute of Technical Physics, University of Saarbriicken, D-66041 Saarbriicken, Germany
Abstract -- Using the inert-gas condensation method we have produced metastable phases of Ta, W and YeOs .The formation of these phases results front both the undercooling achieved during inert-gas condensation as well as the surface-tension-induced increase in pressure with decreasing particle size. For example, we have found that Y203 can be synthesized in the equilibrium cubic phase (a- Y2Oa) as well as in the high-pressure monoclinic phase (?~YeOa)by varying the preparation conditions. The resulting crystal structure is found to be determined by the crystailite size (D). Based on these observations it is shown that by utilizing-inert gas condensation one can prepare metastable structures and alloys that can hardly be synthesized by classical metailurgical processes. © 1997 Acta Metallurgica Inc. INTRODUCTION
The inert-gas condensation (IGC) method for synthesizing nanocrystallinc materials (1) can be used to prepare a range of phases for various materials. In this technique, crystallites of a given material solidify from the vapor phase in an inert-gas atmosphere. The fact that this solidification generally occurs in highly undercooled liquid droplets at moderately fast rates often leads to the nucleation and growth of metastable phases. In this paper we want to focus on the special particle solidification conditions during this process. The theory of particle formation and solidification in IGC is based on the concepts developed by Granqvist and Buhrman (2). Directly above the evaporation source is a region of supersaturated metal vapor, where the nucleation of small 'liquid-like' metal particles takes place. In this narrow region the further growth of the particles occurs mainly by coalescence to larger particles. The gas convection between the hot evaporation source and the cold finger (77 K) transports the particles out of this nucleation and growth region, thus ending the growth process. The particles, which are still 'liquid like', then solidify. The thermodynamic quantities that dominate the solidification process are the driving force for solidification AG, the interfacial energy ,/between the solid and the liquid phase - both strongly dependent on the amount of undercooling AT - the catalytic factor f(0) and the surface tension ~ between the liquid and its vapor phase (3). The kinetics will be strongly affected by the cooling rate. Considering the radiation loss as the cooling mechanism for such small particles should lead to relatively high cooling rates, which may favor metastable phases (4). Various methods are known to influence the cited factors. The closest analogy to the IGC process can be deduced from the dispersion techniques (5) and drop-tube experiments (6). For the dispersion method the highest rcl~rtcd undcrcoolings rclativo to the melting point of the 100
110
W KP~us8AND R BIRFIINGER
Material (AT/TE) are as high as 0.5. This fact leads to interesting effects, such as the formation of metastable phases. The smallest particle size used in this type of experiments is about 4 Inn. By the use of drop-tubes the particle size is usually larger (1 - 5 mm) and the achieved undercooling (AT/TE) is in the range of 0.15 to 0.2 for pure metals. The fact that (AT/TE) is smaller, suggests that the amount of undercooling scales inversely with the particle size. In contrast to the dispersion method, in drop-tube experiments a carrier fluid is no longer necessary and this method can even be applied to highly refractory metals. Taking advantage of the special solidification parameters in drop-tube experiments, a variety of metastable phases can be prepared, though some only as transient states. For the occurrence of metastable structures, similarities between the undercooled melt and the metastable phase are considered to be dominant. This implies a lower interfacial energy )'m~ between the melt and the metastable phases compared to the y, of the melt - stable phase interface. With the IGC technique particles sizes in the range between several to 30 nm can be generated. Since different evaporation sources (resistivity heating, sputtering, e-gun evaporation) are available for the synthesis of nanocrystals by IGC, there are no restrictions on the selection of materials. In addition to facilitating a large amount of undercooling, the extremly small size of the 'liquid-like' particles produced by IGC leads to large effective pressures. According to the Laplace equation, the pressure increase Ap can be a high as in the GPa range (7). Upon solidification, the surface tension term o must be replaced by the surface stress tensor (fij), which can lead to even higher pressures inside the particles (8). Taking all these factors into account, there should be at least two cases in which the IGC leads to new metastable phases. First, metastable phases that are structurally similar to the undercooled melt may be observed due to the undercooling and the cooling rate of small particles in IGC (9). Consequently, the density of these metastable phases should be closer to the liquid ones than the stable modification. Second, the Laplace pressure may encourage the formation of high-pressure phases, which - in contrast to the metastable phases resembling the undercooled melt - will generally have a higher density than the equilibrium crystalline phase. EXPERIMENTAL
The experimental setup for the IGC process used for particle synthesis is described elsewhere (4). As model systems, we have chosen W and Ta on the one hand and Y203 on the other. Because of the high melting points of W and Ta, dc-sputtering was used instead of thermal evaporation to energize the metals into their vapor phase. In earlier work, it has been shown that by adjusting the process parameters, dc-sputtering is a suitable tool for producing nanocrystalline particles similar to those produced by thermal evaporation (10). After collecting the resulting particles on a cold finger, they were scraped into a press and compacted (1.5 GPa) in UHV to a macroscopic sample ( 6 ~ 8 nun, _<0.6 mm thick). These samples can then be analyzed with respect to their structural and physical properties. For the synthesis of Y203, pure Y was evaporated and the Y particles were collected on the cold finger. After stopping the evaporation process the inert gas was removed from the chamber and 02 gas was let in until atmospheric pressure was reached. Subsequently, the powder was taken out of the evaporation chamber and was oxidized in pure 02 atmosphere (350 °C, 35 rain), because the particles were not completely oxidized in the IGC process (11). The resulting powder changed its color from dark grey to white during this last annealing step. Finally, some of this powder was compacted in UHV conditions to pellets of 8 mm diameter.
METASTABLEPHASESSYNTHESlZlEOBYINERT-GAs-CoNDENSATION
1 11
RESULTS AND DISCUSSION Figure 1 shows XRD scans of nanocrystalline W and Ta, respectively. In addition to the effect of peak broadening resulting from the small crystal size of approximatdy 12 nm, it can be seen that the crystallites in the samples do not have the equilibrium structure (SG: Im3m) (12,13). Instead, a crystalline metastable phase could be identified. In the case of W [Fig. l(a)] the sample has the Pm3m structure with a lattice parameter of 5.0548 A, which yields a density of 18.91 g/cm3 (literature value: 19.262 g/cm3). This metastable structure of W is often falsely cited as 'W30' (14), whereas simulations of XRD pattern clearly reveal that the structure is truly W3W, which is called the AI5 modification of W or ~-W. As a control experiment, the experimental setup was changed to lead to film growth on the cold finger instead of particle formation and solidification in the inert gas. This could be proven by checking the morphology of the resulting product by TEM. Similar to the W particles, these films could be scraped off the cold finger and compacted to macroscopic samples. These samples show a crystal size also in the nm range (5 - 20 nm), whereas the crystal structure is now the equilibrium structure with a lattice parameter of 3.165 A (not shown here). This observation clearly confirms that the nucleation conditions determine the crystal structure. In Fig l(b) the diffraction pattern ofa Ta sample is shown. Just as with W, the Ta samples do not have the equilibrium structure. This structure can be identified as O-Ta (SG: P42/mnn, literature density: 16.326 g/cm3) (15); the density of the equilibrium structure is 16.634 g/cm 3 (13). For Y203, there is no difference in the XRD pattern whether the powder or the compacted material was analyzed. In contrast to the white color of the powder, the compacted material is transparent, revealing that the particles are completely oxidized and that there are no scattering centres (e.g. pores) with diameters comparable to the wavelength of visible light. A strong dependence of the resulting structure on the particle size could be detected. In Fig. 2(a) a diffraction pattern of a sample with a crystalline diameter _> 20 nm is shown. The sample consists almost completely (approx. 95 %) of a phase with the equilibrium crystal structure of Y203 (SG: Ia3), which has a density of 5.032 g/cm3 (16). If the parameters of the IGC process are adjusted in such a manner that the powder has a grain size of _< 10nm after the postoxidation, the XRD pattern shows that the sample no longer has the equilibrium structure [fig. 2 Co)]. The new structure of the sample could be identified as the high pressure phase of Y203 (SG: C2/m), which has a density of 5.468 g/cm3 (17).
a)
15
25
35
45
55
65
2 0 [degreesJ
75
85
95
5
I0
15
20
25
30
35
40
45
50
2 0 [degrees]
Figure 1: XRD patterns of (a) W (Cu-Ka) and (b) Ta (Mo-Ka), intensities as arb. units. The vertical lines show the corresponding fiterature values.
112
W KRAUSSAND R BIRRINGER
a)
10
15
20
25
30
10
15
20
25
30
2 0 [degrees] 2 0 [degrees] Figure 2." X R D patterns o f Y203 with particle sizes of(a) ~ 20 nm and (b) ~ 9 nm (Mo-K~), intensities as arb. units. The vertical lines show the corresponding literature values.
CONCLUSIONS The presented work shows that by applying IGC it is possible to synthesize metastablc crystal structures in macroscopic samples stable at room temperature. Candidate systems arc materials that have a high-pressure modification, as Y203, or materials that can form a metastable phase with structural similarities to the undercooled melt, as in the case with both W and Ta.
ACKNOWLEDGEMENTS This work was supported by SFB 277.
REFERENCES 1. 2. 3. 4. 5.
H. Gleiter, Materials Science Forum 189-190, 67 (1995). C.G. Granqvist and R.A. Buhnnan, J. Appl. Phys. 47, 2200 (1976). B. Feuerbacher, Mater. Sei. Rep. 4, 1 (1989). R.F. Cochrane, Euromat 91 Vol 1, p. 18, Inst. of Mat., London (1992). J.H. Perepezko and L.E. Anderson, Synthesis and properties of metastable phases, p.31,Warrendale, PA, Metall. Soc. of AIME (1980) 6. T.J. Rathz, M.B. Robinson, W.H. Holineister and R.J. Bayuziek, Rev. Sci. Instrum. 31, 3846 (1990). 7. Phase Transformations in Metals and Alloys, ed. by D.A. Porter and K.E. Easterling, Chapman and Hall, London (1991). 8. J.P. Borel andA. Chfitelain, Surf. Sci. 156, 572 (1985). 9. L. Cortella, B. Vinet, P.J. Desr6, A. Pasturel, A.T. Paxton and M. van Schilfgaarde, Phys. Rev. Lett. 70, 1469 (1993). 10. V. Haas and R. Birringer, Nanostruct. Mat. 11,491 (1991). 11. G. Skandan, C.M.Foster, H.Frase, M.N,AIi, J.C. Parker and H. Hahn, Nanostruct. Mat.l, 313 (1992) 12. ASTM No.: 4-806. 13. ASTM No.: 4-788. 14. B. Krebs, C. Brendel and H. Sch~ifer: Z. Anorg. Allg. Chem. 553, 127 (1987). 15. ASTM No.: 25-1280. 16. ASTM No.: 43- 1036. 17. ASTM No." 44-399
Pergamon
pII S09~-9773(97)00030.5
METASTABLE PHASES SYNTHESIZED BY INERT-GASCONDENSATION W. Krauss and R. Birringer Institute of Technical Physics, University of Saarbriicken, D-66041 Saarbriicken, Germany
Abstract -- Using the inert-gas condensation method we have produced metastable phases of Ta, W and YeOs .The formation of these phases results front both the undercooling achieved during inert-gas condensation as well as the surface-tension-induced increase in pressure with decreasing particle size. For example, we have found that Y203 can be synthesized in the equilibrium cubic phase (a- Y2Oa) as well as in the high-pressure monoclinic phase (?~YeOa)by varying the preparation conditions. The resulting crystal structure is found to be determined by the crystailite size (D). Based on these observations it is shown that by utilizing-inert gas condensation one can prepare metastable structures and alloys that can hardly be synthesized by classical metailurgical processes. © 1997 Acta Metallurgica Inc. INTRODUCTION
The inert-gas condensation (IGC) method for synthesizing nanocrystallinc materials (1) can be used to prepare a range of phases for various materials. In this technique, crystallites of a given material solidify from the vapor phase in an inert-gas atmosphere. The fact that this solidification generally occurs in highly undercooled liquid droplets at moderately fast rates often leads to the nucleation and growth of metastable phases. In this paper we want to focus on the special particle solidification conditions during this process. The theory of particle formation and solidification in IGC is based on the concepts developed by Granqvist and Buhrman (2). Directly above the evaporation source is a region of supersaturated metal vapor, where the nucleation of small 'liquid-like' metal particles takes place. In this narrow region the further growth of the particles occurs mainly by coalescence to larger particles. The gas convection between the hot evaporation source and the cold finger (77 K) transports the particles out of this nucleation and growth region, thus ending the growth process. The particles, which are still 'liquid like', then solidify. The thermodynamic quantities that dominate the solidification process are the driving force for solidification AG, the interfacial energy ,/between the solid and the liquid phase - both strongly dependent on the amount of undercooling AT - the catalytic factor f(0) and the surface tension ~ between the liquid and its vapor phase (3). The kinetics will be strongly affected by the cooling rate. Considering the radiation loss as the cooling mechanism for such small particles should lead to relatively high cooling rates, which may favor metastable phases (4). Various methods are known to influence the cited factors. The closest analogy to the IGC process can be deduced from the dispersion techniques (5) and drop-tube experiments (6). For the dispersion method the highest rcl~rtcd undcrcoolings rclativo to the melting point of the 100
110
W KP~us8AND R BIRFIINGER
Material (AT/TE) are as high as 0.5. This fact leads to interesting effects, such as the formation of metastable phases. The smallest particle size used in this type of experiments is about 4 Inn. By the use of drop-tubes the particle size is usually larger (1 - 5 mm) and the achieved undercooling (AT/TE) is in the range of 0.15 to 0.2 for pure metals. The fact that (AT/TE) is smaller, suggests that the amount of undercooling scales inversely with the particle size. In contrast to the dispersion method, in drop-tube experiments a carrier fluid is no longer necessary and this method can even be applied to highly refractory metals. Taking advantage of the special solidification parameters in drop-tube experiments, a variety of metastable phases can be prepared, though some only as transient states. For the occurrence of metastable structures, similarities between the undercooled melt and the metastable phase are considered to be dominant. This implies a lower interfacial energy )'m~ between the melt and the metastable phases compared to the y, of the melt - stable phase interface. With the IGC technique particles sizes in the range between several to 30 nm can be generated. Since different evaporation sources (resistivity heating, sputtering, e-gun evaporation) are available for the synthesis of nanocrystals by IGC, there are no restrictions on the selection of materials. In addition to facilitating a large amount of undercooling, the extremly small size of the 'liquid-like' particles produced by IGC leads to large effective pressures. According to the Laplace equation, the pressure increase Ap can be a high as in the GPa range (7). Upon solidification, the surface tension term o must be replaced by the surface stress tensor (fij), which can lead to even higher pressures inside the particles (8). Taking all these factors into account, there should be at least two cases in which the IGC leads to new metastable phases. First, metastable phases that are structurally similar to the undercooled melt may be observed due to the undercooling and the cooling rate of small particles in IGC (9). Consequently, the density of these metastable phases should be closer to the liquid ones than the stable modification. Second, the Laplace pressure may encourage the formation of high-pressure phases, which - in contrast to the metastable phases resembling the undercooled melt - will generally have a higher density than the equilibrium crystalline phase. EXPERIMENTAL
The experimental setup for the IGC process used for particle synthesis is described elsewhere (4). As model systems, we have chosen W and Ta on the one hand and Y203 on the other. Because of the high melting points of W and Ta, dc-sputtering was used instead of thermal evaporation to energize the metals into their vapor phase. In earlier work, it has been shown that by adjusting the process parameters, dc-sputtering is a suitable tool for producing nanocrystalline particles similar to those produced by thermal evaporation (10). After collecting the resulting particles on a cold finger, they were scraped into a press and compacted (1.5 GPa) in UHV to a macroscopic sample ( 6 ~ 8 nun, _<0.6 mm thick). These samples can then be analyzed with respect to their structural and physical properties. For the synthesis of Y203, pure Y was evaporated and the Y particles were collected on the cold finger. After stopping the evaporation process the inert gas was removed from the chamber and 02 gas was let in until atmospheric pressure was reached. Subsequently, the powder was taken out of the evaporation chamber and was oxidized in pure 02 atmosphere (350 °C, 35 rain), because the particles were not completely oxidized in the IGC process (11). The resulting powder changed its color from dark grey to white during this last annealing step. Finally, some of this powder was compacted in UHV conditions to pellets of 8 mm diameter.
METASTABLEPHASESSYNTHESlZlEOBYINERT-GAs-CoNDENSATION
1 11
RESULTS AND DISCUSSION Figure 1 shows XRD scans of nanocrystalline W and Ta, respectively. In addition to the effect of peak broadening resulting from the small crystal size of approximatdy 12 nm, it can be seen that the crystallites in the samples do not have the equilibrium structure (SG: Im3m) (12,13). Instead, a crystalline metastable phase could be identified. In the case of W [Fig. l(a)] the sample has the Pm3m structure with a lattice parameter of 5.0548 A, which yields a density of 18.91 g/cm3 (literature value: 19.262 g/cm3). This metastable structure of W is often falsely cited as 'W30' (14), whereas simulations of XRD pattern clearly reveal that the structure is truly W3W, which is called the AI5 modification of W or ~-W. As a control experiment, the experimental setup was changed to lead to film growth on the cold finger instead of particle formation and solidification in the inert gas. This could be proven by checking the morphology of the resulting product by TEM. Similar to the W particles, these films could be scraped off the cold finger and compacted to macroscopic samples. These samples show a crystal size also in the nm range (5 - 20 nm), whereas the crystal structure is now the equilibrium structure with a lattice parameter of 3.165 A (not shown here). This observation clearly confirms that the nucleation conditions determine the crystal structure. In Fig l(b) the diffraction pattern ofa Ta sample is shown. Just as with W, the Ta samples do not have the equilibrium structure. This structure can be identified as O-Ta (SG: P42/mnn, literature density: 16.326 g/cm3) (15); the density of the equilibrium structure is 16.634 g/cm 3 (13). For Y203, there is no difference in the XRD pattern whether the powder or the compacted material was analyzed. In contrast to the white color of the powder, the compacted material is transparent, revealing that the particles are completely oxidized and that there are no scattering centres (e.g. pores) with diameters comparable to the wavelength of visible light. A strong dependence of the resulting structure on the particle size could be detected. In Fig. 2(a) a diffraction pattern of a sample with a crystalline diameter _> 20 nm is shown. The sample consists almost completely (approx. 95 %) of a phase with the equilibrium crystal structure of Y203 (SG: Ia3), which has a density of 5.032 g/cm3 (16). If the parameters of the IGC process are adjusted in such a manner that the powder has a grain size of _< 10nm after the postoxidation, the XRD pattern shows that the sample no longer has the equilibrium structure [fig. 2 Co)]. The new structure of the sample could be identified as the high pressure phase of Y203 (SG: C2/m), which has a density of 5.468 g/cm3 (17).
a)
15
25
35
45
55
65
2 0 [degreesJ
75
85
95
5
I0
15
20
25
30
35
40
45
50
2 0 [degrees]
Figure 1: XRD patterns of (a) W (Cu-Ka) and (b) Ta (Mo-Ka), intensities as arb. units. The vertical lines show the corresponding fiterature values.
112
W KRAUSSAND R BIRRINGER
a)
10
15
20
25
30
10
15
20
25
30
2 0 [degrees] 2 0 [degrees] Figure 2." X R D patterns o f Y203 with particle sizes of(a) ~ 20 nm and (b) ~ 9 nm (Mo-K~), intensities as arb. units. The vertical lines show the corresponding literature values.
CONCLUSIONS The presented work shows that by applying IGC it is possible to synthesize metastablc crystal structures in macroscopic samples stable at room temperature. Candidate systems arc materials that have a high-pressure modification, as Y203, or materials that can form a metastable phase with structural similarities to the undercooled melt, as in the case with both W and Ta.
ACKNOWLEDGEMENTS This work was supported by SFB 277.
REFERENCES 1. 2. 3. 4. 5.
H. Gleiter, Materials Science Forum 189-190, 67 (1995). C.G. Granqvist and R.A. Buhnnan, J. Appl. Phys. 47, 2200 (1976). B. Feuerbacher, Mater. Sei. Rep. 4, 1 (1989). R.F. Cochrane, Euromat 91 Vol 1, p. 18, Inst. of Mat., London (1992). J.H. Perepezko and L.E. Anderson, Synthesis and properties of metastable phases, p.31,Warrendale, PA, Metall. Soc. of AIME (1980) 6. T.J. Rathz, M.B. Robinson, W.H. Holineister and R.J. Bayuziek, Rev. Sci. Instrum. 31, 3846 (1990). 7. Phase Transformations in Metals and Alloys, ed. by D.A. Porter and K.E. Easterling, Chapman and Hall, London (1991). 8. J.P. Borel andA. Chfitelain, Surf. Sci. 156, 572 (1985). 9. L. Cortella, B. Vinet, P.J. Desr6, A. Pasturel, A.T. Paxton and M. van Schilfgaarde, Phys. Rev. Lett. 70, 1469 (1993). 10. V. Haas and R. Birringer, Nanostruct. Mat. 11,491 (1991). 11. G. Skandan, C.M.Foster, H.Frase, M.N,AIi, J.C. Parker and H. Hahn, Nanostruct. Mat.l, 313 (1992) 12. ASTM No.: 4-806. 13. ASTM No.: 4-788. 14. B. Krebs, C. Brendel and H. Sch~ifer: Z. Anorg. Allg. Chem. 553, 127 (1987). 15. ASTM No.: 25-1280. 16. ASTM No.: 43- 1036. 17. ASTM No." 44-399