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In situ observation of solidification behavior in undercooled Pd–Cu–Ni–P alloy by using a confocal scanning laser microscope
Acta mater. 49 (2001) 615–622 www.elsevier.com/locate/actamat
IN SITU OBSERVATION OF SOLIDIFICATION BEHAVIOR IN UNDERCOOLED Pd–Cu–Ni–P ALLOY BY USING A CONFOCAL SCANNING LASER MICROSCOPE J. -H. KIM1*, S. -G. KIM1 and A. INOUE1, 2 1
Inoue Superliquid Glass Project, ERATO, Japan Science and Technology Corporation (JST), Sendai 9820807, Japan and 2Institute for Materials Research, Tohoku University, Sendai 980-77, Japan ( Received 2 May 2000; received in revised form 16 October 2000; accepted 18 October 2000 )
Abstract—In the undercooled melt of Pd40Cu30Ni10P20 alloy, the solidification behavior including the nucleation and growth of crystals at the micrometer level has been observed in situ by use of a confocal scanning laser microscope combined with an infrared image furnace. The Pd40Cu30Ni10P20 alloy specimens were cooled from the liquid state to various undercooled states under a helium gas flow. Images of solidification progress were obtained by the charge-coupled device image sensor of the confocal scanning laser microscope. Depending on the degree of undercooling, the morphology of the solidification front changed among various types: faceted front, columnar dendritic front, cellular grain and equiaxed grain, etc. The velocities of the solid–liquid interface were measured to be 10⫺5–10⫺7 m/s, which are at least two orders of magnitude higher than the theoretical crystal growth rates. Combining the morphologies observed in the three undercooling regimes and their solidification behaviors, we conclude that phase separation takes place in the undercooled molten Pd40Cu30Ni10P20 alloy. The continuous-cooling–transformation (CCT) diagram was derived from the experimental time–temperature-transformation diagram constructed from solidification onset times under various isothermal annealing conditions. The CCT diagram suggests that the critical cooling rate for glassy solidification is about 1.8 K/s, which is in agreement with previous calorimetric findings. 2001 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. Keywords: Metallic glasses; Phase transformations; Microstructure; Kinetics; Dynamic phenomena
1. INTRODUCTION
In order to investigate directly the growth of crystals in situ in the undercooled liquid of bulk metallic glasses, it is necessary to overcome the effect of thermal radiation from the melts on microscope viewing and to distinguish between the two phases at elevated temperatures. A confocal scanning laser microscope (CSLM) was used in this study to overcome such difficulties for in situ observation of metallic melts. Actually, a Japanese research group has previously reported successful observation of the growth of crystals in Fe–C alloy melts by use of a confocal scanning laser microscope [1]. In addition, the inherent thermal stability of bulk metallic glasses against crystallization and oxidation, which are other important parameters along with the technical difficulties described above for the in situ observation of metallic melts, were considered. Multicomponent alloys recently discovered exhibit excellent glass-forming ability. The critical cooling rate for
* To whom all correspondence should be addressed. Fax: ⫹81-22-243-7616.
glass formation in Pd40Cu30Ni10P20 alloy in a B2O3 flux, for example, is as low as 0.1 K/s [2]. The thermal stability of these bulk metallic glasses against crystallization leads to an experimentally accessible time and temperature window to investigate the transformation behavior from the undercooled liquid into crystalline equilibrium phases. The aim of this paper is to present an in situ observation in a confocal scanning laser microscope of the solidification behavior and a microstructure analysis of Pd40Cu30Ni10P20 alloy solidified during different isothermal annealing treatments at various levels of undercooling. Further, the glass-forming ability (GFA) for this metallic glass is evaluated on the basis of the experimental findings and classical nucleation theory (CNT), and is compared with that obtained in previous reports [2, 3].
2. EXPERIMENTAL PROCEDURES
The master ingot was prepared by arc-melting a mixture of pure metallic Pd, Cu and Ni (purity ⱖ 99.99%) and pre-alloyed Pd60P40 in a purified argon atmosphere. Details of the preparation of bulk glassy
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samples have been reported previously [2]. In the undercooled melt of the Pd40Cu30Ni10P20 alloy, the solidification behavior including nucleation and growth of crystals on the micrometer level was observed in situ by use of a confocal scanning laser microscope combined with an infrared image furnace. A confocal scanning laser microscope (1LM21 from Lasertec Co.; with 0.2 µm resolution at 3500⫻ magnification) was employed with a 1.5 mW He–Ne laser of 632.8 nm wavelength. The laser beam was reflected and scanned by an acoustic optic deflector in the horizontal direction at a rate of 15.7 kHz and by a glavano mirror in the vertical direction at 60 Hz. The observation mechanism and schematic construction of this microscope have been given previously [1]. The Pd40Cu30Ni10P20 alloy specimens, which were cut into cubes 19±2 mg in weight and mirror polished, were melted in an alumina crucible and cooled from the liquid state at about 1.1TL to various undercooled states, as shown in Table 1, under a highpurity helium gas flow. Here, TL is the liquidus temperature, 804 K, of this alloy. The undercooled melt specimens were annealed isothermally at different undercooling regimes ⌬T ⫽ TL⫺Tiso ⫽ 39 to 173 K, where Tiso is the isothermal annealing temperature. The images of solidification progress during annealing were obtained by the charge-coupled device (CCD) image sensor, monitored on the cathode ray tube (CRT) of the confocal scanning laser microscope and recorded on videotape at a rate of 30 frames per second. The solidified specimens were quenched to room temperature after the isothermal annealing treatments. The velocities of the solid–liquid interface at the surface of the undercooled melt were derived from these observations. These experimental velocities were compared with theoretical crystal growth rates, which were calculated on the basis of Turnbull’s approximation and classical nucleation and growth theories [4–7]. In order to evaluate the GFA, the time–temperature–transformation (TTT) diagram was constructed from solidification onset times at which the primary phases were observed on the surface of undercooled melts under various annealing conditions.
The surface morphology and microstructures of solidified samples were observed by optical and scanning electron microscopy (SEM) (JXM-8621MX, Jeol Co.). The latter apparatus was equipped with wavelength-dispersive spectrometers (WDS) for determination of the element concentrations in various phases. Various phases in the solidified samples were analyzed with a micro-area X-ray diffractor (micro-area XRD) (M18XHF22, MAC Sci. Co.; Cu Kα radiation). 3. RESULTS
Figure 1 shows real-time solidification morphologies, taken from the surface of undercooled Pd40Cu30Ni10P20 alloy melts during isothermal annealing treatment at various levels of undercooling, obtained by using a confocal scanning laser microscope. Scanning electron micrographs of the cross-sections of the solidified samples are shown in Fig. 2. In Fig. 1, several kinds of morphology are seen in the undercooled Pd40Cu30Ni10P20 alloy, and they are dominated by the level of undercooling ⌬T. In the small undercooling regime—i.e., for ⌬Tⱕ 68 K, a complex microstructure consisting of two or three subnetworks is seen, as shown in Fig. 2(a), (b), (c) and (d). When the Pd40Cu30Ni10P20 alloy melt is annealed at 765 K (⌬T is 39 K), no obvious solidification front is seen on the surface of the undercooled melt except for the appearance of jag as shown in Fig. 1(a). In Fig. 2(a) and (a⬘) the microstructure of the cross-section of the sample is composed of two irregular phases (a1, a2) surrounded by a single matrix (a3). The matrix is estimated to be a glassy single phase or a glassy phase including fine crystals as evidenced by the existence of some Bragg peaks on the typical broad glassy X-ray pattern taken from the cross-sectional surface of the sample annealed at 765 K, as shown in Fig. 3(a). The approximate composition of all three phases was determined by WDS analysis on at least three different crystals in each case. The measured compositions are given in Table 2. According to Table 2, two irregular phases (a1, a2) are estimated to be nickel palladium phosphide and a copper-rich palladium compound including the other elements of the
Table 1. Conditions of isothermal annealing and morphologies of the solidification front in undercooled Pd40Cu30Ni10P20 alloy meltsa Isothermal annealing temperature, Tiso (K) 765 756 746 741 736 717 708 679 650 631 a
Degree of undercooling, ⌬T ⫽ 804⫺T⬑ (K) 39 48 58 63 68 87 96 125 154 173
Morphology of solidification on surface Jag Jag ⫹ faceted front Faceted front ⫹ jag Columnar dendrite Columnar dendrite Cellular grain Cellular grain Equiaxed grain Fine equiaxed grain Fine spot
Glass transition temperature, Tg, is 575 K and onset temperature for crystallization, Tx, is 670 K at a heating rate of 40 K/min.
Fig. 1. Real-time solidification morphologies, taken from the surface of undercooled Pd40Cu30Ni10P20 alloy melts during isothermal annealing treatment at various levels of undercooling, obtained by using a confocal scanning laser microscope. The degrees of undercooling were: (a) 39 K; (b) 48 K; (c) 58 K; (d) 63 K; (e) 68 K; (f) 87 K; (g) 96 K; (h) 125 K; (i) 154 K; (j) 173 K.
KIM et al.: OBSERVATION OF BEHAVIOR OF Pd–Cu–Ni–P ALLOY 617
Fig. 2. Scanning electron and optical micrographs of cross-sections of the Pd40Cu30Ni10P20 alloy samples solidified during isothermal annealing treatment at various levels of undercooling. The degrees of undercooling were: (a), (a⬘) 39 K; (b) 48 K; (c) 58 K; (d), (d⬘) 63 K; (e), (f) 96 K; (g), (h) 125 K; (i), (j) 154 K. Here, (a⬘) and (d⬘) are the microstructures obtained by using an optical microscope with a magnification of less than 200⫻. (f), (h) and (j) are the microstructures observed in the center region of cross-sections of the samples.
618 KIM et al.: OBSERVATION OF BEHAVIOR OF Pd–Cu–Ni–P ALLOY
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619
shown in Fig. 2(b), (c) and (d). In Fig. 2(b) and (d), one can see light gray faceted dendritic (b3) and white coarse (b4) phases, which are identified to be copperrich and palladium-rich quaternary phases by the WDS analysis measurement given in Table 2. In the intermediate undercooling regime—i.e., for 68 K⬍⌬T⬍125 K, growth and coarsening of grains which have directional solidification fronts on a submillimeter level are seen as shown in Fig. 1(f) and (g). A coarse irregular grain, approximately 100 µm wide, is seen together with other smaller grains in the outer region of the cross-sectional surface as shown in Fig. 2(e). In the interior of the samples, however, acicular phases (c1) surrounded by another black phase (c2), which are located at the center of each irregular grain background (c3, c4), are seen in Fig. 2(f). The acicular phase (c1) is estimated to be nickel palladium phosphide, including copper element below 4 at% (Table 2), as evidenced by the existence of Bragg peaks of Pd2Ni2P on the X-ray patterns in Fig. 3(c). The composition of the irregular grain background is similar to that of the master ingot as given in Table 2. It appears that the solidification progress of the interior crystals is different from that of the grains near the surface of the sample. If a thermal gradient exists in the sample during solidification and the undercooled melt has activated nuclei, the interior primary crystals can originate from the quenched interior nuclei and lead to solidification of the irregular grains at a slightly higher temperature than the surface temperature. On the other hand, grain growth near the surface of the sample may originate predominantly from the surface. Figure 4 shows the relationship between the final distance of movement of the solid–liquid interface after solidification and the annealing temperature, together with three undercooling regimes classified on the basis of the morphology of the solidification front. In the large undercooling regime, which is just above and below the crystallization onset temperature (670 K [2]) for ⌬Tⱖ125 K, spherical or equiaxed polyhedral grain morphologies, which arise from nucleation and growth, are identified and a sharp
Fig. 3. X-ray diffraction patterns taken from the cross-sections of the Pd40Cu30Ni10P20 alloy samples solidified during isothermal annealing treatment at various levels of undercooling. The degrees of undercooling were: (a) 39 K; (b) 58 K; (c) 96 K; (d) 125 K; (e) 154 K.
alloy below 2 at% level. From the compositional point of view, the nickel palladium phosphide identified in the present work is different from the ternary primary phase Ni45Pd34P21, which was described by Donovan et al. in a crystallized Pd40Ni40P20 glass [8]. Recently, crystallized phases in the Pd40Cu30Ni10P20 alloy have been identified as Ni3P, Pd15P2, Pd2Ni2P and Cu3Pd [9, 10]. In Fig. 3, the Bragg peaks of Pd2Ni2P and Cu3Pd are seen on the X-ray patterns taken from the cross-sectional surface of the annealed samples in the present work, along with shifting of the Bragg peaks. In the cases of the samples annealed at 756 K (⌬T is 48 K) and 746 K (⌬T is 58 K), faceted solidification fronts are observed with networks of complex jags on these surfaces as shown in Fig. 1(b) and (c). For the samples annealed at 741 K (⌬T is 63 K) and 736 K (⌬T is 68 K), solidification fronts of columnar dendritic shape are observed as shown in Fig. 1(d) and (e). The microstructures are composed typically of a metallic compound and/or a cellular dendritic precipitation with a eutectic-like background as
Table 2. Measured compositions of the solidified phases in the Pd40Cu30Ni10P20 alloy Isothermal annealing temperature (degree of undercooling)
765 K (39 K)
756 K (48 K)
708 K (96 K)
Phase a1 Phase a2 Matrix a3 Phase b1 Phase b2 Phase b3 Phase b4 Phase c1 Phase c2 Phase c3 Phase c4
Composition (at%)
Remarks
Pd
Cu
Ni
P
38.9 20.7 41.3 40.0 20.0 40.0 46.3 41.1 27.5 41.4 38.5
1.8 77.6 34.4 2.2 78.4 34.7 18.7 3.9 56.7 28.6 32.5
31.2 0.2 5.0 29.7 0.3 8.5 10.2 27.5 7.4 9.5 10.0
28.1 1.5 19.3 28.1 1.3 17.8 24.8 27.4 8.3 20.5 19.0
In Fig. 2(a)
In Fig. 2(b)
In Fig. 2(f)
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KIM et al.: OBSERVATION OF BEHAVIOR OF Pd–Cu–Ni–P ALLOY 4. DISCUSSION
Figure 5 shows the relationship between the experimental velocities of the solid–liquid interface and the annealing temperature together with crystal growth rates calculated on the basis of Turnbull’s approximation. The crystal growth rate can be expressed as [3, 7] uc ⫽
Fig. 4. Relationship between the final distance of movement of the solid–liquid interface and the annealing temperature, together with the three undercooling regimes classified on the basis of the morphology of the solidification front.
decrease in the dimensions of the microstructure occurs, as shown in Fig. 1(h), (i) and (j) and quantified in Fig. 4. Acicular phases with the irregular grain background, which were observed in the samples annealed at 708 K (⌬T is 96 K), are also seen in the interior of these samples, as given in Fig. 2(h) and (i). Considering the morphologies observed in the three undercooling regimes and their solidification behaviors, for which two or more solidification events are detected during isothermal annealing at various undercooling states, it is concluded that a phase separation takes place in the undercooled molten Pd40Cu30Ni10P20 alloy. The first solidification event can be identified as the primary solidification of nickel palladium phosphide competing with other crystals, such as copper-rich palladium compound and quaternary crystals, whereas the final event corresponds to a eutectic-like reaction of the above identified crystals or unidentified crystals. At annealing temperatures lower than 736 K (⌬T is 68 K), however, the primary solidification of nickel palladium phosphide disappears at the melt surface and in the outer regions of the samples. This phase separation in the undercooled molten Pd40Cu30Ni10P20 alloy is evident from the X-ray diffraction patterns taken from the cross-section of the isothermal annealed samples at various temperatures between TL and Tg, as shown in Fig. 3. It appears that the intensities of the Bragg peaks of nickel palladium phosphide and the copper-rich palladium compound decrease and the intensities of other Bragg peaks of unidentified crystals and binary phases, Pd15P2, Ni3P, etc., increase with an increase in the degree of undercooling. It is suggested that undercooling to the temperature region below the critical point leads to a reduction of heterogeneous nucleation and growth for the primary crystal in the undercooled melt. The remaining eutectic-like solidification leads to a microstructure that becomes more and more refined with increasing degree of undercooling.
冋 冉 冊册
fDg ⌬Gm 1⫺exp ⫺ , a0 RT
(1)
where Dg is a diffusion coefficient for atomic motion required for growth, a0 is the mean atomic or ionic diameter for the diffusive jump, ⌬Gm is the molar free energy change and f is the fraction of sites at the interface where growth occurs. For rough interfaces, as for closely packed crystal structures having low entropy of fusion, ⌬Sm ⫽ ⌬Hm/TL⬍2R, f⫽1 was chosen. Here, ⌬Hm and R are the molar heat of fusion and the ideal gas constant. The thermodynamic parameters of equation (1) in this work were set equal to the values used in Ref. [3] as given in Appendix A. The velocities of the solid–liquid interface, which were measured to be 10⫺5–10⫺7 m/s, tend to increase and have more scattered values with an increase in the degree of undercooling below 68 K, but decrease with an increase in the degree of undercooling between 68 K and 174 K. The experimental velocities of the solid–liquid interface are at least two orders of magnitude higher than the theoretical crystal growth rates. This underestimation of the calculated crystal growth rates suggests that there is an error in the parameters, e.g., Dg, f and ⌬Gfm, in equation (1). It is known that Dg was used with the assumption of D ⫽ Di ⫽ Dg in the Stokes–Einstein relation in order to simplify equation (1), where D is the bulk diffusion coefficient and Di is the diffusion coefficient for
Fig. 5. Relationship between the experimental velocities of the solid–liquid interface and the annealing temperature, together with the crystal growth rates calculated on the basis of Turnbull’s approximation in classical nucleation and growth theories [4–7].
KIM et al.: OBSERVATION OF BEHAVIOR OF Pd–Cu–Ni–P ALLOY
atomic transport across the nucleus–matrix interface. Turnbull has already pointed out that Dg is approximately two to three orders of magnitude larger than the corresponding D in Fe–Ni alloys [11]. Therefore the underestimation of the calculated crystal growth rate seems to be caused predominantly by the underestimation of Dg, although there are errors in both parameters. Figure 6 shows the time–temperature–transformation (TTT) diagram for the Pd40Cu30Ni10P20 alloy. In this diagram, the time to form a detectable fraction of primary crystals in the undercooled melt, indicated by the onset of crystal appearance, is shown as a function of temperature. In Fig. 6, curve 1, the so-called time–temperature–nucleation (TTN) diagram was derived from best fitting of the experimental onset times by using equation (2). Under isothermal conditions the number N of nucleation events in a melt of volume Vp formed at a constant undercooling after time t can be expressed by [12] IssVpt ⫽ N.
(2a)
Taking into account the expression for the steadystate nucleation rate Iss,
冉
Iss ⫽ AvNosn exp ⫺
冊
⌬G∗f(q) 16πs3 , ⌬G∗ ⫽ , RT 3⌬G2v
冉
冊
VpAvNosn 16πs3f(q) ⫹ , N 3RT⌬G2v
where Av is the area of nucleating substrate per unit volume of the melt, Nos is the number of molecules per unit area of substrate, f(q) is a wetting factor, n is the frequency of attempts by atoms to cross the phase boundary, s is interface energy, and ⌬Gv is the Gibbs free-energy difference between solid and melt per unit volume. According to equation (2b) a linear relation is expected, provided the temperature dependence of the prefactor AvNosn is neglected. This is a reasonable assumption for the temperature range below the melting temperature at not too large an undercooling [12]. Thus there can be an overestimation of the factor AvNosn at the low temperature above and in the vicinity of the glass temperature, because the temperature dependence of the factor AvNosn is much more pronounced than that of ⌬Gv. Continuous-cooling–transformation (CCT) diagrams have been constructed from the TTT diagrams by the method of Grange and Keifer [13] in order to evaluate GFA. Thus the critical cooling rate Rc was calculated for the cooling curve that just avoids interception of the nose of the CCT curve; i.e., Rc ⫽ (TL⫺Tn)/tn where Tn and tn are the temperature and the time at the nose, respectively. The CCT plot (curve 2 in Fig. 6) suggests that the critical cooling rate for glassy solidification is about 1.8 K/s, which is in agreement with previous experimental findings determined by calorimetric techniques in the corresponding condition [3].
5. CONCLUSIONS
Equation (2a) can be rewritten as ln t ⫽ ⫺ln
621
(2b)
Fig. 6. TTT curve (curve 1) and CCT curve (curve 2) for the Pd40Cu30Ni10P20 alloy constructed by using equation (2b), together with the onset for appearance of solid obtained from the in situ observation findings of solidification progress. The CCT diagram (curve 2) was constructed from the TTT curve (curve 1) following the method of Grange and Keifer [13].
In the undercooled melt of Pd40Cu30Ni10P20 alloy, the solidification behavior, including the nucleation and growth of crystals on a micrometer scale, has been observed in situ by use of a confocal scanning laser microscope combined with an infrared image furnace. The results obtained can be summarized as follows. 1. Depending on the degree of undercooling, the morphology of the solidification front changes among various types, such as faceted front, columnar dendritic, cellular grain and equiaxed grain, etc. 2. The velocities of the solid–liquid interface are measured to be 10⫺5–10⫺7 m/s, which are at least two orders of magnitude higher than the theoretical crystal growth rates. 3. Considering the morphologies observed in the three undercooling regimes and their solidification behaviors, we conclude that a phase separation takes place in the undercooled molten Pd40Cu30Ni10P20 alloy. 4. The continuous-cooling–transformation (CCT) diagram, which was derived from the experimental time–temperature–transformation (TTT) diagram, indicates a critical cooling rate for glassy solidifi-
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cation of about 1.8 K/s for this alloy, which agrees with previous calorimetric findings.
Dg ⫽ a20n ⫽
kT , 3πa0h
(A1)
REFERENCES 1. Chikama, H., Shibata, H., Emi, T. and Suzuki, M., Mater. Trans. JIM, 1996, 37(4), 620. 2. Nishiyama, N. and Inoue, A., Mater. Trans. JIM, 1997, 38, 464. 3. Nishiyama, N. and Inoue, A., Acta mater., 1999, 47(5), 1487. 4. Turnbull, D., Contemp. Phys., 1969, 10, 473. 5. Uhlmann, D. R., J. Non-Cryst. Solids, 1972, 7, 337. 6. Davies, H. A., Phys. Chem. Glasses, 1976, 17, 159. 7. Peker, A., Ph.D. Thesis, California Institute of Technology, 1994. 8. Donovan, P. E., Evans, P. V. and Greer, A. L., J. Mater. Sci. Lett., 1986, 5, 951. 9. Inoue, A. and Nishiyama, N., Mater. Sci. Eng., 1997, 401, 226. 10. Lu, I. R., Wilde, G., Grler, G. P. and Willnecker, R., J. Non-Cryst. Solids, 1999, 250-252, 577. 11. Turnbull, D., Trans. Metall. Soc. AIME, 1961, 221, 422. 12. Herlach, D. M., Mater. Sci. Eng., 1994, R12, 177. 13. Grange, R. A. and Keifer, J. M., Trans. Am. Soc. Met., 1941, 29, 85. 14. Fulcher, G. S., J. Am. Ceram. Soc., 1925, 6, 339. APPENDIX A
A.1. Thermodynamic parameters used in equation (1) 앫 Dg, the diffusion coefficient for atomic motion required for growth, is calculated from
where h ⫽ 9.34⫻10⫺3
冉
冊
4135 , (T⫺447)
(A2)
k and h are the Boltzmann constant and the viscosity, respectively. 앫 Equation (A2) is the temperature dependence of h for the undercooled Pd40Cu30Ni10P20 melt, which is derived from experimental results obtained by using a thermal-mechanical analyzer (TMA) on the assumptions of h ⫽ 1012 Pa s at Tg and h ⫽ 103 Pa s at Tm in a Fulcher-type equation [3, 7, 14]. 앫 a0 is the mean atomic or ionic diameter for the diffusive jump: 2.55⫻10⫺10 m. 앫 ⌬Gm is the molar free-energy change: ⌬Hm (TL⫺T)/TL. 앫 ⌬Hm is the heat of fusion: 4840 J/mol. 앫 TL is the melting temperature: 804 K. 앫 f is the fraction of sites at the interface where growth occurs: 1.
IN SITU OBSERVATION OF SOLIDIFICATION BEHAVIOR IN UNDERCOOLED Pd–Cu–Ni–P ALLOY BY USING A CONFOCAL SCANNING LASER MICROSCOPE J. -H. KIM1*, S. -G. KIM1 and A. INOUE1, 2 1
Inoue Superliquid Glass Project, ERATO, Japan Science and Technology Corporation (JST), Sendai 9820807, Japan and 2Institute for Materials Research, Tohoku University, Sendai 980-77, Japan ( Received 2 May 2000; received in revised form 16 October 2000; accepted 18 October 2000 )
Abstract—In the undercooled melt of Pd40Cu30Ni10P20 alloy, the solidification behavior including the nucleation and growth of crystals at the micrometer level has been observed in situ by use of a confocal scanning laser microscope combined with an infrared image furnace. The Pd40Cu30Ni10P20 alloy specimens were cooled from the liquid state to various undercooled states under a helium gas flow. Images of solidification progress were obtained by the charge-coupled device image sensor of the confocal scanning laser microscope. Depending on the degree of undercooling, the morphology of the solidification front changed among various types: faceted front, columnar dendritic front, cellular grain and equiaxed grain, etc. The velocities of the solid–liquid interface were measured to be 10⫺5–10⫺7 m/s, which are at least two orders of magnitude higher than the theoretical crystal growth rates. Combining the morphologies observed in the three undercooling regimes and their solidification behaviors, we conclude that phase separation takes place in the undercooled molten Pd40Cu30Ni10P20 alloy. The continuous-cooling–transformation (CCT) diagram was derived from the experimental time–temperature-transformation diagram constructed from solidification onset times under various isothermal annealing conditions. The CCT diagram suggests that the critical cooling rate for glassy solidification is about 1.8 K/s, which is in agreement with previous calorimetric findings. 2001 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. Keywords: Metallic glasses; Phase transformations; Microstructure; Kinetics; Dynamic phenomena
1. INTRODUCTION
In order to investigate directly the growth of crystals in situ in the undercooled liquid of bulk metallic glasses, it is necessary to overcome the effect of thermal radiation from the melts on microscope viewing and to distinguish between the two phases at elevated temperatures. A confocal scanning laser microscope (CSLM) was used in this study to overcome such difficulties for in situ observation of metallic melts. Actually, a Japanese research group has previously reported successful observation of the growth of crystals in Fe–C alloy melts by use of a confocal scanning laser microscope [1]. In addition, the inherent thermal stability of bulk metallic glasses against crystallization and oxidation, which are other important parameters along with the technical difficulties described above for the in situ observation of metallic melts, were considered. Multicomponent alloys recently discovered exhibit excellent glass-forming ability. The critical cooling rate for
* To whom all correspondence should be addressed. Fax: ⫹81-22-243-7616.
glass formation in Pd40Cu30Ni10P20 alloy in a B2O3 flux, for example, is as low as 0.1 K/s [2]. The thermal stability of these bulk metallic glasses against crystallization leads to an experimentally accessible time and temperature window to investigate the transformation behavior from the undercooled liquid into crystalline equilibrium phases. The aim of this paper is to present an in situ observation in a confocal scanning laser microscope of the solidification behavior and a microstructure analysis of Pd40Cu30Ni10P20 alloy solidified during different isothermal annealing treatments at various levels of undercooling. Further, the glass-forming ability (GFA) for this metallic glass is evaluated on the basis of the experimental findings and classical nucleation theory (CNT), and is compared with that obtained in previous reports [2, 3].
2. EXPERIMENTAL PROCEDURES
The master ingot was prepared by arc-melting a mixture of pure metallic Pd, Cu and Ni (purity ⱖ 99.99%) and pre-alloyed Pd60P40 in a purified argon atmosphere. Details of the preparation of bulk glassy
1359-6454/01/$20.00 2001 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. PII: S 1 3 5 9 - 6 4 5 4 ( 0 0 ) 0 0 3 5 3 - 0
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KIM et al.: OBSERVATION OF BEHAVIOR OF Pd–Cu–Ni–P ALLOY
samples have been reported previously [2]. In the undercooled melt of the Pd40Cu30Ni10P20 alloy, the solidification behavior including nucleation and growth of crystals on the micrometer level was observed in situ by use of a confocal scanning laser microscope combined with an infrared image furnace. A confocal scanning laser microscope (1LM21 from Lasertec Co.; with 0.2 µm resolution at 3500⫻ magnification) was employed with a 1.5 mW He–Ne laser of 632.8 nm wavelength. The laser beam was reflected and scanned by an acoustic optic deflector in the horizontal direction at a rate of 15.7 kHz and by a glavano mirror in the vertical direction at 60 Hz. The observation mechanism and schematic construction of this microscope have been given previously [1]. The Pd40Cu30Ni10P20 alloy specimens, which were cut into cubes 19±2 mg in weight and mirror polished, were melted in an alumina crucible and cooled from the liquid state at about 1.1TL to various undercooled states, as shown in Table 1, under a highpurity helium gas flow. Here, TL is the liquidus temperature, 804 K, of this alloy. The undercooled melt specimens were annealed isothermally at different undercooling regimes ⌬T ⫽ TL⫺Tiso ⫽ 39 to 173 K, where Tiso is the isothermal annealing temperature. The images of solidification progress during annealing were obtained by the charge-coupled device (CCD) image sensor, monitored on the cathode ray tube (CRT) of the confocal scanning laser microscope and recorded on videotape at a rate of 30 frames per second. The solidified specimens were quenched to room temperature after the isothermal annealing treatments. The velocities of the solid–liquid interface at the surface of the undercooled melt were derived from these observations. These experimental velocities were compared with theoretical crystal growth rates, which were calculated on the basis of Turnbull’s approximation and classical nucleation and growth theories [4–7]. In order to evaluate the GFA, the time–temperature–transformation (TTT) diagram was constructed from solidification onset times at which the primary phases were observed on the surface of undercooled melts under various annealing conditions.
The surface morphology and microstructures of solidified samples were observed by optical and scanning electron microscopy (SEM) (JXM-8621MX, Jeol Co.). The latter apparatus was equipped with wavelength-dispersive spectrometers (WDS) for determination of the element concentrations in various phases. Various phases in the solidified samples were analyzed with a micro-area X-ray diffractor (micro-area XRD) (M18XHF22, MAC Sci. Co.; Cu Kα radiation). 3. RESULTS
Figure 1 shows real-time solidification morphologies, taken from the surface of undercooled Pd40Cu30Ni10P20 alloy melts during isothermal annealing treatment at various levels of undercooling, obtained by using a confocal scanning laser microscope. Scanning electron micrographs of the cross-sections of the solidified samples are shown in Fig. 2. In Fig. 1, several kinds of morphology are seen in the undercooled Pd40Cu30Ni10P20 alloy, and they are dominated by the level of undercooling ⌬T. In the small undercooling regime—i.e., for ⌬Tⱕ 68 K, a complex microstructure consisting of two or three subnetworks is seen, as shown in Fig. 2(a), (b), (c) and (d). When the Pd40Cu30Ni10P20 alloy melt is annealed at 765 K (⌬T is 39 K), no obvious solidification front is seen on the surface of the undercooled melt except for the appearance of jag as shown in Fig. 1(a). In Fig. 2(a) and (a⬘) the microstructure of the cross-section of the sample is composed of two irregular phases (a1, a2) surrounded by a single matrix (a3). The matrix is estimated to be a glassy single phase or a glassy phase including fine crystals as evidenced by the existence of some Bragg peaks on the typical broad glassy X-ray pattern taken from the cross-sectional surface of the sample annealed at 765 K, as shown in Fig. 3(a). The approximate composition of all three phases was determined by WDS analysis on at least three different crystals in each case. The measured compositions are given in Table 2. According to Table 2, two irregular phases (a1, a2) are estimated to be nickel palladium phosphide and a copper-rich palladium compound including the other elements of the
Table 1. Conditions of isothermal annealing and morphologies of the solidification front in undercooled Pd40Cu30Ni10P20 alloy meltsa Isothermal annealing temperature, Tiso (K) 765 756 746 741 736 717 708 679 650 631 a
Degree of undercooling, ⌬T ⫽ 804⫺T⬑ (K) 39 48 58 63 68 87 96 125 154 173
Morphology of solidification on surface Jag Jag ⫹ faceted front Faceted front ⫹ jag Columnar dendrite Columnar dendrite Cellular grain Cellular grain Equiaxed grain Fine equiaxed grain Fine spot
Glass transition temperature, Tg, is 575 K and onset temperature for crystallization, Tx, is 670 K at a heating rate of 40 K/min.
Fig. 1. Real-time solidification morphologies, taken from the surface of undercooled Pd40Cu30Ni10P20 alloy melts during isothermal annealing treatment at various levels of undercooling, obtained by using a confocal scanning laser microscope. The degrees of undercooling were: (a) 39 K; (b) 48 K; (c) 58 K; (d) 63 K; (e) 68 K; (f) 87 K; (g) 96 K; (h) 125 K; (i) 154 K; (j) 173 K.
KIM et al.: OBSERVATION OF BEHAVIOR OF Pd–Cu–Ni–P ALLOY 617
Fig. 2. Scanning electron and optical micrographs of cross-sections of the Pd40Cu30Ni10P20 alloy samples solidified during isothermal annealing treatment at various levels of undercooling. The degrees of undercooling were: (a), (a⬘) 39 K; (b) 48 K; (c) 58 K; (d), (d⬘) 63 K; (e), (f) 96 K; (g), (h) 125 K; (i), (j) 154 K. Here, (a⬘) and (d⬘) are the microstructures obtained by using an optical microscope with a magnification of less than 200⫻. (f), (h) and (j) are the microstructures observed in the center region of cross-sections of the samples.
618 KIM et al.: OBSERVATION OF BEHAVIOR OF Pd–Cu–Ni–P ALLOY
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619
shown in Fig. 2(b), (c) and (d). In Fig. 2(b) and (d), one can see light gray faceted dendritic (b3) and white coarse (b4) phases, which are identified to be copperrich and palladium-rich quaternary phases by the WDS analysis measurement given in Table 2. In the intermediate undercooling regime—i.e., for 68 K⬍⌬T⬍125 K, growth and coarsening of grains which have directional solidification fronts on a submillimeter level are seen as shown in Fig. 1(f) and (g). A coarse irregular grain, approximately 100 µm wide, is seen together with other smaller grains in the outer region of the cross-sectional surface as shown in Fig. 2(e). In the interior of the samples, however, acicular phases (c1) surrounded by another black phase (c2), which are located at the center of each irregular grain background (c3, c4), are seen in Fig. 2(f). The acicular phase (c1) is estimated to be nickel palladium phosphide, including copper element below 4 at% (Table 2), as evidenced by the existence of Bragg peaks of Pd2Ni2P on the X-ray patterns in Fig. 3(c). The composition of the irregular grain background is similar to that of the master ingot as given in Table 2. It appears that the solidification progress of the interior crystals is different from that of the grains near the surface of the sample. If a thermal gradient exists in the sample during solidification and the undercooled melt has activated nuclei, the interior primary crystals can originate from the quenched interior nuclei and lead to solidification of the irregular grains at a slightly higher temperature than the surface temperature. On the other hand, grain growth near the surface of the sample may originate predominantly from the surface. Figure 4 shows the relationship between the final distance of movement of the solid–liquid interface after solidification and the annealing temperature, together with three undercooling regimes classified on the basis of the morphology of the solidification front. In the large undercooling regime, which is just above and below the crystallization onset temperature (670 K [2]) for ⌬Tⱖ125 K, spherical or equiaxed polyhedral grain morphologies, which arise from nucleation and growth, are identified and a sharp
Fig. 3. X-ray diffraction patterns taken from the cross-sections of the Pd40Cu30Ni10P20 alloy samples solidified during isothermal annealing treatment at various levels of undercooling. The degrees of undercooling were: (a) 39 K; (b) 58 K; (c) 96 K; (d) 125 K; (e) 154 K.
alloy below 2 at% level. From the compositional point of view, the nickel palladium phosphide identified in the present work is different from the ternary primary phase Ni45Pd34P21, which was described by Donovan et al. in a crystallized Pd40Ni40P20 glass [8]. Recently, crystallized phases in the Pd40Cu30Ni10P20 alloy have been identified as Ni3P, Pd15P2, Pd2Ni2P and Cu3Pd [9, 10]. In Fig. 3, the Bragg peaks of Pd2Ni2P and Cu3Pd are seen on the X-ray patterns taken from the cross-sectional surface of the annealed samples in the present work, along with shifting of the Bragg peaks. In the cases of the samples annealed at 756 K (⌬T is 48 K) and 746 K (⌬T is 58 K), faceted solidification fronts are observed with networks of complex jags on these surfaces as shown in Fig. 1(b) and (c). For the samples annealed at 741 K (⌬T is 63 K) and 736 K (⌬T is 68 K), solidification fronts of columnar dendritic shape are observed as shown in Fig. 1(d) and (e). The microstructures are composed typically of a metallic compound and/or a cellular dendritic precipitation with a eutectic-like background as
Table 2. Measured compositions of the solidified phases in the Pd40Cu30Ni10P20 alloy Isothermal annealing temperature (degree of undercooling)
765 K (39 K)
756 K (48 K)
708 K (96 K)
Phase a1 Phase a2 Matrix a3 Phase b1 Phase b2 Phase b3 Phase b4 Phase c1 Phase c2 Phase c3 Phase c4
Composition (at%)
Remarks
Pd
Cu
Ni
P
38.9 20.7 41.3 40.0 20.0 40.0 46.3 41.1 27.5 41.4 38.5
1.8 77.6 34.4 2.2 78.4 34.7 18.7 3.9 56.7 28.6 32.5
31.2 0.2 5.0 29.7 0.3 8.5 10.2 27.5 7.4 9.5 10.0
28.1 1.5 19.3 28.1 1.3 17.8 24.8 27.4 8.3 20.5 19.0
In Fig. 2(a)
In Fig. 2(b)
In Fig. 2(f)
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Figure 5 shows the relationship between the experimental velocities of the solid–liquid interface and the annealing temperature together with crystal growth rates calculated on the basis of Turnbull’s approximation. The crystal growth rate can be expressed as [3, 7] uc ⫽
Fig. 4. Relationship between the final distance of movement of the solid–liquid interface and the annealing temperature, together with the three undercooling regimes classified on the basis of the morphology of the solidification front.
decrease in the dimensions of the microstructure occurs, as shown in Fig. 1(h), (i) and (j) and quantified in Fig. 4. Acicular phases with the irregular grain background, which were observed in the samples annealed at 708 K (⌬T is 96 K), are also seen in the interior of these samples, as given in Fig. 2(h) and (i). Considering the morphologies observed in the three undercooling regimes and their solidification behaviors, for which two or more solidification events are detected during isothermal annealing at various undercooling states, it is concluded that a phase separation takes place in the undercooled molten Pd40Cu30Ni10P20 alloy. The first solidification event can be identified as the primary solidification of nickel palladium phosphide competing with other crystals, such as copper-rich palladium compound and quaternary crystals, whereas the final event corresponds to a eutectic-like reaction of the above identified crystals or unidentified crystals. At annealing temperatures lower than 736 K (⌬T is 68 K), however, the primary solidification of nickel palladium phosphide disappears at the melt surface and in the outer regions of the samples. This phase separation in the undercooled molten Pd40Cu30Ni10P20 alloy is evident from the X-ray diffraction patterns taken from the cross-section of the isothermal annealed samples at various temperatures between TL and Tg, as shown in Fig. 3. It appears that the intensities of the Bragg peaks of nickel palladium phosphide and the copper-rich palladium compound decrease and the intensities of other Bragg peaks of unidentified crystals and binary phases, Pd15P2, Ni3P, etc., increase with an increase in the degree of undercooling. It is suggested that undercooling to the temperature region below the critical point leads to a reduction of heterogeneous nucleation and growth for the primary crystal in the undercooled melt. The remaining eutectic-like solidification leads to a microstructure that becomes more and more refined with increasing degree of undercooling.
冋 冉 冊册
fDg ⌬Gm 1⫺exp ⫺ , a0 RT
(1)
where Dg is a diffusion coefficient for atomic motion required for growth, a0 is the mean atomic or ionic diameter for the diffusive jump, ⌬Gm is the molar free energy change and f is the fraction of sites at the interface where growth occurs. For rough interfaces, as for closely packed crystal structures having low entropy of fusion, ⌬Sm ⫽ ⌬Hm/TL⬍2R, f⫽1 was chosen. Here, ⌬Hm and R are the molar heat of fusion and the ideal gas constant. The thermodynamic parameters of equation (1) in this work were set equal to the values used in Ref. [3] as given in Appendix A. The velocities of the solid–liquid interface, which were measured to be 10⫺5–10⫺7 m/s, tend to increase and have more scattered values with an increase in the degree of undercooling below 68 K, but decrease with an increase in the degree of undercooling between 68 K and 174 K. The experimental velocities of the solid–liquid interface are at least two orders of magnitude higher than the theoretical crystal growth rates. This underestimation of the calculated crystal growth rates suggests that there is an error in the parameters, e.g., Dg, f and ⌬Gfm, in equation (1). It is known that Dg was used with the assumption of D ⫽ Di ⫽ Dg in the Stokes–Einstein relation in order to simplify equation (1), where D is the bulk diffusion coefficient and Di is the diffusion coefficient for
Fig. 5. Relationship between the experimental velocities of the solid–liquid interface and the annealing temperature, together with the crystal growth rates calculated on the basis of Turnbull’s approximation in classical nucleation and growth theories [4–7].
KIM et al.: OBSERVATION OF BEHAVIOR OF Pd–Cu–Ni–P ALLOY
atomic transport across the nucleus–matrix interface. Turnbull has already pointed out that Dg is approximately two to three orders of magnitude larger than the corresponding D in Fe–Ni alloys [11]. Therefore the underestimation of the calculated crystal growth rate seems to be caused predominantly by the underestimation of Dg, although there are errors in both parameters. Figure 6 shows the time–temperature–transformation (TTT) diagram for the Pd40Cu30Ni10P20 alloy. In this diagram, the time to form a detectable fraction of primary crystals in the undercooled melt, indicated by the onset of crystal appearance, is shown as a function of temperature. In Fig. 6, curve 1, the so-called time–temperature–nucleation (TTN) diagram was derived from best fitting of the experimental onset times by using equation (2). Under isothermal conditions the number N of nucleation events in a melt of volume Vp formed at a constant undercooling after time t can be expressed by [12] IssVpt ⫽ N.
(2a)
Taking into account the expression for the steadystate nucleation rate Iss,
冉
Iss ⫽ AvNosn exp ⫺
冊
⌬G∗f(q) 16πs3 , ⌬G∗ ⫽ , RT 3⌬G2v
冉
冊
VpAvNosn 16πs3f(q) ⫹ , N 3RT⌬G2v
where Av is the area of nucleating substrate per unit volume of the melt, Nos is the number of molecules per unit area of substrate, f(q) is a wetting factor, n is the frequency of attempts by atoms to cross the phase boundary, s is interface energy, and ⌬Gv is the Gibbs free-energy difference between solid and melt per unit volume. According to equation (2b) a linear relation is expected, provided the temperature dependence of the prefactor AvNosn is neglected. This is a reasonable assumption for the temperature range below the melting temperature at not too large an undercooling [12]. Thus there can be an overestimation of the factor AvNosn at the low temperature above and in the vicinity of the glass temperature, because the temperature dependence of the factor AvNosn is much more pronounced than that of ⌬Gv. Continuous-cooling–transformation (CCT) diagrams have been constructed from the TTT diagrams by the method of Grange and Keifer [13] in order to evaluate GFA. Thus the critical cooling rate Rc was calculated for the cooling curve that just avoids interception of the nose of the CCT curve; i.e., Rc ⫽ (TL⫺Tn)/tn where Tn and tn are the temperature and the time at the nose, respectively. The CCT plot (curve 2 in Fig. 6) suggests that the critical cooling rate for glassy solidification is about 1.8 K/s, which is in agreement with previous experimental findings determined by calorimetric techniques in the corresponding condition [3].
5. CONCLUSIONS
Equation (2a) can be rewritten as ln t ⫽ ⫺ln
621
(2b)
Fig. 6. TTT curve (curve 1) and CCT curve (curve 2) for the Pd40Cu30Ni10P20 alloy constructed by using equation (2b), together with the onset for appearance of solid obtained from the in situ observation findings of solidification progress. The CCT diagram (curve 2) was constructed from the TTT curve (curve 1) following the method of Grange and Keifer [13].
In the undercooled melt of Pd40Cu30Ni10P20 alloy, the solidification behavior, including the nucleation and growth of crystals on a micrometer scale, has been observed in situ by use of a confocal scanning laser microscope combined with an infrared image furnace. The results obtained can be summarized as follows. 1. Depending on the degree of undercooling, the morphology of the solidification front changes among various types, such as faceted front, columnar dendritic, cellular grain and equiaxed grain, etc. 2. The velocities of the solid–liquid interface are measured to be 10⫺5–10⫺7 m/s, which are at least two orders of magnitude higher than the theoretical crystal growth rates. 3. Considering the morphologies observed in the three undercooling regimes and their solidification behaviors, we conclude that a phase separation takes place in the undercooled molten Pd40Cu30Ni10P20 alloy. 4. The continuous-cooling–transformation (CCT) diagram, which was derived from the experimental time–temperature–transformation (TTT) diagram, indicates a critical cooling rate for glassy solidifi-
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KIM et al.: OBSERVATION OF BEHAVIOR OF Pd–Cu–Ni–P ALLOY
cation of about 1.8 K/s for this alloy, which agrees with previous calorimetric findings.
Dg ⫽ a20n ⫽
kT , 3πa0h
(A1)
REFERENCES 1. Chikama, H., Shibata, H., Emi, T. and Suzuki, M., Mater. Trans. JIM, 1996, 37(4), 620. 2. Nishiyama, N. and Inoue, A., Mater. Trans. JIM, 1997, 38, 464. 3. Nishiyama, N. and Inoue, A., Acta mater., 1999, 47(5), 1487. 4. Turnbull, D., Contemp. Phys., 1969, 10, 473. 5. Uhlmann, D. R., J. Non-Cryst. Solids, 1972, 7, 337. 6. Davies, H. A., Phys. Chem. Glasses, 1976, 17, 159. 7. Peker, A., Ph.D. Thesis, California Institute of Technology, 1994. 8. Donovan, P. E., Evans, P. V. and Greer, A. L., J. Mater. Sci. Lett., 1986, 5, 951. 9. Inoue, A. and Nishiyama, N., Mater. Sci. Eng., 1997, 401, 226. 10. Lu, I. R., Wilde, G., Grler, G. P. and Willnecker, R., J. Non-Cryst. Solids, 1999, 250-252, 577. 11. Turnbull, D., Trans. Metall. Soc. AIME, 1961, 221, 422. 12. Herlach, D. M., Mater. Sci. Eng., 1994, R12, 177. 13. Grange, R. A. and Keifer, J. M., Trans. Am. Soc. Met., 1941, 29, 85. 14. Fulcher, G. S., J. Am. Ceram. Soc., 1925, 6, 339. APPENDIX A
A.1. Thermodynamic parameters used in equation (1) 앫 Dg, the diffusion coefficient for atomic motion required for growth, is calculated from
where h ⫽ 9.34⫻10⫺3
冉
冊
4135 , (T⫺447)
(A2)
k and h are the Boltzmann constant and the viscosity, respectively. 앫 Equation (A2) is the temperature dependence of h for the undercooled Pd40Cu30Ni10P20 melt, which is derived from experimental results obtained by using a thermal-mechanical analyzer (TMA) on the assumptions of h ⫽ 1012 Pa s at Tg and h ⫽ 103 Pa s at Tm in a Fulcher-type equation [3, 7, 14]. 앫 a0 is the mean atomic or ionic diameter for the diffusive jump: 2.55⫻10⫺10 m. 앫 ⌬Gm is the molar free-energy change: ⌬Hm (TL⫺T)/TL. 앫 ⌬Hm is the heat of fusion: 4840 J/mol. 앫 TL is the melting temperature: 804 K. 앫 f is the fraction of sites at the interface where growth occurs: 1.