- Email: [email protected]
Rapid solidification of cobalt-titanium alloys induced by nanosecond laser pulses
Materials Science and Engineering, A179/A 180 (1994) 243-248
243
Rapid solidification of cobalt-titanium alloys induced by nanosecond laser pulses S. Vitta Department of Metallurgical Engineering, Indian Institute of Technology, Powai, Bombay 400076 (India) A. L. Greer and R. E. Somekh Department of Materials Science and Metallurgy, University of Cambridge, Pembroke Street, Cambridge CB2 3QZ (UK)
Abstract Phase formation is studied in thin films of Co-Ti alloys, both as sputtered and after quenching following melting by a 5 ns laser pulse. The composition ranges for glass formation are established, and the significance of crystal growth kinetics is examined. The C o - T i system exhibits many types of intermetallic compounds and a wide variety of behaviours is found around their compositions.
1. Introduction
Partitionless crystal growth in an undercooled liquid, as may arise in congruent freezing or with complete solute trapping, occurs at a rate determined by thermal diffusion and by the kinetics of the interfacial processes. The focus of this study is the kinetics which may be diffusion limited or collision limited [1]. In the former case (expected for compounds and ordered solid solutions), even in the absence of solute partitioning, the migration of the crystal-liquid interface requires diffusive-type atomic jumps to effect changes in chemical order; the maximum growth rate (i.e. if not limited by thermal diffusion) is D/it (where D is the atomic diffusivity in the liquid and it the atomic diameter), i.e. about 10 m s- ~. In the latter case (expected for pure metals and disordered solid solutions), diffusivetype jumps are not required and the growth rate is limited only by the collisional frequency; the maximum rate is the speed of sound in the liquid, typically about 4000 m s-t. The distinction between the two regimes is important in rapid solidification, particularly concerning glass formation, which would be expected to be impossible by rapid liquid quenching if collisionlimited growth were to operate. The C o - T i system is suitable for the study of glass formation at compound-forming compositions as it has compounds exhibiting a variety of homogeneity ranges and structures [2]. Glass formation has been reported at 21-23 at.% Ti and about 77 at.% Ti by melt spinning [3], and at 25 at.% Ti and 80 at.% Ti by 30 ns 0921-5093/94/$7.00 SSDI 0921-5093(93)05520-Y
pulse laser quenching [4]. However, Co-Ti exhibits a lower glass-forming ability than for example Nb-Ti in which laser quenching gave only amorphous phases at compound compositions [5]. Of melt quenching methods, pulsed lasers (10 ps,10 ns) permit the most rapid solidification: cooling rates of 101°-1013 K s-1 and isotherm velocities of about 100 m s-l. At such rates solute partitioning is not possible [6], resulting in either a crystalline phase of the same composition as the melt or a glass. While measurement of crystal regrowth velocity is possible after laser melting (for example by monitoring the surface reflectivity or resistivity of the sample during and immediately after the laser pulse [7, 8], most studies of alloys, as here, have relied on microstructural characterization after melting and solidification. In completely crystalline films, the grain size and the total solidification time give a lower bound estimate of growth velocity; for equiaxed structures, the average grain radius is taken as the growth distance; for growth from surrounding unmelted material the distance is taken as the length of the columnar grains. When isolated crystals are found in a glassy matrix, the crystal radius gives the actual growth velocity, albeit only approximately. 2. Experimental methods
The experimental design followed closely that of Lin and Spaepen [5]. As shown in Fig. l(b), a copper substrate was coated with aluminium (about 1 /zm) and © 1994 - Elsevier Sequoia. All rights reserved
S. Vitta et aL
244
/Rapid sofidification of Co-Ti alloys
then the alloy of interest (less than 100 nm, to permit transmission electron microscopy (TEM) characterization without thinning). After laser irradiation, the alloy film was floated off the substrate by dissolution of the underlying aluminium in sodium hydroxide solution. In contrast to ref. 5, the as-deposited alloy films were homogeneous, not layered; they were crystalline or amorphous, depending on composition, but always single phase. The Co-Ti films were deposited by ultrahigh vacuum magnetron sputtering in argon [9]. Elemental targets (Co, less than 25 ppm metal impurity; Ti, 99.7% minimum purity) were used with the substrates rotated underneath; the films correspondingly had "layers" about 0.3 nm thick and are therefore considered homogeneous. The substrates were at floating temperature, estimated to be below 200°C. An Nd-YAG Q-switched mode-locked laser was used, giving 5 ns, 2 m J, pulses after frequency quadrupling to 266 nm wavelength. The beam was focused to 100/zm diameter on the specimen, and had an approximately annular (but very inhomogeneous) intensity distribution (Fig. l(a)). Laser energy was adjusted to give small holes in the deposited films, ensuring melting around the holes. Structural characterization was by TEM (Philips EM300), and compositions were determined (from simultaneous deposits on to silicon substrates) by energy-dispersive spectroscopy. Detailed thermal modelling is ruled out by the homogeneity of the laser beam used in this work. Instead, thermal parameters are estimated using the dimensional arguments presented by Spaepen and Lin
[
(i)
(Ji)
copper
substrote
(iii)
1 lam
Fig. 1. Schematic diagram of a laser-quenched sample: (a) the laser-irradiated area, showing the annular intensity distribution; (b) a section through the sample, showing the material removal by the laser pulse and the melt depth. Three characteristic regions of the Co-Ti thin film are marked. The different horizontal and vertical scales in (b) should be noted.
[10]. The melt depth d varies as shown in Fig. l(b) from about 1/zm (where the radiation was most intense, but with the melting not quite reaching the copper substrate) to about 100 nm (near the edge of the irradiated area, where only the top sputtered film was melted). The temperature gradient, [V I ~ Tm/d ( w h e r e T m is the congruent alloy melting temperature) is then calculated to be in the range 109-10 TM K m - I. The cooling rate 7"= D,hlv r I/d, (where Dth is the thermal diffusivity of the solid alloy) is 101°-10 ~2 K s-l. The total solidification time r ~ TmT is 10- 7_ 10 - 9 s, and the isotherm velocity Uth= i/~/I V T[ is 10-100 m s ' At the edges of the irradiated area there is an unmelted, but heat-affected, zone. In the part of the Co-Ti samples melted and resolidified, three areas may be distinguished (Fig. l(b)). In region (i) at the edge of the hole, the thickness of the alloy layer is reduced. The cooling rate may therefore be somewhat higher than in region (ii) which has the lowest cooling rate (i. e. about 101° K s-l). Region (iii), near the edge of the irradiated area, has a much lower melt depth and the highest quenching rate (i.e. about 10 ~2 K s-~). As discussed below, the quenched microstructures are often different in the three regions.
3. Results
Figure 2 shows the Co-Ti phase diagram and summarizes the results. As-deposited Co-Ti films were all single phase. The samples with 26.8 at.% Ti and 76.7 at.% Ti were completely amorphous, as were all samples of intermediate composition. Outside this range, solid solutions were found. As-deposited cobalt is c.c.p. (high temperature allotrope, r-Co) with a grain size of 50 nm. Addition of 4.5 at.% Ti and 9.3 at.% Ti reduces the grain size to 15 nm without changing the crystal structure. As-deposited titanium is h.c.p. (low temperature allotrope, a-Ti) with a grain size of 50 nm. Addition of 5.4 at.% Co changes the structure to c.c.p. and reduces the grain size to 10 nm. After laser quenching only the 26.8 at.% Ti alloy was completely amorphous. The structures of the other samples are described in order of increasing Ti content. Some representative bright field TEM micrographs are in Fig. 3. Laser-quenched cobalt is c.c.p. with a grain size of about 250 nm, much larger than in the as-deposited films which have the same structure. With 4.5 at.% Ti the structure is still c.c.p, with a similar grain size, but there is a high density of stacking faults. As further titanium is added (19.3 at.%) the faults become less evident, but the grain size decreases towards 0.1 /zm (Fig. 3(a)). With 40.3 at.% Ti, 43.7 at.% Ti and 52.6 at.% Ti (Figs. 3(b), 3(c) and 3(d) respectively) the CoTi phase appears, confirmed by the
S. Vitta etaL
/Rapidsolidification of Co-Ti alloys
Weight Percent Titanium 0
18001
10
20
'
'
30
'
50
'
'
60
'
70 '
80
'
90
'
tOO
I
liquid
t_~ 1400
2
40
12oo
E ,ooo p,\ /
o
to
Co
o.-d.pos ,.df tose.qoeochedl
ocp ~cp
I io=l
L J __
P-Ti I
20
30 4o 50 6o 7o 80 Atomic Percent Titanium
°
b / .
\
CoTi
9o
~oo Ti
S cop ScoTi2
+,cp
~>5+bc~.
Fig. 2. The Co-Ti phase diagram (after Murray [2]), showingthe T0 line for the c.c.p, phase (------). The various equilibrium compounds are indicated, includingthe Laves phases with structures C15 and C36. The phase mixtures found in the present work are indicated; the curved lines indicate uncertain boundaries. The arrows at the top of the diagram indicate the compositions tested. presence of superlattice reflections. Throughout the composition range the CoTi is the only phase in regions (iii) and (ii), but there is some amorphous phase in region (i). The grain size is about 40 nm in region (iii) (iii) (e.g. Figs. 3(b) and 3(d)) but much larger in region (ii) (e.g. Fig. 3(c)), ranging from about 1.0/zm in the middle of the composition range (43.7 at.% Ti) to about 0.2/zm at the edges. With 70.2 at.% Ti and 76.7 at.%Ti (Figs. 3(e) and 3(f) respectively), the structure is amorphous in region (i), is a mixture of amorphous phase and the intermetallic compound CoTi: in region (ii) and is a mixture of a-Ti and CoTi: in region (iii). In samples with 70.2 at.% Ti in regions (i) and (ii), the CoTi: phase appears as isolated roughly cylindrical grains in an amorphous matrix (e.g. Fig. 3(e)). In samples with 76.7 at.% Ti, the fine crystalline structure might have been formed by devitrification of an intermediate glassy phase (Fig. 3(f)). Laser-quenched titanium is b.c.c, with a grain size of 150 nm. The quenching is sufficiently rapid to retain at room temperature the high temperature allotrope,/5-Ti, formed from the melt. With 5.4 at.% Co added, the c.c.p, phase of the as-deposited film is retained after quenching, with a much increased grain size of about 0.2/~m. The equiaxed grains have only a low density of defects.
4. D i s c u s s i o n
In the laser-quenched films there is little occurrence of crystal regrowth from surrounding unmelted mate-
245
rial, even for compositions which are expected to display rapid growth. Pure titanium, for example, solidifies as/5-Ti formed by nucleation and growth in the melt, rather than as a-Ti regrown from the surrounding unmelted a-Ti. There is no evidence for mixing during melting of the Co-Ti layers and the underlying aluminium. There is very limited time for mixing and the alumina layer on the aluminium may act as a barrier. Fortunately for the microstructural studies, the alumina seems not to be an active nucleant for the crystalline phases in Co-Ti. The low areal density of nucleation events leads to large grain sizes and permits determination of crystal growth velocities in some cases. The comparative lack of influence of substrate and surrounding unmelted material facilitates the determination of intrinsic behaviour in this system, as discussed further below.
4.1. The Co3Ti phase This phase is not observed in laser-quenched samples; indeed, a completely amorphous phase is formed at 26.8 at.% Ti, close to the ideal stoichiometry. At lower titanium content the amorphous phase is still found, now in combination with the extended solid solution based on fl-Co. It is significant that in this composition range, in marked contrast to the behaviour elsewhere, a greater degree of amorphicity is found on the laser-quenched samples than in the asdeposited samples. This agrees with the results of melt spinning (cooling rate, 105-106 K S -1) which causes glass formation at similar compositions (21-23 at.% Ti) [3]. There seems to be a particular resistance to crystal nucleation and growth in the vicinity of the Co3Ti composition, remarkable since the phase has the relatively simple Au3Cu structure (ordered c.c.p.). If solidification is too rapid for formation of chemically ordered Co3Ti, then the chemically disordered form might be expected. (For an analysis of disorder trapping, see ref. 11). The undercooling necessary for this to form is easily estimated by extrapolating the TO line for the c.c.p, r - C o solid solution (broken line on Fig. 2). Partitionless solidification to the disordered Co3Ti solid solution should be possible at a moderate undercooling of about 50 K, yet even this is not observed. It may be that the liquid is particularly stable at this composition, perhaps because of the formation of associates; Murray [2] has commented on the particular difficulties of modelling the Co3Ti-liquid equilibrium. However, the main reason for the non-appearance of Co3Ti must be kinetic. This is found also in devitrification experiments. When Inoue et al. [3] annealed amorphous Co-22 at.% Ti samples they found that the first crystalline phase was the Laves phase at 34 at.% Ti, and not Co3Ti; the selection of a phase with a composition substantially different from that of the amorphous alloy must be a consequence of a
S. Vitta et al.
246
/Rapid solidification of Co- Ti alloys
"" .,,.4
::,k'
%
~,
~"
"
..
"?
~"
b
O
,;
•
4,
*/
~lr
4¢'* ,-
Q
C
q
0~5 um
.: :.:," :=:
f
O.25pm|
D
J
I
~
') t
f
t :
-
..
'
'~:~"
~ "
"
1.0IJm
Fig. 3. Bright field transmission electron micrographs, showing representative structures in Co-T] tllmS after laser quenching: (a) 19.3 at.% Ti; (b) 40.3 at.% Ti; (c) 43.7 at.%; (d) 52.6 at.% Ti; (e) 70.2 at.% Ti; (f) 76.7 at.% Ti. substantial kinetic barrier to nucleation or growth of Co3Ti.
4.2. Laves phases Two Laves phases appear in equilibrium in the C o - T i system at 31 at.% and 34 at.% Ti [2]. In equilibrium the phase fields for these compounds are only 0.5 at.% wide. Although the Laves phases at 34 at.% Ti can be obtained by devitrification [3], there was no evidence for either phase in the present work. The
laser-quenched compositions were not very close to those of the Laves phases, so their formation may not be expected; at least it is demonstrated that these phases do not substantially enlarge their composition ranges on rapid quenching.
4.3. Growth of the Co Ti phase This phase has the relatively simple CsCl structure and is observed over a wide composition range (40.3 at.%-52.6 at.% Ti) in laser-quenched samples. The
S. Vittaet al. /Rapidsolidification of Co-Ti alloys compound has previously been reported at compositions as far from stoichiometry as 43.9 at.% and 52 at.% Ti [2]. In the present work the grain size is largest at 43.7 at.% Ti, closest to the centre of the homogeneity range in the equilibrium diagram. Nucleation is expected to be easier at this composition, occurring at lesser undercooling and permitting rapid grain growth. Estimating from the grain size in Fig. 3(c), the growth velocity is about 10 m s- 1. 4.4. Growth of the CoTi phase This phase has a complex structure (of Fe3W3Ctype, E93) with 96 atoms in the unit cell [2], and is therefore surprising that it forms in rapid laser quenching. The isolated crystals in region (ii) of the 70.2 at.% Ti sample (Fig. 3(e)) enable the growth rate to be estimated as 4 m s- 1. This is fast for a complex structure, but within the diffusion-limited regime. The CoTi 2 phase is found also in laser-quenched Co-76.7 at.% Ti. Although not tested directly, it seems likely that the phase has formed without partitioning. Its composition then reflects a considerable extension of its homogeneity range. Although a range has been reported [2], such a large extension has not previously been observed. 4.5. Amorphous phase formation Some previous work on laser quenching, for example that on Nb-Ni using 30 ps pulses [5], suggested that the composition ranges for glass formation are similar for sputter deposition and for ultrarapid liquid quenching. The present work does not bear this out, as the glassy or amorphous phase is much less prevalent after quenching. The present results are, however, consistent with earlier work using melt spinning [3] and 30 ns laser pulses [4]. Comparison of the phase selection in as-deposited and quenched samples shows that the key difference is the presence, or not, of intermetallic compounds. No compounds are found in the as-deposited films, although the CoTi and C o T i 2 compounds appear to nucleate and grow readily during laser quenching. This suggests that the nucleation frequency for intermetallic compounds must be particularly temperature dependent, being significant in undercooled liquids during quenching, but essentially zero at the effective surface temperature at thin film deposition. It may also be that the chemical order in the melt aids compound formation. These effects are related to the diffusion limitation of the growth of the compounds. Similar considerations may apply in solid state amorphization by interfacial reaction; that the amorphous phase forms at all suggests a difficulty in nucleating intermetallic compounds at low annealing temperatures. As already discussed in Section 4.1, special considerations appear to apply for the Co3Ti which is difficult to form at any temperature.
247
Away from the range of compound formation, glass formability is determined by the ease of formation of the terminal solid solutions. For these conditions, Egami and Waseda [12] have proposed a criterion for glass formation based on the influence of atomic size difference on the topographical stability of terminal solid solutions. Taking their 2 parameter to have a value in the range 0.035-0.045 as found for the other systems, the glass-forming range in Co-Ti is predicted to be from 21.5-27.6 at.% Ti to 81.5-85.5 at.% Ti. The competition between the formation of glass and the terminal solid solution has been investigated in detail only on the Co-rich side in the present work. Complete amorphous phase formation is found for Ti contents greater than about 23 at.% in both as-deposited and laser-quenched samples. This limit of the glass formation agrees well with the Egami and Waseda prediction and with the results of melt spinning. It appears that when the competition is with the formation of a terminal solid solution, the limit of glass formability is not strongly dependent on the details of the quenching procedure; even ultrarapid quenching using a pulsed laser or thin film deposition do not significantly extend glass formation beyond the range achievable by melt spinning. Presumably the terminal solid solution can exhibit partitionless collision-limited growth, and in this regime the effect of solute content on thermodynamic parameters is dominant.
5. Conclusions
Apart from the specific findings of phase formation at particular compositions in the Co-Ti system, some general conclusions can be drawn. Pulsed laser quenching of alloys is effective in forming glasses and other metastable microstructures. The solidification is partitionless. When the competition is between glass formation and the formation of a terminal solid solution, the transition to collision-limited growth of the solid solution sets a sharp boundary to glass formation; boundary of the glass-forming composition range (for partitionless freezing) is not very dependent on details of the quenching method. On the contrary, at the compositions of intermetallic compounds where growth is diffusion limited, glass formation is dependent on the quenching method. In some cases it is possible to estimate growth velocities for compounds from the quenched microstructures; they are typically a few metres per second. Some compounds which form readily from the liquid, even on ultrarapid quenching, do not form in thin film deposition; this difficulty of formation at low temperatures may be relevant for solid state amorphization processes. The ease of formation of different compounds does not exhibit a straight-
248
S. Vitta et al.
/Rapid solidification of Co- Ti alloys
forward correlation with their structural complexity; in the present case, C o T i 2 (E93 structure, 96 atoms in unit cell) forms readily, but disordered Co3Ti (c.c.p.) does not.
Acknowledgments S.V. is grateful for the award of a C a m b r i d g e N e h r u Studentship. Deposition facilities s u p p o r t e d by the Science and Engineering Research Council (UK) are acknowledged.
References 1 S.R. Coriell and D. Turnbull, Acta Metall., 30 (1982) 2135. 2 J. L. Murray, Bull. Alloy Phase Diag., 3 (1982) 74.
3 A. Inoue, K. Kobayashi, C. Suryanarayana and T. Masumoto, Scr. Metall., 14 (1980) 119. 4 K. Affolter and M. von Allmen, AppL Phys. A, 23 (1984) 93. 5 C.J. Lin and F. Spaepen, Acta Metall., 34 (1986) 1367. 6 C. J. Lin, F. Spaepen and D. Turnbull, J. Non-Cryst. Solids, 61-62 (1984) 767. 7 C. A. MacDonald, A. M. Malvezzi and E Spaepen, J. Appl. Phys., 65 (1989) 129. 8 J. Y. Tsao, S. T. Picraux, P. S. Peercy and M. O. Thompson, Appl. Phys. Lett., 48 (1986)278. 9 R. E. Somekh, Z. H. Barber, C. S. Baxter, P. E. Donovan, J. E. Evetts and W. M. Stobbs, J. Mater. Sci. Lett., 3 (1984) 217. 10 E Spaepen and C. J. Lin, in M. von Allmen (ed.), Amorphous Metals and Non-Equilibrium Processing, Editions de Physique, Les Ulis, 1984, p. 65. 11 W. J. Boettinger and M. J. Aziz, Acta Metall., 37 (1989) 3379. 12 T. Egami and Y. Waseda, J. Non-Cryst. Solids, 64 (1984) 113.
243
Rapid solidification of cobalt-titanium alloys induced by nanosecond laser pulses S. Vitta Department of Metallurgical Engineering, Indian Institute of Technology, Powai, Bombay 400076 (India) A. L. Greer and R. E. Somekh Department of Materials Science and Metallurgy, University of Cambridge, Pembroke Street, Cambridge CB2 3QZ (UK)
Abstract Phase formation is studied in thin films of Co-Ti alloys, both as sputtered and after quenching following melting by a 5 ns laser pulse. The composition ranges for glass formation are established, and the significance of crystal growth kinetics is examined. The C o - T i system exhibits many types of intermetallic compounds and a wide variety of behaviours is found around their compositions.
1. Introduction
Partitionless crystal growth in an undercooled liquid, as may arise in congruent freezing or with complete solute trapping, occurs at a rate determined by thermal diffusion and by the kinetics of the interfacial processes. The focus of this study is the kinetics which may be diffusion limited or collision limited [1]. In the former case (expected for compounds and ordered solid solutions), even in the absence of solute partitioning, the migration of the crystal-liquid interface requires diffusive-type atomic jumps to effect changes in chemical order; the maximum growth rate (i.e. if not limited by thermal diffusion) is D/it (where D is the atomic diffusivity in the liquid and it the atomic diameter), i.e. about 10 m s- ~. In the latter case (expected for pure metals and disordered solid solutions), diffusivetype jumps are not required and the growth rate is limited only by the collisional frequency; the maximum rate is the speed of sound in the liquid, typically about 4000 m s-t. The distinction between the two regimes is important in rapid solidification, particularly concerning glass formation, which would be expected to be impossible by rapid liquid quenching if collisionlimited growth were to operate. The C o - T i system is suitable for the study of glass formation at compound-forming compositions as it has compounds exhibiting a variety of homogeneity ranges and structures [2]. Glass formation has been reported at 21-23 at.% Ti and about 77 at.% Ti by melt spinning [3], and at 25 at.% Ti and 80 at.% Ti by 30 ns 0921-5093/94/$7.00 SSDI 0921-5093(93)05520-Y
pulse laser quenching [4]. However, Co-Ti exhibits a lower glass-forming ability than for example Nb-Ti in which laser quenching gave only amorphous phases at compound compositions [5]. Of melt quenching methods, pulsed lasers (10 ps,10 ns) permit the most rapid solidification: cooling rates of 101°-1013 K s-1 and isotherm velocities of about 100 m s-l. At such rates solute partitioning is not possible [6], resulting in either a crystalline phase of the same composition as the melt or a glass. While measurement of crystal regrowth velocity is possible after laser melting (for example by monitoring the surface reflectivity or resistivity of the sample during and immediately after the laser pulse [7, 8], most studies of alloys, as here, have relied on microstructural characterization after melting and solidification. In completely crystalline films, the grain size and the total solidification time give a lower bound estimate of growth velocity; for equiaxed structures, the average grain radius is taken as the growth distance; for growth from surrounding unmelted material the distance is taken as the length of the columnar grains. When isolated crystals are found in a glassy matrix, the crystal radius gives the actual growth velocity, albeit only approximately. 2. Experimental methods
The experimental design followed closely that of Lin and Spaepen [5]. As shown in Fig. l(b), a copper substrate was coated with aluminium (about 1 /zm) and © 1994 - Elsevier Sequoia. All rights reserved
S. Vitta et aL
244
/Rapid sofidification of Co-Ti alloys
then the alloy of interest (less than 100 nm, to permit transmission electron microscopy (TEM) characterization without thinning). After laser irradiation, the alloy film was floated off the substrate by dissolution of the underlying aluminium in sodium hydroxide solution. In contrast to ref. 5, the as-deposited alloy films were homogeneous, not layered; they were crystalline or amorphous, depending on composition, but always single phase. The Co-Ti films were deposited by ultrahigh vacuum magnetron sputtering in argon [9]. Elemental targets (Co, less than 25 ppm metal impurity; Ti, 99.7% minimum purity) were used with the substrates rotated underneath; the films correspondingly had "layers" about 0.3 nm thick and are therefore considered homogeneous. The substrates were at floating temperature, estimated to be below 200°C. An Nd-YAG Q-switched mode-locked laser was used, giving 5 ns, 2 m J, pulses after frequency quadrupling to 266 nm wavelength. The beam was focused to 100/zm diameter on the specimen, and had an approximately annular (but very inhomogeneous) intensity distribution (Fig. l(a)). Laser energy was adjusted to give small holes in the deposited films, ensuring melting around the holes. Structural characterization was by TEM (Philips EM300), and compositions were determined (from simultaneous deposits on to silicon substrates) by energy-dispersive spectroscopy. Detailed thermal modelling is ruled out by the homogeneity of the laser beam used in this work. Instead, thermal parameters are estimated using the dimensional arguments presented by Spaepen and Lin
[
(i)
(Ji)
copper
substrote
(iii)
1 lam
Fig. 1. Schematic diagram of a laser-quenched sample: (a) the laser-irradiated area, showing the annular intensity distribution; (b) a section through the sample, showing the material removal by the laser pulse and the melt depth. Three characteristic regions of the Co-Ti thin film are marked. The different horizontal and vertical scales in (b) should be noted.
[10]. The melt depth d varies as shown in Fig. l(b) from about 1/zm (where the radiation was most intense, but with the melting not quite reaching the copper substrate) to about 100 nm (near the edge of the irradiated area, where only the top sputtered film was melted). The temperature gradient, [V I ~ Tm/d ( w h e r e T m is the congruent alloy melting temperature) is then calculated to be in the range 109-10 TM K m - I. The cooling rate 7"= D,hlv r I/d, (where Dth is the thermal diffusivity of the solid alloy) is 101°-10 ~2 K s-l. The total solidification time r ~ TmT is 10- 7_ 10 - 9 s, and the isotherm velocity Uth= i/~/I V T[ is 10-100 m s ' At the edges of the irradiated area there is an unmelted, but heat-affected, zone. In the part of the Co-Ti samples melted and resolidified, three areas may be distinguished (Fig. l(b)). In region (i) at the edge of the hole, the thickness of the alloy layer is reduced. The cooling rate may therefore be somewhat higher than in region (ii) which has the lowest cooling rate (i. e. about 101° K s-l). Region (iii), near the edge of the irradiated area, has a much lower melt depth and the highest quenching rate (i.e. about 10 ~2 K s-~). As discussed below, the quenched microstructures are often different in the three regions.
3. Results
Figure 2 shows the Co-Ti phase diagram and summarizes the results. As-deposited Co-Ti films were all single phase. The samples with 26.8 at.% Ti and 76.7 at.% Ti were completely amorphous, as were all samples of intermediate composition. Outside this range, solid solutions were found. As-deposited cobalt is c.c.p. (high temperature allotrope, r-Co) with a grain size of 50 nm. Addition of 4.5 at.% Ti and 9.3 at.% Ti reduces the grain size to 15 nm without changing the crystal structure. As-deposited titanium is h.c.p. (low temperature allotrope, a-Ti) with a grain size of 50 nm. Addition of 5.4 at.% Co changes the structure to c.c.p. and reduces the grain size to 10 nm. After laser quenching only the 26.8 at.% Ti alloy was completely amorphous. The structures of the other samples are described in order of increasing Ti content. Some representative bright field TEM micrographs are in Fig. 3. Laser-quenched cobalt is c.c.p. with a grain size of about 250 nm, much larger than in the as-deposited films which have the same structure. With 4.5 at.% Ti the structure is still c.c.p, with a similar grain size, but there is a high density of stacking faults. As further titanium is added (19.3 at.%) the faults become less evident, but the grain size decreases towards 0.1 /zm (Fig. 3(a)). With 40.3 at.% Ti, 43.7 at.% Ti and 52.6 at.% Ti (Figs. 3(b), 3(c) and 3(d) respectively) the CoTi phase appears, confirmed by the
S. Vitta etaL
/Rapidsolidification of Co-Ti alloys
Weight Percent Titanium 0
18001
10
20
'
'
30
'
50
'
'
60
'
70 '
80
'
90
'
tOO
I
liquid
t_~ 1400
2
40
12oo
E ,ooo p,\ /
o
to
Co
o.-d.pos ,.df tose.qoeochedl
ocp ~cp
I io=l
L J __
P-Ti I
20
30 4o 50 6o 7o 80 Atomic Percent Titanium
°
b / .
\
CoTi
9o
~oo Ti
S cop ScoTi2
+,cp
~>5+bc~.
Fig. 2. The Co-Ti phase diagram (after Murray [2]), showingthe T0 line for the c.c.p, phase (------). The various equilibrium compounds are indicated, includingthe Laves phases with structures C15 and C36. The phase mixtures found in the present work are indicated; the curved lines indicate uncertain boundaries. The arrows at the top of the diagram indicate the compositions tested. presence of superlattice reflections. Throughout the composition range the CoTi is the only phase in regions (iii) and (ii), but there is some amorphous phase in region (i). The grain size is about 40 nm in region (iii) (iii) (e.g. Figs. 3(b) and 3(d)) but much larger in region (ii) (e.g. Fig. 3(c)), ranging from about 1.0/zm in the middle of the composition range (43.7 at.% Ti) to about 0.2/zm at the edges. With 70.2 at.% Ti and 76.7 at.%Ti (Figs. 3(e) and 3(f) respectively), the structure is amorphous in region (i), is a mixture of amorphous phase and the intermetallic compound CoTi: in region (ii) and is a mixture of a-Ti and CoTi: in region (iii). In samples with 70.2 at.% Ti in regions (i) and (ii), the CoTi: phase appears as isolated roughly cylindrical grains in an amorphous matrix (e.g. Fig. 3(e)). In samples with 76.7 at.% Ti, the fine crystalline structure might have been formed by devitrification of an intermediate glassy phase (Fig. 3(f)). Laser-quenched titanium is b.c.c, with a grain size of 150 nm. The quenching is sufficiently rapid to retain at room temperature the high temperature allotrope,/5-Ti, formed from the melt. With 5.4 at.% Co added, the c.c.p, phase of the as-deposited film is retained after quenching, with a much increased grain size of about 0.2/~m. The equiaxed grains have only a low density of defects.
4. D i s c u s s i o n
In the laser-quenched films there is little occurrence of crystal regrowth from surrounding unmelted mate-
245
rial, even for compositions which are expected to display rapid growth. Pure titanium, for example, solidifies as/5-Ti formed by nucleation and growth in the melt, rather than as a-Ti regrown from the surrounding unmelted a-Ti. There is no evidence for mixing during melting of the Co-Ti layers and the underlying aluminium. There is very limited time for mixing and the alumina layer on the aluminium may act as a barrier. Fortunately for the microstructural studies, the alumina seems not to be an active nucleant for the crystalline phases in Co-Ti. The low areal density of nucleation events leads to large grain sizes and permits determination of crystal growth velocities in some cases. The comparative lack of influence of substrate and surrounding unmelted material facilitates the determination of intrinsic behaviour in this system, as discussed further below.
4.1. The Co3Ti phase This phase is not observed in laser-quenched samples; indeed, a completely amorphous phase is formed at 26.8 at.% Ti, close to the ideal stoichiometry. At lower titanium content the amorphous phase is still found, now in combination with the extended solid solution based on fl-Co. It is significant that in this composition range, in marked contrast to the behaviour elsewhere, a greater degree of amorphicity is found on the laser-quenched samples than in the asdeposited samples. This agrees with the results of melt spinning (cooling rate, 105-106 K S -1) which causes glass formation at similar compositions (21-23 at.% Ti) [3]. There seems to be a particular resistance to crystal nucleation and growth in the vicinity of the Co3Ti composition, remarkable since the phase has the relatively simple Au3Cu structure (ordered c.c.p.). If solidification is too rapid for formation of chemically ordered Co3Ti, then the chemically disordered form might be expected. (For an analysis of disorder trapping, see ref. 11). The undercooling necessary for this to form is easily estimated by extrapolating the TO line for the c.c.p, r - C o solid solution (broken line on Fig. 2). Partitionless solidification to the disordered Co3Ti solid solution should be possible at a moderate undercooling of about 50 K, yet even this is not observed. It may be that the liquid is particularly stable at this composition, perhaps because of the formation of associates; Murray [2] has commented on the particular difficulties of modelling the Co3Ti-liquid equilibrium. However, the main reason for the non-appearance of Co3Ti must be kinetic. This is found also in devitrification experiments. When Inoue et al. [3] annealed amorphous Co-22 at.% Ti samples they found that the first crystalline phase was the Laves phase at 34 at.% Ti, and not Co3Ti; the selection of a phase with a composition substantially different from that of the amorphous alloy must be a consequence of a
S. Vitta et al.
246
/Rapid solidification of Co- Ti alloys
"" .,,.4
::,k'
%
~,
~"
"
..
"?
~"
b
O
,;
•
4,
*/
~lr
4¢'* ,-
Q
C
q
0~5 um
.: :.:," :=:
f
O.25pm|
D
J
I
~
') t
f
t :
-
..
'
'~:~"
~ "
"
1.0IJm
Fig. 3. Bright field transmission electron micrographs, showing representative structures in Co-T] tllmS after laser quenching: (a) 19.3 at.% Ti; (b) 40.3 at.% Ti; (c) 43.7 at.%; (d) 52.6 at.% Ti; (e) 70.2 at.% Ti; (f) 76.7 at.% Ti. substantial kinetic barrier to nucleation or growth of Co3Ti.
4.2. Laves phases Two Laves phases appear in equilibrium in the C o - T i system at 31 at.% and 34 at.% Ti [2]. In equilibrium the phase fields for these compounds are only 0.5 at.% wide. Although the Laves phases at 34 at.% Ti can be obtained by devitrification [3], there was no evidence for either phase in the present work. The
laser-quenched compositions were not very close to those of the Laves phases, so their formation may not be expected; at least it is demonstrated that these phases do not substantially enlarge their composition ranges on rapid quenching.
4.3. Growth of the Co Ti phase This phase has the relatively simple CsCl structure and is observed over a wide composition range (40.3 at.%-52.6 at.% Ti) in laser-quenched samples. The
S. Vittaet al. /Rapidsolidification of Co-Ti alloys compound has previously been reported at compositions as far from stoichiometry as 43.9 at.% and 52 at.% Ti [2]. In the present work the grain size is largest at 43.7 at.% Ti, closest to the centre of the homogeneity range in the equilibrium diagram. Nucleation is expected to be easier at this composition, occurring at lesser undercooling and permitting rapid grain growth. Estimating from the grain size in Fig. 3(c), the growth velocity is about 10 m s- 1. 4.4. Growth of the CoTi phase This phase has a complex structure (of Fe3W3Ctype, E93) with 96 atoms in the unit cell [2], and is therefore surprising that it forms in rapid laser quenching. The isolated crystals in region (ii) of the 70.2 at.% Ti sample (Fig. 3(e)) enable the growth rate to be estimated as 4 m s- 1. This is fast for a complex structure, but within the diffusion-limited regime. The CoTi 2 phase is found also in laser-quenched Co-76.7 at.% Ti. Although not tested directly, it seems likely that the phase has formed without partitioning. Its composition then reflects a considerable extension of its homogeneity range. Although a range has been reported [2], such a large extension has not previously been observed. 4.5. Amorphous phase formation Some previous work on laser quenching, for example that on Nb-Ni using 30 ps pulses [5], suggested that the composition ranges for glass formation are similar for sputter deposition and for ultrarapid liquid quenching. The present work does not bear this out, as the glassy or amorphous phase is much less prevalent after quenching. The present results are, however, consistent with earlier work using melt spinning [3] and 30 ns laser pulses [4]. Comparison of the phase selection in as-deposited and quenched samples shows that the key difference is the presence, or not, of intermetallic compounds. No compounds are found in the as-deposited films, although the CoTi and C o T i 2 compounds appear to nucleate and grow readily during laser quenching. This suggests that the nucleation frequency for intermetallic compounds must be particularly temperature dependent, being significant in undercooled liquids during quenching, but essentially zero at the effective surface temperature at thin film deposition. It may also be that the chemical order in the melt aids compound formation. These effects are related to the diffusion limitation of the growth of the compounds. Similar considerations may apply in solid state amorphization by interfacial reaction; that the amorphous phase forms at all suggests a difficulty in nucleating intermetallic compounds at low annealing temperatures. As already discussed in Section 4.1, special considerations appear to apply for the Co3Ti which is difficult to form at any temperature.
247
Away from the range of compound formation, glass formability is determined by the ease of formation of the terminal solid solutions. For these conditions, Egami and Waseda [12] have proposed a criterion for glass formation based on the influence of atomic size difference on the topographical stability of terminal solid solutions. Taking their 2 parameter to have a value in the range 0.035-0.045 as found for the other systems, the glass-forming range in Co-Ti is predicted to be from 21.5-27.6 at.% Ti to 81.5-85.5 at.% Ti. The competition between the formation of glass and the terminal solid solution has been investigated in detail only on the Co-rich side in the present work. Complete amorphous phase formation is found for Ti contents greater than about 23 at.% in both as-deposited and laser-quenched samples. This limit of the glass formation agrees well with the Egami and Waseda prediction and with the results of melt spinning. It appears that when the competition is with the formation of a terminal solid solution, the limit of glass formability is not strongly dependent on the details of the quenching procedure; even ultrarapid quenching using a pulsed laser or thin film deposition do not significantly extend glass formation beyond the range achievable by melt spinning. Presumably the terminal solid solution can exhibit partitionless collision-limited growth, and in this regime the effect of solute content on thermodynamic parameters is dominant.
5. Conclusions
Apart from the specific findings of phase formation at particular compositions in the Co-Ti system, some general conclusions can be drawn. Pulsed laser quenching of alloys is effective in forming glasses and other metastable microstructures. The solidification is partitionless. When the competition is between glass formation and the formation of a terminal solid solution, the transition to collision-limited growth of the solid solution sets a sharp boundary to glass formation; boundary of the glass-forming composition range (for partitionless freezing) is not very dependent on details of the quenching method. On the contrary, at the compositions of intermetallic compounds where growth is diffusion limited, glass formation is dependent on the quenching method. In some cases it is possible to estimate growth velocities for compounds from the quenched microstructures; they are typically a few metres per second. Some compounds which form readily from the liquid, even on ultrarapid quenching, do not form in thin film deposition; this difficulty of formation at low temperatures may be relevant for solid state amorphization processes. The ease of formation of different compounds does not exhibit a straight-
248
S. Vitta et al.
/Rapid solidification of Co- Ti alloys
forward correlation with their structural complexity; in the present case, C o T i 2 (E93 structure, 96 atoms in unit cell) forms readily, but disordered Co3Ti (c.c.p.) does not.
Acknowledgments S.V. is grateful for the award of a C a m b r i d g e N e h r u Studentship. Deposition facilities s u p p o r t e d by the Science and Engineering Research Council (UK) are acknowledged.
References 1 S.R. Coriell and D. Turnbull, Acta Metall., 30 (1982) 2135. 2 J. L. Murray, Bull. Alloy Phase Diag., 3 (1982) 74.
3 A. Inoue, K. Kobayashi, C. Suryanarayana and T. Masumoto, Scr. Metall., 14 (1980) 119. 4 K. Affolter and M. von Allmen, AppL Phys. A, 23 (1984) 93. 5 C.J. Lin and F. Spaepen, Acta Metall., 34 (1986) 1367. 6 C. J. Lin, F. Spaepen and D. Turnbull, J. Non-Cryst. Solids, 61-62 (1984) 767. 7 C. A. MacDonald, A. M. Malvezzi and E Spaepen, J. Appl. Phys., 65 (1989) 129. 8 J. Y. Tsao, S. T. Picraux, P. S. Peercy and M. O. Thompson, Appl. Phys. Lett., 48 (1986)278. 9 R. E. Somekh, Z. H. Barber, C. S. Baxter, P. E. Donovan, J. E. Evetts and W. M. Stobbs, J. Mater. Sci. Lett., 3 (1984) 217. 10 E Spaepen and C. J. Lin, in M. von Allmen (ed.), Amorphous Metals and Non-Equilibrium Processing, Editions de Physique, Les Ulis, 1984, p. 65. 11 W. J. Boettinger and M. J. Aziz, Acta Metall., 37 (1989) 3379. 12 T. Egami and Y. Waseda, J. Non-Cryst. Solids, 64 (1984) 113.