Nanocrystallization of Ti-Ni-Cu-Sn amorphous alloy

Scripta mater. 43 (2000) 371–376 www.elsevier.com/locate/scriptamat

NANOCRYSTALLIZATION OF Ti-Ni-Cu-Sn AMORPHOUS ALLOY D.V. Louzguine* and A. Inoue Institute for Materials Research, Tohoku University, Katahira 2-1-1, Aoba-Ku, Sendai 980-8577, Japan (Received March 10, 2000) (Accepted April 4, 2000) Keywords: Transmission electron microscopy; Amorphous materials; Phase transformations Introduction At present Ti-based bulk amorphous alloys are of significant scientific and commercial interest due to their high tensile strength, (1,2) 1800 MPa for example in the case of Ti50Ni25Cu25 alloy (alloy’s compositions are given in at%) and their ability to be produced in a bulk form by mold casting (3). Addition of Sn increases significantly not only tensile strength (2200 MPa for Ti50Ni20Cu23Sn7 alloy for example) (4) but also the critical diameter of the bulk sample, that means the maximum diameter of an as-cast cylindrical sample with an amorphous single phase (4). The largest diameter attained for Ti-Ni-Cu-Sn bulk amorphous alloys was 5 mm (4). The above-mentioned result means that Ti-based amorphous alloys are not more the subject of scientific curiosity but can be widely industrially applicable, for example as structural materials. We should also mention their relatively high corrosion resistance at room temperature (5) and a relatively low density of the main alloying element Ti (4500 kg/m3) that implies higher strength/density ratio compared to Fe- or Zr-based bulk amorphous alloys. Ti-Ni-Cu-Sn alloys as well as Ti-Be-Zr (6), Ti-Ni-Cu-Al (1) and Ti-Zr-Ni-Cu (2) system alloys have a relatively wide supercooled liquid region of several tens degrees Kelvin (depending on composition) before crystallization observed by using differential scanning calorimetry (DSC) testing carried out at the heating rate of 0.67 K/s (4). Crystallization behaviour of Ti50Ni25Cu25 (7) and Ti45Ni20Cu25Sn5Zr5 alloys (8) has been studied recently. As has been reported before (4), the addition of Sn to Ti50Ni25Cu25 alloy causes a change of the shape of DSC traces from the single heat release to a double-step one, however structural changes during crystallization process are unstudied. The Ti50Ni20Cu23Sn7 alloy having the highest tensile strength among the Ti-Ni-Cu-Sn alloys investigated (4) and high glass-forming ability has been chosen for the present study. Experimental Procedure An ingot of the Ti50Ni20Cu23Sn7 alloy was prepared by arc-melting the mixture of Ti 99.95 mass% purity, Ni 99.9 mass% purity, Cu 99.99 mass% purity and Sn 99.9 mass% purity in an argon atmosphere. From this alloy, ribbon samples of about 0.02 mm in thickness and 0.9 mm in width were *Can also be written as Luzgin. 1359-6462/00/$–see front matter. © 2000 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved. PII: S1359-6462(00)00425-5

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Figure 1. X-ray diffraction patterns of the Ti50Ni20Cu23Sn7 alloy (a) in as-solidified state, (b) after isothermal calorimetry for 2220 s at 725 K, (c) after DSC at 0.67 K/s up to 787 K and (d) after isothermal annealing for 1200 s at 800 K.

prepared by rapid solidification of the melt on a single copper roller at the wheel surface velocity of 42 m/s. The structure of ribbon samples was examined by X-ray diffraction with monochromatic CuK␣ radiation. Transmission electron microscopy (TEM) was carried out using a JEM 2010 microscope operating at 200 kV equipped with an energy dispersive X-ray spectrometer (EDX). Crystallization temperature and heat of crystallization were examined by differential scanning calorimetry at heating rates ranging from 0.08 to 0.67 K/s. Heating up to testing temperature during isothermal differential calorimetry was conducted at 1.67 K/s. Cooling rate after DSC and isothermal calorimetry was about 0.6 – 0.8 K/s for the first 100 K. Isothermal annealing at 800 K was carried out in vacuum of 5⫻10⫺6 Torr. Mechanical properties were measured with an Instron-type testing machine at 298 K and a strain rate of 8.3⫻10⫺4 s⫺1. Results Heating of the Ti50Ni20Cu23Sn7 amorphous alloy (Fig. 1(a)) above the crystallization temperature of 750 –772 K, depending upon a heating rate (Fig. 2), results in formation of a nanostructure. After completion of the first heat release (marked as heat release A in Fig. 2) the structure consisted of equiaxed grains of 10 – 40 nm in size as illustrated in Fig. 3(a,b). Average composition of the particles 44 at% Ti, 27 at% Ni, 22 at% Cu and 7 at% Sn is close to the alloy composition. A small volume fraction of the residual amorphous phase was also presented. Analyses of the X-ray diffraction (Fig. 1 (c)) and selected area electron diffraction (Fig. 3 (c)) patterns indicated that this phase is a solid solution of Cu and Sn in a cubic Fd3m Ti2Ni phase (9). Lattice parameter of the solid solution of 1.138 nm is

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Figure 2. DSC traces taken at 0.08 K/s, 0.33 K/s and 0.67 K/s.

0.9% larger than that of the binary phase. Experimental and calculated values of d-spacings and integrated intensities of the diffraction peaks related to Ti2Ni solid solution showed a good agreement with each other. Exactly the same diffraction patterns were obtained for the samples annealed for 3.6 ks at 725 K after the completion of the heat release A in isothermal conditions (Fig. 4). No other phases were found. Isothermal calorimetry curves presented in Fig. 4 show that the first DSC heat effect A has a complicated shape. In order to study the beginning of crystallization the samples were annealed for 2220 and 2600 s at 725 K. At the initial stage of the crystallization (annealing for 2220 s at 725 K) the polyhedral grains of the Ti2Ni solid solution with a shape close to spherical (see Fig. 2 (d,e)) were about 10 –20 nm in size. Although only the strongest (511) peak of Ti2Ni solid solution is clearly seen in Fig. 1 (b), the diffraction pattern shown in Fig. 3 (f) contains sharp rings belonging to other diffraction maximums. As in the case of the diffraction pattern shown in Fig. 3 (c), all these diffraction peaks were indexed as belonging to Ti2Ni solid solution. The data are presented in Table 1. The second heat release (marked as release B in Fig. 2) is related to precipitation of the tetragonal phase that is likely a solid solution of Ni and Sn in ␥ TiCu phase (10) (Fig 1 (d)). Lattice parameters of ␥ TiCu solid solution and their ratio deviated from that of the binary ␥ TiCu phase being a ⫽ 0.314 nm and c ⫽ 0.612 nm. Several weak diffraction peaks remained unindexed. At this stage the crystalline grains size of the Ti2Ni solid solution changes slightly and its lattice parameter remain unchanged. Discussion The present study has shown that crystallization behaviour of the Ti50Ni20Cu23Sn7 alloy is significantly different from that of the Ti50Ni25Cu25 (7) and Ti45Ni20Cu25Sn5Zr5 alloys (8) studied earlier. A multicomponent cubic phase with lattice parameters of 0.3047 and 0.3069 nm, respectively, and a grain size of 300 –900 nm and 100 –200 nm, respectively, was a main structural constituent after the completion of the primary crystallization in these two alloys. Crystallization of the Ti50Ni20Cu23Sn7 alloy begins from the primary precipitation of Ti2Ni solid solution with a lattice parameter of 1.138 nm. Although only Ti2Ni solid solution phase was observed to precipitate during the reaction A (see Fig. 1

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Figure 3. (a,d,g) Bright- and (b,e,h) dark-field TEM images, (c,f) selected area electron diffraction patterns and (i) high-resolution TEM image. The samples (a,b,c) after DSC at 0.67 K/s up to 787 K, (d,e,f) after isothermal calorimetry for 2220 s at 725 K and (g,h,i) after DSC at 0.67 K/s up to 850 K.

Figure 4. Isothermal differential calorimetry traces.

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TABLE 1 Experimental d-Spacings Taken from the Selected-Area Electron Diffraction Patterns (Sharp Diffraction Rings) Shown in Fig. 3 (c and f) and Calculated Spacings for a Cubic Structure with Lattice Parameter of a ⫽ 1.138 nm From Fig. 3 (c) after DSC at 0.67 K/s up to 787 K HKL d exp., nm d cal., nm HKL d exp., nm d cal., nm HKL d exp., nm d cal., nm

(220) (222) (331) (422) 0.403 0.324 0.260 0.233 0.4023 0.3285 0.2611 0.2323 (440) (442) (660) (752) 0.203 0.190 0.134 0.1291 0.2012 0.1897 0.1341 0.1289 From Fig. 3 (f) after isothermal calorimetry for 2220 s at 725 K (331) (511) (440) (660) 0.260 0.219 0.202 0.134 0.2611 0.2190 0.2012 0.1341

(511) 0.219 0.2190

(862) 0.114 0.1159

(b,c) and Fig. 3 (c,f)), equiaxed grains belonging to Ti2Ni solid solution phase only were observed at the initial stage of primary crystallization “release A” (see Fig. 3 (d,e)) and after its completion (see Fig. 3 (a,b))), the trace of heat release A studied by DSC and isothermal calorimetry has a complicated shape (see Fig. 4). Thus, primary crystallization of Ti2Ni solid solution cannot be described from the viewpoint of the Avrami exponent (11) value which is widely used for analysis of different phase transformations including crystallization of amorphous phase. This release demonstrates influence of the nanoscale crystal size and the change of the growth rate during primary crystallization according to Kolmogorov (12)–Johnson (13)–Mehl (13)–Avrami (11)–Kelton (14) general exponential equation for the fraction transformed: 4␲ x e共t兲 ⫽ 3V 0

冕 冋冕 t

I共 ␶ 兲

0

␶

t

g共t⬘兲dt⬘

册

3

d␶

(1)

where I(␶) and g(t⬘) are time-dependent nucleation and growth rates, respectively. This feature has also been observed in Zr-based multicomponent alloys (15). Difficulty of nucleation of the Ti2Ni solid solution having a large lattice parameter and complicated composition is one of the reasons for the elevated glass-forming ability of the Ti-Ni-Cu-Sn amorphous alloy studied. As the ␥ TiCu solid solution precipitates from the residual amorphous matrix and the lattice parameter of the Ti2Ni solid solution after the release B remains unchanged, no redistribution of the alloying elements between Ti2Ni and ␥ TiCu phases takes place. The ␥ TiCu solid solution grains occupy the spaces between the Ti2Ni polyhedrons and have a different morphology as shown in Fig. 3 (i). A carcass of the Ti2Ni solid solution phase created at the first stage of crystallization does not allow ␥ TiCu solid solution grains to reach the size larger than the spaces between crystalline grains and these grains remain to be nanocrystals. One should also mention close correspondence of some d-spacings, for example (511) - (102), (440) - (003) and (842) - (203) of Ti2Ni and ␥ TiCu solid solutions, respectively, and that 2c value of ␥ TiCu solid solution is close to the lattice parameter of the Ti2Ni one (the difference is about 7%). Good correspondence of the lattices, also seen in Fig 3 (i), implies formation of ␥ TiCu on the base of Ti2Ni. A low volume fraction of the third unindexed crystalline phase is also present. Although the structure of the Ti50Ni20Cu23Sn7 alloy in nanocrystalline state is favorable from the viewpoint of the mechanical properties (i.e. small grain size and their equiaxed morphology), Ti2Ni intermetallic compound is brittle and does not allow to attain high mechanical strength in nanocrys-

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talline or mixed nanocrystalline ⫹ amorphous state. For example, tensile strength of 1000 ⫹-50 MPa obtained for the ribbon samples heat treated for 1200 s at 800 K is about two times less than that of the as-solidified amorphous samples (4). Conclusions Addition of 7 at% Sn to the ternary Ti50Ni25Cu25 alloy significantly changes its crystallization behaviour leading to formation of the nanocrystalline structure consisting of Ti2Ni and ␥ TiCu solid solution phases in two steps. Solid solution of Cu and Sn in Ti2Ni phase is formed during primary crystallization whereas ␥ TiCu solid solution crystallizes from the residual amorphous matrix. Good correspondence of the lattices implies formation of ␥ TiCu on the base of Ti2Ni. Isothermal calorimetry data show an important role of the nanoscale crystal size and change of the growth rate during primary crystallization. Nanocrystallization of the Ti50Ni20Cu23Sn7 alloy causes its embrittlement and significant reduction of mechanical properties and should be avoided when used as a structure material. References 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15.

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