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Crystallization in amorphous NiZr2, studied by HRTEM
346
Materials Science and Engineering, A 133 ( 1991 ) 346-352
Crystallization in amorphous NiZr2, studied by HRTEM C. Beeli and H.-U. Nissen Laboratory of Solid State Physics, ETH Ziirich, H6nggerberg, CH-8093 Ziirich (Switzerland)
Q. Jiang and R. Liick Max Planck-lnstitut fiir Metallforschung, Inst. fiir Werkstoffwissenschaften, Seestr. 75, D- 7000Stuttgart i (F.R. G.)
Abstract The crystalline textures of amorphous NiZr 2 heated in a differential scanning calorimeter (DSC) at very low heating rates were studied by high-resolution transmission electron microscopy. Three differently heat-treated specimens were compared. One of these was taken out of the DSC just after a strong endothermal peak had occurred, which is unexpected for amorphous binary alloys. The remaining specimens were taken out of the DSC after termination of crystallization. All specimens have a polycrystalline texture, with an average grain size of 100 nm. The grains have the tetragonal A12Cu-type structure, and each of these grains has a nanocrystalline domain texture consisting of rotation twin domains characterized by the following orientation relation: [00111//[111]z; [11011//[11012. The boundary between any two nanocrystals intergrown according to this relation is a planar interface with index (110). The typical distance between such interfaces is 1-3 nm. It is concluded that all specimens are fully crystallized after the occurrence of the strong endothermal peak and that the orientation of the twin domains in the nanocrystalline texture results from medium-range bond orientational order assumed to be present in the amorphous starting material.
1. Introduction For many years the atomic structure of glasses has been a major unsolved problem in the materials sciences. To understand the structure of dense liquids, Bernal [1], in 1964, introduced the so-called dense random packing (DRP) of spheres model. Applied to metallic glasses, this model could explain the most prominent properties of glasses, such as the position of the first peak of the pair correlation function g(R) or the damping of the oscillations of g(R) [2]. However, the D R P model of metallic glasses failed to explain the presence of a short-range order which resembles that of their crystalline counterparts, for instance by the fact that metallic glasses have a clearly defined shape as well as population of the first coordination shell around a certain constituent atom [3, 4]. In 1982 Lamparter, Sperl, Steeb and Bletry demonstrated that even medium-range order was present in glassy Nis1B19 [5]. The boron atoms in this glass were shown to be correlated over distances as large as 10 A. Lee et al. [6] as well as Mizoguchi et aL [7] have shown that the 0921-5093/91/$3.50
short-range order present in glassy N i Z r 2 closely resembles that of its crystalline counterpart. A general model that is able to take into account this medium-range order has been developed by Dubois [8]. This model has successfully been applied to glassy Ni-B. A review of the presently available models explaining the formation and relative stability of metallic glasses has recently been given by Dubois and Janot [9]. Using a refined method to measure the specific heat capacity Cp, L/ick, Jiang and Predel [10, 11] have recently determined the activation energy of the glass transition and of the crystallization for an amorphous Ni34Zr66 alloy, applying heating rates in the range of 0.3125-80 K min- 1. At heating rates below 5 K min-1 the endothermic peak of the glass transition was found to be rather large. Figure 1 shows a schematic drawing of a differential scanning calorimeter (DSC) measurement for low heating rates. It has been suggested [10] that the amount of endothermic enthalpy is a measure of the stability of the amorphous state. From DSC measurements Liick et al. © Elsevier Sequoia/Printed in The Netherlands
347
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TEMPERATURE
Fig. 1. Schematic drawing of a DSC measurement of NiZr 2 for low heating rates.
[10] also concluded that crystallization occurs in a single step. To clarify the origin of the strong endothermal peak, the crystalline textures were studied by electron microscopy at different stages of the heating process for very low heating rates. For high-resolution transmission electron microscopy (HRTEM) both a 300 kV as well as a 200 kV transmission electron microscope (TEM) were used.
2. Experimental Thin tapes (22/~m in thickness) of amorphous stoichiometric NiZr2, obtained by the melt-spin-
Fig. 2. (a) Polycrystalline grain arrangement with a region of amorphous material near the specimen boundaries marked by letter "a'. (b) Diffraction pattern of a single grain of a NiZr z crystal.
348 ning technique were subjected to a stepwise heat treatment (Cp method) using a step height of 1 K and applying an average heating rate of 0.3125 K min-1. The heating process was monitored by a DSC. Three different specimens were used to compare the crystalline textures forming in amorphous NiZr 2 when heated at very low heating rates. Specimen A was taken out of the DSC just after the strong endothermal peak, marked T 3 in Fig. 1, had occurred [10], and specimen B was taken out of the DSC after termination of the reaction at a point marked T 4 in Fig. 1. Specimen C was heated by applying a much larger average heating rate of 10 K min- 1 and was also taken out of the DSC at point T4, after termination of the reaction. Characterization by X-ray diffraction indicated that specimens B and C had become totally crystalline, while X-ray diffraction patterns of specimen A had broadened peaks, which may indicate the presence of amorphous material. The specimens were ion-etched for 1.5-2 h at 5.5 kV, using an argon beam. In addition, thin edges of broken specimens were investigated directly by TEM to check for any artifacts possibly introduced in the ion-etching process. For measurements of the lattice spacings, the absolute camera length of several selected area electron diffraction (SAD) patterns was determined by reference to the SAD patterns of a polycrystalline gold film. Using this camera length, which gave absolute d-values, indexing of the diffraction patterns was possible.
3. Results
Specimens A, B and C invariably consist of a polycrystalline grain arrangement. The crystal grains have the tetragonal A12Cu-type structure, with space group I4/mcm and cell parameters a = 0.6483 nm and c= 0.5267 nm. The average size of the grains is approximately 100 nm (Fig. 2). In specimen A, a small proportion (less than 5 vol.%) of the material was found to be amorphous (Figs. 2 and 3). The nanocrystalline texture is characterized by the following orientation relation: [00111//[11112; [11011//[11012 (Fig. 4). The boundary between any two nanocrystals intergrown according to this relation is a planar interface with index (110). Both equivalent (110)- and (110)-planes have frequently been observed in a single grain. The typical distance between such interfaces is 1-3 nm.
Fig. 3. (a) Bright-field electron micrograph of an amorphous grain included by polycrystallinegrains. The contrast in the polycrystallinegrains indicates the presence of many defects. (b) Diffraction pattern of the amorphous grain shown in the center of (a).
In specimen B no amorphous material was found except in very small domains (1-6 nm in diameter) with a highly disordered structure occurring within the polycrystalline grains (Fig. 4(b)). All polycrystalline grains showed the same nanocrystalline texture. In addition, a polysynthetic twin texture consisting of lamellar twins of NiZr 2 was observed. The entire grain shown in
349
Fig. 4. (a) Diffraction pattern of a single polycrystalline grain. The diffraction pattern indicates the presence of a specific orientation relation (see text). (b) Electron micrograph of a grain with nanocrystalline domain texture. Both types of(110)-planar defects can be recognized. Some amorphous regions are marked by the letter "a".
Fig. 5 is typical for this texture and is an arrangement of polysynthetic twin lamellae. T h e lamellar domains have widths of 0 . 1 - 0 . 4 / ~ m and lengths of 0.8-5 ~ m (Fig. 5). In specimen C the same polycrystalline grains with the nanocrystalline domain texture were found, but no polysynthetic twins such as those occurring in specimen B have been observed. This suggests that in specimens B and C most of the material (probably more than 70%) is poly-
crystalline and has the nanocrystalline domain texture. T h e complicated diffraction patterns shown in Figs. 2(b) and 4(a) cannot be explained as single crystal diffraction patterns of NiZr 2. In Fig. 2(b) additional reflections and strong streaks parallel to the [110J-direction can be recognized. In Fig. 4 there are even two sets of parallel streaks oriented normal to each other. T h e diffraction pattern in Fig. 2(b) can be explained as the super-
350
Fig. 5. Bright-fieldelectronmicrographof a polysynthetictwintexturein NiZr~with manylamellartwin domains.
position of two different diffraction patterns of NiZr 2. In Fig. 6(a) a diffraction pattern in the [111] orientation is superimposed onto a [001Jdiffraction pattern with the [ll0]-directions in parallel orientation. There are additional spots at the positions with l = _+ 1, _+3 .... (such as [101], [211], [213], [ 3 0 3 ] ) i n the [111] diffraction pattern. In single crystal diffraction patterns of NiZr 2 these reflections are kinematically extinct. In Fig. 6(b) these additional spots are plotted as unfilled circles. The high-resolution electron micrograph (Fig. 6(c)) shows domains oriented in the [111J-direction, separated from each other by domains oriented in the [001]-direction. Depending on the thickness of the [001]-domains, there is a relative shift of two adjacent [111]-domains by 0.5 [110]. This situation can be compared with antiphase domains separated by a boundary domain (a [001]-domain) of a certain thickness. Because of these shifts additional reflections are present in the diffraction pattern at the positions [101], [211], etc., i.e. exactly at the positions indicated by the unfilled circles in Fig. 6(b). The effect of similar displacements on the diffraction patterns of the silicate mineral tobermorite has been discussed by Bollmann [12]. The presence of strong streaks in the observed SAD patterns can be explained by the irregularly arranged domains with different widths as well as by atomic disorder at the interfaces. In the grain shown in Fig. 4(b), interfaces of both orientations, (110) as well as (110), are present. Therefore streaks nor-
mal to both sets of interfaces occur in the corresponding diffraction pattern (Fig. 4(a)). 4. Discussion
The occurrence of the specific orientation relation of the nanocrystalline texture can be explained by the geometric theory of interfaces [13, 14]. Two domains with the [lll]-direction and the [001J-direction parallel to each other having a common interface (110) are rotated against each other around the [ll0]-direction, which is perpendicular to the interface. This kind of boundary between two domains is called a twist boundary. The rotation angle is equal to the angle between the [220]-direction and the [112Jdirection, which is 60.1 °. An (ll0)-interface rotated by 60.1 ° around the [ll0]-direction shows a good geometric fit with the unrotated (110)-interface, since the d-values d1220] and dill2 ] are almost equal; i.e. di220]=0.2292 nm and dfll21 = 0.2284 nm. There are at least two possible orientations between the two interfaces. In the first case nickel atoms coincide at the interface, and in the second case, zirconium atoms coincide at the interface. Probably, both these cases occur. The periodicity of the coincidence site lattice in the [110]-direction is doubled. We now briefly discuss the relation between the structure of the original amorphous solid and the resulting nanocrystalline twin domain texture.
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Fig. 6. (a) Superimposition of two diffraction patterns in [lll]-projection and in [001j-projection, respectively, with the (110)-directions in parallel orientation. (b) Diffraction pattern similar to (a) but showing additional diffraction spots, indicated by unfilled circles. (c) High-resolution electron micrograph showing shifts between domains oriented parallel to the [ 111 ]-direction. The amount of shift is 0.5 [ 110]. For a clear recognition of the shifts, the micrograph should be observed at a grazing angle.
T h e results of the electron microscopic investigation show that all three differently heat-treated specimens showed the same nanocrystalline texture as the main structure. This suggests that between point T 2 and point T 3 in the heating process, the a m o r p h o u s N i Z r 2 material has converted into the crystalline texture. T h e differences in the D S C diagrams for small vs. large heating rates do not give any direct information about the structural features of the forming crystalline texture or the a m o r p h o u s starting material.
Since glassy NiZr 2 has a short-range order that closely resembles that of the crystalline structure of N i Z r 2 [6, 7], it is assumed that the model discussed by Dubois [8] can be applied to a m o r phous N i Z r 2. It may be assumed that in small domains of the a m o r p h o u s material a preferred orientation of bonds has existed, which is parallel to that in the nanocrystalline domains. This can be concluded f r o m the extremely small width (1-3 nm) of the twin domains in the crystalline state.
352
The stability of this nanocrystalline domain texture during the heating process from T 3 to T 4 (Fig. 1), which inhibits further grain growth, can possibly be explained as follows. The periodicity of the coincidence site lattice in the [110J-direction is only twice as large as the periodicity of the (110)-plane in the single crystal structure. Therefore the interface energy of these (ll0)-twist boundaries may be very low.
5. Conclusion The transformation of the amorphous material into the nanocrystalline texture takes place between the points T 2 and T3, i.e. when the strong endothermal peak in the DSC measurement occurs. The orientation of the twin domains in the nanocrystalline grains after the transformation into a crystalline solid is a result of a more incipient bond orientational order in the form of small irregularly bounded domains in the amorphous starting material. We hypothesize that the strong endothermal peak, unexpected for binary amorphous alloys, which is observed in the DSC plot for low heating rates [10], is not in the first place due to the glass transition but to rearrangement of atoms in the formation of the nanocrystalline texture. This suggestion is not meant to exclude the possibility that the glass transition makes a minor contribution to the endothermal peak.
Acknowledgments The original melt-spun specimens have been supplied by Prof. S. Steeb and M. Heckele for which we are very grateful. We acknowledge inspiring discussions with Prof. J.-M. Dubois on the structure of metallic glasses.
References 1 J.D. Bernal, Proc. Roy. Soc. A, 280(1964) 299. 2 G.S. Cargill, SolidState Phys., 30(1975) 227. 3 J.-M. Dubois and G. LeCa~r, Acta Metall., 32 (1984) 2101. 4 P. H. Gaskell, Glassy Metals 11, Topics in Applied Physics, 53, Springer, Berlin, 1983, p. 5. 5 P. Lamparter, W. Sperl, S. Steeb and J. Bletry, Z. Naturf., (a)37(1982) 1223. 6 A. Lee, G. Etherington and C. N. J. Wagner, J. NonCrystalline Solids, 61 & 62 (1984) 349. 7 T. Mizoguchi, S. Yada, N. Akutsu, S. Yamada, J. Nishioka, T. Suematsu and N. Watanabe, in S. Steeb and H. Warlimont (eds.), Proc. 5th Int. Conf. on Rapidly Quenched Metals, North-Holland, Amsterdam, 1985, p. 483. 8 J.-M. Dubois, J. Less-Common Metals, 145 (1988) 309. 9 J.-M. Dubois and Chr. Janot, Phil. Mag., B61 (1990) 649. 10 R. Liick, Q. Jiang and B. Predel, in H. Endo (ed.), Proc. Conf. Liquid and Amorphous Metals, LAM 7, Kyoto, 1989, in J. Non-Cryst. Mater., 911 (1990) 117-118. 11 Q. Jiang, R. Liick and B. Predel, Z. Metallk., 81 (1990) 94. 12 W. Bollmann, Z. Kristallogr., 126(1968) 1. 13 W. Bollmann, Phil Mag., 16(1967) 363 and 383. 14 W. Bollmann, Crystal Defects and Crystalline Interfaces, Springer, Berlin, 1970.
Materials Science and Engineering, A 133 ( 1991 ) 346-352
Crystallization in amorphous NiZr2, studied by HRTEM C. Beeli and H.-U. Nissen Laboratory of Solid State Physics, ETH Ziirich, H6nggerberg, CH-8093 Ziirich (Switzerland)
Q. Jiang and R. Liick Max Planck-lnstitut fiir Metallforschung, Inst. fiir Werkstoffwissenschaften, Seestr. 75, D- 7000Stuttgart i (F.R. G.)
Abstract The crystalline textures of amorphous NiZr 2 heated in a differential scanning calorimeter (DSC) at very low heating rates were studied by high-resolution transmission electron microscopy. Three differently heat-treated specimens were compared. One of these was taken out of the DSC just after a strong endothermal peak had occurred, which is unexpected for amorphous binary alloys. The remaining specimens were taken out of the DSC after termination of crystallization. All specimens have a polycrystalline texture, with an average grain size of 100 nm. The grains have the tetragonal A12Cu-type structure, and each of these grains has a nanocrystalline domain texture consisting of rotation twin domains characterized by the following orientation relation: [00111//[111]z; [11011//[11012. The boundary between any two nanocrystals intergrown according to this relation is a planar interface with index (110). The typical distance between such interfaces is 1-3 nm. It is concluded that all specimens are fully crystallized after the occurrence of the strong endothermal peak and that the orientation of the twin domains in the nanocrystalline texture results from medium-range bond orientational order assumed to be present in the amorphous starting material.
1. Introduction For many years the atomic structure of glasses has been a major unsolved problem in the materials sciences. To understand the structure of dense liquids, Bernal [1], in 1964, introduced the so-called dense random packing (DRP) of spheres model. Applied to metallic glasses, this model could explain the most prominent properties of glasses, such as the position of the first peak of the pair correlation function g(R) or the damping of the oscillations of g(R) [2]. However, the D R P model of metallic glasses failed to explain the presence of a short-range order which resembles that of their crystalline counterparts, for instance by the fact that metallic glasses have a clearly defined shape as well as population of the first coordination shell around a certain constituent atom [3, 4]. In 1982 Lamparter, Sperl, Steeb and Bletry demonstrated that even medium-range order was present in glassy Nis1B19 [5]. The boron atoms in this glass were shown to be correlated over distances as large as 10 A. Lee et al. [6] as well as Mizoguchi et aL [7] have shown that the 0921-5093/91/$3.50
short-range order present in glassy N i Z r 2 closely resembles that of its crystalline counterpart. A general model that is able to take into account this medium-range order has been developed by Dubois [8]. This model has successfully been applied to glassy Ni-B. A review of the presently available models explaining the formation and relative stability of metallic glasses has recently been given by Dubois and Janot [9]. Using a refined method to measure the specific heat capacity Cp, L/ick, Jiang and Predel [10, 11] have recently determined the activation energy of the glass transition and of the crystallization for an amorphous Ni34Zr66 alloy, applying heating rates in the range of 0.3125-80 K min- 1. At heating rates below 5 K min-1 the endothermic peak of the glass transition was found to be rather large. Figure 1 shows a schematic drawing of a differential scanning calorimeter (DSC) measurement for low heating rates. It has been suggested [10] that the amount of endothermic enthalpy is a measure of the stability of the amorphous state. From DSC measurements Liick et al. © Elsevier Sequoia/Printed in The Netherlands
347
1 f
o W
TEMPERATURE
Fig. 1. Schematic drawing of a DSC measurement of NiZr 2 for low heating rates.
[10] also concluded that crystallization occurs in a single step. To clarify the origin of the strong endothermal peak, the crystalline textures were studied by electron microscopy at different stages of the heating process for very low heating rates. For high-resolution transmission electron microscopy (HRTEM) both a 300 kV as well as a 200 kV transmission electron microscope (TEM) were used.
2. Experimental Thin tapes (22/~m in thickness) of amorphous stoichiometric NiZr2, obtained by the melt-spin-
Fig. 2. (a) Polycrystalline grain arrangement with a region of amorphous material near the specimen boundaries marked by letter "a'. (b) Diffraction pattern of a single grain of a NiZr z crystal.
348 ning technique were subjected to a stepwise heat treatment (Cp method) using a step height of 1 K and applying an average heating rate of 0.3125 K min-1. The heating process was monitored by a DSC. Three different specimens were used to compare the crystalline textures forming in amorphous NiZr 2 when heated at very low heating rates. Specimen A was taken out of the DSC just after the strong endothermal peak, marked T 3 in Fig. 1, had occurred [10], and specimen B was taken out of the DSC after termination of the reaction at a point marked T 4 in Fig. 1. Specimen C was heated by applying a much larger average heating rate of 10 K min- 1 and was also taken out of the DSC at point T4, after termination of the reaction. Characterization by X-ray diffraction indicated that specimens B and C had become totally crystalline, while X-ray diffraction patterns of specimen A had broadened peaks, which may indicate the presence of amorphous material. The specimens were ion-etched for 1.5-2 h at 5.5 kV, using an argon beam. In addition, thin edges of broken specimens were investigated directly by TEM to check for any artifacts possibly introduced in the ion-etching process. For measurements of the lattice spacings, the absolute camera length of several selected area electron diffraction (SAD) patterns was determined by reference to the SAD patterns of a polycrystalline gold film. Using this camera length, which gave absolute d-values, indexing of the diffraction patterns was possible.
3. Results
Specimens A, B and C invariably consist of a polycrystalline grain arrangement. The crystal grains have the tetragonal A12Cu-type structure, with space group I4/mcm and cell parameters a = 0.6483 nm and c= 0.5267 nm. The average size of the grains is approximately 100 nm (Fig. 2). In specimen A, a small proportion (less than 5 vol.%) of the material was found to be amorphous (Figs. 2 and 3). The nanocrystalline texture is characterized by the following orientation relation: [00111//[11112; [11011//[11012 (Fig. 4). The boundary between any two nanocrystals intergrown according to this relation is a planar interface with index (110). Both equivalent (110)- and (110)-planes have frequently been observed in a single grain. The typical distance between such interfaces is 1-3 nm.
Fig. 3. (a) Bright-field electron micrograph of an amorphous grain included by polycrystallinegrains. The contrast in the polycrystallinegrains indicates the presence of many defects. (b) Diffraction pattern of the amorphous grain shown in the center of (a).
In specimen B no amorphous material was found except in very small domains (1-6 nm in diameter) with a highly disordered structure occurring within the polycrystalline grains (Fig. 4(b)). All polycrystalline grains showed the same nanocrystalline texture. In addition, a polysynthetic twin texture consisting of lamellar twins of NiZr 2 was observed. The entire grain shown in
349
Fig. 4. (a) Diffraction pattern of a single polycrystalline grain. The diffraction pattern indicates the presence of a specific orientation relation (see text). (b) Electron micrograph of a grain with nanocrystalline domain texture. Both types of(110)-planar defects can be recognized. Some amorphous regions are marked by the letter "a".
Fig. 5 is typical for this texture and is an arrangement of polysynthetic twin lamellae. T h e lamellar domains have widths of 0 . 1 - 0 . 4 / ~ m and lengths of 0.8-5 ~ m (Fig. 5). In specimen C the same polycrystalline grains with the nanocrystalline domain texture were found, but no polysynthetic twins such as those occurring in specimen B have been observed. This suggests that in specimens B and C most of the material (probably more than 70%) is poly-
crystalline and has the nanocrystalline domain texture. T h e complicated diffraction patterns shown in Figs. 2(b) and 4(a) cannot be explained as single crystal diffraction patterns of NiZr 2. In Fig. 2(b) additional reflections and strong streaks parallel to the [110J-direction can be recognized. In Fig. 4 there are even two sets of parallel streaks oriented normal to each other. T h e diffraction pattern in Fig. 2(b) can be explained as the super-
350
Fig. 5. Bright-fieldelectronmicrographof a polysynthetictwintexturein NiZr~with manylamellartwin domains.
position of two different diffraction patterns of NiZr 2. In Fig. 6(a) a diffraction pattern in the [111] orientation is superimposed onto a [001Jdiffraction pattern with the [ll0]-directions in parallel orientation. There are additional spots at the positions with l = _+ 1, _+3 .... (such as [101], [211], [213], [ 3 0 3 ] ) i n the [111] diffraction pattern. In single crystal diffraction patterns of NiZr 2 these reflections are kinematically extinct. In Fig. 6(b) these additional spots are plotted as unfilled circles. The high-resolution electron micrograph (Fig. 6(c)) shows domains oriented in the [111J-direction, separated from each other by domains oriented in the [001]-direction. Depending on the thickness of the [001]-domains, there is a relative shift of two adjacent [111]-domains by 0.5 [110]. This situation can be compared with antiphase domains separated by a boundary domain (a [001]-domain) of a certain thickness. Because of these shifts additional reflections are present in the diffraction pattern at the positions [101], [211], etc., i.e. exactly at the positions indicated by the unfilled circles in Fig. 6(b). The effect of similar displacements on the diffraction patterns of the silicate mineral tobermorite has been discussed by Bollmann [12]. The presence of strong streaks in the observed SAD patterns can be explained by the irregularly arranged domains with different widths as well as by atomic disorder at the interfaces. In the grain shown in Fig. 4(b), interfaces of both orientations, (110) as well as (110), are present. Therefore streaks nor-
mal to both sets of interfaces occur in the corresponding diffraction pattern (Fig. 4(a)). 4. Discussion
The occurrence of the specific orientation relation of the nanocrystalline texture can be explained by the geometric theory of interfaces [13, 14]. Two domains with the [lll]-direction and the [001J-direction parallel to each other having a common interface (110) are rotated against each other around the [ll0]-direction, which is perpendicular to the interface. This kind of boundary between two domains is called a twist boundary. The rotation angle is equal to the angle between the [220]-direction and the [112Jdirection, which is 60.1 °. An (ll0)-interface rotated by 60.1 ° around the [ll0]-direction shows a good geometric fit with the unrotated (110)-interface, since the d-values d1220] and dill2 ] are almost equal; i.e. di220]=0.2292 nm and dfll21 = 0.2284 nm. There are at least two possible orientations between the two interfaces. In the first case nickel atoms coincide at the interface, and in the second case, zirconium atoms coincide at the interface. Probably, both these cases occur. The periodicity of the coincidence site lattice in the [110]-direction is doubled. We now briefly discuss the relation between the structure of the original amorphous solid and the resulting nanocrystalline twin domain texture.
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Fig. 6. (a) Superimposition of two diffraction patterns in [lll]-projection and in [001j-projection, respectively, with the (110)-directions in parallel orientation. (b) Diffraction pattern similar to (a) but showing additional diffraction spots, indicated by unfilled circles. (c) High-resolution electron micrograph showing shifts between domains oriented parallel to the [ 111 ]-direction. The amount of shift is 0.5 [ 110]. For a clear recognition of the shifts, the micrograph should be observed at a grazing angle.
T h e results of the electron microscopic investigation show that all three differently heat-treated specimens showed the same nanocrystalline texture as the main structure. This suggests that between point T 2 and point T 3 in the heating process, the a m o r p h o u s N i Z r 2 material has converted into the crystalline texture. T h e differences in the D S C diagrams for small vs. large heating rates do not give any direct information about the structural features of the forming crystalline texture or the a m o r p h o u s starting material.
Since glassy NiZr 2 has a short-range order that closely resembles that of the crystalline structure of N i Z r 2 [6, 7], it is assumed that the model discussed by Dubois [8] can be applied to a m o r phous N i Z r 2. It may be assumed that in small domains of the a m o r p h o u s material a preferred orientation of bonds has existed, which is parallel to that in the nanocrystalline domains. This can be concluded f r o m the extremely small width (1-3 nm) of the twin domains in the crystalline state.
352
The stability of this nanocrystalline domain texture during the heating process from T 3 to T 4 (Fig. 1), which inhibits further grain growth, can possibly be explained as follows. The periodicity of the coincidence site lattice in the [110J-direction is only twice as large as the periodicity of the (110)-plane in the single crystal structure. Therefore the interface energy of these (ll0)-twist boundaries may be very low.
5. Conclusion The transformation of the amorphous material into the nanocrystalline texture takes place between the points T 2 and T3, i.e. when the strong endothermal peak in the DSC measurement occurs. The orientation of the twin domains in the nanocrystalline grains after the transformation into a crystalline solid is a result of a more incipient bond orientational order in the form of small irregularly bounded domains in the amorphous starting material. We hypothesize that the strong endothermal peak, unexpected for binary amorphous alloys, which is observed in the DSC plot for low heating rates [10], is not in the first place due to the glass transition but to rearrangement of atoms in the formation of the nanocrystalline texture. This suggestion is not meant to exclude the possibility that the glass transition makes a minor contribution to the endothermal peak.
Acknowledgments The original melt-spun specimens have been supplied by Prof. S. Steeb and M. Heckele for which we are very grateful. We acknowledge inspiring discussions with Prof. J.-M. Dubois on the structure of metallic glasses.
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