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Local structure in marginal glass forming Al–Sm alloy
Intermetallics 18 (2010) 1676e1682
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Intermetallics journal homepage: www.elsevier.com/locate/intermet
Local structure in marginal glass forming AleSm alloy Y.E. Kalay a, b, *, L.S. Chumbley a, b, M.J. Kramer a, b, I.E. Anderson a, b a b
Ames Laboratory (DOE), Iowa State University, Ames, IA 50011, USA Department of Materials Science and Engineering., Iowa State University, Ames, IA 50011, USA
a r t i c l e i n f o
a b s t r a c t
Article history: Received 11 January 2010 Received in revised form 30 April 2010 Accepted 5 May 2010
The local structure in rapidly quenched Al(100x)Smx (x ¼ 8, 10, 11, and 12) and liquid Al89Sm11 has been investigated using a combination of transmission electron microscopy (TEM) and high-energy synchrotron X-ray diffraction (HEXRD). TEM analysis showed a featureless microstructure with diffuse scattering in rapidly quenched Al(100x)Smx (x ¼ 8, 10, 11, and 12) within the glass formation composition range under the bright field (BF) conditions. Total structure factor analysis of the liquid and as-quenched alloys revealed a pre-peak located well below the main amorphous peak and a distinct side peak. The presence of the pre-peak and the side peak is related to the formation of Sm rich medium range order (MRO) clusters in the liquid that is retained in the as-quenched alloys. Atomic structure models constructed using Reverse Monte Carlo (RMC) simulations from experimentally determined total structure factors and coupled with Voronoi Tessellation analysis indicated icosahedral and deformed bcc-like Voronoi polyhedron (VP) surrounding Al and Sm atoms, respectively. Sm atoms were found to be highly coordinated with Al atoms in the first shell neighborhood. The structural unit sizes corresponding to the extra broad peaks and the first shell neighborhood around Sm atoms have remarkable similarities with the high temperature metastable Al11Sm3 tetragonal phase. The existence of the Sm rich MRO clusters in the as-quenched state is believed to promote the high nucleation density of fcc-Al nanocrystals that form when the material is devitrified by acting as catalyst sites. Ó 2010 Elsevier Ltd. All rights reserved.
Keywords: A. Nanostructured intermetallics B. Glasses, metallic C. Rapid solidification processing E. Simulations, Monte Carlo F. Diffraction
1. Introduction In marginal glass forming amorphous alloys, the first crystallization to occur is often related to ease of nucleation of metastable phases rather than enhanced driving force for the thermodynamically favored phases. This results in the formation of a very high nucleation density of nanocrystals, on the order of 1020e1023 m3, in the amorphous matrix [1e4]. An exact mechanism to explain the presence of such a high nanocrystal density after complete or partial devitrification has not been identified to date. Classical nucleation theory fails to predict the observed number density of nanocrystals [5]. Several mechanisms have been proposed, including a heterogeneous nucleation model [6]; a high density of quenched-in nuclei [3,7]; phase separation in the amorphous state [8,9]; and time-dependent homogeneous nucleation model [10,11]. However, there is still no agreement on the mechanism of highdensity nanocrystal or nanocluster formation. Thus, in order to develop a better understanding of the unusual behaviour of such
* Corresponding author. Ames Laboratory (DOE), Iowa State University, Ames, IA 50011, USA. E-mail address: [email protected] (Y.E. Kalay). 0966-9795/$ e see front matter Ó 2010 Elsevier Ltd. All rights reserved. doi:10.1016/j.intermet.2010.05.005
marginal glass formers, it is important to analyze the as-quenched and liquid structures. Among the marginal glass forming alloys, Al based amorphous alloys [1,7,12e14] have attracted much attention due to their high fracture strengths, exceeding 1 GPa [15,16] and their low densities. In these alloys, an amorphous structure is often observed for Al rich (between 80 and 92 at% Al) AleRE and AleTMeRE (TM: transition metals; RE: rare earth elements) by either mechanical deformation or rapid quenching [1, 12]. The AleSm binary alloy has the widest glass formation range (from 8 to 16 at. % Sm) of all similar AleRE systems [1]. In accordance with this, the binary AleSm system has been chosen as a model to investigate the liquid and the amorphous structures. In a previous study [17] it has been shown that amorphous samples of cold-rolled (transformed in the solid state) Al92Sm8 exhibit a calorimetric signal of glass transition (Tg) that is separated from the first crystallization peak. However, alloys of the same composition quenched from the liquid state using melt spinning do not show any distinct Tg [17]. Fluctuation electron microscopy (FEM) studies identified fcc-Al-like medium range order structures for the as-quenched melt-spun alloy but not for the solid state coldrolled sample [18]. The size of the Al clusters in the as-quenched melt-spun state is between 1 and 2 nm, which makes it difficult to
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detect them using conventional TEM and XRD. Upon heating, these sub-critical size nuclei showed a restricted growth and no distinct glass transition temperature was measured for as-quenched amorphous alloys. Further studies using atom probe tomography (APT) of as-quenched Al90Sm10 revealed clusters of Al atoms in a size range of 2e5 nm [19], indicative of compositional fluctuations in the amorphous state. The APT results also indicated that amorphous AleSm forms a skeleton-like structure, similar to crosslinking in a polymer, and that fcc-Al nucleates between the links of this structure as shown by TEM [20]. Partial devitrification of asquenched melt-spun ribbons [19,20], as well as high pressure gas atomized [21] powders, yields a high nucleation density (population) of fcc-Al, similar to Al92Sm8 as-quenched samples [17]. An interesting observation noted in Al90Sm10 as-quenched meltspun samples [19] is the presence of a pre-peak located well below the primary diffuse scattering peak in HEXRD patterns. Similar types of pre-peaks have previously been observed in other systems [22e24], often being interpreted as topological or chemical ordering occurring to a certain degree in the amorphous phase. However, the structural origin of this pre-peak has never been fully understood or fully described. In this study, as-quenched samples of Al100xSmx (x ¼ 8, 10, 11 and 12) and the liquid structure of Al89Sm11 alloy were analyzed using a combination of high-energy X-ray diffraction (HEXRD) and transmission electron microscopy (TEM). Three-dimensional atomic configurations of amorphous and liquid structures were modeled using the Reverse Monte Carlo (RMC) technique [25], coupled with Voronoi Tessellation analyses [26], based on the HEXRD experiments. The existence of the pre-peak and possible implications that may apply to the observed fcc-Al nanocrystallization with extremely high nucleation density seen in AleSm alloys are discussed.
2. Experimental procedure Ingots of Al(10x)Smx (x ¼ 8, 10, 11, and 12) were prepared by electric arc melting under Ar atmosphere from highly pure Al (99.99 wt%) and Sm (99.9 wt%) elements [27]. Amorphous ribbons with a thickness of 20e30 mm and a width of 1.0e1.2 mm were produced from bulk alloy by Cu block single melt spinning technique under Ar atmosphere at a tangential speed of 30 m/s [27]. X-ray diffraction studies were carried out using high-energy transmission synchrotron X-ray diffraction (HEXRD) at the Advanced Photon Source at Argonne National Laboratory in collaboration with the Midwest Universities Collaborative Access Team (MUCAT). For liquid structure analysis, samples were cast into rods initially then pieces of the rod were inserted into 2 mm diameter, carbon-lined quartz capillaries and sealed in Ar. The sealed quartz capillaries were exposed to 99.586 keV of X-rays corresponding to a wavelength of 0.0124(5) nm. The diffraction data were collected in DebyeeScherrer geometry by a MAR charge coupled device (CCD), with up to 60 s of exposure time. No reaction between the liquid melt and the carbon-lined quartz capillary was observed within the experimental time duration. A diffraction pattern from a similar carbon-lined empty quartz capillary was collected and subtracted from the liquid data sets for the background corrections. The diffraction data from the solid amorphous samples were collected without using any sample holder at room temperature. The raw HEXRD data were corrected for the background and converted to the total structure factor function, S(Q), according to [28,29]
SðQ Þ ¼ 1 þ
P I c ðQ Þ ni¼1 ai jfi ðQ Þj2 Pn 2 j i¼1 ai fi ðQ Þj
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(1)
where Ic(Q) is the coherent scattering intensity normalized to the atomic concentrations, a, and f(Q) is the atomic structure factors for each component in the system as corrected for polarization, absorption, multiple and Compton scattering [30]. In order to construct 3D structural models of amorphous structures, reverse Monte Carlo (RMC) simulations [25,31] were carried out with an RMC analysis program using the S(Q) data derived from the HEXRD experiments. In these simulations random configurations of 20,000 atoms were distributed in a cubic cell with periodic boundary conditions with the proper stoichiometry, density and nearest neighbor distances, as determined from the experimental data and ab-initio calculations [32]. The cut-off distances for the partial pairs for the first shell distances in the RMC calculations were chosen from direct Fourier transforms and abinitio calculations [32] as 0.230, 0.270, and 0.345 nm for amorphous and 0.210, 0.250, and 0.315 nm for liquid to represent AleAl, AleSm and SmeSm cut-offs, respectively. The difference between the measured S(Q) from HEXRD experiments and the calculated SC(Q) from each RMC modeled configuration is determined by
c2o
¼
h i2 n S0 ðQi Þ SC0 ðQi Þ X i¼1
sðQi Þ2
(2)
In order to minimize the c2, approximately 106 iterations were performed for a constant s(Q) value of 0.002. The local atomic environment was investigated by applying Voronoi Tessellation analyses [33]. In this analysis, different types of polyhedrons are defined with 5 digit indices in the form of Ni (i ¼ 3e7) where N represents the number of faces having “i” edges around the central P atom (Al or Sm). The summation of N ð i NÞ gives the coordination number around a specific atom. TEM analyses were performed using an FEI Inc. Tecnai G2 F20 scanning/transmission electron microscope. Samples for electron microscopy were thinned using double jet polishing at 248 K with a solution of 3 vol% HCl, 36 vol% methanol, and distilled water. 3. Experimental results 3.1. As-quenched and liquid structures The rapidly solidified microstructures of melt-spun ribbons as revealed using TEM are shown in Fig. 1. Bright field (BF) images and selected area electron diffraction (SAED) patterns (inset) of each asquenched Al(100x)Smx (x ¼ 8, 10, 11, and 12) alloy ribbon show a featureless microstructure, while SAED patterns indicate a diffuse ring, typical for an amorphous phase. Fig. 2 shows the total structure factor (S(Q)), determined for amorphous and liquid alloys. Note that S(Q) does not show any sharp crystalline peaks in the diffraction patterns. However, two broad peaks appear approximately at scattering vector magnitudes of 12.6 nm1 and 12.7 nm1 (pre-peak) and 33.4 nm1 for amorphous and 32.5 nm1 for liquid (side peak) where Q ¼ 4psinðqÞ=l. The negative values of the S(Q) in the low-Q region are due to a large difference in Al and Sm atomic scattering factors [30]. Since these data were collected without using any sample holder, these extra peaks should be an intrinsic feature of the AleSm system. Comparison of the peaks to the positions expected for fcc-Al diffraction (dashed lines of Fig. 2) reveals that this broad scattering is not due to fcc-Al. Fig. 2b shows the HEXRD pattern of liquid Al89Sm11 alloy at a temperature of 1313 K in comparison with that
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Y.E. Kalay et al. / Intermetallics 18 (2010) 1676e1682 Table 1 Peak positions, the structural unit sizes, and the correlation lengths for the amorphous and liquid alloys. Structural unit size (nm)
Correlation length (nm)
Side peak (10 nm1)
Amorphous Al92Sm8 1.26 Al90Sm10 1.26 Al89Sm11 1.27 Al88Sm12 1.27
0.499 0.498 0.495 0.495
2.4 2.5 2.6 2.5
3.35 3.33 3.34 3.34
Liquid Al89Sm11 1.28
0.493
1.5
3.25
Alloy
Pre-peak (10 nm1)
compared to the amorphous material. Table 1 summarizes peak positions, structural unit sizes, and correlation lengths determined for the pre-peak and the side peak observed in amorphous and liquid samples.
3.2. Reverse Monte Carlo (RMC) simulations
Fig. 1. BF images with SAED (inset) showing a featureless matrix and diffuse rings for (a) Al92Sm8, (b) Al90Sm10, (c) Al89Sm11 and (d) Al88Sm12, respectively.
obtained from the amorphous solid. The most intriguing observation is that the pre-peak in the liquid appears in a similar position to that seen for the solid amorphous alloy. Thus, the appearance of the two broad scattering peaks seen in the amorphous solid may be an indication of an ordered structure other than fcc-Al in rapidly solidified ribbon. Of particular interest is the pre-peak, located well below the major amorphous peak (Q ¼ 26.1 nm1) that is also present in the liquid. An expression similar to the Scherrer equation used for crystalline materials is often applied for disordered materials in order to estimate the size of any short or medium range ordered structure. This expression is given as D z 2p/DQpre-peak, where D is the correlation length and DQ is the half-width of the pre-peak on the S (Q)eQ graph. The structural unit size (R) corresponding to the prepeak can be estimated using the relationship R z 2p/Qpre-peak [34]. In order to identify the position and the width of the pre-peak, a Lorentzian function was fit to the low-Q part of the S(Q)eQ data [35]. In doing this, the intensity of the pre-peak is seen to decrease, but its position remains almost constant for the liquid structure as
Among the compositions within the glass formation range, Al89Sm11 was further investigated in solid and molten states using reverse Monte Carlo simulations. Fig. 3 shows comparisons of the total structure factor measured from HEXRD experiments to results obtained when the data are fit using the indirect RMC method for as-quenched and liquid Al89Sm11 alloys to produce the same S(Q). A reasonably good fit is obtained in both the low (including the prepeak) and the high-Q regions. Fig. 4 shows the partial structure factors (SAleAl(Q), SAleSm(Q), and SSmeSm(Q)) determined for as-quenched and molten alloys. In the low-Q region (Q w 12.7 nm1) of the partial structure factors, a strong pre-peak (arrowed) is seen for SSmeSm(Q). Interestingly, the pre-peak is absent from the SAleAl(Q), and SAleSm(Q) indicating that the pre-peak seen in the total structure factor arises due to SmeSm containing correlations. The complex polyhedral structure in the local atomic configurations generated by RMC was further investigated by Voronoi Tessellation analyses, as shown in Fig. 5. Thirteen and 14-coordinated polyhedrons with indices <0,3,6,4,0> and <0,2,8,4,0> are dominant for Al centered structures. Voronoi polyhedrons (VP) with such indices correspond to a deformed bcc structure [36]. It is also found that some icosahedral-like clusters with indices <0,0,12,0,0>, <0,1,10,2,0>, or <0,2,8,2,0> and coordination of 12 and 13 atoms are formed around Al. The Sm centered polyhedrons almost always have higher coordination as compared to Al centered
Fig. 2. Total structure function S(Q) data for (a) as-quenched Al(100x)Smx (x ¼ 8, 10, 11 and 12) and as-quenched amorphous with liquid Al89Sm11 alloy at 1313 K.
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Fig. 3. Total structure function data measured (open circles) and calculated using RMC (solid line) for as-quenched and liquid Al89Sm11 at 1313 K. The line at the bottom shows the difference between measured and calculated S(Q).
polyhedrons. The indices corresponding to deformed bcc structures, such as <0,1,10,4,0>, <0,2,8,6,0>, <0,1,10,5,0> and <0,2,8,5,0>, are the most frequent polyhedral forming around Sm atoms. The distribution of VP types is variable for liquid states, and no specific type of indice stands out. Despite the highly variable distribution of VP indices, particularly for the liquid state, coordination numbers (CN) numbers are quite consistent. Fig. 6 shows the histogram of CN for Al and Sm center clusters in amorphous and liquid states according to Voronoi Tessellation analysis. Sm centered clusters are highly coordinated, on the order of 16 atoms in the first near-neighbor shell. The first near-neighbor shell calculated with Voronoi Tessellation analysis includes both Sm and Al atoms. The fraction of Al atoms surrounding Sm centered atoms is on the order of 0.94 and 0.89 for amorphous and liquid states respectively. Therefore, according to Voronoi Tessellation analysis the Sm atoms are highly coordinated with Al both in liquid and amorphous states. CN for Al center cluster is on the order of 14 and the fraction of Al in the first neighborhood is around 0.87 for both amorphous and liquid states. Table 2 summarizes the fraction of Al atoms in the first neighborhood for Al and Sm center clusters for different CN. 4. Discussion According to previous high resolution electron microscopy studies [19,20], as-quenched Al90Sm10 alloys produced by high
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pressure gas atomization and melt spinning with wheel speeds of 30 and 40 m/s contain clusters of fcc-Al in a size range of 2e5 nm embedded in an amorphous matrix. In another study [18], clusters of Al that are showing fcc like medium range order (MRO) and having sizes in the range of 0.8e1.5 nm were resolved using fluctuation electron microscopy (FEM) in melt-spun Al92Sm8 produced at a wheel speed of 55 m/s. These results clearly show that the size of fcc-Al-like clusters decreases with increasing cooling rate (i.e. wheel speeds of melt spinner) and MRO or nanocrystals of an fccAl-like structure exist in AleSm alloys that are quenched from the liquid state within the glass formation range, at least for the conditions used in these studies. These pre-existing fcc-Al nanoclusters and the others that nucleate [20] show a restricted growth upon annealing below their crystallization temperatures (Tx) and a high number of density of nanocrystals on the order of 1021e1023 m3 results in the amorphous matrix. The small size of clusters in the as-quenched state makes them difficult to detect in conventional transmission electron microscopy and X-ray diffraction analyses. While fcc-Al nanoclusters cannot be detected with HEXRD, as shown in Fig. 2, an intriguing observation is the existence of a prepeak (Q w 12.6 nm1) at a position well below the major diffuse scattering position for all four as-quenched compositions. As shown in Fig. 2, the pre-peak is a distinct feature in the total structure factor function S(Q) calculated for as-quenched Al(100x)Smx (x ¼ 8, 10, 11, and 12) and liquid Al89Sm11. Similar types of pre-peaks have been reported for other alloy systems [22e24]. The physical origin of such pre-peaks is not clear and their existence is usually attributed to the formation of MRO structures [37,38]. These data were collected at room temperature without using any sample holder, thus eliminating any possible effect due to the quartz capillaries that were used at higher temperatures. These observations indicate that the pre-peak is not an anomaly of the manner in which the HEXRD data were obtained. There is a slight increase in the intensity of the pre-peak with increasing Sm content, which is more pronounced for the second broad peak located approximately at 33.4 nm1. A similar pre-peak is also observed in liquid phase. Fig. 2b shows the total structure factor function after background corrections for liquid Al89Sm11 at 1313 K, approximately 80 K above the liquidus temperature. A pre-peak is clearly seen in both graphs and the position is in good agreement with the pre-peaks observed at room temperatures in the asquenched samples. The correlation lengths and structural unit sizes calculated from the position of the pre-peak for as-quenched and liquid structures are summarized in Table 1. Note that the correlation lengths calculated from pre-peak broadening differ for solid
Fig. 4. Partial structure factors (SAleAl(Q), SAleSm(Q), and SSmeSm(Q)) calculated for (a) as-quenched and (b) liquid Al89Sm11 at 1313 K.
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Fig. 5. Most populated VP histograms of (a, c) Al and (b, d) Sm centered structure models for as-quenched (a, b) and liquid (c, d) Al89Sm11 at 1313 K.
and liquid structures. For as-quenched solids the correlation length is on the order of 2.5 nm and is almost constant with respect to composition. However, correlation length is much smaller in the liquid, being on the order of 1.5 nm. The structural unit sizes calculated from the positions of the prepeaks for the as-quenched samples are almost constant with respect to compositions and are similar to that of liquid Al89Sm11. This shows that the clusters formed in the liquid state might be retained in the as-quenched state with an increase in correlation
length. The positions of the extra peaks observed in liquid and asquenched states are close to some of the major diffraction peaks in a high temperature metastable body centered tetragonal phase, Al11Sm3. This compound is stable between 1351 and 1733 K [39]. The interplanar spacing in t-Al11Sm3 for (002) and (211) planes are given as 0.495 and 0.188 nm, respectively. These planar spacings are close in value to those calculated from the positions of the prepeaks and side peaks for as-quenched Al(100x)Smx (x ¼ 8, 10, 11 and 12) and liquid Al89Sm11.Thus, the broad peaks observed in the total
Fig. 6. Coordination numbers calculated for (a) Al and (b) Sm centered VP models for as-quenched and liquid Al89Sm11 at 1313 K.
Y.E. Kalay et al. / Intermetallics 18 (2010) 1676e1682 Table 2 The fraction of Al atoms in the first shell neighborhood around Sm and Al center atoms. CN
Sm centered
Al centered
Amorphous 17 16 15 14 13 12
0.94 0.94 0.93 0.93 0.93 0.94
0.90 0.89 0.88 0.87 0.86 0.85
Liquid 17 16 15 14 13 12
0.91 0.90 0.90 0.89 0.88 0.87
0.88 0.88 0.88 0.87 0.87 0.86
structure factor function are consistent with a pseudo-tetragonal MRO Al11Sm3 structure. Moreover, according to Voronoi Tessellation analysis, the local atomic configuration around Sm atoms calculated using RMC for the as-quenched state have structural similarities with the first shell neighborhood of Sm atoms in high temperature metastable Al11Sm3 tetragonal phase. Fig. 7a shows the atomic configuration of this phase, which has the I4/mmm space group symmetry. This structure consists of Sm atoms surrounded by 8 Al atoms at a distance of 0.325 nm and another 8 Al atoms at a distance of 0.327 nm. Therefore, every Sm atom in this structure is highly coordinated with Al in their first and second shells and 16 Al atoms exist at a distance of 0.327 nm or less (Fig. 7b). Moreover, in a study where the local atomic structure of amorphous Al92Sm8 was investigated using X-ray absorption fine structure (XAFS) [40], the best fit to the experimental XAFS data were obtained by considering a “single shell” structure of 16 Al atoms surrounding Sm atoms within the first coordination shell. The average coordination number (CN) determined for Sm atoms is on the order of 16 according to Voronoi analyses for both liquid and amorphous states having 90e94% Al atoms within the first shell coordination (Table 2). Therefore, a high coordination of Al atoms around Sm atoms is detected in the first shell structure that is similar to tetragonal Al11Sm3 metastable compound (Fig. 7). According to HEXRD observations of liquid and as-quenched structures a pre-peak is resolved well below the major amorphous peak in total structure factors. Further RMC simulations coupled with Voronoi Tessellation analysis reveal highly coordinated Sm centered clusters in the as-quenched structure and their first shell neighborhood is structurally similar to the high temperature metastable Al11Sm3 tetragonal phase. This particular alloy system has been reported to devitrify to a very high number of nucleation
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density of nanocrystalline fcc-Al on the order of 1020e1023 m3 if the alloy is quenched from liquid [3,17]. According to high temperature HEXRD, the pre-peak and side peaks observed in the liquid and as-quenched states appear at atomic distances in close correspondence to the spacing of (002) and (211) planes in t-Al11Sm3 compound. This would argue for a persistence of pseudo-tetragonal Al11Sm3 atomic ordering in the molten state and subtle growth of this MRO during quenching from liquid. The first shell neighborhood around Sm atoms within this compound consists of highly coordinated Al atoms and the clustering in the melt is believed to form due to strong interactions between Al and Sm atoms. However, as shown by the partial structure factor (Fig. 4), the pre-peak is particularly originated by SmeSm pairs, which require a longer length correlation as compared to the first shell neighborhood. The correlation length of this MRO structure observed in as-quenched state is between 2 and 3 nm according to HEXRD results. This corresponds to 4e6 layers of repeating (002) planes, if the structure is assumed to be a pseudotetragonal MRO Al11Sm3. In contrast the pre-peak in the liquid has a correlation length of only 1.5 nm or 2e3 layers. These clusters in the liquid essentially act like distinct “superatoms”, with a broad size distribution and huge diameters as compared to individual Al and Sm atoms in the matrix. Upon rapid solidification it is hypothesized that these “superatoms” cause a local atomic size effect, similar to that seen in multi-component bulk amorphous alloys, preventing the complete crystallization of the Al rich matrix. This results in the as-quenched structure being divided into nanoscale regions by a network of MRO structure. Depending on the size of these regions a certain percentage of fcc-Al clusters will reach the critical nucleus size with the remainder staying below this size. While admittedly speculative, this explanation agrees with results obtained for Al90Sm10 using three-dimensional atom probe tomography (3D APT) [19]. The effect of cooling rate has a dominant role on the percentage of Al that reaches the critical nucleation size. However, it should be noted that cooling rate has possibly a minor effect on the presence of MRO in the as-quenched state since a prepeak due to Sm rich clustering is also observed for the melt. Therefore, ordering in the liquid needs be studied directly with respect to initial melt temperature, and a more accurate representation of the liquid structure is required. When as-quenched alloys are annealed to devitrify to a high number of density nanocrystals, the MRO appears to act as a catalyst site for nucleation of fcc-Al. Previous results within this system [20] showed that long-range diffusion between fcc-Al clusters is blocked by the Sm rich MRO structure in the matrix, resulting in restricted growth of fcc-Al. Al rich regions with sizes below the critical size for homogeneous nucleation can nucleate via a heterogeneous nucleation mechanism in the vicinity of the MRO clusters. A more detailed study that includes atom probe tomography and scanning transmission electron microscopy equipped with a high angle annular dark field detector is in progress. Such a study should be helpful for identifying the spatial distribution of the MRO and fcc-Al in the as-quenched state, to help elucidate the effect of MRO on producing the observed high nucleation density of fcc-Al in these alloys. 5. Conclusion
Fig. 7. (a) Crystal structure and (b) first and second shells neighborhood of Sm atoms in high temperature metastable tetragonal phase Al11Sm3 with I4/mmm space group symmetry. Red and black colors represent Al and Sm atoms, respectively.
Amorphous alloys of Al(100x)Smx (x ¼ 8, 10, 11, and 12) were produced using the melt spinning technique at a wheel speed of 30 m/s. Detailed analysis of the as-quenched structure by HEXRD revealed a pre-peak and a side peak in total structure factor for both amorphous and liquid states. Partial structure factor analysis showed that the pre-peak arises due to SmeSm correlations. The formation of the pre-peak is related to the formation of clusters in
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the liquid. After the liquid alloy is quenched, the position of the prepeak remains almost constant, indicating that the clusters in the liquid are retained and showed subtle growth during quenching from liquid. The Voronoi analysis of local atomic structures based on RMC simulation suggests a deformed bcc-like structure surrounding Sm with CN on the order of 16 atoms for both amorphous and liquid states. These experimental results and simulations indicate the first shell neighborhood of Sm atoms in asquenched states is structurally similar to the high temperature metastable Al11Sm3 tetragonal phase. The pre-peak, indicating a SmeSm correlation, was found to result due to a higher order correlation within the liquid. These Sm rich clusters in the liquid are thought of as creating Sm and Al rich regions. When as-quenched alloys are annealed, the Sm rich clusters act as catalyst sites promoting heterogeneous nucleation of fcc-Al, while inhibiting long-range diffusion processes between fcc-Al clusters. Acknowledgments Appreciation is expressed to Shaogang Hao for his valuable help in ab-initio calculations. The work at Ames Laboratory was supported by the United States Department of Energy (USDOE), Office of Science (OS), Office of Basic Energy Science (BES), under Ames Laboratory Contract No. DE-AC02-07CH11358, the high-energy X-ray work at the MUCAT sector of the APS was supported by the US Department of Energy, Office of Science, and Basic Energy Sciences under Contract No. DE-AC02-06CH11357. The assistance of Materials Preparation Center of the Ames Laboratory is acknowledged for supplying our samples [27]. References [1] Inoue A. Amorphous, nanoquasicrystalline and nanocrystalline alloys in Albased systems. Prog Mater Sci 1998;43:365e520. [2] Foley JC, Perepezko JH. The devitrification of AleYeFe amorphous alloys. J Non-Cryst Solids 1996;205e207:559e62. [3] Foley JC, Allen DR, Perepezko JH. Analysis of nanocrystal development in AleYeFe and AleSm glasses. Scr Mater 1996;35:655e60. [4] Hirata A, Hirotsu Y, Matsubara E, Ohkubo T, Hono K. Mechanism of nanocrystalline microstructure formation in amorphous FeeNbeB Alloys. Phys Rev B 2006;74:184204. [5] Schroers J, Busch R, Masuhr A, Johnson WL. Continuous refinenement of the microstructure during crystallization of supercooled Zr41Ti14Cu12Ni10Be23 melts. Appl Phys Lett 1999;74:2806e8. [6] Hono K, Ping DH, Ohnuma M, Onodera H. Cu clustering and Si partitioning in the early crystallization stage of an Fe73.5Si13.5B9Nb3Cu1 amorphous alloy. Acta Mater 1999;47:997e1006. [7] Allen DR, Foley JC, Perepezko JH. Nanocrystal development during primary crystallization of amorphous alloys. Acta Mater 1998;46:431e40. [8] Wang XL, Almer J, Liu CT, Wang YD, Zhao JK, Stoica AD, et al. In situ synchrotron study of phase transformation behaviors in bulk metallic glass by simultaneous diffraction and small angle scattering. Phys Rev Lett 2003;91:265501e5. [9] Schneider S, Thiyagarajan P, Johnson WL. Formation of nanocrystals based on decomposition in the amorphous Zr41.2Ti13.8Ni10Be22.5 alloy. Appl Phys Lett 1995;68:493e5. [10] Kelton KF. Time-dependent nucleation in partitioning transformations. Acta Mater 2000;48:1967e80. [11] Kelton KF. Kinetic model for nucleation in partitioning systems. J Non-Cryst Solids 2000;274:147e54.
[12] He Y, Poon SJ, Shiflet GJ. Synthesis and properties of metallic glasses that contain aluminum. Science 1988;241:1640e2. [13] Greer AL. Metallic glasses. Science 1995;267:1947e53. [14] Stratton WG, Hamann J, Perepezko JH, Voyles PM. Electron beam induced crystallization of amorphous Al-based alloys in the TEM. Intermetallics 2006;14:1061e5. [15] Kulik T. Nanocrystallization of metallic glasses. J Non-Cryst Solids 2001;287:145e61. [16] Kim HS, Warren PJ, Cantor B, Lee HR. Mechanical properties of partially crystallized aluminum based amorphous alloys. Nanostruct Mater 1999;11:241e7. [17] Wilde G, Sieber H, Perepezko JH. Glass formation versus nanocrystallization in an Al92Sm8 alloy. Scr Mater 1999;40:779e83. [18] Stratton WG, Hamann J, Perepezko JH, Voyles PM, Mao X, Khare SV. Aluminum nanoscale order in amorphous Al92Sm8 measured by fluctuation electron microscopy. Appl Phys Lett 2005;86:141910e3. [19] Kalay YE, Chumbley LS, Anderson IE. Crystallization behavior in a highly driven marginal glass forming alloy. J Non-Cryst Solids 2008;354:3040e8. [20] Kalay YE, Yeager C, Chumbley LS, Kramer MJ, Anderson IE. Initial crystallization in a nanostructured Al-Sm rare earth alloy. J Non-Cryst Solids, in press, doi:10.1016/j.jnoncrysol.2010.05.005. [21] Kalay YE, Chumbley LS, Anderson IE. Characterization of a marginal glass former alloy solidified in gas atomized powders. Mater Sci Eng A 2008;490:72e80. [22] Yang L, Jiang JZ, Saksl K, Franz H. Origin of the pre-peak in Zr70Cu29Pd1 metallic glass. J Phys Condens Matter 2007;19:1e6. [23] Lin Z, Youshi W, Xiaufang B, Hui L, Weimin W, Jinguo L, et al. Origin of the prepeak in structure factors of liquid and amorphous AleFeeCe alloys. J Phys Condens Matter 1999;11:7959e69. [24] Fang J, Youshi W, Yuanchang S, Chuanjiang Z, Zhiqian Z. Stability of the structure corresponding to the prepeak and amorphous Al90Tm5Ce5 alloys. Mater Sci Eng A 2001;340:212e5. [25] McGreevy RL, Howe MA. RMC: modeling disordered structures. Annu Rev Mater Res 1992;22:217e42. [26] Finney JL. Random packings and the structure of simple liquids. II. The molecular geometry of simple liquids. Proc Roy Soc London Ser A Mathem Phys Sci 1970;319:495e507. [27] Materials Preparation Center, A.L., US DOE Basic Energy Sciences, Ames, IA, USA, www.mpc.ameslab.gov. Avaliable from:. [28] Qiu X, Thompson JW, Billinge SJL. PDFgetX2: a GUI-driven program to obtain the pair distribution function from X-ray powder diffraction data. J Appl Crystallogr 2004;37:678. [29] Kramer MJ, Sordelet DJ. Polymorphism in the short-range order of Zr70Pd30 metallic glasses. J Non-Cryst Solids 2005;351:1586e93. [30] Waseda Y. The structure of non-crystalline materials. Mc-Graw-Hill Inc.; 1980. [31] McGreevy RL, Pusztai L. Reverse Monte Carlo simuation: a new technique for the determination of disordered structures. Mol Simul 1988;1:359e67. [32] Shaogang H. Unpublished work, Ames, 2008. [33] Medvedev NN. The algorithm for three-dimensional Voronoi polyhedra. J Comput Sci 1986;67:223e9. [34] Vateva E, Savova E. New medium-range order features in GeeSbeS glasses. J Non-Cryst Solids 1995;145:192e3. [35] Sokolov AP, Kisliuk A, Soltwisch M, Quitmann D. Medium-range order in glasses: comparison of Raman and diffraction measurements. Phys Rev Lett 1992;69:1540e3. [36] Hirata A, Hirotsu Y, Ohkubo T, Hanada T, Bengus VZ. Compositional dependence of local atomic structures in amorphous Fe100xBx(x ¼ 14,17,20) alloys studied by electron diffraction and high resolution electron microscopy. Phys Rev B 2006;74:214206e15. [37] Elliott SR. Medium-range structural order in covalent amorphous solids. Nature 1991;354:445e52. [38] Wang SY, Wang CZ, Li MZ, Huang L, Ott RT, Kramer MJ, et al. Short- and medium-range order in a Zr73Pt27 glass: experimental and simulation studies. Phys Rev B 2008;78:184204e13. [39] Massalski TB. Binary alloy phase diagrams, vol. 1; 1987. 166 p. [40] Zalewski W, Antonowicz J, Bacewicz R, Latuch J. Local atomic order in Albased metallic glasses studied using XAFS method. J Alloys Compd 2009;468:40e6.
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Local structure in marginal glass forming AleSm alloy Y.E. Kalay a, b, *, L.S. Chumbley a, b, M.J. Kramer a, b, I.E. Anderson a, b a b
Ames Laboratory (DOE), Iowa State University, Ames, IA 50011, USA Department of Materials Science and Engineering., Iowa State University, Ames, IA 50011, USA
a r t i c l e i n f o
a b s t r a c t
Article history: Received 11 January 2010 Received in revised form 30 April 2010 Accepted 5 May 2010
The local structure in rapidly quenched Al(100x)Smx (x ¼ 8, 10, 11, and 12) and liquid Al89Sm11 has been investigated using a combination of transmission electron microscopy (TEM) and high-energy synchrotron X-ray diffraction (HEXRD). TEM analysis showed a featureless microstructure with diffuse scattering in rapidly quenched Al(100x)Smx (x ¼ 8, 10, 11, and 12) within the glass formation composition range under the bright field (BF) conditions. Total structure factor analysis of the liquid and as-quenched alloys revealed a pre-peak located well below the main amorphous peak and a distinct side peak. The presence of the pre-peak and the side peak is related to the formation of Sm rich medium range order (MRO) clusters in the liquid that is retained in the as-quenched alloys. Atomic structure models constructed using Reverse Monte Carlo (RMC) simulations from experimentally determined total structure factors and coupled with Voronoi Tessellation analysis indicated icosahedral and deformed bcc-like Voronoi polyhedron (VP) surrounding Al and Sm atoms, respectively. Sm atoms were found to be highly coordinated with Al atoms in the first shell neighborhood. The structural unit sizes corresponding to the extra broad peaks and the first shell neighborhood around Sm atoms have remarkable similarities with the high temperature metastable Al11Sm3 tetragonal phase. The existence of the Sm rich MRO clusters in the as-quenched state is believed to promote the high nucleation density of fcc-Al nanocrystals that form when the material is devitrified by acting as catalyst sites. Ó 2010 Elsevier Ltd. All rights reserved.
Keywords: A. Nanostructured intermetallics B. Glasses, metallic C. Rapid solidification processing E. Simulations, Monte Carlo F. Diffraction
1. Introduction In marginal glass forming amorphous alloys, the first crystallization to occur is often related to ease of nucleation of metastable phases rather than enhanced driving force for the thermodynamically favored phases. This results in the formation of a very high nucleation density of nanocrystals, on the order of 1020e1023 m3, in the amorphous matrix [1e4]. An exact mechanism to explain the presence of such a high nanocrystal density after complete or partial devitrification has not been identified to date. Classical nucleation theory fails to predict the observed number density of nanocrystals [5]. Several mechanisms have been proposed, including a heterogeneous nucleation model [6]; a high density of quenched-in nuclei [3,7]; phase separation in the amorphous state [8,9]; and time-dependent homogeneous nucleation model [10,11]. However, there is still no agreement on the mechanism of highdensity nanocrystal or nanocluster formation. Thus, in order to develop a better understanding of the unusual behaviour of such
* Corresponding author. Ames Laboratory (DOE), Iowa State University, Ames, IA 50011, USA. E-mail address: [email protected] (Y.E. Kalay). 0966-9795/$ e see front matter Ó 2010 Elsevier Ltd. All rights reserved. doi:10.1016/j.intermet.2010.05.005
marginal glass formers, it is important to analyze the as-quenched and liquid structures. Among the marginal glass forming alloys, Al based amorphous alloys [1,7,12e14] have attracted much attention due to their high fracture strengths, exceeding 1 GPa [15,16] and their low densities. In these alloys, an amorphous structure is often observed for Al rich (between 80 and 92 at% Al) AleRE and AleTMeRE (TM: transition metals; RE: rare earth elements) by either mechanical deformation or rapid quenching [1, 12]. The AleSm binary alloy has the widest glass formation range (from 8 to 16 at. % Sm) of all similar AleRE systems [1]. In accordance with this, the binary AleSm system has been chosen as a model to investigate the liquid and the amorphous structures. In a previous study [17] it has been shown that amorphous samples of cold-rolled (transformed in the solid state) Al92Sm8 exhibit a calorimetric signal of glass transition (Tg) that is separated from the first crystallization peak. However, alloys of the same composition quenched from the liquid state using melt spinning do not show any distinct Tg [17]. Fluctuation electron microscopy (FEM) studies identified fcc-Al-like medium range order structures for the as-quenched melt-spun alloy but not for the solid state coldrolled sample [18]. The size of the Al clusters in the as-quenched melt-spun state is between 1 and 2 nm, which makes it difficult to
Y.E. Kalay et al. / Intermetallics 18 (2010) 1676e1682
detect them using conventional TEM and XRD. Upon heating, these sub-critical size nuclei showed a restricted growth and no distinct glass transition temperature was measured for as-quenched amorphous alloys. Further studies using atom probe tomography (APT) of as-quenched Al90Sm10 revealed clusters of Al atoms in a size range of 2e5 nm [19], indicative of compositional fluctuations in the amorphous state. The APT results also indicated that amorphous AleSm forms a skeleton-like structure, similar to crosslinking in a polymer, and that fcc-Al nucleates between the links of this structure as shown by TEM [20]. Partial devitrification of asquenched melt-spun ribbons [19,20], as well as high pressure gas atomized [21] powders, yields a high nucleation density (population) of fcc-Al, similar to Al92Sm8 as-quenched samples [17]. An interesting observation noted in Al90Sm10 as-quenched meltspun samples [19] is the presence of a pre-peak located well below the primary diffuse scattering peak in HEXRD patterns. Similar types of pre-peaks have previously been observed in other systems [22e24], often being interpreted as topological or chemical ordering occurring to a certain degree in the amorphous phase. However, the structural origin of this pre-peak has never been fully understood or fully described. In this study, as-quenched samples of Al100xSmx (x ¼ 8, 10, 11 and 12) and the liquid structure of Al89Sm11 alloy were analyzed using a combination of high-energy X-ray diffraction (HEXRD) and transmission electron microscopy (TEM). Three-dimensional atomic configurations of amorphous and liquid structures were modeled using the Reverse Monte Carlo (RMC) technique [25], coupled with Voronoi Tessellation analyses [26], based on the HEXRD experiments. The existence of the pre-peak and possible implications that may apply to the observed fcc-Al nanocrystallization with extremely high nucleation density seen in AleSm alloys are discussed.
2. Experimental procedure Ingots of Al(10x)Smx (x ¼ 8, 10, 11, and 12) were prepared by electric arc melting under Ar atmosphere from highly pure Al (99.99 wt%) and Sm (99.9 wt%) elements [27]. Amorphous ribbons with a thickness of 20e30 mm and a width of 1.0e1.2 mm were produced from bulk alloy by Cu block single melt spinning technique under Ar atmosphere at a tangential speed of 30 m/s [27]. X-ray diffraction studies were carried out using high-energy transmission synchrotron X-ray diffraction (HEXRD) at the Advanced Photon Source at Argonne National Laboratory in collaboration with the Midwest Universities Collaborative Access Team (MUCAT). For liquid structure analysis, samples were cast into rods initially then pieces of the rod were inserted into 2 mm diameter, carbon-lined quartz capillaries and sealed in Ar. The sealed quartz capillaries were exposed to 99.586 keV of X-rays corresponding to a wavelength of 0.0124(5) nm. The diffraction data were collected in DebyeeScherrer geometry by a MAR charge coupled device (CCD), with up to 60 s of exposure time. No reaction between the liquid melt and the carbon-lined quartz capillary was observed within the experimental time duration. A diffraction pattern from a similar carbon-lined empty quartz capillary was collected and subtracted from the liquid data sets for the background corrections. The diffraction data from the solid amorphous samples were collected without using any sample holder at room temperature. The raw HEXRD data were corrected for the background and converted to the total structure factor function, S(Q), according to [28,29]
SðQ Þ ¼ 1 þ
P I c ðQ Þ ni¼1 ai jfi ðQ Þj2 Pn 2 j i¼1 ai fi ðQ Þj
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(1)
where Ic(Q) is the coherent scattering intensity normalized to the atomic concentrations, a, and f(Q) is the atomic structure factors for each component in the system as corrected for polarization, absorption, multiple and Compton scattering [30]. In order to construct 3D structural models of amorphous structures, reverse Monte Carlo (RMC) simulations [25,31] were carried out with an RMC analysis program using the S(Q) data derived from the HEXRD experiments. In these simulations random configurations of 20,000 atoms were distributed in a cubic cell with periodic boundary conditions with the proper stoichiometry, density and nearest neighbor distances, as determined from the experimental data and ab-initio calculations [32]. The cut-off distances for the partial pairs for the first shell distances in the RMC calculations were chosen from direct Fourier transforms and abinitio calculations [32] as 0.230, 0.270, and 0.345 nm for amorphous and 0.210, 0.250, and 0.315 nm for liquid to represent AleAl, AleSm and SmeSm cut-offs, respectively. The difference between the measured S(Q) from HEXRD experiments and the calculated SC(Q) from each RMC modeled configuration is determined by
c2o
¼
h i2 n S0 ðQi Þ SC0 ðQi Þ X i¼1
sðQi Þ2
(2)
In order to minimize the c2, approximately 106 iterations were performed for a constant s(Q) value of 0.002. The local atomic environment was investigated by applying Voronoi Tessellation analyses [33]. In this analysis, different types of polyhedrons are defined with 5 digit indices in the form of Ni (i ¼ 3e7) where N represents the number of faces having “i” edges around the central P atom (Al or Sm). The summation of N ð i NÞ gives the coordination number around a specific atom. TEM analyses were performed using an FEI Inc. Tecnai G2 F20 scanning/transmission electron microscope. Samples for electron microscopy were thinned using double jet polishing at 248 K with a solution of 3 vol% HCl, 36 vol% methanol, and distilled water. 3. Experimental results 3.1. As-quenched and liquid structures The rapidly solidified microstructures of melt-spun ribbons as revealed using TEM are shown in Fig. 1. Bright field (BF) images and selected area electron diffraction (SAED) patterns (inset) of each asquenched Al(100x)Smx (x ¼ 8, 10, 11, and 12) alloy ribbon show a featureless microstructure, while SAED patterns indicate a diffuse ring, typical for an amorphous phase. Fig. 2 shows the total structure factor (S(Q)), determined for amorphous and liquid alloys. Note that S(Q) does not show any sharp crystalline peaks in the diffraction patterns. However, two broad peaks appear approximately at scattering vector magnitudes of 12.6 nm1 and 12.7 nm1 (pre-peak) and 33.4 nm1 for amorphous and 32.5 nm1 for liquid (side peak) where Q ¼ 4psinðqÞ=l. The negative values of the S(Q) in the low-Q region are due to a large difference in Al and Sm atomic scattering factors [30]. Since these data were collected without using any sample holder, these extra peaks should be an intrinsic feature of the AleSm system. Comparison of the peaks to the positions expected for fcc-Al diffraction (dashed lines of Fig. 2) reveals that this broad scattering is not due to fcc-Al. Fig. 2b shows the HEXRD pattern of liquid Al89Sm11 alloy at a temperature of 1313 K in comparison with that
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Y.E. Kalay et al. / Intermetallics 18 (2010) 1676e1682 Table 1 Peak positions, the structural unit sizes, and the correlation lengths for the amorphous and liquid alloys. Structural unit size (nm)
Correlation length (nm)
Side peak (10 nm1)
Amorphous Al92Sm8 1.26 Al90Sm10 1.26 Al89Sm11 1.27 Al88Sm12 1.27
0.499 0.498 0.495 0.495
2.4 2.5 2.6 2.5
3.35 3.33 3.34 3.34
Liquid Al89Sm11 1.28
0.493
1.5
3.25
Alloy
Pre-peak (10 nm1)
compared to the amorphous material. Table 1 summarizes peak positions, structural unit sizes, and correlation lengths determined for the pre-peak and the side peak observed in amorphous and liquid samples.
3.2. Reverse Monte Carlo (RMC) simulations
Fig. 1. BF images with SAED (inset) showing a featureless matrix and diffuse rings for (a) Al92Sm8, (b) Al90Sm10, (c) Al89Sm11 and (d) Al88Sm12, respectively.
obtained from the amorphous solid. The most intriguing observation is that the pre-peak in the liquid appears in a similar position to that seen for the solid amorphous alloy. Thus, the appearance of the two broad scattering peaks seen in the amorphous solid may be an indication of an ordered structure other than fcc-Al in rapidly solidified ribbon. Of particular interest is the pre-peak, located well below the major amorphous peak (Q ¼ 26.1 nm1) that is also present in the liquid. An expression similar to the Scherrer equation used for crystalline materials is often applied for disordered materials in order to estimate the size of any short or medium range ordered structure. This expression is given as D z 2p/DQpre-peak, where D is the correlation length and DQ is the half-width of the pre-peak on the S (Q)eQ graph. The structural unit size (R) corresponding to the prepeak can be estimated using the relationship R z 2p/Qpre-peak [34]. In order to identify the position and the width of the pre-peak, a Lorentzian function was fit to the low-Q part of the S(Q)eQ data [35]. In doing this, the intensity of the pre-peak is seen to decrease, but its position remains almost constant for the liquid structure as
Among the compositions within the glass formation range, Al89Sm11 was further investigated in solid and molten states using reverse Monte Carlo simulations. Fig. 3 shows comparisons of the total structure factor measured from HEXRD experiments to results obtained when the data are fit using the indirect RMC method for as-quenched and liquid Al89Sm11 alloys to produce the same S(Q). A reasonably good fit is obtained in both the low (including the prepeak) and the high-Q regions. Fig. 4 shows the partial structure factors (SAleAl(Q), SAleSm(Q), and SSmeSm(Q)) determined for as-quenched and molten alloys. In the low-Q region (Q w 12.7 nm1) of the partial structure factors, a strong pre-peak (arrowed) is seen for SSmeSm(Q). Interestingly, the pre-peak is absent from the SAleAl(Q), and SAleSm(Q) indicating that the pre-peak seen in the total structure factor arises due to SmeSm containing correlations. The complex polyhedral structure in the local atomic configurations generated by RMC was further investigated by Voronoi Tessellation analyses, as shown in Fig. 5. Thirteen and 14-coordinated polyhedrons with indices <0,3,6,4,0> and <0,2,8,4,0> are dominant for Al centered structures. Voronoi polyhedrons (VP) with such indices correspond to a deformed bcc structure [36]. It is also found that some icosahedral-like clusters with indices <0,0,12,0,0>, <0,1,10,2,0>, or <0,2,8,2,0> and coordination of 12 and 13 atoms are formed around Al. The Sm centered polyhedrons almost always have higher coordination as compared to Al centered
Fig. 2. Total structure function S(Q) data for (a) as-quenched Al(100x)Smx (x ¼ 8, 10, 11 and 12) and as-quenched amorphous with liquid Al89Sm11 alloy at 1313 K.
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Fig. 3. Total structure function data measured (open circles) and calculated using RMC (solid line) for as-quenched and liquid Al89Sm11 at 1313 K. The line at the bottom shows the difference between measured and calculated S(Q).
polyhedrons. The indices corresponding to deformed bcc structures, such as <0,1,10,4,0>, <0,2,8,6,0>, <0,1,10,5,0> and <0,2,8,5,0>, are the most frequent polyhedral forming around Sm atoms. The distribution of VP types is variable for liquid states, and no specific type of indice stands out. Despite the highly variable distribution of VP indices, particularly for the liquid state, coordination numbers (CN) numbers are quite consistent. Fig. 6 shows the histogram of CN for Al and Sm center clusters in amorphous and liquid states according to Voronoi Tessellation analysis. Sm centered clusters are highly coordinated, on the order of 16 atoms in the first near-neighbor shell. The first near-neighbor shell calculated with Voronoi Tessellation analysis includes both Sm and Al atoms. The fraction of Al atoms surrounding Sm centered atoms is on the order of 0.94 and 0.89 for amorphous and liquid states respectively. Therefore, according to Voronoi Tessellation analysis the Sm atoms are highly coordinated with Al both in liquid and amorphous states. CN for Al center cluster is on the order of 14 and the fraction of Al in the first neighborhood is around 0.87 for both amorphous and liquid states. Table 2 summarizes the fraction of Al atoms in the first neighborhood for Al and Sm center clusters for different CN. 4. Discussion According to previous high resolution electron microscopy studies [19,20], as-quenched Al90Sm10 alloys produced by high
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pressure gas atomization and melt spinning with wheel speeds of 30 and 40 m/s contain clusters of fcc-Al in a size range of 2e5 nm embedded in an amorphous matrix. In another study [18], clusters of Al that are showing fcc like medium range order (MRO) and having sizes in the range of 0.8e1.5 nm were resolved using fluctuation electron microscopy (FEM) in melt-spun Al92Sm8 produced at a wheel speed of 55 m/s. These results clearly show that the size of fcc-Al-like clusters decreases with increasing cooling rate (i.e. wheel speeds of melt spinner) and MRO or nanocrystals of an fccAl-like structure exist in AleSm alloys that are quenched from the liquid state within the glass formation range, at least for the conditions used in these studies. These pre-existing fcc-Al nanoclusters and the others that nucleate [20] show a restricted growth upon annealing below their crystallization temperatures (Tx) and a high number of density of nanocrystals on the order of 1021e1023 m3 results in the amorphous matrix. The small size of clusters in the as-quenched state makes them difficult to detect in conventional transmission electron microscopy and X-ray diffraction analyses. While fcc-Al nanoclusters cannot be detected with HEXRD, as shown in Fig. 2, an intriguing observation is the existence of a prepeak (Q w 12.6 nm1) at a position well below the major diffuse scattering position for all four as-quenched compositions. As shown in Fig. 2, the pre-peak is a distinct feature in the total structure factor function S(Q) calculated for as-quenched Al(100x)Smx (x ¼ 8, 10, 11, and 12) and liquid Al89Sm11. Similar types of pre-peaks have been reported for other alloy systems [22e24]. The physical origin of such pre-peaks is not clear and their existence is usually attributed to the formation of MRO structures [37,38]. These data were collected at room temperature without using any sample holder, thus eliminating any possible effect due to the quartz capillaries that were used at higher temperatures. These observations indicate that the pre-peak is not an anomaly of the manner in which the HEXRD data were obtained. There is a slight increase in the intensity of the pre-peak with increasing Sm content, which is more pronounced for the second broad peak located approximately at 33.4 nm1. A similar pre-peak is also observed in liquid phase. Fig. 2b shows the total structure factor function after background corrections for liquid Al89Sm11 at 1313 K, approximately 80 K above the liquidus temperature. A pre-peak is clearly seen in both graphs and the position is in good agreement with the pre-peaks observed at room temperatures in the asquenched samples. The correlation lengths and structural unit sizes calculated from the position of the pre-peak for as-quenched and liquid structures are summarized in Table 1. Note that the correlation lengths calculated from pre-peak broadening differ for solid
Fig. 4. Partial structure factors (SAleAl(Q), SAleSm(Q), and SSmeSm(Q)) calculated for (a) as-quenched and (b) liquid Al89Sm11 at 1313 K.
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Fig. 5. Most populated VP histograms of (a, c) Al and (b, d) Sm centered structure models for as-quenched (a, b) and liquid (c, d) Al89Sm11 at 1313 K.
and liquid structures. For as-quenched solids the correlation length is on the order of 2.5 nm and is almost constant with respect to composition. However, correlation length is much smaller in the liquid, being on the order of 1.5 nm. The structural unit sizes calculated from the positions of the prepeaks for the as-quenched samples are almost constant with respect to compositions and are similar to that of liquid Al89Sm11. This shows that the clusters formed in the liquid state might be retained in the as-quenched state with an increase in correlation
length. The positions of the extra peaks observed in liquid and asquenched states are close to some of the major diffraction peaks in a high temperature metastable body centered tetragonal phase, Al11Sm3. This compound is stable between 1351 and 1733 K [39]. The interplanar spacing in t-Al11Sm3 for (002) and (211) planes are given as 0.495 and 0.188 nm, respectively. These planar spacings are close in value to those calculated from the positions of the prepeaks and side peaks for as-quenched Al(100x)Smx (x ¼ 8, 10, 11 and 12) and liquid Al89Sm11.Thus, the broad peaks observed in the total
Fig. 6. Coordination numbers calculated for (a) Al and (b) Sm centered VP models for as-quenched and liquid Al89Sm11 at 1313 K.
Y.E. Kalay et al. / Intermetallics 18 (2010) 1676e1682 Table 2 The fraction of Al atoms in the first shell neighborhood around Sm and Al center atoms. CN
Sm centered
Al centered
Amorphous 17 16 15 14 13 12
0.94 0.94 0.93 0.93 0.93 0.94
0.90 0.89 0.88 0.87 0.86 0.85
Liquid 17 16 15 14 13 12
0.91 0.90 0.90 0.89 0.88 0.87
0.88 0.88 0.88 0.87 0.87 0.86
structure factor function are consistent with a pseudo-tetragonal MRO Al11Sm3 structure. Moreover, according to Voronoi Tessellation analysis, the local atomic configuration around Sm atoms calculated using RMC for the as-quenched state have structural similarities with the first shell neighborhood of Sm atoms in high temperature metastable Al11Sm3 tetragonal phase. Fig. 7a shows the atomic configuration of this phase, which has the I4/mmm space group symmetry. This structure consists of Sm atoms surrounded by 8 Al atoms at a distance of 0.325 nm and another 8 Al atoms at a distance of 0.327 nm. Therefore, every Sm atom in this structure is highly coordinated with Al in their first and second shells and 16 Al atoms exist at a distance of 0.327 nm or less (Fig. 7b). Moreover, in a study where the local atomic structure of amorphous Al92Sm8 was investigated using X-ray absorption fine structure (XAFS) [40], the best fit to the experimental XAFS data were obtained by considering a “single shell” structure of 16 Al atoms surrounding Sm atoms within the first coordination shell. The average coordination number (CN) determined for Sm atoms is on the order of 16 according to Voronoi analyses for both liquid and amorphous states having 90e94% Al atoms within the first shell coordination (Table 2). Therefore, a high coordination of Al atoms around Sm atoms is detected in the first shell structure that is similar to tetragonal Al11Sm3 metastable compound (Fig. 7). According to HEXRD observations of liquid and as-quenched structures a pre-peak is resolved well below the major amorphous peak in total structure factors. Further RMC simulations coupled with Voronoi Tessellation analysis reveal highly coordinated Sm centered clusters in the as-quenched structure and their first shell neighborhood is structurally similar to the high temperature metastable Al11Sm3 tetragonal phase. This particular alloy system has been reported to devitrify to a very high number of nucleation
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density of nanocrystalline fcc-Al on the order of 1020e1023 m3 if the alloy is quenched from liquid [3,17]. According to high temperature HEXRD, the pre-peak and side peaks observed in the liquid and as-quenched states appear at atomic distances in close correspondence to the spacing of (002) and (211) planes in t-Al11Sm3 compound. This would argue for a persistence of pseudo-tetragonal Al11Sm3 atomic ordering in the molten state and subtle growth of this MRO during quenching from liquid. The first shell neighborhood around Sm atoms within this compound consists of highly coordinated Al atoms and the clustering in the melt is believed to form due to strong interactions between Al and Sm atoms. However, as shown by the partial structure factor (Fig. 4), the pre-peak is particularly originated by SmeSm pairs, which require a longer length correlation as compared to the first shell neighborhood. The correlation length of this MRO structure observed in as-quenched state is between 2 and 3 nm according to HEXRD results. This corresponds to 4e6 layers of repeating (002) planes, if the structure is assumed to be a pseudotetragonal MRO Al11Sm3. In contrast the pre-peak in the liquid has a correlation length of only 1.5 nm or 2e3 layers. These clusters in the liquid essentially act like distinct “superatoms”, with a broad size distribution and huge diameters as compared to individual Al and Sm atoms in the matrix. Upon rapid solidification it is hypothesized that these “superatoms” cause a local atomic size effect, similar to that seen in multi-component bulk amorphous alloys, preventing the complete crystallization of the Al rich matrix. This results in the as-quenched structure being divided into nanoscale regions by a network of MRO structure. Depending on the size of these regions a certain percentage of fcc-Al clusters will reach the critical nucleus size with the remainder staying below this size. While admittedly speculative, this explanation agrees with results obtained for Al90Sm10 using three-dimensional atom probe tomography (3D APT) [19]. The effect of cooling rate has a dominant role on the percentage of Al that reaches the critical nucleation size. However, it should be noted that cooling rate has possibly a minor effect on the presence of MRO in the as-quenched state since a prepeak due to Sm rich clustering is also observed for the melt. Therefore, ordering in the liquid needs be studied directly with respect to initial melt temperature, and a more accurate representation of the liquid structure is required. When as-quenched alloys are annealed to devitrify to a high number of density nanocrystals, the MRO appears to act as a catalyst site for nucleation of fcc-Al. Previous results within this system [20] showed that long-range diffusion between fcc-Al clusters is blocked by the Sm rich MRO structure in the matrix, resulting in restricted growth of fcc-Al. Al rich regions with sizes below the critical size for homogeneous nucleation can nucleate via a heterogeneous nucleation mechanism in the vicinity of the MRO clusters. A more detailed study that includes atom probe tomography and scanning transmission electron microscopy equipped with a high angle annular dark field detector is in progress. Such a study should be helpful for identifying the spatial distribution of the MRO and fcc-Al in the as-quenched state, to help elucidate the effect of MRO on producing the observed high nucleation density of fcc-Al in these alloys. 5. Conclusion
Fig. 7. (a) Crystal structure and (b) first and second shells neighborhood of Sm atoms in high temperature metastable tetragonal phase Al11Sm3 with I4/mmm space group symmetry. Red and black colors represent Al and Sm atoms, respectively.
Amorphous alloys of Al(100x)Smx (x ¼ 8, 10, 11, and 12) were produced using the melt spinning technique at a wheel speed of 30 m/s. Detailed analysis of the as-quenched structure by HEXRD revealed a pre-peak and a side peak in total structure factor for both amorphous and liquid states. Partial structure factor analysis showed that the pre-peak arises due to SmeSm correlations. The formation of the pre-peak is related to the formation of clusters in
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the liquid. After the liquid alloy is quenched, the position of the prepeak remains almost constant, indicating that the clusters in the liquid are retained and showed subtle growth during quenching from liquid. The Voronoi analysis of local atomic structures based on RMC simulation suggests a deformed bcc-like structure surrounding Sm with CN on the order of 16 atoms for both amorphous and liquid states. These experimental results and simulations indicate the first shell neighborhood of Sm atoms in asquenched states is structurally similar to the high temperature metastable Al11Sm3 tetragonal phase. The pre-peak, indicating a SmeSm correlation, was found to result due to a higher order correlation within the liquid. These Sm rich clusters in the liquid are thought of as creating Sm and Al rich regions. When as-quenched alloys are annealed, the Sm rich clusters act as catalyst sites promoting heterogeneous nucleation of fcc-Al, while inhibiting long-range diffusion processes between fcc-Al clusters. Acknowledgments Appreciation is expressed to Shaogang Hao for his valuable help in ab-initio calculations. The work at Ames Laboratory was supported by the United States Department of Energy (USDOE), Office of Science (OS), Office of Basic Energy Science (BES), under Ames Laboratory Contract No. DE-AC02-07CH11358, the high-energy X-ray work at the MUCAT sector of the APS was supported by the US Department of Energy, Office of Science, and Basic Energy Sciences under Contract No. DE-AC02-06CH11357. The assistance of Materials Preparation Center of the Ames Laboratory is acknowledged for supplying our samples [27]. References [1] Inoue A. Amorphous, nanoquasicrystalline and nanocrystalline alloys in Albased systems. Prog Mater Sci 1998;43:365e520. [2] Foley JC, Perepezko JH. The devitrification of AleYeFe amorphous alloys. J Non-Cryst Solids 1996;205e207:559e62. [3] Foley JC, Allen DR, Perepezko JH. Analysis of nanocrystal development in AleYeFe and AleSm glasses. Scr Mater 1996;35:655e60. [4] Hirata A, Hirotsu Y, Matsubara E, Ohkubo T, Hono K. Mechanism of nanocrystalline microstructure formation in amorphous FeeNbeB Alloys. Phys Rev B 2006;74:184204. [5] Schroers J, Busch R, Masuhr A, Johnson WL. Continuous refinenement of the microstructure during crystallization of supercooled Zr41Ti14Cu12Ni10Be23 melts. Appl Phys Lett 1999;74:2806e8. [6] Hono K, Ping DH, Ohnuma M, Onodera H. Cu clustering and Si partitioning in the early crystallization stage of an Fe73.5Si13.5B9Nb3Cu1 amorphous alloy. Acta Mater 1999;47:997e1006. [7] Allen DR, Foley JC, Perepezko JH. Nanocrystal development during primary crystallization of amorphous alloys. Acta Mater 1998;46:431e40. [8] Wang XL, Almer J, Liu CT, Wang YD, Zhao JK, Stoica AD, et al. In situ synchrotron study of phase transformation behaviors in bulk metallic glass by simultaneous diffraction and small angle scattering. Phys Rev Lett 2003;91:265501e5. [9] Schneider S, Thiyagarajan P, Johnson WL. Formation of nanocrystals based on decomposition in the amorphous Zr41.2Ti13.8Ni10Be22.5 alloy. Appl Phys Lett 1995;68:493e5. [10] Kelton KF. Time-dependent nucleation in partitioning transformations. Acta Mater 2000;48:1967e80. [11] Kelton KF. Kinetic model for nucleation in partitioning systems. J Non-Cryst Solids 2000;274:147e54.
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