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Kinetics of fcc-Al nanocrystallization in Al90Tb10 metallic glass
Journal of Non-Crystalline Solids 378 (2013) 71–78
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Journal of Non-Crystalline Solids journal homepage: www.elsevier.com/ locate/ jnoncrysol
Kinetics of fcc-Al nanocrystallization in Al90Tb10 metallic glass T. Demirtaş, Y.E. Kalay ⁎ Department of Metallurgical and Materials Engineering, METU, Ankara 06800, Turkey
a r t i c l e
i n f o
Article history: Received 5 May 2013 Received in revised form 11 June 2013 Available online xxxx Keywords: Nanocrystallization; Metallic glass; Crystallization kinetics
a b s t r a c t The crystallization kinetics of Al90Tb10 amorphous alloy were investigated by a combined study of differential scanning calorimetry (DSC), transmission electron microscopy (TEM), Cu-Kα X-ray diffraction (XRD) analyses and microhardness measurements. Amorphous to fcc-Al transformation kinetics were descried through Johnson–Mehl–Avrami (JMA) approach based on the isothermal DSC hold at 220 °C. XRD and TEM revealed the formation of highly populated (~ 1021 m−3) fcc-Al nanocrystals after the first crystallization event. JMA indicated a three dimensional diffusional growth with constant nucleation rate. TEM and XRD analyses agree well with the three dimensional diffusional growth of fcc-Al within the amorphous matrix. Nucleation of fcc-Al nanocrystals was attributed to already existing Al rich clusters embedded in the amorphous matrix. The deviation in cooling rate at the free and the wheel sides caused variations in Al rich cluster size and population through the thickness of the ribbon specimen. Microhardness tests captured this discrepancy which results in anomalous Avrami exponent, n(x). © 2013 Elsevier B.V. All rights reserved.
1. Introduction Nanocrystalline materials have been subjected to considerable amount of research due to their great potential of improving various material properties through grain size refinement such as strength, corrosion resistance or improved magnetic properties [1]. Materials with nanocrystalline microstructure can be derived from vapor, liquid and solid states by applying several different techniques e.g. sputtering, rapid solidification, severe plastic deformation, mechanical alloying and devitrification of amorphous phase [1–3]. When the nanocrystalline aluminum alloys are considered, techniques involving the isothermal devitrification of an amorphous precursor constitute a promising manufacturing pathway [4]. Successful preceding studies have shown that production of low-density highstrength Al alloys with UTS values exceeding 1 GPa is achievable by devitrification of Al based metallic glasses [4,5]. A major challenge in constructing such high-strength nanocrystalline embedded amorphous composite structures, is the control of microstructure development during the devitrification sequence. The ideal microstructure design requires the full control of nanocrystalline size, volume fraction and number density. Therefore, a complete understanding of crystallization mechanism and kinetics is crucial to engineer such nanocrystalline Al alloys with desired microstructural features. Most of the previous studies on devitrification of Al-based metallic glass have been performed in high Al content (between 80 and 92 at.% Al) Al–RE and Al–RE–TM (RE: rare-earth element; TM: transition metal) systems [4]. These alloy systems, also known as marginal glass ⁎ Corresponding author. Tel.: +90 312 210 2525; fax: +90 312 210 2518. E-mail address: [email protected] (Y.E. Kalay). 0022-3093/$ – see front matter © 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.jnoncrysol.2013.06.020
formers, show a limited glass-forming ability upon quenching from their molten states. On the other hand, their devitrification product after first crystallization event involves nanocrystals of fcc-Al with a very high number of nucleation density, on the order of 1020– 1023 m− 3 embedded in the amorphous matrix [4]. Several studies have been conducted to reveal the underlying crystallization mechanism yet there is no clear answer for this atypical nucleation behavior of fcc-Al nanocrystals. The previously proposed models have been gathered around two main approaches namely as; “quenched-in nucleation” [6,7] and “phase separation within the amorphous matrix” [8,9]. According to the first approach, considerable amount of crystalline nuclei form during rapid solidification. However, fast rising of the viscosity of undercooled liquid near the glass transition temperature (Tg) suppresses the crystal growth and crystallite nuclei trap within the amorphous matrix [6]. The existence of “quenched-in nuclei” was confirmed through a comparison of melt-spun and cold-rolled specimens of similar chemical compositions [10]. MRO structures with pseudo-fcc atomic arrangement was detected by fluctuation electron microscopy (FEM) in melt-spun ribbons but they were absent in mechanically deformed specimens [11]. This was explained by the lack of quenched-in nuclei formation in cold-rolled alloys as they were not produced by solidification. The second hypothesis approaches the nanocrystallization of fcc-Al by a mechanism similar to spinodal decomposition [8]. This hypothesis argues that a fully amorphous state phase separates into Al-rich and Al-depleted regions before crystallization. Al-rich regions extending to 74–126 nm [9] are hold responsible for the nucleation of fcc-Al nanocrystals upon low temperature annealing. A time-dependent homogeneous nucleation model [12], namely as “coupled-flux nucleation”, was also introduced
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Fig. 1. (a) BF-TEM image with SAED pattern (in-set) and (b) XRD patterns (from wheel and free sides) of as-quenched Al90Tb10.
to explain high number density of nanocrystallization. Both approaches have been questioned in different aspects such as, chemical integrity of the amorphous alloy, chemical dependence of devitrification pathway and two-stage crystallization observed in certain Al–RE and Al–RE–TM systems [13]. There is still no consensus regarding an exact nucleation mechanism yielding to high-number density of crystals. Some recent works on Al–Y–Fe [14] and Al–Tb [15] have shown that phase separation may actually be present at nanoscales. Such nanoscale segregations may cause chemical and topological orders develop in the amorphous states. Kalay et al. have experimentally demonstrated regions of 1 nm pure Al in as-quenched Al90Tb10 alloy by atom probe tomography [15]. It is interesting to note that some of the sub-nano scale orders were found to exist in molten state, as well [15,16]. Therefore, no matter how fast the liquid alloy cools, some topological and chemical orders persist to exist in as-quenched state [16]. In this study, nanocrystallization of fcc-Al from amorphous Al90Tb10 alloy was investigated in details using differential scanning calorimetry, interrupted X-ray diffraction and transmission electron microscopy experiments. Transformation kinetics was modeled by the well-known Johnson–Mehl–Avrami (JMA) approach [17] according to isothermal DSC and the results were compared with the experimental findings. The applicability of JMA approach in nanoscale phase-separated amorphous systems was discussed.
Specimens for differential scanning calorimetry (DSC) experiments were cut into equal size coupons at 10 mg of average weight. They were fed into Al pans and sealed with Al cover. Specimens were isochronally heated with 100 °C/min heating rate under a protective gas atmosphere of pure N2 and isothermally hold at predetermined temperature of 220 °C below the crystallization point. Same routine was repeated with empty pans for baseline correction. The conventional XRD analyses (Cu-Kα) were conducted at Bragg–Brentano geometry between 10 and 80 2θ° with a scan rate of 0.02°/min from free and wheel sides of ribbon specimens. TEM analyses were conducted by a JEOL JEM2100F scanning/transmission electron microscope. Samples for electron microscopy were electropolished with a solution of 25 vol.% HNO3 and 75 vol.% methanol. The ribbon specimens were subjected to microhardness tests for the investigation of hardness changes through the different levels of crystallization. Analyses were held with a Shimadzu HMV Micro Hardness Tester under a load of 980.7 mN (HV 0.1) for 10 s of dwell time. Ribbons were attached on a bulk aluminum plate having average hardness values of 154 HV. Seven significant indentations were taken into account. In order to eliminate the errors potentially coming from the Al substrate, hardness tests were repeated on a steel plate having an average hardness of 354 HV. 3. Experimental results
2. Experimental procedure
3.1. As-quenched state
Ingots of Al90Tb10 were prepared using electric arc melting under Ar atmosphere from highly pure Al (99.99 wt.%) and Tb (99.9 wt.%) elements. Amorphous ribbons with approximate thickness of 30 μm were produced using a Cu block single melt-spinner at a tangential speed of 30 m/s.
Fig. 1(a) shows the bright-field (BF-TEM) image and selected-area electron diffraction (SAED) pattern of the as-quenched Al90Tb10 alloy. TEM and electron diffraction analyses indicate a completely amorphous matrix. The corresponding X-ray diffraction (XRD) pattern from wheel and free sides of the ribbons, shows no crystalline reflections
Fig. 2. (a) Isochronal DSC curve at 40 °C/min heating rate. Isothermal DSC curves indicating (b) the fcc-Al crystallization and (c) the overall first and second exothermic events. The interception points for TEM and XRD are marked as A, B, C, D, E and F.
T. Demirtaş, Y.E. Kalay / Journal of Non-Crystalline Solids 378 (2013) 71–78 Table 1 Critical interruption points for TEM and XRD analyses. Label
Time (seconds)
Condition
A B C D E F
298 365 469 634 785 1270
On-set (initial crystallization) Peak (initial crystallization) Initial crystallization completed Peak (second crystallization)
(Fig. 1(b)), but the formations of a pre-peak and a side-peak before and after the main amorphous reflection, respectively. The positions of these two broad peaks correspond to scattering vectors of 13 nm−1 (2θ = 18.5°) and 34 nm−1 (2θ = 49.3°). 3.2. Crystallization kinetics Fig. 2(a) shows the isochronal DSC curve collected at 40 °C/min heating rate. Four exothermic reactions were detected between room temperature and 500 °C. The current and the previous studies [15] have indicated that the first devitrification event is the nanocrystallization of fcc-Al phase. The on-set point of initial crystallization was measured as 236 °C. The absence of a clear glass transition temperature is related to relaxation of the marginal glass forming system and it was previously discussed in [15]. Isothermal heating experiments were performed 16 °C below the Tx value. Fig. 2(b) and (c) show the DSC isothermal heating curves for the first 22 and the overall 50 minutes crystallization events, respectively. The devitrified TEM specimens were prepared by interrupted annealing experiments at 220 °C. The interruption points are marked as A, B, C, D, E and F in Fig. 2(a) and (b) and the corresponding annealing times are shown in Table 1. The nanocrystallization of fcc-Al is completed during the first exothermic reaction. The second isothermal reaction yields to an Al–Tb intermetallic structure. Fig. 3 shows the bright field (BF), high-resolution TEM (HRTEM) and high-angle annular dark-field (HAADF) images of
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the second phase forming within the amorphous matrix. EDS indicates 15 ± 1 at.% of Tb within this structure. The lighter color contrast in HAADF clearly shows the relative enrichment of Tb within this intermetallic compound. A face-centered orthorhombic structure with a = 17.7 Å, b = 9.36 Å and c = 5.42 Å is in good agreement with the measured diffraction patterns. However, further detailed analysis should be conducted before a solid conclusion could be drawn. XRD patterns of the as-quenched (Fig. 1(b)) and isothermally annealed (Fig. 4) specimens indicate a pre-peak formation well before the main amorphous reflection peak at approximately 18.5° 2θ (or Q = 13 nm−1). The first Bragg reflection of the second forming phase at point F is actually evolving from the pre-peak during isothermal heating. This is in good agreement with the hypothesis on the persistence of the pre-peak due to strong RE–Al interactions [16] and it will be discussed later. The analysis of the second phase formation is left out of scope of the current study. The first crystallization reaction resulting in high number density of nanocrystal was investigated in details by thermal, X-ray diffraction and electron microscopy techniques. The interrupted XRD (Fig. 4) and TEM analyses (Figs. 5 and 6) from selected critical points on the first crystallization evidently show the formation fcc-Al phase within the amorphous matrix. The fcc-Al reflections on XRD become stronger and sharper with respect to annealing time, indicating the growth of nanocrystals embedded in the amorphous matrix. TEM BF analyses reveal the dendritic nature of growing fcc-Al nanocrystals. Fig. 6 shows the high-resolution TEM images (HRTEM) representing the size and interface morphologies between the nanocrystals and amorphous matrix. The average sizes of these dendritic nanocrystals were estimated by measuring the distance between two far edges of at least 10 nanocrystals for 20 different TEM micrographs for each interrupted annealing point. The average size of the nanocrystals increases with annealing time at 220 °C as shown in Fig. 7(a). Another important parameter to be considered is the change in number density of nanocrystals with respect to crystal growth. In order to get significant measurements, 20 sets of data including five TEM BF images of nanocrystals and one from the hole were taken for each stage of crystallization. In total, for each crystallization step (A to E), 120 BF images were investigated with utmost care.
Fig. 3. (a) BF-TEM, (b) HRTEM and (c) HAADF images of fcc-Al nanocrystals and Al–Tb intermetallic phase formed after the second crystallization event.
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key parameter is the thickness of the specimen. The thickness of specimen can be calculated by the electron mean free path formula [20] as ðIBF Þ t ¼ exp − ðI0 Þ Δ
ð2Þ
where IBF is the bright field average intensity, I0 is the bright field average intensity of fully transmitted beam (the hole), t is the thickness, and Δ is the elastic scattering mean free path. Δ value is taken equal to pure Al [21]. The result of change in nanocrystal population is given in Fig. 7(b). It is interesting to note that the population of nanocrystals is almost constant with respect to annealing time during isothermal heating. The average number density measured at interrupted points is between 1.5 × 1021 and 2.5 × 1021 m−3 with a very little deviation. In the course of analytically simulating the isothermal crystallization event at 220 °C, the Johnson–Mehl–Avrami (JMA) approach is used. The JMA equation [17] is given as n x ¼ 1− exp −½kðt−τÞ Fig. 4. XRD patterns collected at different stages of the first and second crystallization events.
The volume number density is given by this following formula [18,19] Nv ¼
NA tþd
ð1Þ
where Nv is volume particle density, t is the thickness, d is average particle size and NA is the particle density per image. In this equation a
ð3Þ
where x is the volume fraction of the crystallized phase, t is the annealing time, n is the Avrami exponent, τ is incubation time and k is the rate constant. Among these parameters n is particularly important in describing the nucleation and growth mechanisms. Fig. 8 shows the crystallized volume fraction with respect to annealing times calculated by integrating the isothermal DSC scan shown in Fig. 2(b) excluding the incubation time. The incubation time at this temperature was measured as 1.2 min. The Avrami exponent (n), was calculated to be 2.5 through the slope of ln(− ln(1 − x)) versus ln (t − τ) over the range 5–85% transformation at 220 °C as shown in Fig. 9(a). Another approach of conducting JMA analysis is to take
Fig. 5. BF-TEM images showing the (a) as-quenched amorphous state and fcc-Al nanocrystals at different interrupted points of (b) A, (c) B, (d) C, (e) D and (f) E. Insets indicate the corresponding SAED patterns.
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Fig. 6. HRTEM images of critical interrupted points (a) A, (b) B, (c) C, (d) D, (e) E and (f) F. Insets show the corresponding FFT patterns.
the derivative of n with respect to crystallized volume fraction [22] given as
nðxÞ ¼
∂ lnð− lnð1−xÞÞ : ∂ lnðt−τÞ
ð4Þ
This technique, also called as local Avrami exponent, represents a delicate approach of calculating a transformation fraction sensitive Avrami exponent [23]. Fig. 9(b) shows the change in local Avrami exponent with respect to crystallized fraction. n(x) starts with an approximate value of 1.6 and during the main crystallization regime
(5–85%), it stays around 2.5 and for the final stage it shows a sharp increase towards to a value of 4 (N 85%). 4. Discussion The as-quenched Al90Tb10 was found to be fully amorphous within the limitation of HRTEM. XRD diffraction patterns (Fig. 1(b)), on the other hand, clearly show the formation of a pre-peak and a side peak. The earlier studies on as-quenched amorphous Al90Tb10 alloys at similar quenching rates have shown parallel trend in pre-peak formation [15]. The origin of the pre-peak in metallic glass systems is contentious [24]. Our previous findings on Al–RE systems have
Fig. 7. Average (a) nanocrystal size (nm) and (b) number density (m−3) calculated at different interruption points.
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Fig. 8. The crystallized fraction versus annealing time curve for fcc-Al crystallization.
indicated that pre-peak may be resulted due to the medium range order scale topological and chemical separations within the amorphous phase [15]. According to this hypothesis, the amorphous matrix is divided into Al-rich and Al-depleted regions at nano-scales. The approximate average size of the Al-rich regions was measured as 1 nm by atom probe tomography (APT) analyses. Fluctuation electron microscopy (FEM) clearly indicated fcc-like MRO within the Al-rich regions [15]. Similar pre-peak formation was also observed in molten state of Al–RE at elevated temperatures [16]. The ab-initio and synchrotron constrained reverse Monte Carlo simulation have indicated a strong interaction between the Al and RE atoms. In as-quenched amorphous and molten alloys, RE atoms is always found to be highly correlated with Al atoms which give rise of the pre-peak in total structure factor function [16]. The selective accumulation of Al atoms around the RE naturally causes the formation of RE free regions from where the nucleation of fcc-Al takes place during isochronal or low-temperature isothermal annealing. The first Bragg reflection of the secondly forming Al–Tb intermetallic phase (Fig. 4) at point F is actually evolving from the pre-peak during isothermal heating. This is in good agreement with the hypothesis on the existence of pre-peak due to strong RE–Al interactions. The strong interaction of RE–Al atoms, naturally causes the formation of pure Al nanodomains as shown by previous APT results [15]. Accordingly, the nucleation of fcc-Al nanocrystals upon annealing is expected to occur within these Al-enriched regions. Detailed analysis fcc-Al nanocrystallization kinetics was carried out by isothermal DSC measurements and interrupted TEM experiments. The Avrami exponent, n, calculated by linear regression was found to be 2.5. This value indicates a three-dimensional diffusional growth with constant nucleation rate. In three dimensional diffusional growths, particle
grows linearly with t1/2, where t is the isothermal annealing time [25]. This statement is in good agreement with our TEM observations. Fig. 10 shows the change in average particle size with respect to t1/2. In this calculation, average cluster size at 0 min is assumed to be 1 nm according to the previous work of Kalay et al. with the same Al–Tb alloy composition [15]. The change in R with respect to t1/2 was almost perfectly fitted by a linear function (Fig. 10) indicating a diffusional growth. It should be noted that although the crystallized volume fraction is up to 100%, there is considerable amount of remaining amorphous phase. Therefore, the first crystallization event only comprises the fcc-Al formations. The EDS results after the first exothermic reaction, clearly reveal the enrichment of amorphous matrix with Tb atoms. Further growth of fcc-Al is restricted by the disability of the long range diffusion of Tb within the remaining amorphous matrix by soft impingement. The average size of fcc-Al nanocrystals saturates to a value of 40 nm before the second crystallization event and it results in high population of fcc-Al nanocrystals embedded in the amorphous matrix. Whether fcc-Al crystals grow from the pre-existing nuclei or nucleation occurs within the amorphous regions is still not clear. JMA predicts a constant nucleation rate. FEM studies indicated fcc-like MRO within the Al-rich nano-domains. These MRO structures may act as the nuclei of the nanocrystalline fcc-Al. Under constant nucleation rate conditions, the change in number density of nuclei should be linear with respect to annealing time. It is difficult to observe this change within the limitations of electron microscopy. Fig. 7(b) indicates that the change in nucleation density of nanocrystals is steady around 2 × 1021 m−3. On the other hand, it is quite possible to have a range of pure Al cluster size. The ones with critical size may act as the pre-nuclei. However, the others with “undercritical” size stay as amorphous and upon annealing they grow and eventually crystallizes to fcc-Al. Previous APT studies revealed the interconnection of Al-rich clusters without crystallization upon annealing before the first exothermic event. Moreover, relaxation of these glassy alloys induces a clear glass-transition temperature, indicating the amorphous nature of the as-quenched alloys. The local Avrami coefficient, n(x) starts from a value of 1.6 indicating a zero nucleation rate thus growth of pre-existing nuclei. This may indicate the growth of MRO present in the as-quenched state. For the main part of crystallization regime, n(x) stays around 2.5, which is in good agreement with the Avrami exponent calculated by linear regression. The sharp increase of n(x) after 85% of crystallization is usually attributed to a combination of structural and compositional effects [25,26]. The inhomogeneous distribution of nuclei in amorphous ribbon specimens is a common reason for the sharp increase in n(x) at higher crystallization fractions [27]. Because of this effect, the local Avrami exponent should be considered carefully otherwise
Fig. 9. (a) The JMA plot of (a) ln(−ln(1 − x)) versus ln (t − τ) for fcc-Al crystallization. The line represents the best fit by linear regression with a correlation coefficient (R2) of 0.99992. (b) The derivative of JMA plot with respect to crystallized fraction.
T. Demirtaş, Y.E. Kalay / Journal of Non-Crystalline Solids 378 (2013) 71–78
Fig. 10. The average nanocrystal size versus t1/2 graph for fcc-Al crystallization. The line represents the best fit by linear regression with a correlation coefficient (R2) of 0.9945.
it may lead to erroneous conclusions on nucleation and growth mechanisms. If possible, JMA type modeling studies on crystallization amorphous materials should be conducted with corresponding electron and X-ray based techniques. In our case dominant factor of observing anomalous Avrami exponents should be the inhomogeneous distribution of Al-rich clusters due to ribbon geometry of the specimens. It is a well known fact that the wheel side solidifies at a higher cooling rate as compared to the free side of the as-quenched ribbon [28]. The average thickness of the ribbon is around 30 μm. It is not an easy task to quantify and compare the Al-rich clusters on the wheel and free sides of the ribbon. On the other hand, microhardness measurements are quite sensitive to the percent nanocrystallization of the amorphous matrix [29]. In order to compare the percent nanocrystallization with respect to annealing time, specimens were subjected to microhardness tests from their wheel and free sides on Al and steel block substrates. The test results are given in Fig. 11. The average hardness values on Al and steel substrates are found to be close to each other. This indicates that ribbons have sufficient thickness to conduct the microhardness tests and the effect of substrate in hardness measurement is minor. In as-quenched amorphous state, ribbons have an average value of 185 HV on both wheel and free sides. However, after the first interrupted quench at point A, the hardness at wheel side is considerably higher as compared to free side up to point E, where hardness of both sides becomes equal. After the fcc-Al crystallization is completed, wheel and free sides reach similar hardness values of 280 HV. The increase in hardness of
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the partially crystalline ribbon specimens is due to the formation of defect-free nanocrystalline fcc-Al and the increase of Tb content within the amorphous matrix with the rejection of RE atoms during the growth process [4]. Hardness measurements clearly show that these effects have occurred faster on the wheel side. Al-rich clusters (potential sites for fcc-Al crystals) are expected to be smaller and highly populated on the wheel side of ribbon due to the higher cooling rate. These small clusters having relatively large surface/volume ratio accelerate the nucleation and growth process and cause excess amount of RE rejection per unit time from the growing fcc-Al nanocrystals on the wheel side of the ribbon. This results in a sharp increase and a fast exhaustion in wheel side hardness. It should be noted that the amount of RE rejection per unit time cannot alter the main growth mechanism, which is a three-dimensional diffusional growth (n = 2.5) in our case. However, it will definitely change the conditions for soft impingement towards to end of the fcc-Al crystallization. We believe that this discrepancy is the dominant factor for the formation of anomalous local Avrami exponent observed at high crystallized fractions. Although mechanical properties are left out of the scope of this current study, it is also interesting to note that the hardness values measured for partially crystalline Al90Tb10 are quite high as compared to most of the commercial Al alloys used in aerospace and defense industries. As an example, the critical aerospace material, AA2024, has hardness values of 150 HV and 195 HV after aging and equal channel angular pressing (ECAP) process, respectively [30]. In our case, nanocrystalline ribbons show much higher hardness values indicating the potential strength of these materials. The presence of highly populated nanocrystals embedded in the amorphous matrix was previously reported to reduce the well-known plasticity problem of the metallic glass systems. Therefore, provided that enough thickness is achieved, Al–RE nanocrystalline-amorphous metallic glass composites are among the strong candidates for the next generation light-weight high strength alloys. 5. Conclusion The initial isothermal devitrification products of amorphous Al90Tb10 ribbon specimens were investigated in details using combined techniques of DSC, XRD and TEM. Detailed analysis of crystallization kinetics was performed through JMA approach using isothermal DSC scan at 220 °C. The first devitrifying crystalline phase was found to be fcc-Al nucleated at a very high population within the amorphous matrix. The change in number density of fcc-Al during the first crystallization was found to be steady between
Fig. 11. Hardness values measured for as-quenched amorphous and the critical crystallization states on (a) Al and (b) steel substrates.
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1.5 × 1021 and 2.5 × 1021 m− 3 according to TEM and HRTEM analyses. JMA indicated a three-dimensional diffusional growth with constant nucleation rate. Microstructural observations agreed well with the three-dimensional diffusional growth. The average size of nanocrystals saturate at 40–50 nm resulting a RE-rich amorphous matrix. The local Avrami exponent agrees well with the Avrami expression for the major part of crystallization process. However, it showed a sharp non-linearity for the high crystallized fractions. This deviation was found to be related to inhomogeneous nature of nanocrystallization through the thickness based on the microhardness measurements from the free and wheel sides of ribbon specimens. TEM and JMA analyses are in good agreement with the possibility of having MRO pre-nuclei of fcc-Al in addition to regions of pure Al clusters with “undercritical” size. The highly populated fcc-Al nanocrystals were found to nucleate within these regions. The mutual volume percentage of these regions is affected by the effective cooling rate of the molten metal. Acknowledgment The work at the Middle East Technical University was supported by BAP-08-11-2011-106. The assistance of the Materials Preparation Center of the Ames Laboratory is acknowledged for supplying the samples. References [1] H. Gleiter, Prog. Mater. Sci. 33 (4) (1989) 223–315. [2] K. Lu, Mater. Sci. Eng. R. 16 (4) (1996) 161–221. [3] M.A. Meyers, A. Mishra, D.J. Benson, Prog. Mater. Sci. 51 (2006) 427–556.
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Journal of Non-Crystalline Solids journal homepage: www.elsevier.com/ locate/ jnoncrysol
Kinetics of fcc-Al nanocrystallization in Al90Tb10 metallic glass T. Demirtaş, Y.E. Kalay ⁎ Department of Metallurgical and Materials Engineering, METU, Ankara 06800, Turkey
a r t i c l e
i n f o
Article history: Received 5 May 2013 Received in revised form 11 June 2013 Available online xxxx Keywords: Nanocrystallization; Metallic glass; Crystallization kinetics
a b s t r a c t The crystallization kinetics of Al90Tb10 amorphous alloy were investigated by a combined study of differential scanning calorimetry (DSC), transmission electron microscopy (TEM), Cu-Kα X-ray diffraction (XRD) analyses and microhardness measurements. Amorphous to fcc-Al transformation kinetics were descried through Johnson–Mehl–Avrami (JMA) approach based on the isothermal DSC hold at 220 °C. XRD and TEM revealed the formation of highly populated (~ 1021 m−3) fcc-Al nanocrystals after the first crystallization event. JMA indicated a three dimensional diffusional growth with constant nucleation rate. TEM and XRD analyses agree well with the three dimensional diffusional growth of fcc-Al within the amorphous matrix. Nucleation of fcc-Al nanocrystals was attributed to already existing Al rich clusters embedded in the amorphous matrix. The deviation in cooling rate at the free and the wheel sides caused variations in Al rich cluster size and population through the thickness of the ribbon specimen. Microhardness tests captured this discrepancy which results in anomalous Avrami exponent, n(x). © 2013 Elsevier B.V. All rights reserved.
1. Introduction Nanocrystalline materials have been subjected to considerable amount of research due to their great potential of improving various material properties through grain size refinement such as strength, corrosion resistance or improved magnetic properties [1]. Materials with nanocrystalline microstructure can be derived from vapor, liquid and solid states by applying several different techniques e.g. sputtering, rapid solidification, severe plastic deformation, mechanical alloying and devitrification of amorphous phase [1–3]. When the nanocrystalline aluminum alloys are considered, techniques involving the isothermal devitrification of an amorphous precursor constitute a promising manufacturing pathway [4]. Successful preceding studies have shown that production of low-density highstrength Al alloys with UTS values exceeding 1 GPa is achievable by devitrification of Al based metallic glasses [4,5]. A major challenge in constructing such high-strength nanocrystalline embedded amorphous composite structures, is the control of microstructure development during the devitrification sequence. The ideal microstructure design requires the full control of nanocrystalline size, volume fraction and number density. Therefore, a complete understanding of crystallization mechanism and kinetics is crucial to engineer such nanocrystalline Al alloys with desired microstructural features. Most of the previous studies on devitrification of Al-based metallic glass have been performed in high Al content (between 80 and 92 at.% Al) Al–RE and Al–RE–TM (RE: rare-earth element; TM: transition metal) systems [4]. These alloy systems, also known as marginal glass ⁎ Corresponding author. Tel.: +90 312 210 2525; fax: +90 312 210 2518. E-mail address: [email protected] (Y.E. Kalay). 0022-3093/$ – see front matter © 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.jnoncrysol.2013.06.020
formers, show a limited glass-forming ability upon quenching from their molten states. On the other hand, their devitrification product after first crystallization event involves nanocrystals of fcc-Al with a very high number of nucleation density, on the order of 1020– 1023 m− 3 embedded in the amorphous matrix [4]. Several studies have been conducted to reveal the underlying crystallization mechanism yet there is no clear answer for this atypical nucleation behavior of fcc-Al nanocrystals. The previously proposed models have been gathered around two main approaches namely as; “quenched-in nucleation” [6,7] and “phase separation within the amorphous matrix” [8,9]. According to the first approach, considerable amount of crystalline nuclei form during rapid solidification. However, fast rising of the viscosity of undercooled liquid near the glass transition temperature (Tg) suppresses the crystal growth and crystallite nuclei trap within the amorphous matrix [6]. The existence of “quenched-in nuclei” was confirmed through a comparison of melt-spun and cold-rolled specimens of similar chemical compositions [10]. MRO structures with pseudo-fcc atomic arrangement was detected by fluctuation electron microscopy (FEM) in melt-spun ribbons but they were absent in mechanically deformed specimens [11]. This was explained by the lack of quenched-in nuclei formation in cold-rolled alloys as they were not produced by solidification. The second hypothesis approaches the nanocrystallization of fcc-Al by a mechanism similar to spinodal decomposition [8]. This hypothesis argues that a fully amorphous state phase separates into Al-rich and Al-depleted regions before crystallization. Al-rich regions extending to 74–126 nm [9] are hold responsible for the nucleation of fcc-Al nanocrystals upon low temperature annealing. A time-dependent homogeneous nucleation model [12], namely as “coupled-flux nucleation”, was also introduced
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Fig. 1. (a) BF-TEM image with SAED pattern (in-set) and (b) XRD patterns (from wheel and free sides) of as-quenched Al90Tb10.
to explain high number density of nanocrystallization. Both approaches have been questioned in different aspects such as, chemical integrity of the amorphous alloy, chemical dependence of devitrification pathway and two-stage crystallization observed in certain Al–RE and Al–RE–TM systems [13]. There is still no consensus regarding an exact nucleation mechanism yielding to high-number density of crystals. Some recent works on Al–Y–Fe [14] and Al–Tb [15] have shown that phase separation may actually be present at nanoscales. Such nanoscale segregations may cause chemical and topological orders develop in the amorphous states. Kalay et al. have experimentally demonstrated regions of 1 nm pure Al in as-quenched Al90Tb10 alloy by atom probe tomography [15]. It is interesting to note that some of the sub-nano scale orders were found to exist in molten state, as well [15,16]. Therefore, no matter how fast the liquid alloy cools, some topological and chemical orders persist to exist in as-quenched state [16]. In this study, nanocrystallization of fcc-Al from amorphous Al90Tb10 alloy was investigated in details using differential scanning calorimetry, interrupted X-ray diffraction and transmission electron microscopy experiments. Transformation kinetics was modeled by the well-known Johnson–Mehl–Avrami (JMA) approach [17] according to isothermal DSC and the results were compared with the experimental findings. The applicability of JMA approach in nanoscale phase-separated amorphous systems was discussed.
Specimens for differential scanning calorimetry (DSC) experiments were cut into equal size coupons at 10 mg of average weight. They were fed into Al pans and sealed with Al cover. Specimens were isochronally heated with 100 °C/min heating rate under a protective gas atmosphere of pure N2 and isothermally hold at predetermined temperature of 220 °C below the crystallization point. Same routine was repeated with empty pans for baseline correction. The conventional XRD analyses (Cu-Kα) were conducted at Bragg–Brentano geometry between 10 and 80 2θ° with a scan rate of 0.02°/min from free and wheel sides of ribbon specimens. TEM analyses were conducted by a JEOL JEM2100F scanning/transmission electron microscope. Samples for electron microscopy were electropolished with a solution of 25 vol.% HNO3 and 75 vol.% methanol. The ribbon specimens were subjected to microhardness tests for the investigation of hardness changes through the different levels of crystallization. Analyses were held with a Shimadzu HMV Micro Hardness Tester under a load of 980.7 mN (HV 0.1) for 10 s of dwell time. Ribbons were attached on a bulk aluminum plate having average hardness values of 154 HV. Seven significant indentations were taken into account. In order to eliminate the errors potentially coming from the Al substrate, hardness tests were repeated on a steel plate having an average hardness of 354 HV. 3. Experimental results
2. Experimental procedure
3.1. As-quenched state
Ingots of Al90Tb10 were prepared using electric arc melting under Ar atmosphere from highly pure Al (99.99 wt.%) and Tb (99.9 wt.%) elements. Amorphous ribbons with approximate thickness of 30 μm were produced using a Cu block single melt-spinner at a tangential speed of 30 m/s.
Fig. 1(a) shows the bright-field (BF-TEM) image and selected-area electron diffraction (SAED) pattern of the as-quenched Al90Tb10 alloy. TEM and electron diffraction analyses indicate a completely amorphous matrix. The corresponding X-ray diffraction (XRD) pattern from wheel and free sides of the ribbons, shows no crystalline reflections
Fig. 2. (a) Isochronal DSC curve at 40 °C/min heating rate. Isothermal DSC curves indicating (b) the fcc-Al crystallization and (c) the overall first and second exothermic events. The interception points for TEM and XRD are marked as A, B, C, D, E and F.
T. Demirtaş, Y.E. Kalay / Journal of Non-Crystalline Solids 378 (2013) 71–78 Table 1 Critical interruption points for TEM and XRD analyses. Label
Time (seconds)
Condition
A B C D E F
298 365 469 634 785 1270
On-set (initial crystallization) Peak (initial crystallization) Initial crystallization completed Peak (second crystallization)
(Fig. 1(b)), but the formations of a pre-peak and a side-peak before and after the main amorphous reflection, respectively. The positions of these two broad peaks correspond to scattering vectors of 13 nm−1 (2θ = 18.5°) and 34 nm−1 (2θ = 49.3°). 3.2. Crystallization kinetics Fig. 2(a) shows the isochronal DSC curve collected at 40 °C/min heating rate. Four exothermic reactions were detected between room temperature and 500 °C. The current and the previous studies [15] have indicated that the first devitrification event is the nanocrystallization of fcc-Al phase. The on-set point of initial crystallization was measured as 236 °C. The absence of a clear glass transition temperature is related to relaxation of the marginal glass forming system and it was previously discussed in [15]. Isothermal heating experiments were performed 16 °C below the Tx value. Fig. 2(b) and (c) show the DSC isothermal heating curves for the first 22 and the overall 50 minutes crystallization events, respectively. The devitrified TEM specimens were prepared by interrupted annealing experiments at 220 °C. The interruption points are marked as A, B, C, D, E and F in Fig. 2(a) and (b) and the corresponding annealing times are shown in Table 1. The nanocrystallization of fcc-Al is completed during the first exothermic reaction. The second isothermal reaction yields to an Al–Tb intermetallic structure. Fig. 3 shows the bright field (BF), high-resolution TEM (HRTEM) and high-angle annular dark-field (HAADF) images of
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the second phase forming within the amorphous matrix. EDS indicates 15 ± 1 at.% of Tb within this structure. The lighter color contrast in HAADF clearly shows the relative enrichment of Tb within this intermetallic compound. A face-centered orthorhombic structure with a = 17.7 Å, b = 9.36 Å and c = 5.42 Å is in good agreement with the measured diffraction patterns. However, further detailed analysis should be conducted before a solid conclusion could be drawn. XRD patterns of the as-quenched (Fig. 1(b)) and isothermally annealed (Fig. 4) specimens indicate a pre-peak formation well before the main amorphous reflection peak at approximately 18.5° 2θ (or Q = 13 nm−1). The first Bragg reflection of the second forming phase at point F is actually evolving from the pre-peak during isothermal heating. This is in good agreement with the hypothesis on the persistence of the pre-peak due to strong RE–Al interactions [16] and it will be discussed later. The analysis of the second phase formation is left out of scope of the current study. The first crystallization reaction resulting in high number density of nanocrystal was investigated in details by thermal, X-ray diffraction and electron microscopy techniques. The interrupted XRD (Fig. 4) and TEM analyses (Figs. 5 and 6) from selected critical points on the first crystallization evidently show the formation fcc-Al phase within the amorphous matrix. The fcc-Al reflections on XRD become stronger and sharper with respect to annealing time, indicating the growth of nanocrystals embedded in the amorphous matrix. TEM BF analyses reveal the dendritic nature of growing fcc-Al nanocrystals. Fig. 6 shows the high-resolution TEM images (HRTEM) representing the size and interface morphologies between the nanocrystals and amorphous matrix. The average sizes of these dendritic nanocrystals were estimated by measuring the distance between two far edges of at least 10 nanocrystals for 20 different TEM micrographs for each interrupted annealing point. The average size of the nanocrystals increases with annealing time at 220 °C as shown in Fig. 7(a). Another important parameter to be considered is the change in number density of nanocrystals with respect to crystal growth. In order to get significant measurements, 20 sets of data including five TEM BF images of nanocrystals and one from the hole were taken for each stage of crystallization. In total, for each crystallization step (A to E), 120 BF images were investigated with utmost care.
Fig. 3. (a) BF-TEM, (b) HRTEM and (c) HAADF images of fcc-Al nanocrystals and Al–Tb intermetallic phase formed after the second crystallization event.
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key parameter is the thickness of the specimen. The thickness of specimen can be calculated by the electron mean free path formula [20] as ðIBF Þ t ¼ exp − ðI0 Þ Δ
ð2Þ
where IBF is the bright field average intensity, I0 is the bright field average intensity of fully transmitted beam (the hole), t is the thickness, and Δ is the elastic scattering mean free path. Δ value is taken equal to pure Al [21]. The result of change in nanocrystal population is given in Fig. 7(b). It is interesting to note that the population of nanocrystals is almost constant with respect to annealing time during isothermal heating. The average number density measured at interrupted points is between 1.5 × 1021 and 2.5 × 1021 m−3 with a very little deviation. In the course of analytically simulating the isothermal crystallization event at 220 °C, the Johnson–Mehl–Avrami (JMA) approach is used. The JMA equation [17] is given as n x ¼ 1− exp −½kðt−τÞ Fig. 4. XRD patterns collected at different stages of the first and second crystallization events.
The volume number density is given by this following formula [18,19] Nv ¼
NA tþd
ð1Þ
where Nv is volume particle density, t is the thickness, d is average particle size and NA is the particle density per image. In this equation a
ð3Þ
where x is the volume fraction of the crystallized phase, t is the annealing time, n is the Avrami exponent, τ is incubation time and k is the rate constant. Among these parameters n is particularly important in describing the nucleation and growth mechanisms. Fig. 8 shows the crystallized volume fraction with respect to annealing times calculated by integrating the isothermal DSC scan shown in Fig. 2(b) excluding the incubation time. The incubation time at this temperature was measured as 1.2 min. The Avrami exponent (n), was calculated to be 2.5 through the slope of ln(− ln(1 − x)) versus ln (t − τ) over the range 5–85% transformation at 220 °C as shown in Fig. 9(a). Another approach of conducting JMA analysis is to take
Fig. 5. BF-TEM images showing the (a) as-quenched amorphous state and fcc-Al nanocrystals at different interrupted points of (b) A, (c) B, (d) C, (e) D and (f) E. Insets indicate the corresponding SAED patterns.
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Fig. 6. HRTEM images of critical interrupted points (a) A, (b) B, (c) C, (d) D, (e) E and (f) F. Insets show the corresponding FFT patterns.
the derivative of n with respect to crystallized volume fraction [22] given as
nðxÞ ¼
∂ lnð− lnð1−xÞÞ : ∂ lnðt−τÞ
ð4Þ
This technique, also called as local Avrami exponent, represents a delicate approach of calculating a transformation fraction sensitive Avrami exponent [23]. Fig. 9(b) shows the change in local Avrami exponent with respect to crystallized fraction. n(x) starts with an approximate value of 1.6 and during the main crystallization regime
(5–85%), it stays around 2.5 and for the final stage it shows a sharp increase towards to a value of 4 (N 85%). 4. Discussion The as-quenched Al90Tb10 was found to be fully amorphous within the limitation of HRTEM. XRD diffraction patterns (Fig. 1(b)), on the other hand, clearly show the formation of a pre-peak and a side peak. The earlier studies on as-quenched amorphous Al90Tb10 alloys at similar quenching rates have shown parallel trend in pre-peak formation [15]. The origin of the pre-peak in metallic glass systems is contentious [24]. Our previous findings on Al–RE systems have
Fig. 7. Average (a) nanocrystal size (nm) and (b) number density (m−3) calculated at different interruption points.
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Fig. 8. The crystallized fraction versus annealing time curve for fcc-Al crystallization.
indicated that pre-peak may be resulted due to the medium range order scale topological and chemical separations within the amorphous phase [15]. According to this hypothesis, the amorphous matrix is divided into Al-rich and Al-depleted regions at nano-scales. The approximate average size of the Al-rich regions was measured as 1 nm by atom probe tomography (APT) analyses. Fluctuation electron microscopy (FEM) clearly indicated fcc-like MRO within the Al-rich regions [15]. Similar pre-peak formation was also observed in molten state of Al–RE at elevated temperatures [16]. The ab-initio and synchrotron constrained reverse Monte Carlo simulation have indicated a strong interaction between the Al and RE atoms. In as-quenched amorphous and molten alloys, RE atoms is always found to be highly correlated with Al atoms which give rise of the pre-peak in total structure factor function [16]. The selective accumulation of Al atoms around the RE naturally causes the formation of RE free regions from where the nucleation of fcc-Al takes place during isochronal or low-temperature isothermal annealing. The first Bragg reflection of the secondly forming Al–Tb intermetallic phase (Fig. 4) at point F is actually evolving from the pre-peak during isothermal heating. This is in good agreement with the hypothesis on the existence of pre-peak due to strong RE–Al interactions. The strong interaction of RE–Al atoms, naturally causes the formation of pure Al nanodomains as shown by previous APT results [15]. Accordingly, the nucleation of fcc-Al nanocrystals upon annealing is expected to occur within these Al-enriched regions. Detailed analysis fcc-Al nanocrystallization kinetics was carried out by isothermal DSC measurements and interrupted TEM experiments. The Avrami exponent, n, calculated by linear regression was found to be 2.5. This value indicates a three-dimensional diffusional growth with constant nucleation rate. In three dimensional diffusional growths, particle
grows linearly with t1/2, where t is the isothermal annealing time [25]. This statement is in good agreement with our TEM observations. Fig. 10 shows the change in average particle size with respect to t1/2. In this calculation, average cluster size at 0 min is assumed to be 1 nm according to the previous work of Kalay et al. with the same Al–Tb alloy composition [15]. The change in R with respect to t1/2 was almost perfectly fitted by a linear function (Fig. 10) indicating a diffusional growth. It should be noted that although the crystallized volume fraction is up to 100%, there is considerable amount of remaining amorphous phase. Therefore, the first crystallization event only comprises the fcc-Al formations. The EDS results after the first exothermic reaction, clearly reveal the enrichment of amorphous matrix with Tb atoms. Further growth of fcc-Al is restricted by the disability of the long range diffusion of Tb within the remaining amorphous matrix by soft impingement. The average size of fcc-Al nanocrystals saturates to a value of 40 nm before the second crystallization event and it results in high population of fcc-Al nanocrystals embedded in the amorphous matrix. Whether fcc-Al crystals grow from the pre-existing nuclei or nucleation occurs within the amorphous regions is still not clear. JMA predicts a constant nucleation rate. FEM studies indicated fcc-like MRO within the Al-rich nano-domains. These MRO structures may act as the nuclei of the nanocrystalline fcc-Al. Under constant nucleation rate conditions, the change in number density of nuclei should be linear with respect to annealing time. It is difficult to observe this change within the limitations of electron microscopy. Fig. 7(b) indicates that the change in nucleation density of nanocrystals is steady around 2 × 1021 m−3. On the other hand, it is quite possible to have a range of pure Al cluster size. The ones with critical size may act as the pre-nuclei. However, the others with “undercritical” size stay as amorphous and upon annealing they grow and eventually crystallizes to fcc-Al. Previous APT studies revealed the interconnection of Al-rich clusters without crystallization upon annealing before the first exothermic event. Moreover, relaxation of these glassy alloys induces a clear glass-transition temperature, indicating the amorphous nature of the as-quenched alloys. The local Avrami coefficient, n(x) starts from a value of 1.6 indicating a zero nucleation rate thus growth of pre-existing nuclei. This may indicate the growth of MRO present in the as-quenched state. For the main part of crystallization regime, n(x) stays around 2.5, which is in good agreement with the Avrami exponent calculated by linear regression. The sharp increase of n(x) after 85% of crystallization is usually attributed to a combination of structural and compositional effects [25,26]. The inhomogeneous distribution of nuclei in amorphous ribbon specimens is a common reason for the sharp increase in n(x) at higher crystallization fractions [27]. Because of this effect, the local Avrami exponent should be considered carefully otherwise
Fig. 9. (a) The JMA plot of (a) ln(−ln(1 − x)) versus ln (t − τ) for fcc-Al crystallization. The line represents the best fit by linear regression with a correlation coefficient (R2) of 0.99992. (b) The derivative of JMA plot with respect to crystallized fraction.
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Fig. 10. The average nanocrystal size versus t1/2 graph for fcc-Al crystallization. The line represents the best fit by linear regression with a correlation coefficient (R2) of 0.9945.
it may lead to erroneous conclusions on nucleation and growth mechanisms. If possible, JMA type modeling studies on crystallization amorphous materials should be conducted with corresponding electron and X-ray based techniques. In our case dominant factor of observing anomalous Avrami exponents should be the inhomogeneous distribution of Al-rich clusters due to ribbon geometry of the specimens. It is a well known fact that the wheel side solidifies at a higher cooling rate as compared to the free side of the as-quenched ribbon [28]. The average thickness of the ribbon is around 30 μm. It is not an easy task to quantify and compare the Al-rich clusters on the wheel and free sides of the ribbon. On the other hand, microhardness measurements are quite sensitive to the percent nanocrystallization of the amorphous matrix [29]. In order to compare the percent nanocrystallization with respect to annealing time, specimens were subjected to microhardness tests from their wheel and free sides on Al and steel block substrates. The test results are given in Fig. 11. The average hardness values on Al and steel substrates are found to be close to each other. This indicates that ribbons have sufficient thickness to conduct the microhardness tests and the effect of substrate in hardness measurement is minor. In as-quenched amorphous state, ribbons have an average value of 185 HV on both wheel and free sides. However, after the first interrupted quench at point A, the hardness at wheel side is considerably higher as compared to free side up to point E, where hardness of both sides becomes equal. After the fcc-Al crystallization is completed, wheel and free sides reach similar hardness values of 280 HV. The increase in hardness of
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the partially crystalline ribbon specimens is due to the formation of defect-free nanocrystalline fcc-Al and the increase of Tb content within the amorphous matrix with the rejection of RE atoms during the growth process [4]. Hardness measurements clearly show that these effects have occurred faster on the wheel side. Al-rich clusters (potential sites for fcc-Al crystals) are expected to be smaller and highly populated on the wheel side of ribbon due to the higher cooling rate. These small clusters having relatively large surface/volume ratio accelerate the nucleation and growth process and cause excess amount of RE rejection per unit time from the growing fcc-Al nanocrystals on the wheel side of the ribbon. This results in a sharp increase and a fast exhaustion in wheel side hardness. It should be noted that the amount of RE rejection per unit time cannot alter the main growth mechanism, which is a three-dimensional diffusional growth (n = 2.5) in our case. However, it will definitely change the conditions for soft impingement towards to end of the fcc-Al crystallization. We believe that this discrepancy is the dominant factor for the formation of anomalous local Avrami exponent observed at high crystallized fractions. Although mechanical properties are left out of the scope of this current study, it is also interesting to note that the hardness values measured for partially crystalline Al90Tb10 are quite high as compared to most of the commercial Al alloys used in aerospace and defense industries. As an example, the critical aerospace material, AA2024, has hardness values of 150 HV and 195 HV after aging and equal channel angular pressing (ECAP) process, respectively [30]. In our case, nanocrystalline ribbons show much higher hardness values indicating the potential strength of these materials. The presence of highly populated nanocrystals embedded in the amorphous matrix was previously reported to reduce the well-known plasticity problem of the metallic glass systems. Therefore, provided that enough thickness is achieved, Al–RE nanocrystalline-amorphous metallic glass composites are among the strong candidates for the next generation light-weight high strength alloys. 5. Conclusion The initial isothermal devitrification products of amorphous Al90Tb10 ribbon specimens were investigated in details using combined techniques of DSC, XRD and TEM. Detailed analysis of crystallization kinetics was performed through JMA approach using isothermal DSC scan at 220 °C. The first devitrifying crystalline phase was found to be fcc-Al nucleated at a very high population within the amorphous matrix. The change in number density of fcc-Al during the first crystallization was found to be steady between
Fig. 11. Hardness values measured for as-quenched amorphous and the critical crystallization states on (a) Al and (b) steel substrates.
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1.5 × 1021 and 2.5 × 1021 m− 3 according to TEM and HRTEM analyses. JMA indicated a three-dimensional diffusional growth with constant nucleation rate. Microstructural observations agreed well with the three-dimensional diffusional growth. The average size of nanocrystals saturate at 40–50 nm resulting a RE-rich amorphous matrix. The local Avrami exponent agrees well with the Avrami expression for the major part of crystallization process. However, it showed a sharp non-linearity for the high crystallized fractions. This deviation was found to be related to inhomogeneous nature of nanocrystallization through the thickness based on the microhardness measurements from the free and wheel sides of ribbon specimens. TEM and JMA analyses are in good agreement with the possibility of having MRO pre-nuclei of fcc-Al in addition to regions of pure Al clusters with “undercritical” size. The highly populated fcc-Al nanocrystals were found to nucleate within these regions. The mutual volume percentage of these regions is affected by the effective cooling rate of the molten metal. Acknowledgment The work at the Middle East Technical University was supported by BAP-08-11-2011-106. The assistance of the Materials Preparation Center of the Ames Laboratory is acknowledged for supplying the samples. References [1] H. Gleiter, Prog. Mater. Sci. 33 (4) (1989) 223–315. [2] K. Lu, Mater. Sci. Eng. R. 16 (4) (1996) 161–221. [3] M.A. Meyers, A. Mishra, D.J. Benson, Prog. Mater. Sci. 51 (2006) 427–556.
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