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Microstructural and nuclear magnetic resonance studies of solid-state amorphization in Al–Ti–Si composites prepared by mechanical alloying
Acta Materialia 52 (2004) 4133–4142 www.actamat-journals.com
Microstructural and nuclear magnetic resonance studies of solid-state amorphization in Al–Ti–Si composites prepared by mechanical alloying I. Manna
a,*
, P. Nandi a, B. Bandyopadhyay b, K. Ghoshray b, A. Ghoshray a
b
Metallurgical and Materials Engineering Department, I.I.T., Kharagpur 721 302, India b Saha Institute of Nuclear Physics, 1/AF, Bidhannagar, Kolkata 700 064, India
Received 28 December 2003; received in revised form 14 May 2004; accepted 17 May 2004 Available online 17 June 2004
Abstract Three Al30 Ti70 x Six (x ¼ 10, 20, 30), along with an Al-rich (Al50 Ti40 Si10 ) and an Al-lean (Al10 Ti60 Si30 ) elemental powder blends were subjected to mechanical alloying by high-energy planetary ball milling to yield a composite microstructure with varying proportions of amorphous and nanocrystalline intermetallic phases. Microstructural characterization at different stages of milling was carried out by X-ray diffraction, high-resolution transmission electron microscopy and energy dispersive X-ray spectroscopy. Furthermore, 27 Al nuclear magnetic resonance (NMR) studies were undertaken to probe the mechanism of solidstate amorphization. Ball milling leads to alloying, nanocrystallization and partial solid-state amorphization followed/accompanied by strain-induced nucleation of nanocrystalline intermetallic phases from an amorphous solid solution. Both these amorphous and nano-intermetallic phases are associated with characteristic NMR peaks at lower frequencies (than that of pure Al). Thus, mechanical alloying of Al–Ti–Si appears a suitable technique for developing nanocrystalline intermetallic phase/compound dispersed amorphous matrix composites. Ó 2004 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Mechanical alloying; Nuclear magnetic resonance; Aluminum alloy; Amorphous; Nanocrystalline
1. Introduction Al-alloys are the most widely used material for highspecific strength structural applications in aviation and transportation industry. While age hardenable Al-alloys can attain a maximum strength level of 500–600 MPa, it is reported that significant improvement in compressive strength up to 1200–1400 MPa is possible in Al-alloys in amorphous or nanocrystal dispersed amorphous condition [1–3]. For this reason, development of amorphous Al-alloys, particularly Al-based bulk amorphous alloys by solid-state processing has received considerable research attention in the recent past [4–6].
*
Corresponding author. Tel.: +91-3222-283266; fax: +91-3222282280. E-mail address: [email protected] (I. Manna).
Mechanical alloying is a versatile solid-state synthesis route to develop partially or completely amorphous/ glassy microstructure from elemental powder blend, alloys or intermetallic phases/compounds [7–9]. Besides mechanical attrition, solid-state amorphization is also feasible by sandwich/pack rolling [10], heavy-duty wear/ erosion [11] and equi-channel angular pressing [12] that involves severe plastic deformation at slow strain rate. In contrast, solid-state amorphization by mechanical alloying involves deformation at much higher strain rate, cold welding, fragmentation and dynamic recrystallization [9]. Mass transport accompanying mechanical alloying occurs at a much faster effective diffusion rate equivalent to volume diffusivity at an effective elevated temperature [13]. Solid-state amorphization is usually favored in systems with large negative heat of mixing and/or large difference in diffusivity of the constituents at the reference temperature [14]. While the
1359-6454/$30.00 Ó 2004 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2004.05.026
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former condition favors formation of intermetallic phases, the latter restricts or prevents nucleation and growth of crystalline phases, and thereby, promotes retention of amorphous structure. Development of amorphous microstructure also in alloys with positive enthalpy change suggests that increase in interfacial and strain energy contributions may provide significant contribution in solid-state amorphization [15]. Besides alloys, mechanical milling of intermetallic compounds generates adequate density of lattice defects, cause disordering and raise the internal energy of the system to provide the driving force for amorphization [16]. Unlike rapid solidification processing, the microstructure produced by mechanical attrition retains the strain and lattice disorder in a highly metastable state [17]. As a result, the composition range for amorphization is usually greater in mechanically driven processes than that by rapid solidification techniques. Besides strain, chemical reaction may be responsible for amorphization in the solid state over a wide composition range. In addition, contributions from increase in grain boundary area due to nanocrystallization [18] and lattice expansion/strain [19] may aid amorphization. In this regard, impurities may also play an important role in determining the level of lattice distortion, strain/strain rate and destruction of long-range order or crystallinity [20]. We have recently synthesized a number of Al-rich amorphous and nanocrystalline alloys by mechanical alloying route [4–6]. Besides characterizing the microstructural evolution, we have noted the strong influence of initial composition and miling parameters on the extent (partial/complete) of solid-state amorphization. However, the precise mechanism of amorphous phase formation in these alloys has not been identified. In the present study, we have carried out a systematic effort of synthesis and microstructural characterization of a number of Al–Ti–Si ternary alloys. In addition, nuclear magnetic resonance (NMR) studies have been undertaken for selected alloys to determine the mechanism of partial/complete solid-state amorphization.
2. Experimental Al–Ti–Si elemental powder blends with each elemental constituent having at least 99.5 wt% purity and about 50–100 lm particle size were subjected to high-energy mechanical attrition in FRITCH P5 planetary ball mill with ball to powder ratio of 10:1 using tungsten carbide (WC) coated vial and balls (10 mm diameter). Milling is done in wet (toluene) medium to prevent agglomeration of powders and to retard oxidation beyond initial stage of milling. It also avoids welding/coating of Al to the milling media (balls/vial) and thereby changing the milling dynamics. Mechanical alloying was carried out with three different ternary powder blends with the same
Al content (30 at.%) but varying proportions of Ti and Si, namely Al30 Ti60 Si10 , Al30 Ti50 Si20 and Al30 Ti40 Si30 . Two more blends, one Al-rich (Al50 Ti40 Si10 ) and another Al-lean (Al10 Ti60 Si30 ) were also selected to study the role of Al-content in solid-state amorphization. It may be noted that Al content is varied over a wide range as the NMR spectra related to Al are studied in this work as a probe to monitor amorphization. The identity and sequence of phase evolution in different stages of mechanical alloying were studied by X-ray diffraction (XRD) analysis using a Philips (PW1710) diffractometer with Cu-Ka (0.1542 nm) radiation. Average grain size (dc ) was determined from broadening of the most intense peak of the concerned phases using Voigt method [21] that allows judicious elimination of the contributions due to instrumental and strain effects in the observed broadening of the peak profile and comparing the latter with that from a standard/annealed sample with identical composition. It may be noted that Voigt analysis is based on Scherrer principle of crystallite size determination using XRD-analysis [22]. The results of the XRD-analysis concerning grain size and amorphization were verified by transmission electron microscopy (TEM) using a Philips CM-20 TEM instrument operated at 200 kV using high-resolution, bright field images, and selected area diffraction (SAD) analysis. The composition of the selected milled product at appropriate stages was analyzed/verified by energy dispersive X-ray spectroscopy (EDS) attached to the TEM during microstructural investigation. 27 Al NMR was measured at room temperature at a magnetic field of 7.04 tesla. The spectra of the samples were spread over a wide frequency region, and were recorded by sweeping the frequency in steps, in which a narrow region of the spectrum was excited with a long radiofrequency (rf) pulse at low power, using the solidecho pulse sequence 90–s1 –90–s2 –echo with phase cycling. The amplitude of the Fourier transformed time domain signal at each step was measured and plotted. The narrow spectra were obtained by exciting in the middle of the spectrum by a short rf-pulse at high power. The reference frequency, i.e., 27 Al signal in a solution of AlCl3 was at 78.1608 MHz. In addition, magnetic susceptibility of Al50 Ti40 S10 was measured at room temperature with a vibrating sample magnetometer (Model 155 of EG&G, Princeton Applied Research) to probe the presence of magnetic impurities that could affect the NMR measurements.
3. Results and discussion 3.1. Magnetic susceptibility studies Fig. 1 shows the magnetic field dependence of magnetization of Al50 Ti40 S10 at different stages of milling.
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0.07 0h 5h 10 h 15 h 20 h 30 h Ti metal
Magnetization (emu/g)
0.06 0.05 0.04 0.03 0.02 0.01 0.00
0
4
8
12
16
20
H (kOe) Fig. 1. Field dependence of magnetization of Al50 Ti40 Si10 at different hours of milling at 300 K.
Among the three constituent elements, only Ti contained a small amount of ferromagnetic impurity. The magnetic susceptibility of most metals that do not have localized electrons, is due to the Pauli paramagnetism of free electrons. The paramagnetism of Ti is one order higher in magnitude than that of Al and Si is diamagnetic. The non-linearity at small field values is due to small amount of magnetic impurity present only in Ti. As the formation of the alloy proceeds with increasing milling time, the susceptibility of Al50 Ti40 S10 decreases. It indicates that the metallic character of the alloy diminishes with increase in milling time. Nevertheless,the magnetic susceptibility of Al50 Ti40 S10 after 30 h of mechanical alloying remains Pauli paramagnetic. 3.2. Studies on microstructural evolution by XRD and TEM Fig. 2 shows the XRD-profiles of Al30 Ti60 Si10 powder blend subjected to different hours of planetary ball milling. It is evident that the characteristic peaks of the constituent elements gradually undergo decrease in intensity and increase in width (full width at half maximum, FWHM) to finally disappear beyond 10 h of milling. Continued milling up to 30 h leads to evolution of at least two new sets of peak due to WC (impurity from milling media) and AlTi3 . The latter increases in volume fraction during milling from 20 to 30 h. Since the profile at 10 h records no characteristic peak of AlTi3 except a broad halo between 2h ¼ 35° to 44°, continued milling of the predominantly amorphous matrix from 10 h onwards seems to induce strain-induced nucleation of AlTi3 from an amorphous solid solution. Similar observation was earlier reported by us as well as others [23,24]. It may be mentioned that no significant change in microstructure is noted beyond 30 h of milling.
Fig. 2. XRD-profiles of Al30 Ti60 Si10 powder blend subjected to different hours of planetary ball milling. Note that nanocrystalline AlTi3 forms from an amorphous matrix on continued milling beyond 10 h.
Fig. 3(a) shows the high-resolution transmission electron microscope (HRTEM) image of Al30 Ti60 Si10 following 20 h of mechanical alloying. It is evident that the microstructure is predominantly nanocrystalline with a small volume fraction of region among the crystallites devoid of lattice fringes. At least two types of interplaner spacings (d) are recorded from the microstructure which correspond well with dð201Þ ¼ 0:2162 nm of AlTi3 [25] and dð101Þ ¼ 0:1881 nm of WC [26]. The corresponding SAD pattern shows diffraction rings characteristic of the AlTi3 crystallites (Fig. 3(b)). The isolated spots are likely due to WC impurity picked up during milling. Similar microstructural analysis at 30 h milled sample reveals identical nanocrystalline aggregate with marginally higher volume fraction and grain size of AlTi3 . Perhaps, continued milling increases the volume fraction and size of AlTi3 crystallites at the expense of amorphous matrix. Fig. 4 reveals the summary of XRD patterns of the Al30 Ti50 Si20 powder blend during high-energy ball milling from 0 to 30 h. The peaks due to the constituent elements undergo reduction in intensity and broadening
Fig. 3. (a) High-resolution transmission electron microscope (HRTEM) image of Al30 Ti60 Si10 following 20 h of mechanical alloying and (b) the corresponding SAD pattern.
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Fig. 4. XRD-profiles of Al30 Ti50 Si20 powder blend subjected to highenergy planetary ball milling for 0–30 h. Note that nanocrystalline Ti5 Si3 forms from an amorphous matrix on continued milling beyond 20 h.
of width until 10 h to finally form a predominantly amorphous aggregate at 20 h of milling. Continued milling up to 30 h appears to induce strain-induced nucleation of nanocrystalline Ti5 Si3 or its extended solid solution. Thus, the microstructural evolution of Al30 Ti50 Si20 during high energy ball milling appears identical as that of Al30 Ti60 Si10 except the identity of the nanocrystalline phase that evolves from the amorphous matrix. This difference could be attributed to the difference in initial blend composition. Fig. 5(a) shows the HRTEM image of Al30 Ti50 Si20 after 24 h of milling. The microstructure reveals a typical amorphous + nanocrystalline composite aggregate. The thinner regions of the powder particles (along the edges) are predominantly amorphous. The crystalline region records a d-spacing that matches well with dð211Þ ¼ 0:2198 nm of Ti5 Si3 [27]. Fig. 5(b) shows the corresponding SAD pattern that records both diffraction rings and spots due to the presence of both nanoand micro-crystallites of Ti5 Si3 respectively. It may be mentioned that the volume fraction of amorphous region is negligible at 30 h of milling. Thus, Al30 Ti60 Si10 seems more prone to amorphization followed by strain-
Fig. 5. (a) HRTEM image of Al30 Ti50 Si20 after mechanical alloying for 24 h and (b) the corresponding SAD pattern.
induced nucleation of nanocrystalline intermetallic phase by continued milling than that of Al30 Ti50 Si20 . Fig. 6 shows the XRD profiles of the Al30 Ti40 Si30 powder blend as a function of milling time during highenergy planetary ball milling. The microstructural evolution seems identical as that in Al30 Ti60 Si10 and Al30 Ti50 Si20 except that the microstructure does not turn predominantly amorphous at an intermediate stage (10– 20 h) of milling. Instead, nanocrystalline Ti5 Si3 (or Ti7 Al5 Si12 ) and Al3 Ti seem to nucleate from the extended solid solution that forms between 5 and 10 h following mutual dissolution of constituent elements by mechanical alloying. It is interesting to note that the Ti5 Si3 (211) peak shrinks in width between 20 and 30 h of milling indicating perhaps an increase in the crystallite size during continued mechanical alloying. Such grain coarsening or crystallite growth may be aided by enhanced mass transport [13] during extended hours of milling when no new phase evolution or grain size reduction is feasible. Fig. 7(a) shows the HRTEM image of Al30 Ti40 Si30 milled for 30 h. The microstructure consists of very fine crystallites (<10–15 nm) embedded in amorphous matrix. Careful observation in other areas reveals similar amorphous + nanocrystalline microstructure with varying degree of relative volume fraction. Fig. 7(b) shows the corresponding SAD pattern confirming that 30 h of mechanical alloying of Al30 Ti40 Si30 produces a nanocrystal dispersed amorphous matrix composite microstructure. In all the three alloys, formation of amorphous or nanocrystalline solid solution seems a prerequisite for strain-induced nucleation of the intermetallic phases during continued milling. Perhaps, grain size reduction or nanocrystallization by mechanical alloying facilitates formation of a nanocrystalline or amorphous extended
Fig. 6. XRD profiles of the Al30 Ti40 Si30 powder blend as a function of milling time. Note that nanocrystalline Ti5 Si3 nucleates at an early stage (10 h) of milling.
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Fig. 7. (a) HRTEM image and (b) SAD pattern of Al30 Ti40 Si30 milled for 30 h.
solid solution in 10–20 h of milling. Subsequently, continued milling enables strain-induced nucleation of a relevant intermetallic phase (Ti5 Si3 , AlTi3 , TiAl3 , etc.) that possesses significantly more negative enthalpy of formation ðDH Þ than the amorphous/nanocrystalline solid solution. The related thermodynamic analysis to predict the phase equilibrium is discussed in the following section. In order to investigate the role of Al in microstructural evolution of the present ternary systems, two more powder blends were subjected to mechanical alloying under identical conditions. Fig. 8 shows the XRD profiles of the Al-lean Al10 Ti60 Si30 powder blend subjected to mechanical alloying up to 30 h. It is not clear whether an amorphous phase forms in this powder blend. However, nanocrystalline Ti5 Si3 and AlTi3 peaks are detected at the final stage of milling (along with WC impurity). On the other hand, XRD profiles in Fig. 9 shows that mechanical alloying of Al-rich Al50 Ti40 Si10 powder blend leads to formation of a two-phase microstructure comprising amorphous and nanocrystalline intermetallic phase (Al3 Ti). Though it is not clear whether solid-state amorphization preceded nucleation of nanocrystalline Al3 Ti or vice versa, the final micro-
Fig. 8. XRD profiles of the Al-lean Al10 Ti60 Si30 powder blend subjected to mechanical alloying up to 30 h.
Fig. 9. XRD profiles of the Al-rich Al50 Ti40 Si10 powder blend subjected to mechanical alloying up to 30 h.
structure (at 30 h) seems predominantly amorphous with the presence of nanocrystalline Al3 Ti in it. Fig. 10(a) shows HRTEM evidence of nanocrystal dispersed amorphous matrix aggregate of Al50 Ti40 Si10 after 30 h of mechanical alloying. The corresponding SAD pattern shows a diffused intensity halo with additional low intensity rings substantiating the XRD and HRTEM evidences of amorphous + nanocrystalline composite microstructure of the present alloy after 30 h of milling. Comparison of microstructures at 20 and 30 h of milling reveals that the volume fraction of the nanocrystalline phase marginally decreases between 20 and 30 h of milling. Perhaps, further microstructural disorder/defect introduced by continued milling would transform the microstructure completely amorphous. It
Fig. 10. (a) High resolution TEM image of the milled product following 30 h of milling of Al50 Ti40 Si10 showing a predominantly amorphous microstructure with a few isolated nanometric (<8 nm) crystalline grains and (b) SAD pattern of the same Al50 Ti40 Si10 sample (from another area) showing amorphous as well as nanocrystalline grains with a few bright spots due to some coarse crystalline grains.
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is worth mentioning that isolated nanocrystalline regions existed in the microstructure even after 50 h of milling. Possibly, milling with higher strain rate, at lower temperature or with a different composition is necessary for complete solid-state amorphization by mechanical alloying. It is predicted that nanocrystalline dispersed amorphous rather than single phase amorphous or nanocrystalline microstructure is more likely to yield greater strength and toughness [1,28]. From the present results, it appears that mechanical alloying is a convenient method of producing in situ nucleated nanocrystalline intermetallic dispersed amorphous matrix composite in Al-rich Al–Ti–Si alloys by appropriate choice of composition and milling parameters. 3.3. Thermodynamic analysis The identity and sequence of phase evolution during mechanical alloying is a complex function of the initial powder blend, thermo-physical properties of the constituents, and employed milling parameters/conditions. In a mechanically driven system like mechanical alloying, the phase equilibrium is based on non-equilibrium thermodynamics involving both equilibrium as well as metastable or non-equilibrium phases including extended and amorphous solid solutions. In this regard, the models by Miedema [29,30] provide a convenient tool to estimate the enthalphy/heat of formation ðDH Þ of different phases in a binary system as follows: 1=3 2 DH ¼ K½P ðD/ Þ2 þ QðDgws Þ R;
ð1Þ
where K is the proportionality constant, Dgws is the difference in electron density at the boundary of Wigner–Seitz (ws) cells and D/ the difference in electronegativity parameters of the concerned elements. Q, P and R are empirical constants, dependent on the type of elements (transition or non-transition). The necessary parameters and their values are documented in the literature [31]. As already pointed out, this approach is valid only for binary systems. Since a similar universally applicable
formulation to calculate DH for ternary or multi-component system is not available, we have calculated DH for the present alloys using a weighted mean approach as follows: DHABC ¼ 0:5ðXAB DHAB þ XBC DHBC þ XAC DHAC Þ:
Here, X and DH represent the mole fraction and enthalpy of formation of the binary or ternary phases comprising elements A, B, C (as shown in the subscripts), respectively. Though XA þ XB þ XC ¼ 1, the relative concentration of each element in any given binary combination for calculating the terms DHAB , DHBC or DHAC needed in Eq. (2), say in Al50 Ti40 Si10 , is taken as XAlTi ¼ 0:9, XAlSi ¼ 0:6 and XTiSi ¼ 0:5, such that, 0.5 ðXAlTi þ XTiSi þ XAlSi Þ ¼ 1. For calculating the respective binary DH terms, say DHAB for A–B, the corresponding relative mole fraction of A in A–B, say XA0 ðin ABÞ ¼ XA =ðXA þ XB Þ. Thus, XA0 in A–B may be different than that in A–C. For instance, XA in A–B and in A–C for Al50 Ti40 Si10 are 0.55 and 0.83, respectively. However, ½XA0 ðin ABÞ þ XA0 ðin ACÞ þ XB0 ðin ABÞ þ XB0 ðin BCÞ þ XC0 ðin ACÞ þ XC0 ðin BCÞ =3 ¼ 1 for the same alloy, i.e., Al50 Ti40 Si10 . These calculations are based on Vegard’s law type assumption that enthalpy is a linear function of mole fraction. Table 1 summarizes the DH values calculated for the present ternary alloys as well as the stable binary intermetallic phases/compounds possible in the concerned systems. It is apparent that the predicted DH values for crystalline and amorphous solid solutions are comparable, while the same for some of the intermetallic compounds are significantly smaller. Hence, it is probable that mutual dissolution during early stage of milling leads to development of crystalline or amorphous solid solution. Enhanced rate of mass transport, significant grain size reduction (nanocrystallization) and high degree of plastic strain could lead to strain-induced nucleation of intermetallic phases with more negative DH from the nanocrystalline or amorphous solid solutions. However, direct nucleation of the intermetallic phase/compound may concur or succeed the formation of the nanocrystalline/amorphous solid solution de-
Table 1 Enthalpy of formation ðDH Þ of the possible phases in the Al–Ti–Si system Possible solid solution or intermetallic phases Al50 Ti40 Si10 Al30 Ti60 Si10 Al30 Ti50 Si20 Al30 Ti40 Si30 Al10 Ti60 Si30 AlTi3 TiAl3 TiAl Ti5 Si3
ð2Þ
Enthalpy of formation ðDH Þ in kJ/mol Crystalline solid solution
Amorphous solid solution
Intermetallic phase/compound
)27.62 )26.09 )32.73 )34.84 )30.83
)28.89 )24.46 )36.04 )39.98 )31.55
– – – – – )38.04 )39.43 )60.44 )78.45
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3.4. Mechanism of amorphization by NMR investigation In the absence of localized moments, the position of the resonance line in a metallic sample is determined by, (a) the Knight shift, K, originating from the Fermi contact interaction ðHhf ¼ 2ð8p=3ÞlB cN hI:SðrÞdðrÞÞ between the nuclei and the conduction electrons, (b) the dipolar interaction between nuclear and electronic spin, and (c) the interaction between nuclear spin with the orbital motion of the electrons related to van Vleck paramagnetism [32,33]. The first interaction produces an isotropic shift, the second interaction is anisotropic in nature, and the contribution of the third interaction is negligible in case of Al resonance. The width and shape of the resonance line are determined by the nuclear dipolar interaction, anisotropic Knight shift in case of non-spherical character of the conduction band, and the broadening produced by the bulk magnetic susceptibility of the sample. Additionally, for nuclei with spin I > 1=2 being studied in a compound or an alloy, the interaction between the nuclear electric quadrupole moment and the electric field gradient present around the nucleus, also causes a broadening of the resonance line in polycrystalline samples. The Knight shift can be written, using the definition of the hyperfine field, Hhf , as K ¼ ðHhf =NA lB Þv, where v is the molar magnetic susceptibility and NA is Avogadro’s number. For simple metals and alloys the magnetic susceptibility is dominated by the Pauli paramagnetism, as discussed in Section 3.1 of magnetic susceptibility; for which v ¼ 2NA l2B N ðEF Þ, where N ðEF Þ is the density of state at the Fermi level. So for the Knight shift in the case of simple metal, K ¼ ð1=2Þðce =cN ÞN ðEF ÞA, where A is the hyperfine coupling constant. Fig. 11 shows the 27 Al NMR spectra of Al50 Ti40 Si10 at various intervals of milling. It is evident that the peak due to pure Al (at 78.258 MHz) is distinct at the initial stage of milling (5 h), but reduces in intensity with increased broadening of the linewidth until 10 h. Upon further milling, the resonance component from pure Al is greatly reduced in intensity and becomes very weak after 15 h. However, a trace of pure Al is obtained in NMR spectrum until 20 h of milling. The concerned XRD-pattern of this alloy at similar stage of milling (Fig. 9) substantiate that elemental Al undergoes mutual dissolution to form amorphous and another crystalline phase beyond 10 h of milling.
Al50Ti40Si10 alloy 5h 10 h 15 h 20 h 30 h
Intensity (arb. unit)
pending on the kinetic factors and nucleation barrier. Thus, the present thermodynamic calculations based on simplified weighted mean DH approach predict that an amorphous + nanocrystalline intermetallic phase aggregate or composite microstructure is anticipated in mechanical alloying of the present alloys, which corroborates the experimental results obtained by XRD and TEM analysis.
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30 h 20 h 15 h
10 h
5h
78.1
78.2
78.3
78.4
78.5
Frequency (MHz) Fig. 11. 27 Al NMR at 300 K in Al50 Ti40 Si10 alloys at different hours of milling. The vertical line on x-axis indicates the reference position of Al2 O3 .
In addition to the pure Al line as mentioned above, the sample after 10 h of milling shows a resonance line of broad structure at a lower frequency compared to the resonance line of pure Al metal. Two peaks can be distinguished in the broad structure, indicating the occurrence of at least two different alloy products, both having Al as a component, at this stage of milling. The resonance line extends to frequencies higher than that covered by the resonance in pure Al. This may be an indication of Al nuclei in one of the intermediate products of milling, experiencing an environment of non-cubic electric field gradient, and producing a quadrupolar interaction. This large broadening may also arise because of strain-induced distortion around the aluminum nucleus as a result of milling. It is to be noted that the X-ray diffraction pattern (Fig. 9) in this stage showed the constituent elements, though with huge broadening of Bragg peaks, but could not indicate the presence of any alloy product. Upon further milling, the resonance component from pure Al is greatly reduced in intensity, and the broad structure tend to become symmetric, now having only one peak, positioned close to the reference frequency. It indicates that of the many alloy components which were formed at intermediate stages of milling, only one or a few are retained. In the sample milled for 30 h, 27 Al resonance shows only a broad asymmetric line positioned near the reference and gives no trace of pure Al. A possible explanation for the broad NMR spectrum at the intermediate stage and its subsequent narrowing is that an amorphous extended solid solution phase is formed at an early stage of milling. Continued milling leads to the strain-induced nucleation of the stable coordination or environments of the constituent elements which are retained in the form of either nanocrystalline or amorphous intermetallic compounds, as products until the final stages of alloying.
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Fig. 12 shows the NMR spectra of Al50 Ti40 Si10 and Al10 Ti60 Si30 at the final milling stage (30 h). The two alloys shown in this figure represent the composition limits of the alloys investigated in this study in terms of Al-content. More precisely, Al50 Ti40 Si10 and Al10 Ti60 Si30 represent the Al-rich and Al-lean composition limits among the present alloys. It may be noted that the signal intensity, in case of Al-lean sample, is much smaller than that in case of Al-rich sample, and in former case, the data acquisition time was four times longer than in the latter. The spectrum obtained in Al50 Ti40 Si10 can be fitted as a sum of two components of resonance lines, both having a Lorentzian distribution, as shown in Fig. 12. The larger or dominant component has a linewidth (full width at half maximum, FWHM) of about 45 kHz and isotropic shift of 0.03%. The smaller component has a similar linewidth, but a slightly larger shift of about 0.045%. The alloy therefore consists of at least two components, which could be very similar to each other in their structure and properties. More significantly, the large linewidth, which is almost three times the width of pure Al resonance (Fig. 11), is definitely not only due to Al–Al nuclear dipolar interaction. The magnetic interactions in these samples should be negligible as is evident from the magnetic susceptibility data and do not contribute to the broadening of the resonance line. On the other hand, the grain size reduction and amorphization result in a large variation in the atomic environment around Al atoms in the alloy products, leading to a distribution in the local field around the Al nuclei, and giving rise to the increased linewidth. Also shown in Fig. 12 is the spectrum of Al10 Ti60 Si30 fitted, in this case too, as a sum of two Lorentzian components. One of the components has FWHM of Al50Ti40Si10 Al10Ti60Si30 Sintered Al50Ti40Si10 Sintered Al10Ti60Si30
Al50Ti40Si10 mechanical alloyed
Al30Ti60Si10 Al30Ti50Si20 Al30Ti40Si30
Al50Ti40Si10 sintered
Intensity (arb. unit)
Intensity (arb.unit)
50 kHz and shift 0.04%, which are similar to one described above. The other component shows a width of 36 kHz and negligible shift of 0.008%. The results therefore indicate the formation of two very different Al containing components in the alloy. Moreover, as discussed in the beginning of this section, the components might have different metallicity. This suggests that Allean samples might have a higher probability of yielding an inhomogeneous mixture in the alloy products. Fig. 13 presents a comparison of NMR spectra of the remaining alloys with same (30 at.%) Al-content at the final milling stage (30 h). The peak due to pure Al is consistently absent. The spectral features are similar to those of Al50 Ti40 Si10 at comparable stage (Fig. 11), though the spectra for Al30 Ti50 Si20 and Al30 Ti40 Si30 are distinctly bell-shaped unlike that of Al30 Ti60 Si10 . In the latter, spectrum seems to have two well-defined alloy components each having some distribution of the environment around Al resulting from amorphization. Suitable comparison of concerned XRD (Fig. 2) and HRTEM (Fig. 3) result reveals that Al30 Ti60 Si10 undergoes solid-state amorphization like Al50 Ti40 Si10 under comparable conditions of milling. Thus, the NMR analysis suggest that solid state amorphization in the present alloys proceeds by developing a new atomic environment or coordination (there may develop more than one Al alloy component as in Al30 Ti60 Si10 ) with the constituent elements during mechanical alloying that can be detected by the NMR spectrum. The bell-shaped spectra, on the other hand, indicate the presence of a single Al-based alloy component having a wider distribution in the surroundings of Al. This might occur due to high silicon content in the systems. This is quite evident from relevant XRD patterns of Al30 Ti50 Si20 (Fig. 4) and Al30 Ti40 Si30 (Fig. 6), respectively, wherein the lines belonging to AlTi3 are
Amorphous
Al10Ti60Si30 mechanical alloyed
Al10Ti60Si30 sintered
78.05
78.10
78.15
78.20
78.25
78.30
Frequency (MHz) Fig. 12. 27 Al NMR at 300 K in the Al–Ti–Si system after 30 h milling for Al lean (Al ¼ 10%) and Al rich (Al ¼ 50%) samples, before and after sintering. The continuous lines are theoretical fit to the spectra. The broken lines in both the spectra indicate the two constituent components used in the fit as described in the text. The vertical line indicates the reference position of Al2 O3 .
78.0
78.1
78.2 Frequency (MHz)
78.3
78.4
Fig. 13. 27 Al NMR at 300 K in the Al–Ti–Si systems after 30 h of milling with 30% Al in all three samples. The vertical line indicates the reference position of Al2 O3 .
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either much broader or totally absent as compared to that for Al30 Ti60 Si10 shown in Fig. 2. 3.5. Effect of sintering on phase equilibrium In order to identify the amorphous phases, Al50 Ti40 Si10 and Al10 Ti60 Si30 milled for 30 h, were heated at 20 °C/min to 700 °C, kept at that temperature for 1 h, and then rapidly cooled to room temperature and studied by XRD and NMR. The XRD pattern shown in Fig. 14(a), for Al50 Ti40 Si10 , exhibits reasonably welldefined diffraction peaks, characteristic of crystalline/ nanocrystalline phases, belonging to the intermetallics Al3 Ti and Ti5 Si3 , which, therefore, are the main components of the alloy. Fig. 14(b) corresponds to the Allean sample, and shows the formation of the compounds AlTi3 , as expected, and also Ti5 Si3 . The 27 Al NMR spectra for these two sintered samples are shown in Fig. 12 and can be compared with the spectra before sintering. The figure also shows the theoretical fit to the spectra. The spectrum for sintered Al50 Ti40 Si10 can be fitted to one Lorentzian distribution with a FWHM of 18 kHz and isotropic shift of 0.025%. Clearly, the resonance is obtained from crystalline Al3 Ti, as observed in X-ray. Sintering results in a more uniform local environment around Al nuclei compared to the as-milled amorphous samples and the linewidth is therefore reduced. The isotropic shift of the amorphous sample is larger than that of sintered sample and this shows a larger metallicity of amorphous sample. However, the amorphization-induced anisotropic broadening may also result in a shift in the mean position of the spectrum. The spectrum for sintered Al10 Ti60 Si30 can only be fitted as a sum of two components, as in case of asmilled sample. The dominant component corresponds to AlTi3 , and the linewidth of AlTi3 that was 36 kHz in as-
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milled amorphous sample becomes 18 kHz after sintering. The isotropic shift is also reduced from 0.008% to 0.001%. The other component that exists in a very small amount in the NMR spectrum could not be separately identified in XRD pattern. We believe that this unidentified alloy component exists in the amorphous sample as a metastable phase and tends to disappear upon sintering.
4. Summary and conclusion From the present study, it appears that development of nanocrystalline intermetallic dispersed Al-rich amorphous matrix composite is feasible by mechanical alloying of appropriate Al–Ti–Si powder blend. The intermetallic phases/compounds formed include TiAl3 , Ti3 Al, Ti5 Si3 and Ti7 Al5 Si12 depending on the initial composition selected. During mechanical alloying, the phase evolution sequence comprises mutual dissolution or alloying, nanocrystallization, partial/complete amorphization and finally in situ nucleation of nanocrystalline intermetallic phases from the amorphous precursor during continued milling. Enthalpy calculation based on modified Miedema model corroborates the nature and sequence of phase evolution mentioned above. NMR studies also substantiate that mechanical alloying of the present elemental blends leads to alloying, nanocrystallization, amorphization, followed by strain-induced nucleation of nanocrystalline intermetallic phases/compounds to develop a composite microstructure. Furthermore, the NMR results identify characteristic new environment of atomic coordination around Al atoms for both amorphous and nanocrystalline intermetallic phases. The spectra located at lower frequencies compared to that of pure Al metal indicating that the density of electrons at the Fermi level at Al site is reduced as a result of formation of these phases, giving rise to smaller values of Knight shift. Thus, NMR results as well as magnetic susceptibility measurements indicate that metallic properties are reduced following amorphization and strain-induced nucleation of nanocrystalline intermetallics. Further studies are in progress to consolidate the milled product and assess mechanical properties of the composite.
Acknowledgements
Fig. 14. XRD patterns of 30 h milled (a) Al50 Ti40 Si10 and (b) Al10 Ti60 Si30 quenched to room temperature after heating up to 700 °C.
The authors wish to thank Professors H.J. Fecht, F. Banhart and P.M.G. Nambissan for useful technical discussion on this work. Partial financial support from the Humbold Foundation, Bonn and Council of Scientific and Industrial Research, New Delhi (Grant No. 70 (0048) 03-EMRII) to one of the authors (I.M.) is gratefully acknowledged.
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References [1] Inoue A. Acta Mater 2000;48:279. [2] Masumoto T. Mater Sci Eng A 1994;179/180:8. [3] Kiminami CS, Basim ND, Kaufman MJ, Amateau MF, Eden TJ, Galbraith GJ. Key Eng Mater 2001;189–191:503. [4] Chattopadhyay PP, Gannabatula RNR, Pabi SK, Manna I. Scripta Mater 2001;45:1191. [5] Nandi P, Chattopadhyay PP, Pabi SK, Manna I. Mater Sci Eng A 2003;359:11. [6] Manna I, Chattopadhyay PP, Banhart F, Fecht HJ. Zeit Metal 2003;94(7):835. [7] Suryanarayana C. Prog Mater Sci 2001;46:1. [8] Koch CC. In: Cahn RW, Hassan P, Kramer EJ, editors. Processing of metals and alloys, vol. 15. Weinheim: VCH Publication; 1991. p. 193. [9] Fecht HJ. In: Koch CC, editor. Nanostructured materials – processing, properties, and applications. Norwich, NY: Noyes Publication, William Andrew Publishing; 2002. p. 73. [10] Eleskandarany MS. Metal Mater Trans A 1996;27:3267. [11] Schwartz RB, Petrich RR, Shaw CK. J Non-Cryst Solids 1985;76:281. [12] Oleszak D, Burzynska-Szyszko M, Matyja H. J Mater Sci Lett 1993;12:3. [13] Pabi SK, Das D, Mahapatra TK, Manna I. Acta Mater 1998;46:3501. [14] Hellstern E, Schultz L. Mater Sci Eng 1987;93:213. [15] Bai HY, Michaelsen C, Gente C, Bormann R. Phys Rev B 2001;63:064202. [16] Jang JSC, Koch CC. Mater Res 1990;5:498.
[17] Birringer R, Altounian Z, Srom Olsen JO, Schultz L. Mater Sci Eng 1988;97:317. [18] Desre PJ. Nanostruct Mater 1994;4(8):957. [19] Gaffet E, Faudot F, Harmelin M. Mater Sci Forum 1992;8890:375. [20] Chakk Y, Berger S, Weiss BZ, Brook Levinson E. Acta Metal Mater 1994;42:3679. [21] Keijser TH, Langford IL, Mittemeijer EJ, Vogel BP. J Appl Cryst 1982;15:308. [22] Klug HP, Alexander LE. X-ray diffraction procedures. New York: Wiley; 1967. p. 491. [23] Liu KW, Zhang JS, Wang JG, Chen GL. J Non-Cryst Solids 1997;215:140. [24] Liu KW, Zhang JS, Wang JG, Chen GJ. J Mater Res 1998;13:1198. [25] JCPDS 14-0451. [26] JCPDS 02-1055. [27] JCPDS 08-0041. [28] Zhou JB, Rao KP, Chung CY. J Mater Proc Tech 2003;139:434. [29] Boom R, deBoer FR, Miedema AR. J Less-common Met 1976;45:237. [30] Boom R, deBoer FR, Miedema AR. J Less-common Met 1976;46:271. [31] DeBoer FR, Boom R, Mattens WCM, Miedema AR, Niessen AK. In: DeBoer FR, Pettifor DG, editors. Cohesion in metals. 2nd ed. Amsterdam, Holland: North-Holland Physics Publishing; 1989. [32] Slichter CP. Principles of magnetic resonance. 3rd ed. SpringerVerlag; 1990. [33] Narath A. In: Freeman AJ, Frankle RB, editors. Hyperfine interactions. Academic Press; 1967.
Microstructural and nuclear magnetic resonance studies of solid-state amorphization in Al–Ti–Si composites prepared by mechanical alloying I. Manna
a,*
, P. Nandi a, B. Bandyopadhyay b, K. Ghoshray b, A. Ghoshray a
b
Metallurgical and Materials Engineering Department, I.I.T., Kharagpur 721 302, India b Saha Institute of Nuclear Physics, 1/AF, Bidhannagar, Kolkata 700 064, India
Received 28 December 2003; received in revised form 14 May 2004; accepted 17 May 2004 Available online 17 June 2004
Abstract Three Al30 Ti70 x Six (x ¼ 10, 20, 30), along with an Al-rich (Al50 Ti40 Si10 ) and an Al-lean (Al10 Ti60 Si30 ) elemental powder blends were subjected to mechanical alloying by high-energy planetary ball milling to yield a composite microstructure with varying proportions of amorphous and nanocrystalline intermetallic phases. Microstructural characterization at different stages of milling was carried out by X-ray diffraction, high-resolution transmission electron microscopy and energy dispersive X-ray spectroscopy. Furthermore, 27 Al nuclear magnetic resonance (NMR) studies were undertaken to probe the mechanism of solidstate amorphization. Ball milling leads to alloying, nanocrystallization and partial solid-state amorphization followed/accompanied by strain-induced nucleation of nanocrystalline intermetallic phases from an amorphous solid solution. Both these amorphous and nano-intermetallic phases are associated with characteristic NMR peaks at lower frequencies (than that of pure Al). Thus, mechanical alloying of Al–Ti–Si appears a suitable technique for developing nanocrystalline intermetallic phase/compound dispersed amorphous matrix composites. Ó 2004 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Mechanical alloying; Nuclear magnetic resonance; Aluminum alloy; Amorphous; Nanocrystalline
1. Introduction Al-alloys are the most widely used material for highspecific strength structural applications in aviation and transportation industry. While age hardenable Al-alloys can attain a maximum strength level of 500–600 MPa, it is reported that significant improvement in compressive strength up to 1200–1400 MPa is possible in Al-alloys in amorphous or nanocrystal dispersed amorphous condition [1–3]. For this reason, development of amorphous Al-alloys, particularly Al-based bulk amorphous alloys by solid-state processing has received considerable research attention in the recent past [4–6].
*
Corresponding author. Tel.: +91-3222-283266; fax: +91-3222282280. E-mail address: [email protected] (I. Manna).
Mechanical alloying is a versatile solid-state synthesis route to develop partially or completely amorphous/ glassy microstructure from elemental powder blend, alloys or intermetallic phases/compounds [7–9]. Besides mechanical attrition, solid-state amorphization is also feasible by sandwich/pack rolling [10], heavy-duty wear/ erosion [11] and equi-channel angular pressing [12] that involves severe plastic deformation at slow strain rate. In contrast, solid-state amorphization by mechanical alloying involves deformation at much higher strain rate, cold welding, fragmentation and dynamic recrystallization [9]. Mass transport accompanying mechanical alloying occurs at a much faster effective diffusion rate equivalent to volume diffusivity at an effective elevated temperature [13]. Solid-state amorphization is usually favored in systems with large negative heat of mixing and/or large difference in diffusivity of the constituents at the reference temperature [14]. While the
1359-6454/$30.00 Ó 2004 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2004.05.026
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former condition favors formation of intermetallic phases, the latter restricts or prevents nucleation and growth of crystalline phases, and thereby, promotes retention of amorphous structure. Development of amorphous microstructure also in alloys with positive enthalpy change suggests that increase in interfacial and strain energy contributions may provide significant contribution in solid-state amorphization [15]. Besides alloys, mechanical milling of intermetallic compounds generates adequate density of lattice defects, cause disordering and raise the internal energy of the system to provide the driving force for amorphization [16]. Unlike rapid solidification processing, the microstructure produced by mechanical attrition retains the strain and lattice disorder in a highly metastable state [17]. As a result, the composition range for amorphization is usually greater in mechanically driven processes than that by rapid solidification techniques. Besides strain, chemical reaction may be responsible for amorphization in the solid state over a wide composition range. In addition, contributions from increase in grain boundary area due to nanocrystallization [18] and lattice expansion/strain [19] may aid amorphization. In this regard, impurities may also play an important role in determining the level of lattice distortion, strain/strain rate and destruction of long-range order or crystallinity [20]. We have recently synthesized a number of Al-rich amorphous and nanocrystalline alloys by mechanical alloying route [4–6]. Besides characterizing the microstructural evolution, we have noted the strong influence of initial composition and miling parameters on the extent (partial/complete) of solid-state amorphization. However, the precise mechanism of amorphous phase formation in these alloys has not been identified. In the present study, we have carried out a systematic effort of synthesis and microstructural characterization of a number of Al–Ti–Si ternary alloys. In addition, nuclear magnetic resonance (NMR) studies have been undertaken for selected alloys to determine the mechanism of partial/complete solid-state amorphization.
2. Experimental Al–Ti–Si elemental powder blends with each elemental constituent having at least 99.5 wt% purity and about 50–100 lm particle size were subjected to high-energy mechanical attrition in FRITCH P5 planetary ball mill with ball to powder ratio of 10:1 using tungsten carbide (WC) coated vial and balls (10 mm diameter). Milling is done in wet (toluene) medium to prevent agglomeration of powders and to retard oxidation beyond initial stage of milling. It also avoids welding/coating of Al to the milling media (balls/vial) and thereby changing the milling dynamics. Mechanical alloying was carried out with three different ternary powder blends with the same
Al content (30 at.%) but varying proportions of Ti and Si, namely Al30 Ti60 Si10 , Al30 Ti50 Si20 and Al30 Ti40 Si30 . Two more blends, one Al-rich (Al50 Ti40 Si10 ) and another Al-lean (Al10 Ti60 Si30 ) were also selected to study the role of Al-content in solid-state amorphization. It may be noted that Al content is varied over a wide range as the NMR spectra related to Al are studied in this work as a probe to monitor amorphization. The identity and sequence of phase evolution in different stages of mechanical alloying were studied by X-ray diffraction (XRD) analysis using a Philips (PW1710) diffractometer with Cu-Ka (0.1542 nm) radiation. Average grain size (dc ) was determined from broadening of the most intense peak of the concerned phases using Voigt method [21] that allows judicious elimination of the contributions due to instrumental and strain effects in the observed broadening of the peak profile and comparing the latter with that from a standard/annealed sample with identical composition. It may be noted that Voigt analysis is based on Scherrer principle of crystallite size determination using XRD-analysis [22]. The results of the XRD-analysis concerning grain size and amorphization were verified by transmission electron microscopy (TEM) using a Philips CM-20 TEM instrument operated at 200 kV using high-resolution, bright field images, and selected area diffraction (SAD) analysis. The composition of the selected milled product at appropriate stages was analyzed/verified by energy dispersive X-ray spectroscopy (EDS) attached to the TEM during microstructural investigation. 27 Al NMR was measured at room temperature at a magnetic field of 7.04 tesla. The spectra of the samples were spread over a wide frequency region, and were recorded by sweeping the frequency in steps, in which a narrow region of the spectrum was excited with a long radiofrequency (rf) pulse at low power, using the solidecho pulse sequence 90–s1 –90–s2 –echo with phase cycling. The amplitude of the Fourier transformed time domain signal at each step was measured and plotted. The narrow spectra were obtained by exciting in the middle of the spectrum by a short rf-pulse at high power. The reference frequency, i.e., 27 Al signal in a solution of AlCl3 was at 78.1608 MHz. In addition, magnetic susceptibility of Al50 Ti40 S10 was measured at room temperature with a vibrating sample magnetometer (Model 155 of EG&G, Princeton Applied Research) to probe the presence of magnetic impurities that could affect the NMR measurements.
3. Results and discussion 3.1. Magnetic susceptibility studies Fig. 1 shows the magnetic field dependence of magnetization of Al50 Ti40 S10 at different stages of milling.
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0.07 0h 5h 10 h 15 h 20 h 30 h Ti metal
Magnetization (emu/g)
0.06 0.05 0.04 0.03 0.02 0.01 0.00
0
4
8
12
16
20
H (kOe) Fig. 1. Field dependence of magnetization of Al50 Ti40 Si10 at different hours of milling at 300 K.
Among the three constituent elements, only Ti contained a small amount of ferromagnetic impurity. The magnetic susceptibility of most metals that do not have localized electrons, is due to the Pauli paramagnetism of free electrons. The paramagnetism of Ti is one order higher in magnitude than that of Al and Si is diamagnetic. The non-linearity at small field values is due to small amount of magnetic impurity present only in Ti. As the formation of the alloy proceeds with increasing milling time, the susceptibility of Al50 Ti40 S10 decreases. It indicates that the metallic character of the alloy diminishes with increase in milling time. Nevertheless,the magnetic susceptibility of Al50 Ti40 S10 after 30 h of mechanical alloying remains Pauli paramagnetic. 3.2. Studies on microstructural evolution by XRD and TEM Fig. 2 shows the XRD-profiles of Al30 Ti60 Si10 powder blend subjected to different hours of planetary ball milling. It is evident that the characteristic peaks of the constituent elements gradually undergo decrease in intensity and increase in width (full width at half maximum, FWHM) to finally disappear beyond 10 h of milling. Continued milling up to 30 h leads to evolution of at least two new sets of peak due to WC (impurity from milling media) and AlTi3 . The latter increases in volume fraction during milling from 20 to 30 h. Since the profile at 10 h records no characteristic peak of AlTi3 except a broad halo between 2h ¼ 35° to 44°, continued milling of the predominantly amorphous matrix from 10 h onwards seems to induce strain-induced nucleation of AlTi3 from an amorphous solid solution. Similar observation was earlier reported by us as well as others [23,24]. It may be mentioned that no significant change in microstructure is noted beyond 30 h of milling.
Fig. 2. XRD-profiles of Al30 Ti60 Si10 powder blend subjected to different hours of planetary ball milling. Note that nanocrystalline AlTi3 forms from an amorphous matrix on continued milling beyond 10 h.
Fig. 3(a) shows the high-resolution transmission electron microscope (HRTEM) image of Al30 Ti60 Si10 following 20 h of mechanical alloying. It is evident that the microstructure is predominantly nanocrystalline with a small volume fraction of region among the crystallites devoid of lattice fringes. At least two types of interplaner spacings (d) are recorded from the microstructure which correspond well with dð201Þ ¼ 0:2162 nm of AlTi3 [25] and dð101Þ ¼ 0:1881 nm of WC [26]. The corresponding SAD pattern shows diffraction rings characteristic of the AlTi3 crystallites (Fig. 3(b)). The isolated spots are likely due to WC impurity picked up during milling. Similar microstructural analysis at 30 h milled sample reveals identical nanocrystalline aggregate with marginally higher volume fraction and grain size of AlTi3 . Perhaps, continued milling increases the volume fraction and size of AlTi3 crystallites at the expense of amorphous matrix. Fig. 4 reveals the summary of XRD patterns of the Al30 Ti50 Si20 powder blend during high-energy ball milling from 0 to 30 h. The peaks due to the constituent elements undergo reduction in intensity and broadening
Fig. 3. (a) High-resolution transmission electron microscope (HRTEM) image of Al30 Ti60 Si10 following 20 h of mechanical alloying and (b) the corresponding SAD pattern.
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Fig. 4. XRD-profiles of Al30 Ti50 Si20 powder blend subjected to highenergy planetary ball milling for 0–30 h. Note that nanocrystalline Ti5 Si3 forms from an amorphous matrix on continued milling beyond 20 h.
of width until 10 h to finally form a predominantly amorphous aggregate at 20 h of milling. Continued milling up to 30 h appears to induce strain-induced nucleation of nanocrystalline Ti5 Si3 or its extended solid solution. Thus, the microstructural evolution of Al30 Ti50 Si20 during high energy ball milling appears identical as that of Al30 Ti60 Si10 except the identity of the nanocrystalline phase that evolves from the amorphous matrix. This difference could be attributed to the difference in initial blend composition. Fig. 5(a) shows the HRTEM image of Al30 Ti50 Si20 after 24 h of milling. The microstructure reveals a typical amorphous + nanocrystalline composite aggregate. The thinner regions of the powder particles (along the edges) are predominantly amorphous. The crystalline region records a d-spacing that matches well with dð211Þ ¼ 0:2198 nm of Ti5 Si3 [27]. Fig. 5(b) shows the corresponding SAD pattern that records both diffraction rings and spots due to the presence of both nanoand micro-crystallites of Ti5 Si3 respectively. It may be mentioned that the volume fraction of amorphous region is negligible at 30 h of milling. Thus, Al30 Ti60 Si10 seems more prone to amorphization followed by strain-
Fig. 5. (a) HRTEM image of Al30 Ti50 Si20 after mechanical alloying for 24 h and (b) the corresponding SAD pattern.
induced nucleation of nanocrystalline intermetallic phase by continued milling than that of Al30 Ti50 Si20 . Fig. 6 shows the XRD profiles of the Al30 Ti40 Si30 powder blend as a function of milling time during highenergy planetary ball milling. The microstructural evolution seems identical as that in Al30 Ti60 Si10 and Al30 Ti50 Si20 except that the microstructure does not turn predominantly amorphous at an intermediate stage (10– 20 h) of milling. Instead, nanocrystalline Ti5 Si3 (or Ti7 Al5 Si12 ) and Al3 Ti seem to nucleate from the extended solid solution that forms between 5 and 10 h following mutual dissolution of constituent elements by mechanical alloying. It is interesting to note that the Ti5 Si3 (211) peak shrinks in width between 20 and 30 h of milling indicating perhaps an increase in the crystallite size during continued mechanical alloying. Such grain coarsening or crystallite growth may be aided by enhanced mass transport [13] during extended hours of milling when no new phase evolution or grain size reduction is feasible. Fig. 7(a) shows the HRTEM image of Al30 Ti40 Si30 milled for 30 h. The microstructure consists of very fine crystallites (<10–15 nm) embedded in amorphous matrix. Careful observation in other areas reveals similar amorphous + nanocrystalline microstructure with varying degree of relative volume fraction. Fig. 7(b) shows the corresponding SAD pattern confirming that 30 h of mechanical alloying of Al30 Ti40 Si30 produces a nanocrystal dispersed amorphous matrix composite microstructure. In all the three alloys, formation of amorphous or nanocrystalline solid solution seems a prerequisite for strain-induced nucleation of the intermetallic phases during continued milling. Perhaps, grain size reduction or nanocrystallization by mechanical alloying facilitates formation of a nanocrystalline or amorphous extended
Fig. 6. XRD profiles of the Al30 Ti40 Si30 powder blend as a function of milling time. Note that nanocrystalline Ti5 Si3 nucleates at an early stage (10 h) of milling.
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Fig. 7. (a) HRTEM image and (b) SAD pattern of Al30 Ti40 Si30 milled for 30 h.
solid solution in 10–20 h of milling. Subsequently, continued milling enables strain-induced nucleation of a relevant intermetallic phase (Ti5 Si3 , AlTi3 , TiAl3 , etc.) that possesses significantly more negative enthalpy of formation ðDH Þ than the amorphous/nanocrystalline solid solution. The related thermodynamic analysis to predict the phase equilibrium is discussed in the following section. In order to investigate the role of Al in microstructural evolution of the present ternary systems, two more powder blends were subjected to mechanical alloying under identical conditions. Fig. 8 shows the XRD profiles of the Al-lean Al10 Ti60 Si30 powder blend subjected to mechanical alloying up to 30 h. It is not clear whether an amorphous phase forms in this powder blend. However, nanocrystalline Ti5 Si3 and AlTi3 peaks are detected at the final stage of milling (along with WC impurity). On the other hand, XRD profiles in Fig. 9 shows that mechanical alloying of Al-rich Al50 Ti40 Si10 powder blend leads to formation of a two-phase microstructure comprising amorphous and nanocrystalline intermetallic phase (Al3 Ti). Though it is not clear whether solid-state amorphization preceded nucleation of nanocrystalline Al3 Ti or vice versa, the final micro-
Fig. 8. XRD profiles of the Al-lean Al10 Ti60 Si30 powder blend subjected to mechanical alloying up to 30 h.
Fig. 9. XRD profiles of the Al-rich Al50 Ti40 Si10 powder blend subjected to mechanical alloying up to 30 h.
structure (at 30 h) seems predominantly amorphous with the presence of nanocrystalline Al3 Ti in it. Fig. 10(a) shows HRTEM evidence of nanocrystal dispersed amorphous matrix aggregate of Al50 Ti40 Si10 after 30 h of mechanical alloying. The corresponding SAD pattern shows a diffused intensity halo with additional low intensity rings substantiating the XRD and HRTEM evidences of amorphous + nanocrystalline composite microstructure of the present alloy after 30 h of milling. Comparison of microstructures at 20 and 30 h of milling reveals that the volume fraction of the nanocrystalline phase marginally decreases between 20 and 30 h of milling. Perhaps, further microstructural disorder/defect introduced by continued milling would transform the microstructure completely amorphous. It
Fig. 10. (a) High resolution TEM image of the milled product following 30 h of milling of Al50 Ti40 Si10 showing a predominantly amorphous microstructure with a few isolated nanometric (<8 nm) crystalline grains and (b) SAD pattern of the same Al50 Ti40 Si10 sample (from another area) showing amorphous as well as nanocrystalline grains with a few bright spots due to some coarse crystalline grains.
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is worth mentioning that isolated nanocrystalline regions existed in the microstructure even after 50 h of milling. Possibly, milling with higher strain rate, at lower temperature or with a different composition is necessary for complete solid-state amorphization by mechanical alloying. It is predicted that nanocrystalline dispersed amorphous rather than single phase amorphous or nanocrystalline microstructure is more likely to yield greater strength and toughness [1,28]. From the present results, it appears that mechanical alloying is a convenient method of producing in situ nucleated nanocrystalline intermetallic dispersed amorphous matrix composite in Al-rich Al–Ti–Si alloys by appropriate choice of composition and milling parameters. 3.3. Thermodynamic analysis The identity and sequence of phase evolution during mechanical alloying is a complex function of the initial powder blend, thermo-physical properties of the constituents, and employed milling parameters/conditions. In a mechanically driven system like mechanical alloying, the phase equilibrium is based on non-equilibrium thermodynamics involving both equilibrium as well as metastable or non-equilibrium phases including extended and amorphous solid solutions. In this regard, the models by Miedema [29,30] provide a convenient tool to estimate the enthalphy/heat of formation ðDH Þ of different phases in a binary system as follows: 1=3 2 DH ¼ K½P ðD/ Þ2 þ QðDgws Þ R;
ð1Þ
where K is the proportionality constant, Dgws is the difference in electron density at the boundary of Wigner–Seitz (ws) cells and D/ the difference in electronegativity parameters of the concerned elements. Q, P and R are empirical constants, dependent on the type of elements (transition or non-transition). The necessary parameters and their values are documented in the literature [31]. As already pointed out, this approach is valid only for binary systems. Since a similar universally applicable
formulation to calculate DH for ternary or multi-component system is not available, we have calculated DH for the present alloys using a weighted mean approach as follows: DHABC ¼ 0:5ðXAB DHAB þ XBC DHBC þ XAC DHAC Þ:
Here, X and DH represent the mole fraction and enthalpy of formation of the binary or ternary phases comprising elements A, B, C (as shown in the subscripts), respectively. Though XA þ XB þ XC ¼ 1, the relative concentration of each element in any given binary combination for calculating the terms DHAB , DHBC or DHAC needed in Eq. (2), say in Al50 Ti40 Si10 , is taken as XAlTi ¼ 0:9, XAlSi ¼ 0:6 and XTiSi ¼ 0:5, such that, 0.5 ðXAlTi þ XTiSi þ XAlSi Þ ¼ 1. For calculating the respective binary DH terms, say DHAB for A–B, the corresponding relative mole fraction of A in A–B, say XA0 ðin ABÞ ¼ XA =ðXA þ XB Þ. Thus, XA0 in A–B may be different than that in A–C. For instance, XA in A–B and in A–C for Al50 Ti40 Si10 are 0.55 and 0.83, respectively. However, ½XA0 ðin ABÞ þ XA0 ðin ACÞ þ XB0 ðin ABÞ þ XB0 ðin BCÞ þ XC0 ðin ACÞ þ XC0 ðin BCÞ =3 ¼ 1 for the same alloy, i.e., Al50 Ti40 Si10 . These calculations are based on Vegard’s law type assumption that enthalpy is a linear function of mole fraction. Table 1 summarizes the DH values calculated for the present ternary alloys as well as the stable binary intermetallic phases/compounds possible in the concerned systems. It is apparent that the predicted DH values for crystalline and amorphous solid solutions are comparable, while the same for some of the intermetallic compounds are significantly smaller. Hence, it is probable that mutual dissolution during early stage of milling leads to development of crystalline or amorphous solid solution. Enhanced rate of mass transport, significant grain size reduction (nanocrystallization) and high degree of plastic strain could lead to strain-induced nucleation of intermetallic phases with more negative DH from the nanocrystalline or amorphous solid solutions. However, direct nucleation of the intermetallic phase/compound may concur or succeed the formation of the nanocrystalline/amorphous solid solution de-
Table 1 Enthalpy of formation ðDH Þ of the possible phases in the Al–Ti–Si system Possible solid solution or intermetallic phases Al50 Ti40 Si10 Al30 Ti60 Si10 Al30 Ti50 Si20 Al30 Ti40 Si30 Al10 Ti60 Si30 AlTi3 TiAl3 TiAl Ti5 Si3
ð2Þ
Enthalpy of formation ðDH Þ in kJ/mol Crystalline solid solution
Amorphous solid solution
Intermetallic phase/compound
)27.62 )26.09 )32.73 )34.84 )30.83
)28.89 )24.46 )36.04 )39.98 )31.55
– – – – – )38.04 )39.43 )60.44 )78.45
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3.4. Mechanism of amorphization by NMR investigation In the absence of localized moments, the position of the resonance line in a metallic sample is determined by, (a) the Knight shift, K, originating from the Fermi contact interaction ðHhf ¼ 2ð8p=3ÞlB cN hI:SðrÞdðrÞÞ between the nuclei and the conduction electrons, (b) the dipolar interaction between nuclear and electronic spin, and (c) the interaction between nuclear spin with the orbital motion of the electrons related to van Vleck paramagnetism [32,33]. The first interaction produces an isotropic shift, the second interaction is anisotropic in nature, and the contribution of the third interaction is negligible in case of Al resonance. The width and shape of the resonance line are determined by the nuclear dipolar interaction, anisotropic Knight shift in case of non-spherical character of the conduction band, and the broadening produced by the bulk magnetic susceptibility of the sample. Additionally, for nuclei with spin I > 1=2 being studied in a compound or an alloy, the interaction between the nuclear electric quadrupole moment and the electric field gradient present around the nucleus, also causes a broadening of the resonance line in polycrystalline samples. The Knight shift can be written, using the definition of the hyperfine field, Hhf , as K ¼ ðHhf =NA lB Þv, where v is the molar magnetic susceptibility and NA is Avogadro’s number. For simple metals and alloys the magnetic susceptibility is dominated by the Pauli paramagnetism, as discussed in Section 3.1 of magnetic susceptibility; for which v ¼ 2NA l2B N ðEF Þ, where N ðEF Þ is the density of state at the Fermi level. So for the Knight shift in the case of simple metal, K ¼ ð1=2Þðce =cN ÞN ðEF ÞA, where A is the hyperfine coupling constant. Fig. 11 shows the 27 Al NMR spectra of Al50 Ti40 Si10 at various intervals of milling. It is evident that the peak due to pure Al (at 78.258 MHz) is distinct at the initial stage of milling (5 h), but reduces in intensity with increased broadening of the linewidth until 10 h. Upon further milling, the resonance component from pure Al is greatly reduced in intensity and becomes very weak after 15 h. However, a trace of pure Al is obtained in NMR spectrum until 20 h of milling. The concerned XRD-pattern of this alloy at similar stage of milling (Fig. 9) substantiate that elemental Al undergoes mutual dissolution to form amorphous and another crystalline phase beyond 10 h of milling.
Al50Ti40Si10 alloy 5h 10 h 15 h 20 h 30 h
Intensity (arb. unit)
pending on the kinetic factors and nucleation barrier. Thus, the present thermodynamic calculations based on simplified weighted mean DH approach predict that an amorphous + nanocrystalline intermetallic phase aggregate or composite microstructure is anticipated in mechanical alloying of the present alloys, which corroborates the experimental results obtained by XRD and TEM analysis.
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30 h 20 h 15 h
10 h
5h
78.1
78.2
78.3
78.4
78.5
Frequency (MHz) Fig. 11. 27 Al NMR at 300 K in Al50 Ti40 Si10 alloys at different hours of milling. The vertical line on x-axis indicates the reference position of Al2 O3 .
In addition to the pure Al line as mentioned above, the sample after 10 h of milling shows a resonance line of broad structure at a lower frequency compared to the resonance line of pure Al metal. Two peaks can be distinguished in the broad structure, indicating the occurrence of at least two different alloy products, both having Al as a component, at this stage of milling. The resonance line extends to frequencies higher than that covered by the resonance in pure Al. This may be an indication of Al nuclei in one of the intermediate products of milling, experiencing an environment of non-cubic electric field gradient, and producing a quadrupolar interaction. This large broadening may also arise because of strain-induced distortion around the aluminum nucleus as a result of milling. It is to be noted that the X-ray diffraction pattern (Fig. 9) in this stage showed the constituent elements, though with huge broadening of Bragg peaks, but could not indicate the presence of any alloy product. Upon further milling, the resonance component from pure Al is greatly reduced in intensity, and the broad structure tend to become symmetric, now having only one peak, positioned close to the reference frequency. It indicates that of the many alloy components which were formed at intermediate stages of milling, only one or a few are retained. In the sample milled for 30 h, 27 Al resonance shows only a broad asymmetric line positioned near the reference and gives no trace of pure Al. A possible explanation for the broad NMR spectrum at the intermediate stage and its subsequent narrowing is that an amorphous extended solid solution phase is formed at an early stage of milling. Continued milling leads to the strain-induced nucleation of the stable coordination or environments of the constituent elements which are retained in the form of either nanocrystalline or amorphous intermetallic compounds, as products until the final stages of alloying.
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Fig. 12 shows the NMR spectra of Al50 Ti40 Si10 and Al10 Ti60 Si30 at the final milling stage (30 h). The two alloys shown in this figure represent the composition limits of the alloys investigated in this study in terms of Al-content. More precisely, Al50 Ti40 Si10 and Al10 Ti60 Si30 represent the Al-rich and Al-lean composition limits among the present alloys. It may be noted that the signal intensity, in case of Al-lean sample, is much smaller than that in case of Al-rich sample, and in former case, the data acquisition time was four times longer than in the latter. The spectrum obtained in Al50 Ti40 Si10 can be fitted as a sum of two components of resonance lines, both having a Lorentzian distribution, as shown in Fig. 12. The larger or dominant component has a linewidth (full width at half maximum, FWHM) of about 45 kHz and isotropic shift of 0.03%. The smaller component has a similar linewidth, but a slightly larger shift of about 0.045%. The alloy therefore consists of at least two components, which could be very similar to each other in their structure and properties. More significantly, the large linewidth, which is almost three times the width of pure Al resonance (Fig. 11), is definitely not only due to Al–Al nuclear dipolar interaction. The magnetic interactions in these samples should be negligible as is evident from the magnetic susceptibility data and do not contribute to the broadening of the resonance line. On the other hand, the grain size reduction and amorphization result in a large variation in the atomic environment around Al atoms in the alloy products, leading to a distribution in the local field around the Al nuclei, and giving rise to the increased linewidth. Also shown in Fig. 12 is the spectrum of Al10 Ti60 Si30 fitted, in this case too, as a sum of two Lorentzian components. One of the components has FWHM of Al50Ti40Si10 Al10Ti60Si30 Sintered Al50Ti40Si10 Sintered Al10Ti60Si30
Al50Ti40Si10 mechanical alloyed
Al30Ti60Si10 Al30Ti50Si20 Al30Ti40Si30
Al50Ti40Si10 sintered
Intensity (arb. unit)
Intensity (arb.unit)
50 kHz and shift 0.04%, which are similar to one described above. The other component shows a width of 36 kHz and negligible shift of 0.008%. The results therefore indicate the formation of two very different Al containing components in the alloy. Moreover, as discussed in the beginning of this section, the components might have different metallicity. This suggests that Allean samples might have a higher probability of yielding an inhomogeneous mixture in the alloy products. Fig. 13 presents a comparison of NMR spectra of the remaining alloys with same (30 at.%) Al-content at the final milling stage (30 h). The peak due to pure Al is consistently absent. The spectral features are similar to those of Al50 Ti40 Si10 at comparable stage (Fig. 11), though the spectra for Al30 Ti50 Si20 and Al30 Ti40 Si30 are distinctly bell-shaped unlike that of Al30 Ti60 Si10 . In the latter, spectrum seems to have two well-defined alloy components each having some distribution of the environment around Al resulting from amorphization. Suitable comparison of concerned XRD (Fig. 2) and HRTEM (Fig. 3) result reveals that Al30 Ti60 Si10 undergoes solid-state amorphization like Al50 Ti40 Si10 under comparable conditions of milling. Thus, the NMR analysis suggest that solid state amorphization in the present alloys proceeds by developing a new atomic environment or coordination (there may develop more than one Al alloy component as in Al30 Ti60 Si10 ) with the constituent elements during mechanical alloying that can be detected by the NMR spectrum. The bell-shaped spectra, on the other hand, indicate the presence of a single Al-based alloy component having a wider distribution in the surroundings of Al. This might occur due to high silicon content in the systems. This is quite evident from relevant XRD patterns of Al30 Ti50 Si20 (Fig. 4) and Al30 Ti40 Si30 (Fig. 6), respectively, wherein the lines belonging to AlTi3 are
Amorphous
Al10Ti60Si30 mechanical alloyed
Al10Ti60Si30 sintered
78.05
78.10
78.15
78.20
78.25
78.30
Frequency (MHz) Fig. 12. 27 Al NMR at 300 K in the Al–Ti–Si system after 30 h milling for Al lean (Al ¼ 10%) and Al rich (Al ¼ 50%) samples, before and after sintering. The continuous lines are theoretical fit to the spectra. The broken lines in both the spectra indicate the two constituent components used in the fit as described in the text. The vertical line indicates the reference position of Al2 O3 .
78.0
78.1
78.2 Frequency (MHz)
78.3
78.4
Fig. 13. 27 Al NMR at 300 K in the Al–Ti–Si systems after 30 h of milling with 30% Al in all three samples. The vertical line indicates the reference position of Al2 O3 .
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either much broader or totally absent as compared to that for Al30 Ti60 Si10 shown in Fig. 2. 3.5. Effect of sintering on phase equilibrium In order to identify the amorphous phases, Al50 Ti40 Si10 and Al10 Ti60 Si30 milled for 30 h, were heated at 20 °C/min to 700 °C, kept at that temperature for 1 h, and then rapidly cooled to room temperature and studied by XRD and NMR. The XRD pattern shown in Fig. 14(a), for Al50 Ti40 Si10 , exhibits reasonably welldefined diffraction peaks, characteristic of crystalline/ nanocrystalline phases, belonging to the intermetallics Al3 Ti and Ti5 Si3 , which, therefore, are the main components of the alloy. Fig. 14(b) corresponds to the Allean sample, and shows the formation of the compounds AlTi3 , as expected, and also Ti5 Si3 . The 27 Al NMR spectra for these two sintered samples are shown in Fig. 12 and can be compared with the spectra before sintering. The figure also shows the theoretical fit to the spectra. The spectrum for sintered Al50 Ti40 Si10 can be fitted to one Lorentzian distribution with a FWHM of 18 kHz and isotropic shift of 0.025%. Clearly, the resonance is obtained from crystalline Al3 Ti, as observed in X-ray. Sintering results in a more uniform local environment around Al nuclei compared to the as-milled amorphous samples and the linewidth is therefore reduced. The isotropic shift of the amorphous sample is larger than that of sintered sample and this shows a larger metallicity of amorphous sample. However, the amorphization-induced anisotropic broadening may also result in a shift in the mean position of the spectrum. The spectrum for sintered Al10 Ti60 Si30 can only be fitted as a sum of two components, as in case of asmilled sample. The dominant component corresponds to AlTi3 , and the linewidth of AlTi3 that was 36 kHz in as-
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milled amorphous sample becomes 18 kHz after sintering. The isotropic shift is also reduced from 0.008% to 0.001%. The other component that exists in a very small amount in the NMR spectrum could not be separately identified in XRD pattern. We believe that this unidentified alloy component exists in the amorphous sample as a metastable phase and tends to disappear upon sintering.
4. Summary and conclusion From the present study, it appears that development of nanocrystalline intermetallic dispersed Al-rich amorphous matrix composite is feasible by mechanical alloying of appropriate Al–Ti–Si powder blend. The intermetallic phases/compounds formed include TiAl3 , Ti3 Al, Ti5 Si3 and Ti7 Al5 Si12 depending on the initial composition selected. During mechanical alloying, the phase evolution sequence comprises mutual dissolution or alloying, nanocrystallization, partial/complete amorphization and finally in situ nucleation of nanocrystalline intermetallic phases from the amorphous precursor during continued milling. Enthalpy calculation based on modified Miedema model corroborates the nature and sequence of phase evolution mentioned above. NMR studies also substantiate that mechanical alloying of the present elemental blends leads to alloying, nanocrystallization, amorphization, followed by strain-induced nucleation of nanocrystalline intermetallic phases/compounds to develop a composite microstructure. Furthermore, the NMR results identify characteristic new environment of atomic coordination around Al atoms for both amorphous and nanocrystalline intermetallic phases. The spectra located at lower frequencies compared to that of pure Al metal indicating that the density of electrons at the Fermi level at Al site is reduced as a result of formation of these phases, giving rise to smaller values of Knight shift. Thus, NMR results as well as magnetic susceptibility measurements indicate that metallic properties are reduced following amorphization and strain-induced nucleation of nanocrystalline intermetallics. Further studies are in progress to consolidate the milled product and assess mechanical properties of the composite.
Acknowledgements
Fig. 14. XRD patterns of 30 h milled (a) Al50 Ti40 Si10 and (b) Al10 Ti60 Si30 quenched to room temperature after heating up to 700 °C.
The authors wish to thank Professors H.J. Fecht, F. Banhart and P.M.G. Nambissan for useful technical discussion on this work. Partial financial support from the Humbold Foundation, Bonn and Council of Scientific and Industrial Research, New Delhi (Grant No. 70 (0048) 03-EMRII) to one of the authors (I.M.) is gratefully acknowledged.
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