Effect of zirconium on thermal stability of nanocrystalline aluminium alloy prepared by mechanical alloying

Effect of zirconium on thermal stability of nanocrystalline aluminium alloy prepared by mechanical alloying

Journal of Alloys and Compounds 688 (2016) 571e580 Contents lists available at ScienceDirect Journal of Alloys and Compounds journal homepage: http:...

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Journal of Alloys and Compounds 688 (2016) 571e580

Contents lists available at ScienceDirect

Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom

Effect of zirconium on thermal stability of nanocrystalline aluminium alloy prepared by mechanical alloying V.M. Suntharavel Muthaiah, Suhrit Mula* Department of Metallurgical and Materials Engineering, Indian Institute of Technology Roorkee, Roorkee 247667, Uttarakhand, India

a r t i c l e i n f o

a b s t r a c t

Article history: Received 18 February 2016 Received in revised form 29 June 2016 Accepted 4 July 2016 Available online 7 July 2016

In the present study, a series of Al-x% Zr (x ¼ 1e10 at.%) compositions were prepared by mechanical alloying (MA) to investigate the solid solubility extension of Zr in Al and its thermal stability. The elemental powder blends were mechanically alloyed under high purity argon atmosphere in stainless steel grinding media using SPEX 8000 M high energy ball mill. The milling was carried out for 8 h at room temperature and the ball to powder ratio was maintained at 10:1. Formation of disordered solid solutions is validated using Miedema’s semi-empirical model. X-ray diffraction (XRD) analysis confirms the formation of disordered solid solution up to 1at.% Zr; whereas, Al3Zr and Al9.83Zr0.17 intermetallic phases were found to form as per the XRD pattern of 2e10% Zr alloys. Variation of lattice parameter confirmed the formation of Al-1% Zr solid solution. Crystallite size was estimated to be 41 and 30 nm, respectively, for the as-milled 1 and 10% Zr alloys. The matrix grains were found to be stabilized after annealing at 550  C, and the XRD crystallite size of Al-10% Zr retained at ~58 nm. TEM and AFM analysis confirmed that the grain size of the as-milled as well as annealed samples was retained in the nanometre range (<100 nm), which corroborates well with the XRD crystallite size. Vickers microhardness of the asmilled 10% Zr alloy was found to be 2.4 GPa, which decreased to 1.7 GPa after annealing at 550  C. This demonstrated that the dissolution of Zr in Al has a larger strengthening effect and Zr plausibly played a pivotal role in retaining the matrix grains size in the nanoscale range at high temperature. © 2016 Elsevier B.V. All rights reserved.

Keywords: Nanocrystalline Al-Zr alloys Mechanical alloying Thermal stability Miedema’s semi-empirical model Grain growth

1. Introduction Aluminium (Al)-based alloys have a wide range of industrial applications, namely, aerospace, automobile, electronics, domestic construction and power distribution due to their high strength to weight ratio, excellent corrosion resistance, formability and thermal conductivity. In the recent years, nanocrystalline Al-based alloys have received a great attention due to their unique mechanical properties such as high specific yield strength and appreciable formability and good corrosion resistance [1]. However, the nanostructured Al alloys are often considered unsuitable for high temperature engineering applications because of instability of the grain size. Reduction of strength at elevated temperature is accounted to be the major reason behind this incompatibility [2,3]. Thus, a mechanism to strengthen the nanocrystalline Al-based alloys to make them applicable at relatively higher temperature appears to be very important.

* Corresponding author. E-mail address: [email protected] (S. Mula). http://dx.doi.org/10.1016/j.jallcom.2016.07.038 0925-8388/© 2016 Elsevier B.V. All rights reserved.

In pure nanocrystalline metals, excess grain boundary free energy provides a huge driving force for grain growth at relatively low temperatures [4,5]. The grain growth leads to coarsening of the nanocrystalline grains and resulted in a poor mechanical strength. The grain growth becomes very serious for nanocrystalline alloys of low melting materials like Al and its alloys [6]. It is proposed that the grain size can be retained within sub-micron level by the formation of fine and homogeneous intermetallic compounds [7,8], such as in Al-Co [9] and Al-Ru [10] alloys. Various studies have been carried out to achieve grain size stabilization by the addition of small amounts of thermodynamically insoluble solute elements (oversize solute atoms) such as W (in Al) [7], which acts as a precipitation or dispersion hardening element and effectively stabilizes the grain size at high temperature [11]. Two basic approaches are used to suppress the grain growth of the nanocrystalline materials at relatively high temperatures [12]. Kinetic stabilization mechanism involves restricting the grain boundary movement by solute drag, chemical ordering and precipitation of second phase particle(s) (Zener pinning) as reported in Fe-Zr [13], Cu-Zr [14] and Pd-Zr [15] systems. The other approach to

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retain nanocrystalline grain size at high temperature is the thermodynamic stabilization. It is accomplished by reducing the excess grain boundary energy of metastable (e.g. nanocrystalline) material almost to zero by solute segregation at the grain boundary [16]. Thus, a metastable equilibrium can be achieved and no driving force would be available for further grain coarsening [16]. Zr is found to act as a good stabilizing element at high temperature as reported for Fe-Zr [13] and Pd-Zr alloys [15]. The equilibrium solubility of Zr in Al is very limited and it is only 0.083 at.%1 at room temperature. On the other hand, nonequilibrium processing methods like mechanical alloying (MA) [17], inert gas condensation (IGC) [18] and rapid solidification processing (RSP) [19] can be used to bring thermodynamically insoluble alloying elements into the lattice of solvent up to some extent to produce disordered solid solutions [20]. Rittner et al. [18] reported a high hardness of 3 GPa of Al-Zr alloy produced through inert gas condensation technique. Al-Zr-Fe ultrafine grains (grain size ~260 nm) synthesised by physical vapour deposition technique showed a high tensile strength of 800 MPa [21]. Therefore, it is possible to produce high strength Al-Zr alloys, if the metastable nanocrystalline Al grains are stabilized by the thermodynamic and/ or kinetic mechanisms. The MA is preferred among the nonequilibrium processes, because it is very effective to produce highly supersaturated solid solutions easily. The MA is also reported as a feasible solid state processing route for synthesis of large quantities of nanostructured materials [17]. Thus, it is reasonable to develop Al-Zr nanocrystalline alloys by MA and study their thermal stability. The main aim of the study is to investigate the feasibility of formation of the Al-Zr disordered solid solutions by MA and effectiveness of Zr additions on the thermal stability of such disordered alloys. 2. Experimental Elemental powders blends of aluminium (Alfa Aesar 325 mesh, 99.9% purity) and zirconium (Hi-Media, 200 mesh, 99.7% purity) were mechanically alloyed in 440 stainless steel grinding media using SPEX 8000 M high energy ball mill. The grinding media consists of 8 mm (17 in numbers) and 6 mm (16 in numbers) diameter balls, and a ball-to-powder weight ratio of 10:1 was maintained throughout the milling time. The vial was sealed in high purity argon atmosphere (purity <10 ppm O2) prior to milling and the milling was carried out for 8 h at room temperature. The asmilled powder samples were compacted as disks specimens in a tungsten carbide die-punch using a pressure of 300 MPa by a hydraulic press. Then, the disk samples were annealed in batches at different temperatures from 150 to 550  C for 1 h under (Ar þ 2% H2) atmosphere. X-ray diffraction (XRD) study was performed using a Cu Ka radiation (l ¼ 0.154 nm) at a scan rate of 1 /min in a Rigaku X-ray diffractometer. The average crystallite size was calculated from the broadening of 3 major peaks after eliminating the broadening effects due to strain by using the plot between BrCosq vs. Sinq [22]. The precise lattice parameter of the Al-based alloys (aAl) was calculated from 3 major XRD peaks by the extrapolation of aAl vs. (cos2q/sin2q) plot to cosq ¼ 0 [23]. The samples were polished to achieve mirror finished surface and Buehler Microhardness (Model: UHL Technische Mikroskopie VMHT) tester was employed to perform Vickers microhardness measurements. The microhardness test was carried out using 50 g load at a speed of 15 mm per second with a dwell time of 15 s for each indentation. Atomic force microscopy (AFM) analysis was performed on the disks samples using a silicon nitrate probe in VECCO di Innova

1

Hereafter, at.% will be written as %, unless it is indicated.

atomic force microscope (Model: TS-150). Transmission electron microscopy (TEM) analysis was carried out for some selected asmilled and annealed samples using a JEOL 2000FX electron microscope at a beam energy of 200 keV. Sample preparation for TEM analysis was carried out by drop cast technique using carbon coated copper (Cu) grid. 3. Results and discussion 3.1. Effect of Zr concentration Fig. 1a shows the XRD patterns of the as-milled Al-Zr alloys. It can be noticed that the Zr peak is not detectable from the XRD patterns of any compositions. The XRD phase analysis also reveals that plausibly a disordered solid solution is formed up to 1% of Zr alloy. This is so as only Al peaks are detected from the relevant XRD pattern (if any dilute quantity of intermetallic phase(s) formed in 1% Zr alloy, it is not possible to detect by XRD technique due to its limited detectability for the second phase. Beyond 1% of Zr, probability of formation of intermetallic phases such as Al3Zr and Al9.83Zr0.17 increased with increase in Zr concentration (2e10% Zr alloys). This clearly indicates that the solid solubility of Zr in Al plausibly could not be enhanced beyond 1 at.% by MA due to the formation of the intermetallic phases (as detected from the XRD patterns of 2e10% Zr alloys). It can be noticed, especially from the high angle peaks, that the peak intensity gradually decreased and peak width at FWHM increased with the increase in the Zr content. This is due to the combined effect of grain size refinement (finer crystallite size) and increase in the residual strain [24]. As the sensitivity of XRD to determine the presence of secondary phases has a limit of 1e2 wt%, the limit of formation of disordered solid solution has been analyzed by precise lattice parameter measurement of Al-based solid solutions. The magnified view of (111) peak in Fig. 1b indicates that there is a peak shift towards lower 2q position with respect to the standard peak position of pure Al (111). It is to be noted that the peak shift is very small and it can be detected only for the 1% Zr alloy. The peak shift for the alloys containing more than 1% Zr is not detectable as per the XRD analysis. It should be remembered that the peak shift towards the lower diffraction angles is an indication in the increase in the lattice parameter [23], which is possible only due to the formation of solid solution, i.e. Zr atoms dissolved in Al lattices. The variation of lattice parameter (calculated from the XRD analysis) of the Al-rich solid solution (aAl) can be a very useful tool to predict the solid solubility of a solute [14,23,25]. Fig. 2 shows the variation in the aAl for all the compositions. The aAl was calculated from 3 major XRD peaks to reveal the dissolution of Zr in Al during MA. The refined values of aAl were calculated from the plot of aAl vs. (cos2q/sin q) after extrapolation of cosq ¼ 0 [22]. It is observed that initially the aAl increased slightly from 4.0495 Å (pure Al) to 4.0592 Å corresponding to Al-1% Zr alloy. After that there is almost no significant change in the aAl even after addition of more Zr (Fig. 2), which indicates no further dissolution of Zr in Al occurred in the same conditions of MA. The initial increase in the aAl indicates the dissolution of larger in size Zr atoms (atomic radius ¼ 0.158 nm) into the lattice of Al (atomic radius ¼ 0.143 nm) [26]. And thereafter, the added Zr mostly converted to intermetallic phases as indicated in the XRD plots (Fig. 1a). The temperature rise during milling possibly played an important role leading to the formation of intermetallic compounds, especially in the high stacking fault energy Al alloys [17]. It is well known that the alloying elements which are forming intermetallic phases do not play any role in the variation of lattice parameter of matrix phase; only the alloying content which dissolves in the lattice of solvent can change the lattice parameter of solvent [27].

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Fig. 1. (a) XRD patterns of 8 h ball-milled samples of Al containing 1e10 at.% Zr, (b) Close examination of Al (111) major peak.

Fig. 3 shows the variation of crystallite size of the 8 h milled samples of Al-Zr alloys. The average crystallite size was calculated from 3 major XRD peaks using Williamson and Hall analysis [22]. The broadening effect due to the strain has been eliminated by using the plot between BrCosq vs. Sinq. The peak broadening was evaluated by measuring the full width at half maximum (FWHM) intensity. The crystallite size estimated to be ~97 nm for the asmilled pure Al. The crystallite size decreased with increase in the Zr concentration. The average crystallite size decreased from 97 nm, which corresponds to pure Al, to 41 nm for the as-milled 1% Zr alloy. Thereafter, the decrease in the crystallite size was found to be limited with increase in the Zr content, and the lowest value of the crystallite size was found to be ~30 nm which corresponds to the as-milled Al-10% Zr sample. The gradual decrease in the crystallite size is attributed to the profound plastic deformation,

Fig. 2. Variation of lattice parameters of as-milled Al-Zr samples with increasing Zr content.

Mula et al. [25] reported that the aAl in Al-Mn-Ce alloy initially decreased due to dissolution of smaller in size Mn (0.126 nm) [26] atoms into the solid solution of Al (0.143 nm) [26] during the early stage of MA; whereas, the increase in aAl beyond 15 h of MA was attributed to the delayed dissolution of larger size Ce (0.182 nm) atoms [26] in Al. Azimi et al. [14] also reported that the lattice parameter of Cu-3% Zr alloy initially increased up to 50 h of milling, and then started to decrease due to recovery and recrystallization because of increase in the temperature for long time milling (>50 h). The supersaturated dissolved Zr atoms came out of the solution due to recrystallization leading to decrease in the lattice parameter of Cu-alloy for higher milling time [14]. The formation of Al-Zr solid solutions by MA is validated by Miedema’s semiempirical model discussed later in this section.

Fig. 3. Variation of crystallite sizes & lattice micro strains of as-milled Al-Zr samples.

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repeated fracturing and cold welding of the powder particles during MA [17]. Fig. 3 also shows the changes in lattice microstrain (%) of the as-milled alloys as a function of Zr content. The lattice microstrain was found to increase from 0.156% (for pure Al) to 0.218% for Al-1% Zr. This is further enhanced to 0.3442% corresponding to the as-milled sample of Al-10% Zr. The work hardening rate increases due to increase in the dislocation density and leads to the formation of fine grains during MA [24]. It can be noticed that there is no significant change observed in the crystallite size and lattice microstrain of the alloys containing beyond 1% Zr. This clearly indicates that steady state (saturation) level was achieved for 2 or more at.% of Zr alloys. Fig. 4a and b shows AFM micrographs of the 8 h milled samples of Al-1% Zr and Al-10% Zr compositions, respectively. Presence of uniformly distributed grains clearly indicates that the nanocrystalline structures were developed by MA for both the compositions. The average grain size obtained from AFM analysis was found to be ~59 nm for Al-1% Zr and 49 nm for Al-10% Zr alloys, respectively. TEM analysis of the ball milled sample of Al-10% Zr was carried out to evaluate the grain size and crystallinity of the microstructure. Fig. 5a shows a representative dark field TEM image of the same specimen. The dark field TEM image shows the presence of nanocrystalline grains (size <100 nm). The grain size (43 nm) obtained from the TEM analysis is consistent with the grain size calculated from the AFM analysis (49 nm) and XRD crystallite size (~30 nm) as discussed earlier. Fig. 5b shows selected area electron diffraction (SAED) pattern of the same sample recorded from the encircled region of Fig. 5a. The spotty continuous rings of Al (as indicated) in the SAED pattern confirms that the alloy grain size is in the nanometer range; while the faint rings, as indicated in Fig. 5b, are detected to be from Al3Zr (101), Al3Zr (114), Al3Zr (118) and Al3Zr (200) confirming the presence of Al3Zr in the ball milled sample. Besides the rings of Al and Al3Zr, some partially revealed rings/spots are also visible in the SAED pattern. And these rings/ spots are due to the presence of fine grains of the other Al-Zr intermetallic compounds [28,29]. It should be noted that the formation of Al3Zr intermetallic compounds were also confirmed from the XRD analysis. Therefore, it can be concluded that the added 10% Zr was partially dissolved in Al and remaining amount formed

various Al-Zr intermetallic compounds after 8 h of MA.

3.2. Thermodynamic analysis The Gibbs free energy changes required for the formation of a metastable supersaturated solid solution of Zr in Al was calculated as per the Miedema’s semi-empirical model using the parameters given in Table 1. The change in Gibbs free energy DGm can be obtained from the following equation as follows:

DGm ¼ DHm  T DSm

(1)

where, DHm and DSm are the enthalpy and entropy of mixing for the formation of disordered solid solution from a mixture of pure solvent and solute atoms, T is the temperature at which the solid solution forms. According to Miedema’s semi-empirical model, the enthalpy of formation of the solid solution can be calculated from the following relation [30].

DHm ¼ DHelastic þ DHchemical þ DHstructural

(2)

where, DHchemical is the chemical contribution of engenderment and breaking of atomic bonds, DHstructural is the contribution due to structural changes that occur due to crystal structure of solvent and solute atoms and difference in valence electron. DHelastic is the elastic contribution atomic size mismatch.

  2=3 2=3  2 S  2pf CS ðXA VA þXB VB Þ Q DHchemical ¼  1=3  1=3 ðDfÞ2 þ n1=3  ws P P þ nBws nAws (3) where, 4 is the work function, V is the molar volume and nws is the electron density of the constituents. P, Q and S are the empirical constants related to constituent elements. f(Cs) is the concentration function, which can be calculated for solid solutions as follows:

. P ¼ 14:1 kJ V2 cm; S=P ¼ 0; Q =P ¼ 9:4:

Fig. 4. AFM micrographs of 8 h milled samples (a) 1% Zr and (b) 10% Zr.

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Fig. 5. TEM analysis of the as-milled sample of Al-10% Zr composition: (a) Dark field (b) SAED pattern taken from the region shown in (a).

Table 1 Required parameters for thermodynamic analysis according to Miedema’s semiempirical model in Al-Zr system [30,31]. Parameters

Al

Zr

1 n1/3 ) ws (cm Vm (cm3/mol)

1.61 10.0 4.2 7.52 2.62

1.39 14.02 4.12 9.11 3.3

4*(V)

K (*1010 N/m2) G (*1010 N/m2)

2=3

CAS ¼

xA VA 2=3

xA VA

2=3

; CBS ¼ 2=3

þ xB V B

xB VB 2=3

xA VA

2=3

þ xB VB

(4)

DHstructural is a very small value as it is related to number of valance electrons per atom, which can be neglected in the estimation of DHm. The elastic part of enthalpy can be expressed as [30].

DHelastic ¼ XA XB ðXA DEAinB þ XB DEBinA Þ DEAinB ¼

2KA GB ðDVÞ2 3KA VB þ 4GB VA

and

DEBinA ¼

(5) 2KB GA ðDVÞ2 3KB VA þ 4GA VB

Fig. 6 shows the Gibbs free energy, enthalpies and entropy change, which are required for the formation of disordered Al-Zr solid solution, as a function of XZr at 298 K. Here, the elastic and chemical enthalpy dominantly contributes for the formation of AlZr solid solution. The elastic contribution of enthalpy, DHelastic is primarily due to atomic size difference between Al (0.143 nm) solvent and Zr (0.158 nm) solute atoms. On the other hand, the chemical contribution of enthalpy, DHchemical is due to the large difference in the electron density and comparatively minute difference in the work function between the A and B elements. Total Gibb’s free energy change required (as shown in Fig. 6) for the formation of solid solution spontaneously for all the proposed alloys is found to be positive as per Miedema’s model. Therefore, formation of Al-Zr solid solution is never possible by any equilibrium methods. The extension of solid solubility (increase in solid solubility) can only be achieved for the Al-Zr system (which exhibits positive heat of mixing) using a non-equilibrium MA technique. The Gibbs free energy change to form solid solution in Al-Zr system can be achieved by the excess energy stored in the milled product in 2 ways: (i) in the form of surface energy due to crystallite size reduction and (ii) lattice microstrain energy due to increase in

(6)

where, DEAinB and DEBinA are the change in elastic energy due to the element A dissolved in element B and the element B dissolved in element A. K and G are the bulk and shear modulus. The configurational entropy of mixing, DS can be calculated from the following equation:

DSm ¼ RðXA InXA þ XB InXB Þ

(7)

R (¼ 8.314 J/mol K) is the universal gas constant, XA and XB denotes the mole fraction of element A and B, respectively.

DHamorphous ¼ DHChemical þ aTfuse

(8)

where, a ¼ 3.5 Jmol1 K1 and T fuse is defined as: A B T fuse ¼ ðXA Tm þ XB Tm Þ

(9)

where, Tm is the melting temperatures of the A and B elements.

Fig. 6. Gibbs free energy change, enthalpies and entropy for the formation of solid solution in Al-Zr system at room temperature.

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the dislocations density. The Gibbs free energy change due to reduction in crystallite size, DGb, can be calculated from the following eq. (10) [32,33]:

DGb ¼ g

 A Vm V

(10)

where, g ¼ 324 mJ/m2 is the grain boundary free energy for Al [32], (A/V) is the surface area to volume ratio of the grain and Vm is the molar volume. Fig. 7 shows experimentally obtained values of DGb as a function of Zr concentration for the investigated compositions. The grain size was considered to be spherical for the mechanically alloyed samples [22]. DGb was found to increase with decrease in crystallite size of Al-based alloys. It is known that the elastic strain energy of the metals increases with the increase in the dislocation density [34]. The Gibbs free energy change in the presence of dislocations can be calculated from the following eqs. (11) and (12):

DGS ¼ xrVm x¼



(11)



DGb2 re  In 4p b

(12)

where, r is the dislocation density, Vm is the molar volume and x is the elastic strain energy per unit length of dislocation. G is the shear modulus, b is the burgers vector and re is the outer cut off radius. In nanocrystalline materials re can be taken as the crystallite size of the matrix [35]. r for the mechanically alloyed sample can be estimated from the following eq. (13) [36,37]:

pffiffiffi 2 3  2 1=2 ε r¼ Db

(13)

where, D is the average crystallite size and ε is the lattice microstrain, which can be measured from the XRD data. Fig. 8 shows the increase in the experimentally estimated Gibbs free energy change, DGs, due to the increase in the dislocation density. It can be observed from the Fig. 8 that DGb due to the reduction in the crystallite size is very large compared to that of DGs. The experimentally obtained total Gibbs free energy change, DGT, due to reduction in the crystallite size (DGb) and due to increase in the dislocations density (DGs) is summarized and compared in Table 2 with that of the theoretically calculated values (DGMiedema) using Miedema’s model.

Fig. 7. Gibbs free energy change, DGb, due to the reduction in crystallite size as a function Zr concentration.

Fig. 8. Variation Gibbs free energy change due to increase in dislocation density, DGs, as a function of Zr concentration.

It can be noticed (Table 2) that the experimentally calculated stored energy, DGT (¼DGb þ DGs) was found to be bit higher (0.48 kJ/mol) only for the as-milled Al-1% Zr alloy than that theoretically required energy (0.473 kJ/mol) for the formation of disordered solid solution. Therefore, again it confirms that only 1 at.% of Zr can be dissolved in Al under the present experimental conditions. Thermodynamic calculation shows that total experimental Gibbs free energy change did not reach the theoretically required label for the formation of disordered solid solution for 2e10% Zr alloys. This is so as the alloying content forming intermetallic compounds does not decrease the crystallite size of the solid solution. The disorderness (entropy) of powder materials increases as more heat is generated (enthalpy) during milling. It dominates to cause lattice disorder of Al alloy which can be designated as reversibly ordered intermetallic. Hence, the reversibly ordered structure tends to form intermetallic compound [38].

3.3. Thermal stability and mechanical properties The as-milled samples were annealed in batches at 150, 250, 350, 450 and 550  C for 1 h under high purity (Arþ2% H2) atmosphere and then XRD data were recorded for all the samples. The annealed samples were examined for microhardness values, and these are correlated with the crystallite size and phase evolution of the annealed samples to evaluate their thermal stability. Fig. 9a and b, respectively, show the XRD patterns along with phase analysis of the annealed 1 and 10% Zr alloys. The phase analysis revealed that the formation of intermetallic phases such as Al3Zr and Al9.83Zr0.17 was increased with increase in the annealing temperature. The XRD analysis clearly shows that the width (FWHM) of (e.g. Al 111 reflection) the peaks was decreased with increase in the intensity as the annealing temperature increases (Fig. 9a and b). It can be noticed that after annealing, more number of peaks (less intensity) corresponding to Al3Zr and Al9.83Zr0.17 are appeared especially in the XRD patterns of Al-10% Zr sample. These intermetallic phases plausibly played an important role to stabilize the matrix grains by Zenner pining at higher temperature. The thermal stability of the nanocrystalline Al-Zr alloys was ascertained to the presence intermetallic compounds such as (L12 - Al3Zr) [39]. Fig. 10 shows the variation of crystallite size (calculated from XRD data) as a function of annealing temperature. The crystallite size was found to stabilize in a smaller size range (i.e. less grain growth) with increase in Zr content at constant annealing temperature. The crystallite size of the Al-1% Zr sample annealed at 150  C was found to be ~95 nm and the same was estimated to be

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Table 2 Total Gibbs free energy change due to reduction in crystallite size and due to increase in the dislocation density is compared with that obtained from Miedema’s model. Zr (at.%)

DGb (kJ/mol)

DGs (kJ/mol)

DGT ¼ DGb þ DGs (kJ/mol)

DGMiedema (kJ/mol)

1 2 3 4 5 10

0.474 0.498 0.540 0.571 0.627 0.648

0.00642 0.00791 0.00945 0.01120 0.012633 0.01295

0.4805 0.5063 0.5494 0.5829 0.6397 0.6609

0.473 0.971 1.471 1.969 2.463 4.832

Fig. 9. Comparison of XRD pattern of Al-Zr alloys annealed at various temperatures.

Fig. 10. Variation of crystallite size of the annealed samples as a function of annealing temperatures and Zr concentration.

44 nm for 10% Zr sample, which was annealed under same conditions. After annealing at 550  C, the crystallite size of Al-1% Zr sample was increased to 165 nm, while it was retained at 58 nm for

10% Zr sample. The same was estimated to be much >200 nm for pure Al. It is also clearly visible that the crystallite size was increased with decrease in Zr concentration at a constant annealing temperature (Fig. 10). The increase of Zr concentration suppressed the grain growth at higher temperature possibly due to the segregation of Zr atoms as well as due to the formation of intermetallic compounds (Al3Zr and Al9.83Zr0.17). The strong stabilizing effect of Zr was also observed by VanLeeuwen et al. [15] in Pd-Zr alloys. They described that the grain growth was continued up to 1200  C, then the grain size was stabilized at value of 80 and 50 nm, respectively, for 19 and 20 at.% of Zr after 24 h of annealing at 1500  C. Darling et al. [13] also observed a strong stabilizing effect of Zr in FeeZr alloy. They reported a grain size of 50 nm of Fe-4 at.% Zr alloy after annealing at 1400  C for 1 h [13]. Fig. 11a and b shows AFM micrographs of Al-1% Zr alloy annealed at (a) 250  C and (b) 550  C respectively; whereas, Fig. 11c shows AFM micrograph of Al-10% Zr alloy annealed at 550  C. The average grain size obtained from AFM analysis was found to be 88 nm for Al-1% Zr annealed at 250  C. It can be noted from Figs. 11 and 4 that the grain size obtained from AFM analysis is corroborated well with the crystallite size of the annealed as well as the asmilled samples as discussed earlier (Fig. 4a and b). The kinetic as well as thermodynamics mechanisms played the pivotal role to suppress the grain growth and hinder the dislocations movement at high temperature. The thermodynamics mechanism is known to

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Fig. 11. AFM micrograph of annealed samples: (a) 250  C Al-1%Zr (b) 550  C Al-1%Zr (c) 550  C Al-10%Zr.

be more effective in the grain size stabilization as the excess grain boundary free energy can be reduced almost to zero by the solute segregation along the grain boundaries [12,16,40]. The AFM micrographs (Fig. 11) clearly indicate that the grain growth was restricted at the high temperatures by the kinetic (Zener pinning by intermetallic precipitates) and thermodynamics mechanisms (segregation of Zr atoms along matrix grain boundary). Fig. 12a and b, respectively, show dark field TEM image and corresponding SAED pattern of Al-10% Zr sample, which was annealed at 550  C for 1 h. The SAED pattern was recorded from the encircled region as shown in Fig. 12a. The dark field TEM image shows that the nanocrystalline grains are well-within 100 nm. The grain size (obtained by TEM analysis) was found to be consistent with the XRD crystallite size (58 nm) and AFM grain size (65 nm) as discussed earlier. As shown in Fig. 12b, the corresponding SAED pattern shows discontinuous rings, which is due to the coarsening of Al-matrix grains as well as intermetallic particles during

annealing. The formation of intermetallic phase(s) (L12- Al3Zr) was verified from the extra rings of the SAED pattern other than the Al (as indicated). The presence of the same intermetallic phase was also identified in the XRD pattern (Fig. 9b) of the annealed sample of the same composition. The intermetallic compound enhanced the strength of the Al-Zr alloys and also stabilized the matrix grains well within the nanometer level at a high temperature. Similar kind of strengthening effects was also reported by Mula et al. in Cu-Y alloys [41]. Fig. 13 shows the microhardness values as a function of annealing temperatures. Indentation was made on individual particles within the compact (pellet). To avoid the influences of the indentation plastic zone with particle boundaries, a particle diameter-to-indent depth of at least 10:1 was maintained for the hardness tests. Each reported hardness value is the average of 10 such measurements. The microhardness value was found to be 1.57 GPa for the as-milled sample of Al-1% Zr and the same is

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Fig. 12. TEM analysis of Al-10%Zr sample annealed at 550  C: (a) Dark field image (b) SAED pattern.

(such as L12-Al3Zr) are responsible to obtain such high thermal stability. At higher annealing temperatures (beyond 350  C), the fast decrease in the hardness is attributed to the coarsening of the matrix grains as well as the intermetallic precipitates. 4. Conclusions The Al-Zr nanocrystalline alloys were prepared by mechanical alloying of the elemental powder blends for high strength aerospace and structural applications. Their thermal stability, mechanical properties and microstructural features were investigated in detail. The following findings are summarized from the analysis of the experimental results and discussion.

Fig. 13. Microhardness of as-milled and annealed samples of Al-Zr alloys at various temperatures.

estimated to be 2.35 GPa for Al-10% Zr alloy. It is important to note that initially the hardness value was increased up to 150  C and then it showed a decreasing trend. The maximum hardness of 2.56 GPa was measured for the annealed (at 150  C) sample of the Al-10% Zr alloy. It can be noted that (from Fig. 10) there is bit grain coarsening occurred during the annealing and the crystallite size was increased to ~45 nm compared to 30 nm for the as-milled sample of the same composition. Therefore, higher hardness is mainly due to the formation of more amount of harder intermetallic compounds during annealing at low temperature (150  C). Moreover, the segregation of Zr atoms along the grain boundaries also stabilized the matrix grains to retain the hardness. The microhardness of the annealed samples (at 550  C) was found to be only 0.75 GPa for Al-1% Zr and 1.65 GPa for Al-10% Zr alloy. It is significant to note that the hardness values were decreased slowly with increase in the annealing temperature up to 350  C, especially for the 10% Zr alloy. The bit decrease in the hardness is attributed to the coarsening of the matrix grains. Thermodynamic mechanism of the solute segregation along the grain boundaries as well as the kinetic mechanism of Zener pinning by second phase particles

(1) The as-milled sample of Al-1 at.% Zr showed a complete dissolution of solute after MA for 8 h. This is confirmed from the calculation of the excess stored energy as per the Miedema’s semi empirical model, lattice parameter variation and XRD phase analysis. The as-milled samples containing more than 1 at.% Zr showed the formation of very fine nanocrystalline intermetallic compounds such as Al3Zr and Al9.83 Zr0.17 along with the formation of Al-based disordered solid solution. This is confirmed by the TEM analysis. (2) The Al-Zr solid solutions with intermetallic compounds showed an excellent thermal stability, especially for 10% Zr alloy. The Al-Zr solid solution contained up to 1 at.% Zr and the excess amount of added Zr resulted in a formation of the intermetallic compounds, e.g. L12-Al3Zr, which played an important role in the stabilization of the nanocrystalline grains. The corresponding hardness and crystallite size of the Al-10% Zr were found to be 1.65 GPa and 58 nm even after annealing at 550  C. (3) The kinetic mechanism, i.e., Zener pinning by second phase particles (e.g., L12-Al3Zr) was found to be responsible for achieving such a high thermal stability. Acknowledgement The authors are highly acknowledged Department of Metallurgical and Materials Engineering and Institute Instrumentation Centre, IIT Roorkee for providing the facilities and support to carry out the research work.

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