Bulk Nanostructured Metals and Alloys : Processing , Structure , and Thermal Stability

1Department of Chemical Engineering and Materials Science, University of California, Davis, CA 95616, USA 2Department of Materials Science and Engineering, North Carolina State University, Raleigh, NC 27695-7907, USA 3Department of Materials Science and Engineering, University of North Texas, Denton, TX 76203, USA 4Department of Metals and Materials Engineering, University of British Columbia, Vancouver, BC, Canada V6T 1Z4

Nanostructured metals and alloys have been the focus of research during the past few decades due to their interesting microstructures and promising mechanical properties. Despite the fact that this class of materials shows intriguing properties, there are roadblocks that prevent large-scale production and industrial application of these materials that include but not limited to low ductility and poor thermal stability. In this regard, tailoring mechanical and microstructural properties by developing new alloys and inducing different microstructural features via various processing routes have been the major point of interest for researchers to overcome the difficulties in industrialization of nanomaterials. High tendency of nanograins to reduce their energy by grain coalescence and growth brings the thermodynamics and kinetics of grain growth to attention in order to understand the phenomena and develop thermally stable nanostructures. Attempts in this regard include reducing the mobility of a grain boundary by means of mechanisms such as solute drag, second-phase particle pinning, and chemical ordering. On the other hand, the demand for producing large-sized bulk parts, especially in structural applications, has led to the emergence of new processing techniques and further application-based approaches to the issue of the nanostructured metals and alloys. This special issue addresses the different challenges for up scaling the production of nanostructured metals and alloys with emphasis on processing, microstructure, and thermal stability. We would like to express our appreciation to all authors in this special issue for their help and efforts in addressing these current issues in the field. Furthermore, thanks are extended to all reviewers for enhancing the quality of these papers.
This special issue contains six papers related to synthesis and characterization of nanostructured metals and alloys. In "Thermal stability of neodymium aluminates high-k dielectric deposited by liquid injection MOCVD using single-source heterometallic alkoxide precursors," P. Taechakumput et al. studied the effects of high-temperature post deposition annealing (PDA) on the properties of the NdAlO x thin films, deposited by metalorganic chemical vapor deposition (MOCVD) using single-source precursor. These NdAlO x thin films were shown to remain amorphous up to 50 • C as indicated by XRD. No significant level of crystallinity or movement of metal ions was evident after annealing at 950 • C as indicated by MEIS energy spectra. Good electrical integrity was maintained even after 950 • C PDA showing the extracted dielectric permittivity of 12, a low leakage density of 7 × 10 −7 Acm −2 at 2 MV cm −1 , and a density of interface states at flat band D it of 4.01 × 10 11 cm −2 eV −1 . These features make the neodymium aluminate a potential candidate for the dielectric replacement.
In their paper, Z. S. Hu et al. studied "Align Ag nanorods via oxidation reduction growth using RF-sputtering" and demonstrated an oxidation reduction growth (ORG) technique with mixed-gas sputtering to create Ag nanorod arrays via oxide-assisted growth without any chemical solutions or contamination from aqueous solution. The ORG methodology is used to deposit an Ag buffer layer with silver oxide nanoclusters to obtain Ag nanorod arrays using a two-step mixed gaseous process. The Ag nanorods grew in the original locations of reduced metal nuclei after thermal composition of silver oxide nanoclusters. The Ag nanorods stood vertically on the Ag buffer interlayer and grew from the interface between the Ag grains and the Ag buffer interlayer. The success of the technique provides support for the oxidation reduction growth (ORG) mechanism and proves suitable for fabrication of Ag nanorods in the semiconductor industry. Due to photoactivation caused by radiative recombination of Fermi level electrons and dband holes, the observed photoluminescence spectra of the Ag nanorods were observed to be 2.17 eV in both air and vacuum.
In "Synthesis of bulk nanostructured DO 22 superlattice of Ni 3 (Mo, Nb) with high strength, high ductility, and high thermal stability," H. M. Tawancy showed that bulk nanostructured DO 22 -type superlattice with high strength, high ductility, and high thermal stability can be synthesized in a Ni-Mo-Nb alloy with a composition approaching Ni 3 (Mo, Nb) by a simple aging heat treatment at 700 • C. Upon thermal aging, the grains of the high temperature fcc phase are subdivided into ordered crystals on the nanoscale (10-20 nm) with room-temperature yield strength of about 820 MPa and tensile ductility of 35%. Plastic deformation in the ordered state is found to predominantly occur by twinning on (111) planes of the parent fcc structure indicating that the superlattice preserves the twinning systems of the parent phase leading to the observed high ductility.
In their paper, E. Huerta et al. presented "Elastic modulus determination of Al-Cu film alloys prepared by thermal diffusion." In this study, the elastic modulus of Al-50 at.% Cu alloy films with 50-250 nm thickness, prepared by thermal evaporation on Kapton 50HN flexible substrates and postthermal diffusion, was investigated. The morphology and mechanical properties of the Kapton foil substrate and metallic alloys were investigated to distinguish features and properties of the formed Al-Cu alloy from the bare substrate. Pure Al and Cu films with thickness of 50 nm were analyzed for comparison. The elastic modulus of 50-250 nm thick Al-50 at.% Cu alloys thermally evaporated on to Kapton substrates and postformed by thermal diffusion was investigated. Al 2 Cu phase was the dominant crystalline phase formed as determined by XPS. Force-strain curves of the Al-Cu alloys were obtained by subtracting the force-strain curve of the Kapton substrate from the force-strain curve of the Al-Cu/Kapton material system. The elastic modulus was obtained from the slope of the corresponding stress-strain curves. Elastic modulus of the Al-Cu alloys decreased as the film thickness increased, and their values were determined to be in the range of 106.1 to 77.8 GPa for 50 to 250 nm thick alloys, respectively. The elastic modulus of the studied Al-Cu alloys was found to lie between the corresponding bulk values of the Al and Cu films. The elastic modulus measured for 50 nm thick Al was higher than its corresponding bulk value, while the elastic modulus of 50 nm thick Cu was smaller than its bulk value. The highest values of the mean grain size and rms-roughness were found for the 150 nm thick Al-Cu alloy. The methodology used to obtain the elastic modulus does not yield alloy microfractures because of the small strains (<1%) applied during tensile testing.
In their paper, M. Tavoosi and coworkers studied "Consolidation of amorphous Al 80 Fe 10 Ti 5 Ni 5 powders by hot pressing." The purpose of their study was to investigate the feasibility of fabricating amorphous Al 80 Fe 10 Ti 5 Ni 5 powders by mechanical alloying and consolidation into bulk samples by a hot-pressing technique. The crystallization process of this amorphous alloy is a one-stage mode of the Al 13 (Fe,Ti) 4 and Al 3 Ti intermetallic compounds. The results showed that the as-milled amorphous Al 80 Fe 10 Ti 5 Ni 5 powders were consolidated successfully into bulk metallic glasses by a hot-pressing technique. The temperature and pressure for successful condensation of amorphous powders in the hotpressing method were 550 • C and 600 MPa, respectively. During the consolidation, the amorphous phase does not remain, and an AlTi intermetallic phase precipitates in the amorphous matrix.
In their paper, K. S. Lin et al. investigated "Synthesis and characterization of metal hydride/carbon aerogel composites for hydrogen storage." The synthesis, characterization, and H 2 adsorption capacity of carbon aerogels (CAs)/metallic hydride nanocomposites as a catalyst were studied. Experimentally, The H 2 storage capacity of metallic samples was measured by a TGA microbalance method. In addition, fine structure and crystallinity of metallic hydride was identified by BET nitrogen adsorption isotherms, HR-TEM, FE-SEM/EDS, XRD, and XANES/EXAFS. The carbon aerogels were very effective in improving the hydrogen storage capacity of the Fe-, Ti-doped MgH 2 samples with the "hydrogen spillover" route. Additionally, metallic dispersion of Pd nanoparticles onto the CAs might improve the hydrogen adsorption abilities. Higher surface area from CAs and the defect criteria of the MgH 2 surface due to the doping of the metallic particle cause improvement in the hydrogen adsorption capacities. Conversely, Ti-doping onto the sodium aluminum hydride only improves slightly the hydrogen storage capabilities.

Introduction
Hydrogen is recognized as a clean fuel because of its almost complete combustion in air with notably high energy delivery ability. Recently, owning to the need for efficient and safe H 2 storage, enormous attention has been paid on the advantages of synthesizing composite metal hydrides [1][2][3]. Complex metal hydride composites in the form of ABH 4 , where A is an alkali metal and B is a group III element, have been widely studied in solution as proton acceptors to enhance H 2 adsorption abilities [4][5][6][7]. NaAlH 4 is different from any other metal hydrides and borohydrides with similar structures since it is capable to reversibly store H 2 after doping with transition metals (e.g., Ti, Fe, or Zr) [8][9][10]. Ti-doping also improves significantly the kinetic and cycling performance of NaAlH 4 at moderate temperatures around 373 K and ambient pressures. All these properties make it suitable as mobile, lightweight H 2 storage materials for potential application on a hydrogen fuel cell [11][12][13][14][15][16][17][18][19][20][21][22][23]. Practically, achievable H 2 storage capacity for Ti-NaAlH 4 is only 3.7 wt%, falls short of the theoretical value of 5.6 wt% for NaAlH 4 [20][21][22][23].
Magnesium and its alloys are potential hydrogen storage material because of their very high hydrogen storage abilities. MgH 2 is also relatively light and cheaper when compared with other metal hydrides. Magnesium is often alloyed with transition metals (e.g., Fe, Ti, Zr, Zn, or Mn) to increase the adsorption and desorption rates [24][25][26][27][28][29]. Similarly, FeTi is a viable compound with a hydrogen adsorption capacity of around 1.90 wt%. However, activation process of FeTi is necessary to get rid of the TiO 2 layer, which would otherwise hinder the hydrogenation/dehydrogenation. Consequently, higher temperatures and pressures are frequently required to achieve reproducible, maximum hydrogen adsorption/desorption capacities in the compounds [30].
Palladium-doped monolithic carbon aerogels (CAs) show a high surface area and pore volume. These characteristics are strongly influenced by the chemical nature, distribution and dispersion of the palladium particles [31][32][33]. Furthermore, addition of 1 wt% Pd can be effective for the hydrogen adsorption capacity of FeTi to achieve its theoretical hydrogen storage capacity of 1.9 wt% [31][32][33].
As carbon materials such as graphite or CAs have been shown to reduce the hydrogen adsorption and desorption 2 Journal of Nanomaterials temperatures of Mg [34], it is not unreasonable to expect that CA addition could also be able to improve the hydrogen adsorption behavior of the FeTi system. There is a further advantage of using Pd-doped CAs as it has been shown that an addition of 1 wt% Pd in the FeTi can effectively enhance its hydrogen adsorption capacity [31][32][33].
X-ray absorption near edge structure (XANES)/extended X-ray absorption fine structure (EXAFS) spectroscopy offers a powerful tool to understand the oxidation states and fine structures of Ti atoms in the microstructure of metal hydride composites [35][36][37][38]. Thus, a thorough study of the location of Ti species in metal hydrides may be helpful to explain the enhancement of hydrogen adsorptive kinetics by identifying the compounds in the hydrides. XANES/EXAFS spectroscopy is also an excellent technique for characterizing the valency and local structure of Ti species in a complex Timetal hydride with short-range orders [37][38][39][40][41]. These studies are conducted to identify the optimum hydrogen storage capacity for hydrogen storage system. In this work, synthesis and characterization, and advantages/disadvantages of different metal-hydride systems such as Mg, TiFe, or Nabased metal hydrides, complex hydrides compounds, were investigated by field-emission scanning electron microscopy (FE-SEM), X-ray diffraction (XRD), and XANES/EXAFS spectroscopy. In addition, the adsorptive H 2 storage capacity was also studied using a high-pressure thermogravimetric analyzer.

Preparation of Metal-Doped Carbon Aerogels.
Metalloaded organic aerogels were prepared following the method developed by Pekala [42] for resorcinol-formaldehyde aerogels. In a typical procedure, CA was synthesized by first stirring vigorously resorcinol (20.70 g), Na 2 CO 3 (0.04 g), and equal amount of deionized water in a polypropylene jar at room temperature until resorcinol dissolved. Formaldehyde (30.94 g of purity 36.5%) was then added to the solution, followed by 0.01 M PdCl 2 (Reagent ACS R.D.H., Germany) to acquire the doping of palladium. The solution was left at room temperature for 24 h, then at 323 K for the next 24 h, and finally at 365 K for another 72 h to obtain a dark black monolith. The product was rinsed with acetone three times and soaked overnight at room temperature, yielding the hydrogel materials. After drying at room temperature, the hydrogels were isolated by heating in an oven at temperatures from 273 to 1,100 K under an Ar atmosphere for several hours to obtain the Pd-doped CAs.

Preparation of Metal
Hydrides. As received TiCl 3 (Aldrich, 99.99%, anhydrous) was used as the catalyst precursor. Crystalline NaAlH 4 (Fluka, 99.9%) was purified by tetrahydrofuran (THF) (Aldrich, 99.9%, anhydrous) solution and vacuum-dried while being filtered through 0.7 μm filter paper to remove any residual Al powder and other solids from the solution. The purified NaAlH 4 was mixed with the catalyst precursor in THF to produce a Ti-doped sample containing up to 20 mol% metal (relative to either Na or Al). The THF was evaporated while the NaAlH 4 and catalyst precursor were mixed manually for about 30 min, using a mortar and pestle, or until the sample appeared to be completely dry. A Ti-NaAlH 4 sample was then ball-milled for 2 h using a SPEX 8000 high-energy mixer/mill at a milling speed of 1,000 rpm, and the weight ratio of ball to powder was 10 : 1. 0.5-1.0 gram of sample was subjected to ball-milling each time in a 65 cm 3 SUS 304 vial by ZrO 2 balls of 1.3 cm in diameter. Fe-and Ti-doped MgH 2 was prepared following Yao et al. [43], where MgH 2 was obtained by heating Mg powders (Alfa Aesar) to 400 • C in a hydrogen storage rig. 90 wt% conversion to MgH 2 was achieved according to the volumetric calculations. 1.5-2.0 grams of nanophase iron powder (100 nm), Ti (Aldrich Co), and MgH 2 samples were mixed in the composition of MgH 2 -5.0 wt% (2Fe + Ti, molar ratio) and mechanically ball-milled by using a ZrO 2 ball mixer/mill (model SPEX 8000). The samples were subjected to ball-milling as described above for 2 h, using a weight ratio of ball to powder of 8 : 1 The powder was then ballmilled again with 5.0 wt% of CAs for an additional 10 h at a milling speed of 1,000 rpm. All procedures described above were carried out in a flowing argon gas glove box to prevent the metal hydride samples from oxidation. The samples were then used for hydrogen adsorption property tests and microstructure characterization.

Hydrogen Adsorption Measurements.
The hydrogen isotherms were measured gravimetrically at different temperatures using a method previously described by Eddaoudi et al. [44]. A Cahn Thermax 500 microgravimetric balance with a sensitivity of 1 μg was used to measure the change in mass of samples suspended within a glass enclosure under a certain atmosphere. A pressure sensor, with a range of 0 to 68 atm (at 1,000 • C) and sensitivity of 0.011 atm, was used to measure the hydrogen pressure in the chamber. Samples were outgassed overnight until a constant mass was attained; these varied from 0.2 to 2.0 g. Prior to admittance of argon gas, the entire chamber and manifold were evacuated overnight. The system was purged at room temperature three times with the argon gas before cooling to 77 K, and gases were passed through a molecular sieve trap immersed in liquid nitrogen to remove any condensable impurities or moisture before being exposed to the sample. Pressures were measured with the range covering 1 to 30 atm. Hydrogen was added incrementally, and data points were recorded when no further change in mass was observed. The dynamic hydrogen adsorption data on Pd-, Ti-, Fe-doped CAs, NaAlH 4 , or MgH 2 batch adsorption experiments were also conducted at room temperature (298 K) and different pressures. The adsorbed amount of hydrogen was calculated after the buoyancy correction.

Characterization of Carbon Aerogels and Metal Hydrides.
The morphologies, crystallinity, particle size distribution, and the compositions of as-synthesized metal hydrides and metal-doped CAs were determined by FE-SEM equipped with EDS (Hitachi, S-4700 Type II) and HR-TEM (Zeiss 10C). XRD (MAC Science, MXP18) was performed to identify the phases and crystallinities in the samples. Complex of metallic hydride and CAs was scanned from 20 to 80 • (2θ) with a scan rate of 4 • (2θ) min −1 , and monochromatic CuK α radiation was used. The recorded specific peak intensities and 2θ values can be further identified by a computer database system (JCPDS). The surface areas of the CAs were measured by BET (Brunauer-Emmett-Teller) nitrogen adsorption (Micromeritics ASAP 2010 Instrument). For the BET surface area measurement, the samples were scraped from the sample tube substrate and powdered so as to avoid any influence from the steel tube. Prior to measurement, all samples were degassed at 423 K for 1 h. For the calculation of the BET surface areas, the relative pressure range P/P 0 from 0.05 to 0.2 was used. The pore radius distribution was determined by Barrett, Joyner, and Halenda (BJH) method.
The XANES/EXAFS spectra were collected at BL16A1 and BL01C1 (SWLS) beamlines at the National Synchrotron Radiation Research Center (NSRRC) of Taiwan. The electron storage ring was operated with the energy of 1.5 GeV and a current of 100-200 mA. A Si(111) double crystal monochromator (DCM) was used for providing highly monochromatized photon beams with energies of 1 to 15 keV (BL16A1) and 5 to 30 keV (BL01C1) and resolving power (E/ΔE) of up to 7,000. Data were collected in fluorescence or transmission mode with a Lytle ionization detector [39] for Ti (4,966 eV) and Pd (24,350 eV) K-edge experiments at room temperature. The photon energy was calibrated by characteristic preedge peaks in the absorption spectra of Ti standards. The raw absorption data in the region of 50 to 200 eV below the edge position were fit to a straight line using the least-square algorithms [37][38][39][40]. The XANES was extended to energy of the order of 50 eV above the edge. The k 2 -weighted and EXAFS spectra were Fourier transformed to R space over the range between 2.5 and 12.5Å −1 . The EXAFS data were analyzed by using the

Morphology of Metal-Doped Carbon Aerogels and Metal
Hydrides. The HR-TEM image and FE-SEM micrographs of metal hydride and CAs complexes synthesized are shown in Figure 1. Figure 1(a) is an HR-TEM observation of the Pd-doped aerogel structure consisting of spherical primary Pd(0) nanoparticles. It can be seen that palladium particles of size 3 to 10 nm were uniformly distributed throughout the gel. FE-SEM micrograph in Figure 1(b) reveals the pellet-like or irregularly shaped Ti-NaAlH 4 crystalline nanoparticle, with diameters approximately 50-80 nm. Correspondingly, Fe and Ti nanoparticles with an average diameter of 50-80 nm were dispersed uniformly on the surface of the doped MgH 2 surface as shown in Figure 1(c).

X-Ray Powder Diffraction.
The structure of the compound was confirmed by structural refinement of X-ray diffraction patterns shown in Figure 2. The intensive peaks appearing at small 2θ angles were characteristics of porous materials which possess numerous pores or cavities. Whereas Figure 2  correction for instrumental broadening) applied to the halfheight of the maximum intensity diffraction peak. Metallic Na 3 AlH 6 and Al were consistent with the characteristic peak at 32.92 and 38.6 • , respectively [24]. From the FE-SEM measurement and XRD patterns, the titanium compound distributed on the surface of the NaAlH 4 can be confirmed.
The X-ray diffraction patterns of the Pd-doped CAs indicate that the Pd nanoparticles have a crystalline structure (Figure 2(b)). According to FE-SEM micrographs, the well dispersion of metallic Pd nanoparticles in the CAs apparently facilitated the porous structure. The XRD patterns of Pddoped CAs show those two characteristic 2θ sharp peaks with higher intensities around 40.12 and 46.77 • , indicating a largely crystalline structure. The mean particle size was about 5-8 nm from HR-TEM observation that may help the mesoporous structure of CAs to adsorb hydrogen. The diffraction angles at around 20 and 44 • , corresponding to the (002) and (101) diffraction peaks of graphite, respectively, are represented in Figure 2(b). In addition, hexagonal structure of the CAs started to become disordered with the heat treatment, owning to the progressive graphitization of the sample. Figure 2(

XANES and EXAFS. Figures 3(a)-3(e)
show the preedge XANES spectra of Ti K-edge (4,966 eV) in nanophase Ti-NaAlH 4 nanoparticles and anatase-typed TiO 2 standard [45]. These spectra exhibit an absorbance feature of 4,966 eV for the 1s to 3d transition, which is forbidden by the selection rule in case of perfect octahedral symmetry. The sharp feature at 4,987 eV, due to the dipole allowed of 1s to 3p xy electron transition, indicates the existence of Ti(IV). The The reversible hydrogenation of NaAlH 4 and dehydrogenation to NaH and Al occurred in two steps summarized as follows [1][2][3][4][5][6]. 3NaAlH 4 Na 3 AlH 6 + 2Al + 3H 2 3.7 wt% H 2 , 2Na 3 AlH 6 6NaH + 2Al + 3H 2 1.9 wt% H 2 . (2) Remarkably, the addition of the Ti species lowers the activation energy of the adsorption/desorption processes and enhanced the kinetic rates of these reactions by several orders of magnitude [46][47][48][49][50]. However, the exact mechanism still required further investigations. Notably, the hydrogendepleted multicomponent products resulting from (2) were found to form large crystallites of the complex metal hydride upon the reverse reaction in the presence of Ti during hydrogenation [45][46][47]. This is an important clue to understand the role of Ti where it implies a concerted migration of metal atoms over large distances. Local structural arrangements without the involvement of large-scale atomic migration would instead result in an amorphous or nanocrystalline structure.
The main scientific issues concerning the chemical forms (or speciation) of active species ultimately depend on molecular-scale structure and properties. Basic understanding at this scale is essential for further understanding the catalytic behaviors of Ti-NaAlH 4 complexes synthesized at different mol% values of Ti. Generally, Ti K-edge EXAFS spectroscopy in Figure 4 can provide the information on the Ti atomic arrangement of catalysts in terms of bond distance, bond angle, coordination number, kind of near neighbors, and thermal or static disorders. Experimentally, the Ti species with a central Ti atom have a Ti-O bond distance of 1.95 ± 0.02Å and a coordination number of 4.07, in a Ti-NaAlH 4 composite after ball-milling for 2 h using a high-energy mixer/mill. Comparatively, however, a rutiletyped TiO 2 has a central Ti atom with a Ti-Ti and a Ti-(O)-Ti bonds, and their bond distances are 2.96 ± 0.02Å and 3.57 ± 0.02Å, respectively [29][30][31][32]40]. Since the differences of mass density (ρ) and band structures between anatase-and rutile-typed TiO 2 , bond distances of Ti-O are 1.95 ± 0.02Å and 1.93 ± 0.02Å, respectively [9][10][11][12][13][14].
The EXAFS data shown in Figure 4 reveals that the peak intensity of first-shell Ti-O bond increased with decreasing mol% values of Ti-dopants for a Ti-NaAlH 4 composite after ball-milling for 2 h using a high-energy mixer/mill. These results may offer a further explanation on the higher oxidation states of Ti species for a Ti-NaAlH 4 complex after ball-milling at higher rates and temperatures. Moreover, each Ti(IV) ion in the framework of nanophase Ti-NaAlH 4 complex crystals is surrounded by six O 2− ions [9][10][11][12][13][14]. Since the bond lengths of the first and second shells for Ti-O bonding are 1.94 ± 0.02Å and 1.97 ± 0.02Å, similarly with the coordination numbers of four and two, respectively, an octahedral structure of the Ti-NaAlH 4 composites was found using EXAFS measurement [45]. As shown in Table 1 and Figure 5, the Ti species with a central Ti atom have a Ti-O bond distance of 2.81 ± 0.02Å and a coordination number of 1.83, in a Ti-NaAlH 4 composite in 8 mol% Ti-NaAlH 4 metal hydride composites (more than 2 mol%) after the hydrogen adsorption processes. It can be seen that all Ti species in Ti-NaAlH 4 metal hydride composites can be reduced in the hydrogen adsorption processes and be kept in the metallic Ti form.

N 2 Gas Adsorption Isotherms Analyses.
Nonspecific physical adsorption of the CAs was carried out to measure the total surface area and pore size distribution for Pddoped CAs, as shown in Figure 6. The surface area, pore size distribution, and crystalline diameter of CAs were calculated according to the adsorption data summarized in Table 2. A large surface area was generally observed for Pd-doped CAs and beneficial characteristic for a great variety of applications. The microcrystalline structure of Pd on the CAs is confirmed from the HR-TEM analysis (Figure 1(a)), resulting in an increase of the surface area. The nitrogen adsorption-desorption isotherms exhibit a hysteresis behavior, indicating the microstructural nature of the porous specimens. Moreover, the type IV hysteresis isotherm obtained is represented in Figure 6(a). The adsorption hysteresis was observed in the region of a relative pressure P/P 0 above 0.4. The specific surface areas of CAs and Pd-doped CAs are 687 and 725 m 2 g −1 , with the pore volumes of 0.62 and 0.64 cm 3 g −1 , respectively, as shown in Table 2. It also indicates that these resulting factors of surface enhancement depended on the doping of the Pd on the CAs and consisted of higher porous structure [49,50]. These results may be caused by the dispersion of the Pd on the CAs surface and that indicated the metallic dispersion would be helpful for adsorption criteria.
The intensive peak of the Pd(0) K-edge region around 24,368 eV and the stronger peak at 24,392 eV were due to the orbital 3d → 5s or 4p hybridization caused by unoccupied d bands, as shown in Figure 7. These quantified the doping of Pd on the surface of carbon aerogels. Therefore, in the absorption where the orbital 3d → 5s or 4p hybridization was intensive, the p-like density of states was notably enhanced and the absorption increased like the second absorption peak in the XANES/EXAFS spectra of the Pd standards (Figures 7(b) and 8(a)). The structural parameters of the Pd powder and the palladium doped CAs, obtained from the best fit to the EXAFS data, are shown in Table 1. The coordination numbers and bond lengths also were well compared with the Pd metal and palladium-doped CAs crystal structure data. The Debye-Waller factors (Δσ 2 ) were less than 0.015 (Å 2 ), which indicated that the palladium confirmatively has center Pd atoms coordinated by Pd-Pd bonding. The bulk Pd metal possessed Pd-Pd bond distance of 2.74Å ± 0.02Å and a coordination number of 9.59. For comparison, fresh CAs-Pd clusters were found to have a Pd-Pd bond distance of 2.12Å with a coordination number of 6.18. This shortening of the bond distance with a lower coordination number was caused by random motion of surface atom on the small Pd particles and was responsible for the increase in the surface area of the palladium doped CAs samples, whereas hydrogen reduced bulk CA-doped Pd cluster consisted with the Pd-Pd bond distance of 2.73Å and a coordination number of 9.23. This revealed that the Pd nanoparticles were notably activated in the reduction process [48]. Moreover, the shortening of the bond distance with coordination number probably was caused by the random motion of surface atom on the small Pd particles and therefore increases the surface area of the Pd-doped CAs samples. Furthermore, this result also revealed that the Pd nanoparticles were well dispersed on the surface of CAs, which might improve the amount of hydrogen storage efficiency significantly.

Hydrogen Gas Adsorption Isotherms Analyses.
Comparisons of the hydrogen adsorption of metal hydride and   Mesoporous structure of the CAs, which has a very high surface area as confirmed from the BET nitrogen isotherm analysis, improved the hydrogen adsorption behavior. This is inline with the finding of Imamura et al. [34], who reported that the uptake of hydrogen in a C-Mg composite took place at a significantly lower temperature than that in pure magnesium. Similarly, Jung et al. [51] suggested that the faster adsorption kinetics of MgH 2 systems may result from the very high defect densities of the damaged MgH 2 surfaces caused by the ball-milling of Mg with hard and small oxide particles. Defects provide hydrogen an easy path to go into Mg powders. Therefore, the hydrogen adsorption capacity of 2.7 wt% was obtained with the addition of FeTi nanoparticles into the MgH 2 samples by ball-milling ( Figure 9(b)). Thus, minute particles of FeTi were found to be the most useful catalyst to improve the hydrogen adsorption properties of MgH 2 . Nitrogen adsorption by the  CAs or Pd-doped CAs material clearly shows reversible type IV isotherms at each of the activation stages, indicative of permanent porosity in Figure 6(a). The selectivity of H 2 adsorption over Pd-doped CAs was 1.6 wt%, as shown in Figure 9(c). Due to the doping of Pd on CAs, access of hydrogen was easier in the vacancies created by the mesoporous structure with a higher surface area that was confirmed with the BET nitrogen isotherms measurement. In addition, this suggests that an antidispersive force plays an important role in case of Pd-doped CAs samples and the expected specific interaction with Pd active sites. Thus, this work confirmed that metal doping by Pd on CAs may improve the hydrogen adsorption criteria. An intriguing observation concerns the dependence of the rehydrogenation Journal of Nanomaterials  Hydrogen adsorption (wt%) kinetics on the amount of Ti doping. Moreover, as shown in Figure 9(a), it was observed that the hydrogen uptakes of FeTi-CAs-MgH 2 systems could be enhanced markedly by the phenomenon of "hydrogen spillover" via a simple technique for building carbon bridges [52]. One of the possible reasons is that the hydrogen molecules do not readily dissociate on Mg surface. Experimentally, the catalytic effect was found on hydrogen adsorption of mixing transition metal of Ti atoms into Mg hydride powder during ball-milling. The likely mechanism proposed for the FeTi-CAs-MgH 2 hetero-structures can be also summarized concisely: (1) adsorption of hydrogen on the Ti surface, (2) dissociation of hydrogen and chemisorption of atomic hydrogen on the surface, (3) migration of atomic hydrogen onto the CAs support, and finally (4) chemisorptive spillover onto the MgH 2 substrate. In addition, spillover is facilitated through the use of a carbon bridging compound between the Ti/CAs complex and the MgH 2 substrate. Typically, an 8 mol% Tidop onto the NaAlH4 (Figure 9(d)) revealed the maximum hydrogen adsorption capacity of 1.2 wt% within this class of adsorptive materials. In addition, perfect single crystal TiO 2 was inert toward reaction with H 2 [51]. However, H 2 was absorbed by TiO 2 -surfaces that contained a higher density of defects in the crystal structure. Therefore, surface of MgH 2 in the presence of the metal-doped Fe and Ti with CAs created a high defect density that introduced a higher hydrogen adsorption behavior in the adsorption or desorption processes.

Conclusions
The synthesis, characterization, and H 2 adsorption capacity of CAs/metallic hydride nanocomposites as a catalyst were investigated in the present work. Experimentally, the H 2 storage capacity of metallic samples was conducted and measured by a TGA microbalance method. In addition, fine structures and crystallinity of metallic hydride were identified by BET nitrogen adsorption isotherms, HR-TEM, FE-SEM/EDS, XRD, and XANES/EXAFS. The CAs were very effective in improving the hydrogen storage capacity of the Fe-, Ti-doped MgH 2 samples with the "hydrogen spillover" route. Metallic dispersion such as Pd nanoparticles onto the CAs may improve the hydrogen adsorption abilities. Higher surface area from CAs and the defect criteria of the MgH 2 surface due to the doping of the metallic particle caused to improve the hydrogen adsorption capacities. On the other hand, Ti-doping onto the sodium aluminum hydride can only improve slightly the hydrogen storage capabilities.

Introduction
Alloys with Al content of 80 to 90 at.% have attracted considerable attention due to their combination of good ductility and high strength [1][2][3][4][5][6][7]. Among the various aluminum alloys, Al-Fe systems are of technological interest due to their advantageous properties, including high specific strength, high specific stiffness, good strength at intermediate temperatures, and excellent corrosion resistance at elevated temperatures under oxidizing, carburizing, and sulfidizing atmospheres [8][9][10][11][12]. A new class of Al-based alloys, which has received a great deal of attention in recent years, is amorphous as well as nanocrystalline [13][14][15][16][17][18][19][20][21][22][23]. Indeed, amorphous alloys exhibit several superior properties that cannot be obtained in crystalline materials. In general, bulk amorphous alloys can be fabricated by two main processes: direct solidification from the melt and consolidation of an amorphous powder or ribbons. Generally, very high cooling rates are required for the formation of an amorphous phase from the liquid state and only a few amorphous alloys with high glass-forming ability can be applied to the fabrication of bulk amorphous alloys by a direct solidification process. Due to the requirement of a high cooling rate, amorphous alloys have largely been fabricated in the form of powders, ribbons, and wires with small thickness or diameter [14][15][16][17], and hence the application of amorphous alloys as a structural material has been limited.
The glass-forming ability of Al-based alloys is so low and there is not any report about formation of bulk amorphous structure for these alloys by the use of direct solidification methods [11][12][13]. Nevertheless, a great number of amorphous alloys in Al-based alloys have been synthesized by various preparation methods of rapid quenching from liquid or vapor and solid-state reactions [13]. However, the maximum thicknesses of the resulting amorphous alloys have usually been limited. The limitation of the maximum sample thickness has prevented a wide extension of application fields of amorphous alloys. Consequently, great efforts have been devoted to preparation of a bulk amorphous alloy from amorphous alloy powder by using various techniques of warm pressing, warm extrusion, explosive compaction, hot pressing, and spark plasma sintering [13,24]. The powder metallurgy process promotes high strength bonding between particles and a microstructure similar to that of a wrought product [25][26][27][28][29][30].  The purpose of this study is the investigation of the feasibility of fabricating amorphous Al 80 Fe 10 Ti 5 Ni 5 powders by mechanical alloying and consolidating them into bulk samples by a hot-pressing technique.

Experimental Procedure
Commercial elemental powders of Al (99%), Fe (99.9%), Ti (99%), and Ni (99.99%) were used as raw materials. The mixture of elemental powders with a composition of Al 80 Fe 10 Ti 5 Ni 5 (in atomic percentage) was mechanically alloyed in a planetary ball mill under an argon atmosphere in a steel container at room temperature. A rotation speed of 250 rpm and a ball to powder ratio of 10 : 1 were employed. 1 wt% stearic acid powder supplied by Merck was used as the process control agent (PCA).
The produced amorphous powders were consolidated in a vacuum hot-pressing machine to prepare bulk amorphous discs with 10 mm diameter and 1 mm thickness. Hot pressing was performed at different temperatures and pressures. The hot-pressing chamber was evacuated and pressure was applied during the entire consolidating process.
The as-milled powders and consolidated samples were characterized by X-ray diffraction (XRD) analysis, scanning electron microscopy (SEM), transition electron microscopy (TEM), and differential scanning calorimetry. The XRD analysis was performed using a Philips diffractometer (40 kV) with Cu Kα radiation (λ = 0.15406 nm). The XRD patterns were recorded in the 2θ range of 30-70 • (step size 0.03 • and time per step 1 s). The morphology, fracture surface, and cross-sections of prepared samples were examined using SEM. The microstructure of the produced amorphous powder was investigated by TEM and a selected area diffraction (SAD) pattern analysis carried out at an accelerating voltage of 200 kV resolution of 0.19 nm. A differential thermal analysis (DTA) was conducted to study the thermal stability of the produced amorphous alloy using a Reometric STA 1500 differential thermal analyzer. The samples were placed in Al 2 O 3 pans and heated in a dynamic Ar atmosphere up to 1200 • C at a constant heating rate of 40 • C/min. accompanied by a remarkable decrease in their intensities occurred as a result of refinement of crystallite size and enhancement of lattice strain. Increasing milling time up to 10 h led to the disappearance of Fe, Ti, and Ni peaks. This may be due to dissolution of these elements in the Al lattice. By increasing the milling time, two narrow peaks corresponding to the decagonal phase (D-phase) are observed in the XRD patterns. Analysis of the XRD patterns in Figure 1 reveals that the D-phase, which formed from the solid solution during milling, could be amorphized by further milling (after 40 h) and there is no significant change as the milling time is prolonged to 100 h. A TEM micrograph and SAD pattern of powders ball milled for 40 h are also presented in Figure 2. As can be seen in this figure, the SAD pattern of this sample consists of only amorphous halo and there is no evidence of the presence of any crystalline phase, which is in agreement with the XRD result. This indicates that the product of mechanical alloying of the Al-10%Fe-5%Ni-5%Ti powder mixture is an amorphous phase. SEM cross-sectional micrographs of powder particles after different milling times are presented in Figure 3. According to the micrographs, in the early stages of mechanical alloying (MA), the ductile components are flattened to platelet shapes, cold welded together, and form a composite , the interlamellar spacing decreases and true alloying at the atomic level occurs, resulting in the formation of solid solutions, intermetallics, or even amorphous phases. In this stage, the interlayer spacing disappears or becomes so fine that it is no longer visible using SEM (Figure 3(e)). The morphology of Al 80 Fe 10 Ti 5 Ni 5 powders after 40 h of ball milling is also shown in Figure 4. As seen here, the powders milled for 40 h are spherical and have an average diameter of 2 μm. Intensity (cps) study the thermal stability of the produced amorphous phase, the sample was examined by DSC under continuous heating conditions. Figure 5 shows the DSC heating traces of the Al 80 Fe 10 Ti 5 Ni 5 amorphous alloy at a constant heating rate of 20 • C/min. As seen in this figure, two peaks appear in the DSC curve (an exothermic peak at 950 • C and an endothermic peak at 1100 • C). To analyze the crystallization process responsible for the exothermic peak, the as-blended sample was annealed under an Ar atmosphere at 1000 • C for 20 min. The XRD pattern and a TEM micrograph of the Al 80 Fe 10 Ti 5 Ni 5 amorphous phase after annealing at 1000 • C are presented in Figures 6 and 7, respectively. As can be seen here, the annealed sample consists mainly of an Al 13 (Fe,Ni) 4 [31] and Al 3 Ti [32] intermetallic phases. Therefore, the exothermic peak in Figure 5 is attributed to precipitation of these phases from the amorphous phase. Meanwhile, Journal of Nanomaterials the endothermic peak at 1100 • C in Figure 5 is for the melting of the produced alloy. These results indicate that the total transformation sequence of the Al 80 Fe 10 Ti 5 Ni 5 amorphous alloy is a one-stage process in the temperature range from 930 • C to 980 • C. In fact, according to the DSC results, the crystallization in Al 80 Fe 10 Ti 5 Ni 5 proceeds in the same manner as in the Al 80 Fe 10 Ti 10 [11] system and is different from that in Al 80 Fe 20 [23] and Al 83 Fe 17 [10] amorphous systems. In contrast to the Al 80 Fe 10 Ti 5 Ni 5 amorphous alloy (which exhibits onestage crystallization on heating), the Al 80 Fe 20 and Al 83 Fe 17 amorphous phases exhibit three-stage crystallization during heating.

Consolidation of the Ball-Milled Amorphous Powders.
Fully amorphous powder particles with high crystallization temperature (950 • C) in Al 80 Fe 10 Ti 5 Ni 5 alloy provide good conditions for the fabrication of a bulk amorphous material. In this study, the as-milled powders were consolidated by hot pressing into a disk shape. To optimize the condensation parameter, the ball-milled amorphous powders were hot pressed at different temperatures (400, 450, 500, and 550 • C) under various pressures (200, 300, 400, 500, and 600 MPa) for 30 min. In order to investigate the condensed samples, the fracture surface morphology after compression tests was examined using SEM. Fracture surface photographs of samples hot pressed at different temperatures under a constant pressure of 600 MPa and under different pressure values at a constant temperature of 550 • C are shown in Figures 8 and 9, respectively. As seen in these figures, the trace of the interparticle boundaries between the powders increases and the porosity in the hot pressed samples decreases as the pressing temperature and pressure increase. These results indicate that the optimum temperature and pressure for consolidation of the amorphous powders via the hot-pressing method are 550 • C and 600 MPa, respectively, in our experimental conditions. The polished cross-sectional view of the consolidated sample at 550 • C under 600 MPa in Figure 10 shows no remaining pores. Figure 11 shows the XRD pattern of the consolidated at 550 • C under 600 MPa. From the peaks corresponding to AlTi, it can be concluded that the amorphous phase  the consolidation processes and the produced bulk material is not fully amorphous.

Thermal Behavior of Produced Bulk Material.
In order to investigate the thermal behavior and microstructural changes of the consolidated bulk amorphous alloy (below crystallization temperature), pressed samples were annealed at several temperatures for 3 h. XRD patterns of as-hot pressed and annealed samples at 550 • C, 600 • C, 650 • C, and 700 • C are presented in Figure 14. During the consolidation of the amorphous powders at 550 • C for 30 min, the AlTi phase precipitates in the amorphous matrix. Analysis of the XRD patterns in Figure 14 reveals that by annealing the samples at temperature above 550 • C (below crystallization temperature), the AlTi diffraction peaks disappeared and several peaks corresponding to Al 13 (Fe,Ni) 4 and Al 3 Ti intermetallic

Conclusions
In the present work, we fabricated amorphous Al 80 Fe 10

Introduction
The current tendency in the electronic industry is to use flexible polymers as substrates to fabricate electronic circuits, rendering portable, low-cost, low-weight, and flexible electronic devices. For this reason, it is necessary to determine and understand the behavior of the physical properties of nanofilms deposited on flexible substrates [1,2]. Particular attention needs to be paid to the understanding of the mechanical properties, especially the elastic modulus. Among the different methods used to estimate the elastic modulus of thin film materials are those based on vibrations [3], tensile [4], and bending testing [5]. In the electronic industry, the aluminum-copper alloys (Al-Cu) are used as metallic interconnectors [6], because of the reduction of the electromigration effects in aluminum [7]. The inclusion of polymeric materials as substrates in the electronic devices fabrication makes the interconnectors and electronic components suffer strains by Joule effect during their operation. Thus, it is important to study the mechanical behavior of these devices by knowing the elastic modulus of the interconnector materials at nanothicknesses for improving their performance. Figure 1 shows the Al-Cu alloy phase diagram in bulk [8,9] where different phases can be observed. Al-Cu alloy phases β 0 , β, γ 0 , and ε 2 are formed from 1373 to 823 K, while θ (Al 2 Cu) [10], η 1 (AlCu), ξ 2 (Al 3 Cu 4 ), δ (Al 2 Cu 3 ), and γ 1 (Al 4 Cu 9 ) appear as stable phases below 823 K. As can be seen, the formation of the different phases is mainly determinated by the concentration of each element and the cooling rate during solidification [11,12]. An important crystallographic phase with adequate electrical characteristics for the microelectronic industry is the AlCu phase (η 1 ), which is a stoichiometric phase formed below 823 K. In this work, the elastic modulus of Al-50 at. % Cu alloy films with 50-250 nm thickness, prepared by thermal evaporation on Kapton 50HN flexible substrates and post thermal diffusion, is investigated. The morphology and mechanical properties of the Kapton foil substrate and metallic alloys were investigated in order to determine features and properties of the formed Al-Cu alloy from differences with the bare substrate. Pure Al and Cu 50 nm thickfilms were also analyzed for comparison.

Alloys Preparation and Substrate Selection.
For Al-Cu alloys preparation, high-purity Cu (99.999%) from CERAC was first deposited on 12.5 μm thick Kapton (50HN) foil substrates by thermal evaporation into a vacuum chamber with a pressure of 6.0 × 10 −5 Torr and a deposition rate of 0.2 nm/s. Film thickness and deposition rate were measured and monitored in situ with a TM-400 controller with a quartz crystal sensor. High-purity Al (99.999%) was subsequently deposited on the Cu/Kapton system, forming an Al/Cu bilayer on top of the Kapton substrate. The chosen order of deposition is related to the rapid and continuous oxidation of Cu in comparison with the high stability of the Al oxide. The Al-Cu alloys were prepared with a nominal 50 : 50% atomic concentration and deposited with 50, 100, 150, 200, and 250 nm as total thickness. The corresponding individual film thicknesses for this atomic concentration were calculated as t Al = 58.4% and t Cu = 41.6% of the total thickness. Four samples were simultaneously deposited for each batch. Once the bilayers were formed, Al-Cu alloys were post-formed by thermal diffusion at 673 K by a thermal annealing process during 3 h into a vacuum oven with Argon gas to avoid oxidation. The natural cooling process of the oven, from 673 K to 313 K, required 2 additional hours of the films into the Argon atmosphere. Due to the annealing process, the Al-Cu/Kapton alloys suffered thermal stresses as observed by curling of the samples. The residual stress is assumed to arise from shrinkage of the substrate up to 1% at 673 K (as reported by the supplier) and by the thermal diffusion of the material [13][14][15][16]. Additionally, 50 nm thick samples of only Al and only Cu were thermally deposited to prepare Al/Kapton and Cu/Kapton systems to be used as references. Kapton was selected as the substrate material due to its high melting temperature, smooth surface, and relatively low elastic modulus. The surfaces morphology of Kapton as-received and after thermal annealing at 673 K as well as used for the metallic alloys were examined by atomic force microscopy (AFM) in tapping-mode. AFM images of 2 × 2 μm 2 size, 512 × 512 pixels 2 of resolution, and 1 Hz as scanning rate were obtained. Preliminary results of the mechanical properties of Kapton 50HN substrates and Al-Cu/Kapton alloys were reported [17] and determined from strain-stress curves obtained from a homemade universal testing machine [18] with a load cell of 220 N. The sample strain was measured by using the machine cross-head displacement, taking into account the compliance of the universal testing machine. The machine compliance, estimated as 0.16 μm/N, is small enough to avoid affectation on the elastic modulus determination. Tensile tests of samples were performed with a cross-head speed of 1 μm/s (strain rate of 6.0 × 10 −5 s −1 ). The gauge length (distance between clamps) used for the tensile testing machine was 20 mm. Samples of rectangular geometry of 40 mm length and 4 mm wide were used for tensile tests. No damages with the clamps were produced on the sample surfaces after tensile testing during sample holding. This can be observed in Figure 2, where a sample pressed by the clamps is shown (Figure 2(a)); meanwhile in Figures 2(b) and 2(c), micrographs of the nondamaged grip zone of the sample after the tensile testing can be observed with different magnifications. were examined by atomic force microscopy (AFM). Rmsroughness (R rms ) and grain size (D) before and after the tensile tests were measured using the scanning probe image processor (SPIP) software from Image Metrology Co. The microstructure of the Al-Cu alloys was examined by a scanning electron microscopy (SEM), both by secondary electrons (surface) and by backscattered electrons (to identify elemental composition on different zones of the alloy) with 25 keV using a SEM-Philips XL30 ESEM. Additionally, elemental concentration from different zones of the formed alloys was measured by energy dispersive spectroscopy (EDS).

Elastic Modulus of the Al-Cu Alloys.
Two tensile tests were carried out for each system, one with the metallic alloy/substrate and the second one without the metallic film. The first tensile test was carried out to the Al-Cu alloy bonded to the Kapton substrate to obtain the force-strain curve of the complete Al-Cu/Kapton system. The Al-Cu film alloy was then removed from the Kapton substrate by etching the alloy during 5 minutes into a weak sulfuric acid solution. This process etched the metallic film without affecting the elastic modulus of the substrate as was confirmed by the tensile testings of the Kapton substrates before and after film etching. The second tensile test was subsequently performed to the etched substrate to obtain the force-strain curve of the bare substrate. The tensile stress (σ) of the alloy as a function of applied strain (ε) was calculated by subtracting the forcestrain curve of the Kapton substrate F sub (ε), from the forcestrain curve corresponding to the Al-Cu/Kapton foil F tot (ε), according to the relation [19], where t f and w f are the thickness and width of the alloy, respectively. Force-strain curves of the thin film/polymer system, polymer substrate, and the thin film curve, used to obtain the mechanical properties of the thin films, are plotted in Figure 3. The lower dashed-line curve in Figure 3 corresponds to the curve of the thin film, which was obtained by subtracting the measured force-strain curve of the polymer substrate from the force-strain curve of the thin film/polymer system. Additionally, batches of only Al and Cu films of 50 nm thick deposited on similar Kapton substrates were analyzed by this method to estimate thier corresponding elastic modulus.

Substrate Analysis.
The surface morphology of Kapton substrate both before and after the annealing process was measured by AFM, finding no significant changes in the rms-roughness. Rms-roughness values of the Kapton surface were determined in a range of 1.1-2.9 nm, confirming that Kapton foil presents a uniform low-roughness surface which is suitable to use as a substrate for depositing alloy films with uniform thickness and low-roughness. Additionally, the mechanical properties of Kapton substrates were examined before and after thermal annealing. Figure 4 shows two stress-strain curves from Kapton foil obtained from asreceived, and after 3 h of annealing treatment. Both asreceived and annealed Kapton foil present an initial linear elastic behavior for low strain values (ε ≤ 1.3%). The rule of ε = 0.1% offset applied to both stress-strain curves yields deviations from linearity at ε = 1.3% for the as-received Kapton and ε = 0.9% for the annealed Kapton.
The corresponding elastic moduli were estimated by fitting a straight line into this region. The corresponding mean value of the elastic modulus obtained in this way for as-received Kapton is E wa = 3.0 ± 0.1 GPa, meanwhile for annealed Kapton is E a = 2.7 ± 0.1 GPa, that is, a reduction of elastic modulus of ∼10% for the annealed Kapton. This slight reduction of the elastic modulus was taken into account to estimate the elastic modulus of the Al-Cu alloys, such that the annealed curve was used for subsequent calculations. For the case of pure metallic films (Cu and Al), the elastic modulus was estimated using the mechanical properties of as-received Kapton since no annealing process is involved. Figure 5 shows the stress-strain curves corresponding to the elastic zone of only Kapton obtained after 3 h of thermal annealing and etching during 5 minutes into a weak sulfuric acid solution. The obtained stress-strain curves of the substrate show high reproducibility for the different replicates tested. For all tensile tests, a strain range of 0.9% was used to estimate the elastic modulus. Negligible variations in the estimated elastic modulus of the annealed Kapton (E a = 2.7 GPa) after etching into a weak sulfuric acid solution were observed. At this strain range, Kapton substrate is into the elastic zone and does not suffer damage that can affect the mechanical properties of the films. Once Kapton substrate was characterized, the tensile tests of the Al-Cu/Kapton system were conducted.

Phase Alloy Analysis.
The formed phases of the prepared Al-Cu/Kapton alloys were determined by XPS. XPS spectra in the 1000-0 eV energy range were obtained. Subsequent high-resolution XPS spectra of Al-2p, Cu-2p 1/2 , and Cu-2p 3/2 Journal of Nanomaterials  were determined, as is shown in Figure 6. The main Al peak (Al-2p) appearing at 74 eV (not shown) corresponds to an oxidized state of Al. The Cu spectrum consists of a set of 2p 1/2 and 2p 3/2 peaks, associated to the Al 2 Cu phase [20]. According to the Al-Cu bulk phases diagram for an Al-50 at % Cu, the expected phase is AlCu. However, Al 2 Cu phase may also be produced during the cooling process [21].

Microstructure and Morphology of Al-Cu Alloys.
The microstructure of the Al-Cu alloys was analyzed by EDS-SEM technique using the backscattered secondary electrons mode. Figures 7(a) and 7(b) show two SEM images of the alloy microstructure after the annealing time of 1 h and 3 h, respectively. As can be observed in Figure 7 Figure 7(c) shows an EDS general analysis of the sample annealed for 3 h. Additionally, evidence of diffusion of Cu or Al into the Kapton foil during the annealing process was not found, which was confirmed by the EDS analysis carried out in localized areas of the Kapton where the Al-Cu alloy was removed. Figure 8 shows representative forcestrain curves of the Al-Cu/Kapton system deposited with different film thicknesses (t f = 50 to 250 nm, nominal values) and of neat Kapton substrate as the baseline. Representative force-strain curves from four samples for each Al-Cu/Kapton system and the corresponding curve for Kapton substrate are included.

Alloys Tensile Tests.
Reported thickness values correspond to the deposited thicknesses of the formed Al-Cu alloy as measured by the quartz crystal sensor during thermal evaporation. As can be observed, the slope of the force-strain curve increases as the film thickness increases. The corresponding stress-strain curves of the Al-Cu alloys prepared with different thickness were calculated using Figure 8 and Equation (1) and are shown in Figure 9. From each curve of Figure 9, the elastic moduli of different Al-Cu alloys were estimated.
The elastic modulus values were estimated as the initial slope of the straight line of each curve in the range of 0 ≤ ε ≤ 0.6%, that is, within the linear elastic regime. The mean values and the corresponding standard deviations estimated  As known, the reported elastic modulus for Al and Cu in bulk is 70 and 130 GPa, respectively [22]. The elastic moduli of the Al-Cu film alloys fall between the bulk moduli of Al and Cu, approaching to the bulk modulus of Al as the alloy thickness increases. Additionally, the elastic moduli of 50 nm thick thermally evaporated samples of Al/Kapton and Cu/Kapton were measured for comparison. The corresponding stress-strain curves of 50 nm thick Al and Cu are shown in Figure 10. These curves were obtained from the force-strain curves shown inset. For Al films, a mean elastic modulus of 129.1 ± 4.2 GPa was obtained, a significantly higher value than its bulk value (70 GPa).
The obtained elastic modulus for Cu films was 107.1 ± 4.6 GPa, a lower value as compared with its bulk value (130 GPa). Values of 108.8 ± 11.7 GPa of the elastic modulus Journal of Nanomaterials for Cu films have been reported by other authors [23] by using an optical diffraction technique. They report this elastic modulus value for Cu in a range of 0.1 to 2.0 μm film thickness. The lower elastic modulus value obtained herein for 50 nm-thick Cu film may be explained by the lack of coalescence among grains or clusters for such a small film thickness, which seems to be particularly relevant for copper. This affirmation is also supported by the high electrical resistivity reported for Cu films of 50 nm thickness [24]. After mechanical testing, the elongated surfaces of the tested alloys were examined by SEM. SEM analysis did not show microcracks on the surface of the elongated alloys, confirming that loading was kept within the elastic regime. The surface morphology was also obtained by AFM in order to examine differences with the pristine condition (prior to testing). Figure 11 shows sequential AFM images of the Al-Cu alloys deposited with different thickness, corresponding to scenarios before and after tensile testing (images do not correspond to the same zone). As-grown images obtained from different zones of each alloy showed similar rms-roughness values, suggesting a homogeneous growth condition during evaporation.
From AFM images, the rms-roughness (R rms ) and the mean grain size (D mean ) were measured for each thickness of Al-Cu alloy both before and after tensile testing. Their corresponding mean values and standard deviations are shown in Figure 12. Both R rms and D mean show an initial increment with increased film thickness until maximum values are reached at alloy thickness of 150 nm; after alloy thickness of 150 nm, the R rms and D mean values show a slight reduction. The largest values of the roughness and grain size both before and after tensile testing were found for the 150 nm thick alloy. After deformation, the initial surface roughness value decreases, meanwhile the initial grain size increases.
After annealing (alloy formation), diffusion process and diffusion time affect the final surface of alloys given the dependence of R rms and D mean with the film thickness and formation temperature. In our case, the 150 nm thick Al-Cu film suffered the largest surface morphology changes may be due to the combined action of the alloy thickness, diffusion temperature, and diffusion time. The diffusion process discussed here for Al-Cu film formation was found to be a different behavior than the growth process, where according to the scaling laws, the rms-roughness value and the grain size of the surface increase with increased thickness (or deposition time) until a saturation is reached.

Conclusions
The elastic modulus of 50-250 nm thick Al-50 at.% Cu alloys thermally evaporated onto Kapton substrates and postformed by thermal diffusion has been investigated. Al 2 Cu phase was the dominant crystalline phase formed as determined by XPS. Force-strain curves of the Al-Cu alloys were obtained by subtracting the force-strain curve of the Kapton substrate from the force-strain curve of the Al-Cu/Kapton material system, and elastic modulus was obtained from the slope of the corresponding stress-strain curves. Elastic modulus of the Al-Cu alloys decreased when the film thickness increased, and their values were determined in the range from 106.1 to 77.8 GPa for 50 to 250 nm thick alloys, respectively. The elastic modulus of the studied Al-Cu alloys was found between the corresponding bulk values of the Al and Cu films. The elastic modulus measured for 50 nm thick Al was higher than its corresponding bulk value, while the elastic modulus of 50 nm thick Cu was smaller than its bulk value. The highest values of the mean grain size and rms-roughness were found for the 150 nm thick Al-Cu alloy. Overall, application of a tensile strain/stress yielded a tendency to increase the grain size and decrease the surface roughness of the alloys. The methodology used to obtain the elastic modulus does not yield alloy microfractures because of the small strains (<1%) applied during tensile testing.

8
Journal of Nanomaterials

Introduction
A common feature of bulk nanostructured materials [1][2][3][4][5][6][7][8][9] and intermetallic compounds [10] is relatively low ductility, which limits their usefulness as engineering materials. In the case of nanostructured materials, deformation behavior is more controlled by grain boundary migration and twinning as opposed to dislocation motion and multiplication. To improve ductility by promoting those deformation modes, research efforts have been focusing on synthesis techniques particularly severe plastic deformation [11][12][13][14][15][16][17]. Likewise, there has been an increasing interest in improving the ductility of intermetallic compounds and ordered alloys because of their unique combination of physical and chemical properties [18,19].
Recently, it has been demonstrated that long-range ordering in Ni-Mo-based alloy can be used to synthesize bulk nanostructured superlattices combining high strength and high ductility [20][21][22][23][24]. It has been shown that the closelyrelated Pt 2 Mo-type (Ni 2 Mo), DO 22 (Ni 3 Mo), and D1 a (Ni 4 Mo) superlattices coexist during long-range ordering in these systems [22,23], which form the basis for many commercial alloys with important applications in the chemical process, petrochemical and power generation industries [25]. Each of these superlattices can directly be derived from the parent face-centered cubic structure by minor atoms rearrangement on {420} fcc planes. In the case of the Pt 2 Motype superlattice, every third plane is occupied by Mo atoms and planes in-between contain only Ni atoms. Likewise, for the DO 22 superlattice, every fourth plane is occupied by Mo atoms, and every fifth plane is occupied by Mo atoms in the case of the D1 a superlattice. Long-range ordering in these systems results in considerable strengthening corresponding to room-temperature yield strength of 800 MPa or higher [22]. However, the corresponding tensile ductility ranges from less than 5% to about 40% depending upon the exact chemical composition, which determines the most stable ordered phase and its morphology. Therefore, by balancing the chemical composition it is possible to synthesize bulk nanostructured superlattices with high strength and high ductility as demonstrated in the case of the D1 a superlattice of Ni 4 Mo alloy [21,22] and Ni-Mo-Cr alloys where the metastable Ni 2 Mo with Pt 2 Mo-type superlattice is stabilized by controlled amounts of Cr [23,24]. In the case of binary Ni 3 Mo alloy, large thin platelets of Ni 3 Mo form during the earlier stages of thermal aging rendering the material extremely hard and brittle [26]. The objective of this paper is to show that bulk nanostructured DO 22 -type superlattice with high strength, high ductility, and high thermal stability can be synthesized in a Ni-Mo-Nb alloy with a composition approaching Ni 3 (Mo, Nb).

Experimental Procedure
The alloy studied had a composition of 64. Specimens for light optical metallography were etched in a solution consisting of 80% HCl and 20% of 15 mol % chromic acid to reveal the initial grain structure. Thermal aging experiments were carried out at 700 • C for up to 100 hours followed by air cooling. All tensile tests were conducted at room temperature and the fracture surfaces were examined in the secondary electron mode of a scanning electron microscope operating at 20 keV. Thin foils for transmission electron microscopy were prepared by the jet polishing technique in a solution of 30% nitric acid in methanol. All the foils were examined at an accelerating voltage of 200 keV. Since the parent fcc structure and DO 22 superlattice are closely related, all diffraction patterns are indexed in terms of the fcc structure.   Figure 1(a). Full recrystallization is indicated by the high density of annealing twins, which typifies an fcc material with relatively low stacking fault energy [27]. Similar to the case of binary Ni-Mo alloys, the alloy contained shortrange order as indicated by the appearance of diffuse intensity maxima at {1 1/2 0} positions in the <100> electron diffraction pattern of Figure 1(b). It has been reported in an earlier study that short-range order tends to lower the stacking fault energy [28]. This is found to be consistent with the observed deformation substructure.

Results and Discussion
As an example, Figure 1 dislocation to cross slip are common features of fcc materials having relatively low stacking fault energy [27]. This has an important implication on the preferred deformation mode in the ordered state, which enhances ductility as shown later. During the early stages of thermal aging (15 minutes at 700 • C) very fine precipitates were observed as shown in the example of Figure 2. The bright-field TEM image of Figure 2(a) shows that the precipitates assume the morphology of thin short platelets revealed by elastic strain contrast. At this stage, the precipitates appeared to be too fine to produce characteristic diffraction effects. Figure 2 Figure 3. A schematic illustration of the crystallographic relationship between the parent fcc lattice and DO 22 superlattice is given in Figure 3(a). The corresponding {100} reciprocal lattice intersection is shown in Figure 3(b), which is to be compared with the observed <100> diffraction pattern of Figure 3(c) obtained after 24 hours of aging at 700 • C. As can be seen, the observed diffraction pattern contains the characteristic reflections of the DO 22 Figure 4(a). It is observed that the superlattice assumes the morphology of nanosized particles with an average size of 10-20 nm. Also, it is observed as the aging time was extended from 24 hours (Figure 4(a)) to 100 hours (Figure 4(b)), the change in superlattice morphology was insignificant demonstrating its relatively high thermal stability. The corresponding effect of short-range order → longrange order transformation on the room temperature tensile properties is summarized in Figure 5. After 1 hour of aging at 700 • C, the 0.2% yield strength was nearly doubled reaching about 800 MPa. However, the material retained about 80% of its initial ductility in the short-range ordered state. With continued aging up to 100 hours, the material maintained a combination of high yield strength of about 820 MPa and high tensile ductility of about 35% consistent with the observed high thermal stability of the DO 22 superlattice.
Comparative true tensile stress-strain diagrams in the short-range and long-range ordered states are shown in Figure 6. It is noted that the strain hardening rate is significantly lowered by long-range ordering to the DO 22 superlattice reflecting a change in the deformation behavior. In contrast with the short-range ordered state where deformation behavior typifies an fcc material with relatively low stacking fault energy as demonstrated in Figure 1(c), twinning on {111} planes is found to be the predominant deformation mode in the long-range ordered state as illustrated in the example of Figure 7. This indicates that although long-range ordering could suppress most of the slip systems characteristic of the parent fcc lattice, many of the twinning systems remain active suggesting that the stacking fault energy of the material is further lowered by long-range ordering enhancing deformation by twinning. A similar behavior is found to occur in ordered Ni-Mo [21,22] and Ni-Mo-Cr [23,24] alloys. Figure 8 shows secondary electron SEM images illustrating the morphology of tensile fractures surfaces in the annealed (short-range ordered) and aged (long-range ordered) states. Typical of ductile materials, in both cases, fracture occurred by dimple-type rupture. However, in the long-range ordered state, the dimples can be seen to be finer and shallower, which could be related to higher density of crack nucleation sites as well as relatively lower ductility in the long-range ordered state.

Conclusion
It is concluded that a bulk nanostructured DO 22 superlattice with high strength, high ductility, and high thermal stability can be synthesized in a Ni-Mo-Nb alloy with composition approaching Ni 3 (Mo, Nb) by a simple aging heat treatment at 700 • C. Upon thermal aging, the grains of the hightemperature fcc phase are subdivided into ordered crystals on the nanoscale (10-20 nm) with room-temperature yield strength of about 820 MPa and tensile ductility of 35%. Plastic deformation in the ordered state is found to predominantly occur by twinning on {111} planes of the parent fcc structure indicating that the superlattice preserves the twinning systems of the parent phase leading to the observed high ductility.

Introduction
One-dimensional (1D) metal nanostructures, including rods, wires, and tubes, have recently attracted interest and been applied to optical-electronic and sensing nanodevices due to their unique optical and conductive properties and thermal conductivity [1][2][3][4][5]. Among the various metals, silver exhibits effective thermal conductivity and the highest electrical conductivity and can be synthesized into Agbased compounds with different compositions. Silver is more reactive than gold or platinum; therefore, silver is the most favorable candidate for various applications [6,7]. For the silver oxide system (Ag x O), silver can be reduced by a thermal decomposition reaction due to its low melting point and active energy [8].
One-dimensional (1D) silver nanostructures are used as the substrates to detect surface enhanced Raman scattering (SERS) and surface plasmon resonance (SPR) and for enhancing the photoluminescence of the ZnO system [9][10][11][12][13][14]. Over the last decade, many methods have been used to fabricate Ag nanowires and nanorods. Xu et al. used porous alumina templates to carry out chemical deposition for the fabrication of silver nanorod arrays [15]. Chen et al.
fabricated Ag nanowires via a polybasic process [16]. Murphy and Jana demonstrated a seed-mediated method for obtaining Ag nanowires [17]. All the aforementioned methods require a chemical solution or template to fabricate the Ag 1D nanostructure. To obviate the chemical solutions and reduce complexity, Zhao et al. grew Ag nanorod arrays using oblique angle deposition [18]. Mohanty et al. used vapor-solid phase synthesis to produce Ag nanowires [19]. Besides the aforementioned fabrications, a few literatures have demonstrated the formation of the various nanostructures by sputtering [20][21][22][23], and oxide-assisted growth has been developed for one-dimensional nanostructures in previous studies [24][25][26]. Notably, oxidation reduction growth (ORG) of 1D metal structures prepared using a sputtering method has never been reported. Additionally, when the bulk metal reduces to nanoscale, the luminescence will affect the electron-hole recombination due to the larger aspect ratio [27,28].
This study has demonstrated an oxidation reduction growth (ORG) technique with mixed-gas sputtering to create Ag nanorod arrays without any chemical solutions or contamination from aqueous solution via oxide-assisted growth [29]. The ORG methodology is used to deposit an Ag buffer layer with silver oxide nanoclusters to obtain Ag nanorods arrays using a two-step mixed gaseous process. The success of the technique provides support for the oxidation reduction growth (ORG) mechanism and proves suitable for fabrication of Ag nanorods in the semiconductor industry.

Experiment
Prior to sputtering, a p-Si (100) substrate was first cleaned in acetone and then cleaned in isopropanol for 300 s. The p-Si substrate was then dipped into hydrofluoric acid to remove the native oxide and finally rinsed with deionized water. For RF sputtering, the cleaned p-Si was loaded into the sputter chamber to perform sputtering, and the distance between the sample surface and silver target was approximately 3 cm. The RF sputter system was evacuated to below 1.8 × 10 −4 Torr, and argon gas was introduced with the flow rate controlled at 15 sccm. The gas pressure was fixed at 15 × 10 −3 Torr during the sputtering process, and the RF power was set at 80 W. The sputtering times are 20 and 30 min for silver oxide clusters and silver nanorods, respectively. The deposition rates of silver, silver oxide, and nanorods are 3, 1, and 10 nm/min, respectively. The amount of the introduced oxygen is 1 sccm during reactive sputtering.
A schematic diagram for the oxidation reduction growth of silver nanorods is illustrated in Figure 1. The whole sputtering process was continuous without venting until Ag nanorods had grown. In Figure 1(a), at the beginning of sputtering, a thin silver film was predeposited on the Si surface with a constant flow of 15 sccm argon gas. Next, a trace of oxygen was introduced and mixed with argon as the source of plasma to form the silver oxide seed during sputtering, as shown in Figure 1(b). The silver oxide seed randomly dispersed over the Ag buffer layer and was able to form the contiguous thin film due to the trace amounts of oxygen. Notably, the melting point of silver oxide is normally approximately 300 • C and even lower at the nanoscale [8]. When engaged in the sputtering process, the temperature of the sample will rise to between 200 • C and 300 • C allowing silver oxide to dissolve and release O 2 and return the silver to the original silver oxide site (Figure 1(c)). In Figure 1(d), sputtering is performed without oxygen, and the silver nanorods start to grow from the reduction metal silver site. The oxidation reduction growth (ORG) of silver nanorods is a novel and continual process, which uses sputtering.
The morphology of the Ag nanorods was investigated by scanning electron microscopy (SEM), and the crystallization of the silver nanorods and silver oxide seeds were examined using X-ray diffraction (XRD). The incident angle of Xray was fixed 1 • for the phase identification of the extreme surface. For the detailed structure, the orientation was detected by transmission electron microscopy (TEM), and the selected area electron diffraction (SAED) was observed on the body of the nanorods and the silver buffer layer. For PL measurements, a 405 nm laser was used as an excitation light, and detection was performed in a photomultiplier tube. Figure 2 shows tilted view SEM images and glancing angle XRD for the silver oxide seeds under sputtering 20 min and Ag nanorods under sputtering 30 min. As shown in Figures 2(a) and 2(b), not only were the silver oxide seeds observed but also the Ag nanorods uniformly distributed over the sample surface. Notably, the density of the silver oxide seeds was approximately 150/um 2 , approaching that of Ag nanorods. The range for the diameters of the Ag nanorods was from 40 to 55 nm while the diameter of the silver oxide seeds was approximately 60 nm. The average diameter of the Ag nanorods was slightly smaller than that of the silver oxide seeds.

Results and Discussions
The energy dispersive spectrometer (EDS) data of the Ag nanorods is exhibited in Figure 2(c). A main peak of silver and a weak oxygen signal can be observed on the spectrum. According to the ratio of weight between Ag and oxygen, most of the peaks were Ag, and these were initially identified as pure Ag nanorods. The inset of Figure 2(c) is a cross section view of the Ag nanorods using a focused ion beam and most of the Ag nanorods aligned to the substrate. A silver buffer layer existed beneath the Ag nanorods, and this result is consistent with previous reports (not the present growth method) that one-dimensional aligned nanostructures also require the same material layer to grow due to the lattice match [30][31][32][33]. The thickness of the present Ag buffer layer was about 45 nm, and the length of the Ag nanorods was from 250 to 300 nm.
The crystallization of both the silver oxide nanoseeds and Ag nanorods was investigated using glancing angle XRD. Figure 2(d) shows the main Ag peaks, including (111), (200), (220), and (311), which coincide with the thin silver film [34]. In Ag nanorods case, all signals really came from Ag nanorods due to glancing angle XRD with 1 • incidence, but the peaks of Ag signal in silver oxide nanoseeds resulted from the Ag buffer layer. Notably, the signal of silver oxide (11 − 1) appears at 28 • in some of the silver oxide nanoclusters in addition to the Ag peaks. The silver oxide phase was Ag 3 O 4 and was produced during the mixed-gas sputtering process with the input of oxygen [35,36]. After the thermal Journal of Nanomaterials   decomposition, the oxygen resulting from the reduction reaction of silver oxides was pumped out. In addition, the production of silver oxides reduced metallic silver seeds for the growth of Ag nanorods. However, when the Ag 3 O 4 phase is compared with the AgO phase, the Ag 3 O 4 phase is more thermodynamically unstable (lower free energy) [37,38]; therefore, the Ag 3 O 4 can easily react in the consequent sputtering process to induce the growth of Ag nanorods with a chemical reaction of Ag 3 O 4(s) → 3Ag (s) + 2O 2(g) [39]. This two-step oxidation reduction process is responsible for the growth of the nanorods. Therefore, the present oxidation reduction growth (ORG) mechanism can also be applied to other metal systems with low melting point [40]. Figure 3(a) shows TEM observations with Ag nanorods aligned to the Ag buffer layer, and the corresponding SAED pattern from the white circle zone is shown in Figure 3(b). According to the electron diffraction pattern, the Ag nanorods were ascertained to be the single-crystal structure; the high resolution TEM image is shown in Figure 3(c). Additionally, the Ag nanorods grow from the thermal-reduced Ag nucleus above the Ag buffer layer. A lattice image of the interfacial layer between the Ag nanorods and the grain is shown in the Figure 3(d), which demonstrates a perfect lattice arrangement. Figure 3(b) also shows that the Ag nanorods are crystalline and the interface zones between the Ag nanorods and the thermal-reduced Ag nucleus are coherent. From the cross-section observation, the Ag buffer layer was composed of many silver grains and this result was coincident with the phases of XRD spectra ( Figure  2(d)). Notably, the Ag nanorods of the ORG structure also possessed excellent properties.
A PL spectrum of the Ag nanorod arrays under the irradiation of a 408 nm laser could be seen in Figure 4. The spectra ranges of all samples, which exhibit the strongest intensity near 2.17 eV, were from ∼1.9 to ∼2.45 eV. Basically, PL peaks in the visible wavelength can be regarded as a radiative recombination between Fermi level electrons and sp-or d-band holes [39,41]. However, silver oxide forms easily in air and it may also provide a possible route of luminescence for photoreduced metallic Ag under visible illumination [27,42]. Herein, PL spectrum of Ag nanorods was also measured under the vacuum to avoid the formation of silver oxide during the irradiation of laser in air, and there is no shift for the spectra of Ag nanorods both in air and in vacuum.

Conclusion
In summary, Ag nanorods were fabricated using an oxidation reduction growth (ORG) method by sputtering without catalysts or chemical solutions. The Ag nanorods grew in the original locations of reduced metal nuclei after thermal composition of silver oxide nanoclusters. The Ag nanorods stood vertically on the Ag buffer interlayer and grew from the interface between the Ag grains and the Ag buffer interlayer. PL spectra of Ag nanorods both in air and vacuum were observed at 2.17 eV due to the photoactivation caused by a radiative recombination of Fermi level electrons and d-band holes.

Introduction
Recently, considerable effort has been exerted in developing high-κ rare earth oxide, M 2 O 3 (M = La, Pr, Nd, etc.), as a replacement of the conventional SiO 2 -based gate dielectric material [1]. The incorporation of neodymium (Nd) ions in insulating layers has important applications for solid-state laser materials, luminescent materials, protective coatings, and gate dielectric applications [2,3]. However, Nd 2 O 3 is thermally unstable upon annealing and can be partially transformed to NdO(OH) when exposed to atmospheric conditions [4]. One of many solutions to enhance the thermal stability is the incorporation of aluminium (Al) to develop innovative multifunctional advanced lanthanidealuminates-based ceramics, MAlO 3 (M = La, Pr, Gd, and Nd). The lanthanide aluminates are promising high-κ candidates as they combine the advantages of the high permittivity of the lanthanide oxide with the chemical and thermal stability of Al 2 O 3 . Furthermore, they remain amorphous up to high temperatures, leading to a large reduction in leakage current relative to polycrystalline M 2 O 3 films during CMOS processing [5,6]. Work on NdAlO x was mostly reported as a ceramic material for microwave applications and as a diffusion barrier in solid-oxide fuel cells [7,8]. Growth of NdAlO x thin films have previously been achieved by various deposition methods, including pulsed laser deposition [9], chemical vapor deposition [10,11], e-beam evaporation [8], and atomic layer deposition [4]. To date, however, little is still known about the physical and electronic characteristics of NdAlO x due to a lack of suitable precursors with appropriate stability, volatility, and decomposition characteristics. This has motivated us to further exploit these perovskite thin films for gate dielectric applications using an alternative singlesource precursor. The use of single-source precursor allows better mixing of the components at atomic level, significantly lower decomposition temperature, and free from halide ions contamination [11].
In this work, the effects of high-temperature postdeposition annealing (PDA) on the properties of the NdAlO x thin films, deposited by metalorganic chemical vapor deposition (MOCVD) using single-source precursor, were studied.

Experimental
Near stoichiometric NdAlO x thin films (Nd/Al = 0.87) were deposited on n-type silicon (100) substrates by liquid injection MOCVD at 450 • C on an Aixtron AIX 200FE AVD reactor fitted with the "TriJet" TM liquid injector system [12], utilizing the single-source precursor [NdAl(OPr i ) 6 (Pr i OH)] 2 . Selected films were subjected to high-temperature (750-950 • C) postdeposition annealing (PDA) in pure nitrogen (N 2 ) ambient for 60 seconds. Subsequently, a postmetallization forming gas anneal (FGA) was carried out at 400 • C for 30 minutes using H 2 : N 2 in the ratio 1 : 9, together with a control as-deposited sample.
X-ray diffraction (XRD) was performed on the studied films using nickel-filtered Cu Kα radiation (λ = 1.5405Å) with a 2θ increment of 0.2 • per second. The samples were scanned over a θ/2θ range of 20 • to 40 • . Medium energy ion scattering (MEIS) experiments were carried out using a nominal 200 keV He + ion beam and a 70.5 • scattering angle. Cross-section transmission electron microscopy (TEM) was carried out using a JEOL 2000 FX operated at 500 kV. Capacitance-voltage (C-V) measurements were conducted on the MOS capacitors of the structure (Au/NdAlO x /SiO 2 /n-Si) using a HP4192 impedance analyzer, with 30 mV RMS probe signal, at various frequencies (1 kHz-1 MHz). Leakage current (I-V) measurements were obtained using a Keithley K230 programmable voltage source and a 617 type electrometer.

Results and Discussion
Phase transitions of the gate dielectric, as a function of PDA temperatures, were accessed by X-ray diffraction (XRD). The X-ray diffraction traces of the as-deposited and PDA samples (regardless of PDA temperature) exhibited a diffraction peak consistent with the (200) peak from the silicon substrate. No other diffraction features were observed (Figure 1(a)), suggesting that they were essentially amorphous. In addition, MEIS results indicated that no significant level of crystallinity or movement of metal ions was in evidence in the 950 • C PDA film as demonstrated in Figure 1(b). Figure 2 shows TEM micrographs in which the thicknesses of the high-κ stack, including the interfacial layer, were evaluated. The high-κ thickness and a thin native oxide interlayer, adjacent to the silicon substrate, changed from 11 nm and 1.5 nm, respectively, and to 10.4 nm and 2.5 nm respectively, after 950 • C PDA. This could be due to interdiffusion of oxygen between SiO 2 and NdAlO x . The growth of the SiO 2 layer, following annealing, is also visible in the MEIS energy spectrum (Figure 1(b)). The amorphous nature of the thin films is also observed in the TEM analysis ( Figure 2) confirming the XRD findings. The high-frequency C-V characteristics of as-deposited and PDA samples are shown in Figure 3.  The inset is a plot of capacitance equivalent thickness (CET) versus NdAlO x physical thickness, which shows an increase in the dielectric permittivity (7 to 12) and changes of the interfacial layer (1.5 to 2.5 nm) after PDA treatment.
during annealing at the NdAlO x /SiO 2 interface. Terman analysis [13] yields an interface density of states, D it , of 4.94 × 10 11 cm −2 eV −1 and 4.01 × 10 11 cm −2 eV −1 at midgap for as-deposited and 950 • C PDA samples, respectively. This is lower than D it of other recent high-κ candidates, D it of 2.5 × 10 12 cm −2 eV −1 found in La 1.3 Hf 1.0 O 4.1 [14], but higher than the interface state density of 5 × 10 10 cm −2 eV −1 shown by Y x Hf 1−x O y (x = 0.065) [15]. Despite an increase of the interfacial layer in the PDA samples, the measured capacitance in strong accumulation was found to be higher than that of the as-deposited samples. The inset of Figure 3 shows a plot of the high-κ dielectric thickness against capacitance equivalent thickness (CET). The slope revealed the NdAlO x dielectric permittivity (κ) to be 7 and 12 in the as-deposited and 950 • C PDA films, respectively. The increase of dielectric permittivity (κ), observed after receiving thermal treatment, could be due to a small change of the crystal symmetry, with the formation of some small nanometer scale crystallites with higher permittivity phases of NdAlO x [16,17].
The leakage current densities (J) at 2 MV cm −1 in all samples, even after 950 • C PDA, were below 1 × 10 −7 A cm −2 (Figure 4), which is comparable with other leading edge high-κ dielectrics [18]. The average breakdown, regardless the annealing treatment, also occurs at the equivalent field strength of 7 MV cm −1 .

Conclusions
The NdAlO x thin films deposited by liquid injection MOCVD have been shown to remain amorphous up to 950 • C as shown by XRD analysis. No significant level of crystallinity or movements of metal ions was also in evidence after annealing at 950 • C as indicated in MEIS energy spectra. Electrical properties of NdAlO x samples, after high-temperature annealing, were presented. Good electrical integrity was maintained even after 950 • C PDA as shown by C-Vand I-V results, showing the extracted dielectric permittivity of 12, a low leakage density of 7 × 10 −7 A cm −2 at 2 MV cm −1 , and a density of interface states at flatband D it of 4.01 × 10 11 cm −2 eV −1 . These features 4 Journal of Nanomaterials make the neodymium aluminate a potential candidate for the dielectric replacement.