Structural and mechanical properties of Al―Mg―B films: Experimental study and first-principles calculations

    Structural and mechanical properties of Al-Mg-B films: Experimental study and first-principles calculations V.I. Ivashchenko, P.L. Scrynskyy, S.N. Dub, O.O. Butenko, A.O. Kozak, O.K. Sinelnichenko PII: DOI: Reference:

S0040-6090(15)01309-7 doi: 10.1016/j.tsf.2015.12.059 TSF 34928

To appear in:

Thin Solid Films

Received date: Revised date: Accepted date:

30 September 2015 22 December 2015 24 December 2015

Please cite this article as: V.I. Ivashchenko, P.L. Scrynskyy, S.N. Dub, O.O. Butenko, A.O. Kozak, O.K. Sinelnichenko, Structural and mechanical properties of Al-Mg-B films: Experimental study and first-principles calculations, Thin Solid Films (2015), doi: 10.1016/j.tsf.2015.12.059

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ACCEPTED MANUSCRIPT Structural and mechanical properties of Al-Mg-B films: Experimental study and first-principles calculations V.I. Ivashchenko1, P.L. Scrynskyy1, S.N. Dub2,

Frantsevych Institute for Problems of Material Science, NAS of Ukraine, 3, Krzhyzhanovsky

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O.O. Butenko1, A.O. Kozak1, O.K. Sinelnichenko1

Str., 03142 Kyiv, Ukraine 2

Bakul Institute for Superhard Materials, NAS of Ukraine, 2, Avtozavodska Str., 04074 Kyiv,

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Ukraine

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The Al-Mg-B films were deposited on silicon substrates by direct current magnetron sputtering from the AlMgB14 target at low discharge power and at substrate temperature ranging from 100 to 500 C. The deposited films have been annealed at 1000 °C in vacuum, and

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characterized by X-ray diffraction, atomic force microscopy, Fourier transform infra-red

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spectroscopy, nano- and micro-indentation, and scratch testing. The films exhibit lower hardness than the bulk AlMgB14 material, which is due to their amorphous structure in which the strong

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intra-icosahedron B-B bonds are almost lacking and the weaker B-O bonds are predominant. After the annealing, a reduction of a number of B-O bonds and a formation of crystallites in the films lead to an increase in the nanohardness and elastic modulus. The as-deposited films exhibit

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a low coefficient of friction of 0.08-0.12. First-principles studies show that the icosahedra in amorphous AlMgB14-based materials are not fully developed, which is the reason of their lower mechanical performance.

Keywords: AlMgB14, magnetron sputtering, chemical bonding, hardness, tribological properties, first-principles study.

1.

INTRODUCTION

The hard material AlMgB14 (BAM) was first synthesized and investigated by Matkovich and Economy [1]. Later on, the structure of BAM was refined with more advanced diffraction techniques by Higashi and Ito [2]. They determined the orthorhombic 64-atom unit cell (space group Imma, No. 74) that consisted of four icosahedra (B12 units) and eight inter-icosahedron boron atoms. The metal atoms are located in interstices between the boron icosahedra and are

ACCEPTED MANUSCRIPT bonded to the icosahedra via the inter-icosahedron boron atoms. However, the metal atoms sites were not completely occupied: the occupancy of the aluminum and magnesium sites was about 75% and 78%, respectively. This corresponds to a composition for the material of Al0.75Mg0.78B14 [2]. The BAM materials were mainly synthesized using a flux growth process

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[1,2] and by hot pressing of powder [3].

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The bulk BAM materials have been intensively studied during last decade owing to their high

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hardness of 25-35 GPa [3,4], low friction coefficient and chemical inertness [3,5]. Incorporation of oxygen during and after the processing strongly influences the composition and mechanical properties of the final product.

In contrast to the bulk materials, the investigation of the films based on BAMs is in its

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infancy. In Table 1 we summarized the information on the deposition methods and mechanical properties of the BAM based thin films. One can see that the nanohardness (H) and elastic

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modulus (E) of the films are very sensitive to the deposition conditions and method used, changing in the range of 7-51 GPa and 80-300 GPa, respectively. The films were deposited at substrate temperature, TS, varying from room temperature to 600 °C. A systematic investigation

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of an effect of TS on the structural and mechanical properties of the films has not been carried out

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so far. There is also a contradiction regarding the structure of the films: the studies [5-10] showed that the films were amorphous and the annealing at 1200 °C did not lead to

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crystallization [5], whereas the authors [11] found that the films deposited by using a low discharge power regime (~1.1 W/cm2) at TS=400 °C were nanocrystalline. In this work we aimed to obtain additional information about the effect of substrate

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temperature on the structure and properties of the films. In order to interpret the properties of the films, first-principles investigations of BAM compounds were carried out.

2.

EXPERIMENTAL AND THEORETICAL PROCEDURES USED

2.1 Preparation of the films The films were deposited by magnetron sputtering from an AlMgB14 target (a disk with the diameter and thick of 72 mm and 4 mm, respectively, made in the company “China Material Technology Co., Ltd.”; purity 99.9 %) at different substrate temperatures of 100, 200, 350, 450 and 500 C. The substrates were polished Si (100) wafers. Before the deposition, the silicon wafers were cleaned in the bath with a 5 % HF solution to remove the native oxide. Afterwards, the substrates were rinsed in de-ionized water and dried in nitrogen. Finally, they were sputteretched in argon plasma in the chamber prior to the deposition. The substrate bias was -50 V. The argon flow rate, pressure and DC discharge power were 51 sccm, 0.17 Pa and 1.2 W/cm2,

ACCEPTED MANUSCRIPT respectively. The base pressure was approximately 10-3 Pa. The film deposited at 450 C has been annealed in vacuum (~10-3 Pa) at 1000 C for two hours.

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2.2 Film Characterization

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X-ray diffraction (XRD) investigations of the films have been done using a diffractometer

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“DRON-3M”. Their surface has been studied by means of atomic-force microscope (AFM) “NanoScope IIIa Dimension 3000TM”, and the chemical bonding by Fourier transform infrared spectroscopy (FTIR) using a “FSM 1202” LLC “Infraspek” spectrometer. The chemical states has been studied by an EC 2401 X-ray photoelectron spectroscopy (XPS, USSR) using Mg K

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X-ray radiation (E=1253.6 eV). The Au 4f7/2 and Cu 2p3/2 peaks with binding energy of 84.00.05 eV and 932.660.05 eV, respectively, were used as a reference. The nanoindentation

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was carried out with the help of a device G200 equipped with the Berkovich indenter. Nanohardness and elastic modulus have been determined using the Oliver and Pharr procedure [12]. Knoop hardness (HK) was determined by a device “MICROMET 2103 Microhardness

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Tester” (BUEHLER, USA) at a load of 10 mN. The thickness of the films was estimated by a

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"Micron - alpha" optical profilometer (Ukraine). The film thicknesses approximate 0.7-0.8 m, and slightly decrease with increasing TS. The scratch tests were performed with a “Micron-

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gamma” scratch tester (Ukraine) using a Vickers diamond pyramidal tip with a linear scan rate of 9 m/s and a ramping load from 0 to 0.3 N.

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2.3 Computational aspects To investigate phase stability and to explain the properties of AlMgB14 and Al0.75Mg0.75B14 structures (we would like to remind that the Al0.75Mg0.78B14 phase was detected in the experiment [2]) we singled out the initial orthorhombic 64-atom supercell with the composition of Mg8B56 (space group Imma, No. 74) [13]. The initial structures Al4Mg4B56 and Al3Mg3B56 were constructed by means of the substitution of the Mg atoms by Al atoms. The first-principles calculations were performed with the quantum ESPRESSO code [14] using periodic boundary conditions and the generalized gradient approximation (GGA) of Perdew, Burke and Ernzerhof (PBE) [15] for the exchange-correlation energy and potential, Vanderbilt ultra-soft pseudo-potentials were used to describe the electron-ion interaction [16]. The criterion of convergence for the total energy was 10-6 Ry/formula unit (1.36·10-5 eV/formula unit). In order to speed up the convergence, each eigenvalue was convoluted by a Gaussian with a width σ = 0.025 Ry (0.34 eV). The cut-off energy of 36 Ry (489.6 eV) was used.

ACCEPTED MANUSCRIPT The initial structures were optimized by simultaneously relaxing the supercell basis vectors and the atomic positions inside the supercell using the Broyden-Fletcher-Goldfarb-Shanno (BFGS) algorithm [17]. The relaxation of the atomic coordinates and of the supercell was considered to be complete when atomic forces were less than 1.0 mRy/Bohr (25.7 meV/Å),

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stresses were smaller than 0.05 GPa, and the total energy during the structural optimization

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carried out using the (6 4 3) Monkhorst-Pack mesh [18].

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iterative process was changing by less than 0.1 mRy (1.36 meV). The geometry optimization was

The quantum molecular dynamics (QMD) calculations of the initial relaxed structures were carried out with fixed unit cell parameters and volume (NVT ensemble, constant number of particles-volume-temperature). First, the structures were let to evolve at 4500 K for 4 ps, and

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then the melts were cooled down to 300 K during 20 ps. In the QMD calculations, the time step was 20 atomic units (about 10-15 s). The system temperature was kept constant by rescaling the

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velocity, and the variation of the total energy was controlled during each temperature step. In the large-scale QMD simulation, the chosen reduced k-points mesh (2 2 2) were used in order to save computing time without compromising accuracy. The justification of such an approach was

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validated in Refs. [19,20]. After QMD equilibration, the geometry of the resultant structures

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were optimized by simultaneously relaxing the supercell basis vectors and the positions of atoms inside the supercells using the BFGS algorithm [17].

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The above-described pseudo-potential procedure was used to study the phonon spectra of both the crystalline and amorphous phases. The phonon densities of states (PHDOSs) of the largescale crystalline and amorphous systems were calculated using a procedure based on force

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autocorrelation function (FACF). We modified the procedure based on the Fourier transform of the velocity autocorrelation function [21] to use forces instead of velocities, since in “Quantum ESPRESSO” code the forces are the output information [22]. The FACF-PHDOSs were determined using QMD simulations for 3 ps at 10 K in the NVT ensemble. The shear stress of the structures was calculated as follows: first, an incremental shear distortion was imposed, then the dihedral angle between the supercell vectors that corresponds to the given shear was fixed, and simultaneously the basis supercell vectors and the atomic coordinates were relaxed. For both tensile and shear strains the structural parameters at a previous step were used to calculate the Hellmann-Feynman stress for the next step.

3.

RESULTS AND DISCUSSONS

3.1 Experimental investigations

ACCEPTED MANUSCRIPT We investigated the AFM images of the films deposited at low and high temperatures (not shown here). The roughness RMS determined at the area of 5 m  5 m was 1.6 nm and 1.0 nm for films deposited at 100 C and 500 C, respectively. These data show that the roughness of the films surfaces is low, and it decreases with increasing substrate temperature.

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In Fig. 1 we show the XRD patterns of a film deposited at 450 °C and annealed at 1000 °C. One

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can see that the as-deposited films are X-ray amorphous whereas after annealing it contains a

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crystalline phase. The Bragg reflexes of the annealed film around 2  12° and 25° correspond to the rhombohedral or tetragonal boron [PDF: 011-0618, 012-0377, 031-0206, and 031-0207]. These reflexes can also be attributed to the AlB12 phase [PDF: 012-0640]. Thus, the annealing at 1000 C

related to boron oxides were not observed.

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leads to the formation of the boron and AlB12 crystallites in the amorphous matrix. Crystallites The FTIR spectra of the films deposited at different TS are shown in Fig. 2. The numbers

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denote the wavenumbers of the absorption bands related to the B-O-B bending and B-O stretching vibrations in the tetragonal BO4 and trigonal BO3 units [23]. As will be shown below, the inter-icosahedron vibrations can also contribute to the absorption band at 1055-1100 cm-1

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[8,9]. It is seen that an increase in TS leads to reduction of the intensity of B-O vibrations.

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In Fig. 3 we show the XPS core-level spectra of the films deposited at 450 °C and annealed at 1000 C. To identify the peaks in the measured spectra we used the literature data [5,9,24-29].

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The XPS Al 2p spectra show the one main peak that can be assigned to Al-O bonding [24]. The shift of this peak towards high binding energies (BE) after annealing is considered as the strengthening of the Al-B bonds [25]. In the Mg 2s spectra, the Mg-O bonds are predominant

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[5,26]. There are a peak and a shoulder in the B 1s spectra. The shoulder at about 192 eV originates from B-O bonds [5], whereas the peaks around 188.6-189.1 eV are characteristic for Al-borides [9]. It is seen that the annealing leads to strengthening the borides and B-B bonds. The peak in the O 1s spectrum of the as-deposited film corresponds to boron oxide [27]. After annealing, the part of oxygen released from the B-O bonds forms new Mg-O bonds [28]. As a result, after annealing, the O 1s spectrum shifts towards higher binding energy and its intensity increases. The characteristic feature of the Al 2p, Mg 2s and B 1s spectra is the shift of the mean peaks towards high binding energies after annealing. These findings show that there are Al-B and pure B units, as well the fragments of the Al2O3, MgO, B2O3, MgAl2O4 and possibly other oxide configurations in the amorphous films, and the annealing promotes the formation of new metal-boron, boron-boron and magnesium–oxygen bonds. Using the measured XPS spectra, we estimated the elemental composition of the films. The content of Al, Mg, B and O was approximately 6.1 at.%, 1.9 at.%, 60.0 at.%, 32.0 at.%, respectively. After annealing, the chemical composition changed insignificantly: the boron content decreased by 2.3 at.%, and the

ACCEPTED MANUSCRIPT oxygen content increased by 2.1 at.%. Oxygen is abundant in the deposited films due to absorption of water and oxygen species during sample exposure to air [3,5]. Returning to the results of the XRD and FTIR investigations (cf. Figs. 1 and 2), we see that these results are consistent with the XPS data: all the investigations predict the presence of the

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B-O, Al-O, Mg-O, B-B and Al-B bonds in the as-deposited amorphous films. The B-B, Al-B and

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Mg-O bonds strengthen, and the B-O bonds weaken after annealing.

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The nanohardness and elastic modulus of the deposited and annealed films are presented in Fig. 4 as functions of indentation depth. The dependences of the values of H and E, as well as HK on substrate temperature are shown in Fig. 5. The nanohardness and elastic modulus of the amorphous Al-Mg-B films are smaller as compared to those of the silicon substrate (12.5 GPa

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and 140 GPa, respectively). An exception is found for the high temperature and annealed films. For these films, nanohardness and elastic modulus are comparable with those of the silicon

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substrate. Figures 4 and 5 show that the nanohardness, elastic modulus and Knoop hardness increase with the deposition temperature. The annealing also promotes an increase of these characteristics. The Knoop hardness is higher than the nanohardness by approximately 40-50 %.

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We suppose that the observed difference between H and HK is due to the different manners of

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indentation – the static indentation for the determination of the Knoop hardness and the dynamic indentation for the determination of the nanohardness. It follows that the deposited amorphous

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Al-Mg-B films exhibit lower hardness than that of the bulk BAM. This is due to the amorphous nature of the films in which the strong B-B bonds are almost absent and the weaker B-O bonds dominate (cf. Fig. 2 and 3). The observed increase of the mechanical characteristics with

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substrate temperature and annealing is assumed to be due to strengthening the B-B and Al-B network. After annealing, the appearance of the B and AlB12 crystallites also promotes the improvement of the mechanical properties of the films. The scratch tests showed that the low temperature film delaminated at a load of 0.2 N, whereas other as-deposited films exhibited good adhesion to silicon substrates at a load above 0.3 N. In Fig. 6 we show the friction force as function of load in the adhesion region for the films deposited at TS=200 C and 500 C. One can see that the roughness of the high temperature film is slightly better compared to that of the low temperature films in agreement with the results of the AFM measurements. The friction coefficients extracted from the dependences presented in Fig. 6 is about 0.12 and 0.08 for the films deposited at low and high Ts, respectively.

3.2 Theoretical investigations

We have investigated the crystalline (c) and amorphous (a) AlMgB14 and Al0.75Mg0.75B14

ACCEPTED MANUSCRIPT structures to try to explain the experimental findings presented above. We found that the formation energies of both compounds were negative, and the formation energy for cAl0.75Mg0.75B14 is lower by 0.011 eV/atom as compared with c-AlMgB14. Based on this finding, one can expect that the BAM compounds will be formed with deficiency of the metal atoms in

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agreement with the experimental [2] and theoretical [30] investigations.

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In Fig. 7 we show the atomic configuration of a-Al0.75Mg0.75B14. An analysis of the atomic

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configuration presented in Fig. 7 shows that: i) the icosahedra in the amorphous sample are not fully developed (cf. Fig. 7b) and ii) the number of boron atoms between the incomplete icosahedra in the amorphous phase is higher than in the crystalline phase. These peculiarities of the atomic configuration of the amorphous sample are reflected in the phonon density of states

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(PHDOS) shown in Fig. 8. Firstly we note that for the crystal phases an acceptable agreement between the calculated and experimental spectra is observed, which validates the computational

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procedures. For the amorphous sample, the intense new vibrations related to the intericosahedron boron atoms are observed in the range of 1050-1100 cm-1. One can see that these vibrations are absent in the crystalline structures. We suppose that these vibrations form the

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absorption band at 1050-1100 cm-1 in the infrared spectra of the deposited amorphous films (cf.

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Fig. 2). It should be noted that the same band was detected in the FTIR spectra of other amorphous Al-Mg-B films [8,9].

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In order to estimate the ideal strength of the BAM materials we calculated the stress-strain relation in shear for the weakest shear plane (see Fig. 9). The ideal shear strength decreases in the sequence c-Al0.75Mg0.75B14 – c-AlMgB14 – a-AlMgB14. For c-Al0.75Mg0.75B14, the shear ideal

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strength of 24 GPa is very close the hardness of 26-28 GPa determined for a single crystal of cAl0.75Mg0.78B14 [32]. The amorphous sample has the smallest shear ideal strength, and, therefore it possess also the lowest hardness. Let us comment the experimental results on the hardness of the amorphous films summarized in Table 1. We suppose that the high hardness of the amorphous films can be assigned to the availability of the sub-nanometer icosahedra embedded in the randomized B-B matrix. The icosahedra can aggregate to form larger units [9]. The hardness enhancement in such structures can occur according to the mechanism described by S. Veprek for a variety of nanocomposites published in the last years, where the observed hardness enhancement has been due to smaller crystallite size [33]. The high oxygen content in the films precludes from the formation of the icosahedra due to the formation of B2O3-like units. As a result, the icosahedra are not fully developed, and, according to our calculations, such amorphous structures will possess low hardness.

ACCEPTED MANUSCRIPT 4. CONCLUSIONS

Al-Mg-B films were deposited on silicon wafers by DC magnetron sputtering of an AlMgB14 target at low discharge power of 1.2 W/cm2 and different substrate temperatures between 100 and

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500 C. The as-deposited films were X-ray amorphous. The formation of boron crystallites

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occurred during annealing at 1000 C. An increase in substrate temperature leads to reduction of

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the film surface roughness and friction coefficient, and to an increase of hardness, elastic modulus and film adhesion. The films exhibit a nanohardness of 6-14 GPa and Knoop hardness of 8-18 GPa that are lower than those of the bulk BAM material of 25-35 GPa. The experimental investigations and first-principles calculations show that the observed low strength of the as-

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deposited amorphous films is due to their amorphous network, in which the icosahedra are not fully developed and the weaker B-O bonds dominate. The annealed film exhibits higher hardness

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than the as-deposited films, which is explained by the presence of boron crystallites in the annealed film. The reduction of the B-O bonds with increasing substrate temperature and after annealing promotes an improvement of the mechanical properties of the films. The as-deposited

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films have a relatively low coefficient of friction in the range of 0.08-0.12.

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ACKNOWLEDGEMENTS

This work was supported by the STCU Contract No. 5964. We would like to thank Stan

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Veprek for many critical comments.

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ACCEPTED MANUSCRIPT Table Captions

Table 1. Deposition methods and mechanical properties of the BAM-based films: PLD– pulsed laser deposition; MS-magnetron sputtering; DC-direct current; RF-radio frequency; Ts–

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substrate temperature; D-film thickness; H–nanohardness; E-elastic modulus.

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Fig. 1: XRD patterns of the film deposited at 450 °C (1) and annealed at 1000 °C (2). Fig. 2: FTIR spectra of the films deposited at 100 (1), 200 (2), 350 (3), 450 (4) and 500 C (5).

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Fig. 3: XPS core level spectra from the as-deposited (solid line) and annealed (dashed line) films

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prepared at TS = 450 C. For the sake of comparison, the XPS spectra for pure Al, Al2O3 and Mg

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are also shown. The vertical lines denote the binding energies at: 75.8 eV Al-O bonding in nonstoichiometric Al2O3 [24], 75.01 eV, Al-B bonding [9], 74.1 eV in MgAl2O4 [25] for the Al 2p spectra; 90.85 eV Mg-O bonding [5], 90.1 eV Mg-B bonding in MgB2 [26] for the Mg 2s spectra; 191.88 eV B-O bonding [5], 188.6 eV Borides [9] and 187.62 eV B-B bonding [9] for

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the B 1s spectra; 533.3 eV B-O bonding [27], 532.9 eV Mg-O bonding [28] and 531.8 eV Al-O bonding [29] for the O 1s spectra.

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Fig. 4: Nanohardness (H) and elastic modulus (E) as functions of indentation depth (h) for films deposited at different TS (cf. Fig. 2). The curve 6 corresponds to the sample 4 annealed at 1000 C.

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Fig. 5: Nanohardness (H), Knoop hardness (HK) and elastic modulus (E) of the Al-Mg-B films

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as functions of substrate temperatures (TS). Fig. 6: Friction force as a function of load.

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Fig. 7: Supercells for the a-Al3Mg3B56 (a) and the fragment of an icosahedron observed in amorphous BAM materials (b). The large, meddle and small balls denote Ng, Al and B atoms, respectively.

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Fig. 8: Calculated phonon density of states (PHDOS) for amorphous and crystalline Al0.75Mg0.75B14 in comparison with the Raman spectrum of crystalline Al0.75Mg0.78B14 [31]. Fig. 9: Calculated stress-strain relations for the weakest shear plane (001)[100].

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Figure 1

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Figure 2

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Figure 3

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Figure 4

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Figure 5

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Figure 6

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Figure 7

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Figure 8

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Figure 9

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Method

Target

TS

D

H

E

Ref.

(°C)

(m)

(GPa)

(GPa)

8-12

80-110

[5]

45-51

280-300

[6,7]

7.1-30.7

120-290

[8]

AlMgB14+20%TiB2

500

1.0

PLD

Al0.95MgSi0.05B14

27, 300

0.3-0.4

Al(RF), Mg(DC), B(RF)

25, 600

0.5-0.6

DC-MS

AlMg, B

200

0.5

17-30

N/A

[9]

DC-MS

AlMgB14+20%TiB2

N/A

3.0

17.8*

N/A

[10]

RF-MS

AlMgB14

25, 400,

N/A

N/A

N/A

[11]

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500, 600

*

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Knoop hardness.

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DC-RF-MS

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 Al-Mg-B films were deposited at different substrate temperatures.  The as-deposited films were amorphous, whereas the annealed ones were nanostructured.  Mechanical properties analyzed as functions of substrate and annealing temperatures.  Ab-initio MD simulations of AlMgB14-based materials were carried out.  Both experimental and theoretical investigations enabled one to explain film properties.