Oxidation Mechanism of Aluminum Nanopowders - The Journal of

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Oxidation Mechanism of Aluminum Nanopowders Marie-Vanessa Coulet,*,† Benoit Rufino,† Pierre-Henry Esposito,† Thomas Neisius,‡ Olivier Isnard,§ and Renaud Denoyel† †

Aix Marseille Université, CNRS, MADIREL UMR 7246, Marseille 13397 Cedex 20, France Aix Marseille Université, CNRS, Fédérationde de Recherche en Sciences Chimiques, Marseille 13397 Cedex 20, France § Université Grenoble Alpes, CNRS, Institut Néel, Grenoble 38042 Cedex 9, France ‡

ABSTRACT: Aluminum nanopowders, oxidized at different temperatures using thermogravimetric analyses performed in high resolution mode, are characterized in terms of morphology, structure and microstructure. The particle structure is modeled via geometrical considerations that enable the calculation of the variation of specific surface area during oxidation. A two-step oxidation scenario is proposed. In the early oxidation stage, that is, for temperatures up to 650 °C where a pseudoplateau is reached, the oxidation, which occurs by diffusion of oxygen or aluminum through the alumina layer, leads to a core−shell structure. At higher temperatures, that is, above the melting point of aluminum, outward diffusion of aluminum through the oxide shell is controlling the reaction rate. The reaction interface is then located at the external surface and voids are formed inside the particles. This result is confirmed by energy filtered electron micrographs that allow distinguishing a thin metallic aluminum layer outside the alumina shell. This suggests that the migration of aluminum toward the surface of the particles is faster than the oxidation. Some insights on the nucleation process during the crystallization of liquid aluminum are also proposed which are related to the particle microstructure: heterogeneous nucleation is proposed to govern the crystallization of liquid aluminum and to give a signature of the alumina layer structural state.

1. INTRODUCTION Under standard conditions aluminum is covered by a thin amorphous oxide layer acting as a protective coating against further reaction with oxygen, water, and so forth. Understanding the formation and growth of this alumina layer has been the source of numerous research works. In the field of thin films, studies are focusing on the first stages of formation of this layer and are motivated by the potential use of alumina layer for high-κ gate dielectrics for microelectronic applications.1 Generally, the oxide growth is performed by a thermal oxidation of bare aluminum surface under a partial oxygen pressure.2,3 Two different growth regimes have been proposed in the temperature range between 100 and 500 °C.4 Below 400 °C, a single fast oxidation stage is evidenced which corresponds to the formation of an amorphous oxide layer of limited thickness. The growth rate is realized by the electric field controlled outward migration of Al and leads to the formation of an Al-deficient amorphous layer. At higher temperatures (T > 400 °C), the oxidation proceeds in two steps: an initial fast oxidation stage similar to the one described for low temperatures followed by a second slow oxidation stage. This second oxidation step leads to the thickness increase of the amorphous alumina layer and its subsequent crystallization in γalumina. During this stage, two growth mechanisms are proposed: outward diffusion of Al cations as long as the oxide layer is amorphous and inward diffusion of oxygen along grain boundaries as soon as the oxide crystallizes. The nature of the Al substrate is shown to influence the growth of the oxide © XXXX American Chemical Society

layer. A thermodynamic model proposed that the critical thickness up to which the amorphous oxide is more stable than the corresponding crystalline oxide depends on the orientation of the aluminum substrate.5 In powders, the growth of the alumina layer under oxidative condition is at the center of many investigations because it can give some hints in the understanding of the combustion properties of the powders. Indeed, Al powders are commonly used as solid fuel ingredient in propellant and explosives formulations because of the high oxidation enthalpy of aluminum that confers the necessary energy for ignition. More recently, nanosized powders were proposed to enhance combustion efficiency and to ensure the ignition.6−8 Understanding the oxidation mechanism of Al powders is thus crucial for tailoring reactive powders. Generally, in the case of powders, the previously described “fast oxidation stage” has occurred during the synthesis or postsynthesis so that the particles are covered with a thin amorphous alumina layer.9 For spherical powders, the thickness of this native alumina layer is around 5 nm in micrometric powders,10 3 nm in spherical nanopowders produced by wire-electro-explosion technique10,11 and around 4.5 nm in flakes-like nanopowders.12 Depending on the heating rate and/or the size of the particles, two main mechanisms are proposed to explain the oxidation of Received: July 29, 2015 Revised: September 24, 2015

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DOI: 10.1021/acs.jpcc.5b07321 J. Phys. Chem. C XXXX, XXX, XXX−XXX

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The Journal of Physical Chemistry C

where the image is formed. This filter selects only electrons that have a defined energy spread. Indeed, when a sample is illuminated by an incident electron beam, electrons interact elastically or inelastically with the sample. The inelastically scattered ones lose a certain amount of their initial energy by exciting the sample (phonons, plasmons, inner-shell excitations). Beside the phonons, the excitation of plasmons is the most important inelastic scattering processes. In the case of aluminum, the characteristic bulk plasmon energy for aluminum is 15 eV and the energy loss spectrum shows a very sharp peak. This is in contrast with Al2O3 plasmon peak which presents a less intense and broader maximum around 25 eV. Thus, metallic parts of the powder can easily be distinguished from the oxide by forming the electron micrograph only with the electrons which have suffered an energy loss of around 15 eV. The measurements were performed using a FEI TITAN 80− 300 transmission electron microscope (TEM) equipped with an image corrector and a GATAN TRIDIEM imaging filter for the acquisition of energy filtered (EF) images. The acceleration voltage was 200 kV. The Al-plasmon images were recorded by using a 4 eV slit centered on the Al bulk plasmon peak at 15 eV. Microstructural changes were studied using neutron diffraction experiments and differential scanning calorimetry (DSC). Neutron diffraction experiments were realized on the D1B-CRG two-axis diffractometer at Institut Laue Langevin (Grenoble, France). The powder samples were placed into vanadium containers which contribute only incoherently to the scattered signal. The experiments were performed at a wavelength of 2.52 Å using a pyrolitic graphite (002) monochromator. The scattered neutrons were detected by a 3 He/Xe position-sensitive detector composed of a multielectrodes system with 400 cells which span an angular range 2θ of 80°. The position detector was set in order to cover an angular domain ranging between 50° and 130°. Differential scanning calorimetry (DSC) measurements were performed with a Tian-Calvet type apparatus (DSC111-Setaram) using alumina pans and under Ar flux (30 mL/min). The heating rate was fixed to 5 °C/min between 20 and 400 °C, and to 2 °C/ min between 400 and 700 °C. The calibration (in temperature and in enthalpy) of the calorimeter was done using the melting of pure references.

Al powders. The so-called “melt-dispersion” mechanism13,14 has been proposed for very fast heating rates. In this model, the melting occurs before any transition of the alumina layer and the liquid ejection is the starting point of the oxidation process. A diffusion controlled mechanism is generally assumed for low heating rates.15,16 In that case, the oxidation of aluminum nanopowders is stepwise and the various steps are attributed to polymorphic transitions occurring in the alumina shell which affect the transport properties of the ions and modify the diffusion rate.17 The values delimiting the different stages depend on the heating rate18 and on the particle size distribution.10,19 The location of the reaction interface and the nature of the diffusing ions (Al, O, or both) remain the subject of numerous works.20−22 This paper is devoted to the study of oxidation mechanisms of aluminum nanopowders heated at low rates. Under such experimental conditions, only diffusion based mechanisms will be considered and some of the pending questions will be addressed. Are the oxygen atoms diffusing up to the aluminum core over the whole oxidation temperature range? Is the melting of aluminum responsible for a change in the diffusing species? Can the reaction interfaces be experimentally localized? This Article is organized as follows: a first section is devoted to sample preparation and characterizing techniques. The experimental results are presented and discussed in a second part. The morphology and structural changes of the oxidized powders are first presented and then the role of the alumina shell during the oxidation process and its importance for aluminum phase change are discussed. At last, an oxidation scenario is proposed that includes a geometrical modeling of the evolution of the particle structure.

2. MATERIALS AND METHODS An aluminum nanopowder (Nano-L) provided by SNPE is the sample of interest of this paper. It has been elaborated using the wire-electro-explosion technique followed by a passivation step in a solution of stearic acid.23−25 The powder consists in spherical particles having a mean radius about 100 nm and an initial amorphous alumina layer of 3 nm thickness.10 Thermogravimetric measurements were performed under dynamic dry air atmosphere (60 mL/min) using a TAInstrument apparatus (TG Q 500 series) working in the high resolution scanning mode. This mode is based on sample controlled analysis.26 The heating rate is adapted automatically to the mass change (either in loss or in gain) that cannot exceed a predefined value. This implies that in the domains were the mass change is important; the heating rate may be reduced down to 0.2 °C/min. In order to estimate the initial amount of active aluminum in the initial powder, TGA measurement up to 1200 °C were performed on the Nano-L sample. Those experiments were performed under dry air atmosphere (60 mL/min) using a Setaram apparatus (SETSYS Evo) and using a heating rate of 2 °C/min. The specific surface areas of the oxidized powders were quantified by nitrogen adsorption−desorption isotherms analyses performed at 77 K using a commercial Micrometrics ASAP 2010 apparatus. Before adsorption measurement, the samples were evacuated at 120 °C under vacuum for 5 h. The evolution of morphology and core−shell structure was studied using transmission electron microscopy (TEM). In order to distinguish the metallic core from the oxide shell, energy filtered TEM (EFTEM) images were acquired. EFTEM consists of placing a filter between the sample and the plane

3. RESULTS AND DISCUSSION 3.1. Oxidized Samples. Ex situ oxidation was performed by heating Nano-L sample under air at several temperatures using thermogravimetric analyses (TGA) performed in high resolution mode. Figure 1 presents TGA graphs showing the stepwise oxidation of this aluminum nanopowder up to 900 °C. It is similar to the ones described in the literature.10,19 The step observed around 500 °C is linked to the crystallization of the amorphous alumina shell and the mass gain around 30% is characteristic for nanopowders. Several samples were prepared at different stages of the oxidation process and Figure 1 summarizes the various preparations: Nano-L-dep will be considered as the initial sample: it represents the Nano-L sample after a thermal treatment up to 400 °C that ensures the removal of the stearic acid layer without any oxidation of the sample. Nano-L-ox-1 corresponds to a slight oxidation up to 500 °C leading to a mass increase of 2.6%. Nano-L-ox-2a and Nano-L-ox-2b were prepared by oxidizing the sample up to 512 and 529 °C, respectively. This corresponds to an oxidation of 17% and 19%, respectively. B


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3.2. Morphology. As shown in TEM micrographs (Figure 2), the particles present a spherical morphology over the whole oxidation range and the average size of the particles slightly increases with increasing the oxidation range. Energy filtered images (bottom row in Figure 2) allow a clear distinction between metallic aluminum core and oxidized shell of the particles. From these images, we further observe that the core−shell morphology of the particles is preserved over the whole range of oxidation. EFTEM allows evidencing an increase of the alumina shell: thicknesses around 5 nm are measured for the Nano-L-dep sample (Figure 2a′) while for higher oxidation rates, they reach values up to 50 nm (Figure 2e′). Moreover for samples Nano-L-ox2a and Nano-L-ox3, the fringes observed in the standard bright field images (see arrow in Figure 2c and d) indicate the presence of strain inside the metallic core. At last, hollow shell particles are observed for sample Nano-L-ox5 (Figure 2e and e′). The presence of such particles has been already reported in literature but for higher oxidation temperatures.20 Interestingly, for this sample, EFTEM images allow distinguishing metallic aluminum around the empty shell (see arrows in Figure 2e′). This observation, that suggests that the remaining aluminum has migrated outside the core, comes withstand recent molecular dynamics calculations.22 It has been indeed proposed that, when the Al core melts, the nanoparticle surface becomes Al rich due do the migration of ions. The migration rate of aluminum is then faster than the oxidation process. It is worth noting that a careful inspection of Figure 2e′ reveals the presence of another shell (probably amorphous Al2O3) that surrounds the thin aluminum layer. BET specific surface areas deduced from adsorption isotherms are given in Table 1. One may observe that the values of the specific surface area slightly decrease as the particles are oxidized up to 739 °C. Such a behavior is consistent with TEM observation. It is worth noting that, for the sample oxidized up to 900 °C, the specific surface area increases. This point will be discussed later. The C parameter, which gives an indication on the adsorbate−adsorbent interaction, is also given in Table 1.

Figure 1. High resolution thermogravimetry analyses performed on Nano-L sample in order to prepare the samples investigated in this study. The number in % corresponds to the mass gain associated with each sample preparation.

Nano-L-ox-3 corresponds to an oxidation at the level of the first pseudoplateau at 595 °C and corresponds to a mass increase equal to 32.8%. Nano-L-ox-4 corresponds to an oxidation up to 739 °C leading to a mass increase of 49.3%. Nano-L-ox-5 corresponds to an oxidation up to 900 °C. This is the maximum temperature accessible with the TGA apparatus used for this study. At this temperature the mass increase was equal to 67.3%. Figure 1 gives a graphical representation of the prepared samples. It also shows the very good reproducibility of the preparation since all the TGA recordings (corresponding to the various samples) are superimposable in their common temperature range. It indicates that all the samples have followed the same reaction pathway and can be reasonably compared.

Figure 2. Standard TEM images (top row) and corresponding EFTEM images (bottom row) for Nano-L-dep (a,a′) (oxidized up to 400 °C), NanoL-ox1 (b,b′) (oxidized up to 500 °C), Nano-L-ox2a (c,c′) (oxidized up to 512 °C), Nano-L-ox3 (d,d′) (oxidized up to 595 °C), and Nano-L-ox5 (e,e′) (oxidized up to 900 °C). EFTEM images allow a clear distinction between metallic aluminum core and oxidized shell of the particles. C


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The Journal of Physical Chemistry C Table 1. Specific Surface Area Measured Using Adsorption Isotherm Performed at 77 Ka as (m2 g−1) C a

Nano-L-dep

Nano-L-ox-1

Nano-L-ox-2a

Nano-L-ox-3

Nano-L-ox-4

Nano-L-ox-5

10.06 100

9.51 132

9.44 226

9.17 224

8.36 219

9.26 91

The values of C parameter are also given.

decrease of crystallized aluminum is concomitant with the increase of crystallized alumina as expected from the Trunov model.17,19 In spite of the presence of crystallized nuclei of alumina in the initial powder (∼1%), the major crystallization event occurs for the sample oxidized at step 3. This confirms that the first mass increase around 500 °C is related to the crystallization of alumina. Finally, it is worth noting that the AlN mass fraction decreases with the oxidation temperature and is not anymore detected in the Nano-L-ox-5 sample. This can be explained by the oxidation of AlN which is reported to start around 800 °C.28 DSC analyses also provide information about the changes of microstructure of the nanopowders. Heat flow recordings obtained upon heating and cooling are given in Figures 5 and 6, respectively. As shown in Figure 5, a broad exothermal effect is observed around 580 °C for Nano-L-dep and Nano-L-ox-1 samples. As demonstrated in our previous studies,27 this peak

Figure 3. Experimental neutron diffraction patterns measured at 2.52 Å and at room temperature for the oxidized samples. The vertical markers refer to the calculated Bragg reflexions for Al (first row), AlN (second row), and γ-Al2O3 (third row). The patterns are vertically shifted to ensure legibility.

3.3. On the Role of the Alumina Shell. Neutron diffraction patterns of oxidized samples are given in Figure 3. In the initial sample Nano-L-dep, three phases are evidenced: aluminum (Al), aluminum Nitride (AlN), and alumina (Al2O3). This is consistent with previous finding.27 A strong intensity increase of Al2O3 diffraction peaks is observed for sample Nano-L-ox-3 which is the sample oxidized up to the first plateau of TGA recordings. Using the structures proposed by Rufino et al.,27 systematic Rietveld refinements were performed. The only free parameters were the cell parameters and scale factors for the three phases. The results are summarized in Figure 4 that presents the evolution of the weight percentage for the three phases as a function of the oxidation temperature. It shows that the

Figure 5. DSC recordings upon heating for the various oxidized powders. The inset emphasizes the exothermic peak characteristic of alumina crystallization.

Figure 4. Evolution of the mass percentage for the Al, Al2O3, and AlN phases as a function of the oxidation temperature for Nano-L series of samples obtained from Rietveld refinement. The dashed line represents the mass gain as a function of the oxidation temperature.

Figure 6. DSC recordings upon cooling for the various oxidized powders. D


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The Journal of Physical Chemistry C corresponds to the crystallization of the alumina shell. Its presence confirms that most of the initial particles, in spite of containing a small amount of crystallized Al2O3, are covered by an amorphous alumina layer. The Nano-L-ox-1 sample still contains a large amount of amorphous alumina as shown by the DSC exothermal peak but the heat associated (17 J/gsample) is lower than for the initial sample (23 J/gsample) and the peak maximum is shifted toward lower temperatures (580 °C against 588 °C). These results may be related to the polydispersity of the alumina thickness as illustrated by the EFTEM images in Figure 2. During the oxidation between 0 and 3%, there may be some thicker amorphous alumina regions that reach their critical nucleus for crystallization owing to the slight oxidation. This is confirmed by Rietveld analyses of neutron powder diffractograms that show an increase of the mass percentage of crystallized alumina from 1.3% for sample Nano-L-dep up to 4.5% for sample Nano-L-ox-1 (Figure 4). This is confirmed by TEM images (Figure 2a′ and b′) that show a thickness increase of the alumina layer between Nano-L-dep and Nano-L-ox 1 samples. The shift in crystallization temperature thus could suggest that the crystallization temperature is lowered with increasing thickness of the alumina layer. From sample Nano-L-ox-2a and following, the absence of any exothermal peak on DSC recording suggests that there is no more amorphous alumina in the samples or that the residual amount of the amorphous phase is negligible. Consequently, above an oxidation temperature of ∼510 °C, the thickness increase of the alumina shell occurs by growth of crystallized alumina. At this step, the inflection point of TGA recording that corresponds to the maximum growing speed is overpassed and the rate of oxidation is slowing down until the first pseudoplateau. These data suggest that the first oxidation mechanism, corresponding to the rapid oxidation of bare aluminum domains due to the breaking of the alumina shell upon crystallization, occurs for less than 17% mass gain. The endothermic peaks located around 600 °C in DSC recordings (Figure 5) are related to the melting of confined aluminum. It is well-known that confinement shifts the melting temperature, according to the Gibbs−Thomson equation, in a direction which depends on the wetting properties of the liquid. In the case of surface melting (liquid wets the walls better than solid), a decrease of temperature should be observed, which is the case for aluminum confined in its alumina shell.10,29 Looking at the evolution of the onset temperature with oxidation rate, a slight decreasing trend is observed with the exception of the sample oxidized at 595 °C, that is, up to the first plateau in TGA curves. This is unexpected since the size of aluminum core is supposed to decrease continuously with the oxidation rate. Not only the melting temperature but also the melting enthalpy has an unusual behavior for that sample. Figure 7 shows the mass percentage of aluminum deduced from the melting enthalpy values (assuming that the enthalpy is independent of size effects in the considered range) compared to the ones derived from the Rietveld refinement. Both methods are sensitive to the amount of crystalline phase so that they can be compared. It should be mentionned that both values were normalized to the initial amount of active aluminum deduced from TGA measurement up to 1200 °C. This corresponds to an initial amount of active aluminum equal to 82% in Nano-L-dep sample. As observed in Figure 7, for the sample oxidized up to 595 °C (Nano-L-ox3), the amount of aluminum deduced from enthalpy values is significantly lower

Figure 7. Evolution of the mass percentage of aluminum as a function of the oxidation temperature deduced from DSC measurements (open symbols) and from neutron diffraction experiments (full symbols). Note that the values were normalized to the initial active aluminum content obtained from TGA analyses up to 1200 °C (see text). The dashed line is a guide for the eyes, and the double ended arrow marks the variation in the aluminum percentage estimated using either melting enthalpy or crystallization enthalpy in the case of Nano-L-ox3.

than the expected one (see the double ended arrow in Figure 7). In order to explain this behavior, two hypotheses can be done: (i) additional stress appears when the sample is oxidized up to 600 °C or (ii) the wetting parameters are changing. Both phenomena invalidate the simple behavior predicted by the Gibbs−Thomson equation. Among the stress, one can imagine that after the oxidation the aluminum is submitted at a higher pressure by the alumina layer. Such a pressure effect usually shifts the melting temperature toward higher values. Another possible effect is the change of wetting properties in the system: the surface tension of amorphous alumina/liquid aluminum interface could be different from that of crystalline alumina/ liquid aluminum so causing a different temperature as predicted by the Gibbs−Thomson equation. The behavior of the oxidized nanopowders during solidification allows to test those hypotheses. As shown in the DSC recordings in Figure 6, the crystallization behavior is surprising since it reveals a two-step solidification process of the nanopowders whereas a constant shift of solidification temperature with oxidation rate was rather expected. Similarly to the melting process, the mass percentage of crystallized aluminum has been deduced from the crystallization enthalpy values. The values obtained are given in Figure 7. They are in very good agreement with the ones obtained from the melting enthalpy. Interestingly, the values obtained for sample Nano-L-ox-3 are higher than during the melting cycle and agree fairly well with the mass percentage obtained from Rietveld refinements. Thus, the low melting enthalpy for sample Nano-L-ox-3 can be explained by the presence of strain in the aluminum core which is consistent with the TEM observations of Figure 2d. The melting inside the DSC leads to strain relaxation. Since the cooling is performed at low rate, the liquid can recrystallize without exercising any stress in the particles core. This also proves that the first mass increase in TGA is not correlated to the melting of the aluminum core that would have been stress free in such a case. Even if the presence of strain has been evidenced in sample oxidized up to the first plateau of TGA curve, the wetting E


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The Journal of Physical Chemistry C properties of liquid aluminum also have to be discussed because of the observed two step crystallization process. No strong effect is observed in the variation of the peak onset for the two peaks. The solidification of a confined liquid is generally driven by the nucleation of the solid inside the liquid. This is why there is generally a large hysteresis between melting and solidification, the melting transition itself being metastable.30 The nucleation may be either homogeneous, that is, the nucleus appears inside the liquid, or heterogeneous, that is, the nucleus appears at the interface between the liquid and the solid interface. Interestingly, in the case of aluminum, the undercooled liquid does not exist over a large temperature domain, contrary to other metals. As suggested by Turnbull and Cech,31 this may indicate that the alumina layer does influence the crystallization of aluminum. This implies that a mechanism of heterogeneous crystallization can be evoked in the case of liquid aluminum, the nucleation taking place at the alumina/ aluminum interface. The nature of the surface should then strongly influence the crystallization and one may expect that the alumina shell has not the same surface state if it has crystallized under Ar atmosphere (i.e., with a deficiency in O2) or under air. Such arguments may explain the difference between the samples Nano-L-dep and Nano-L-ox1 and the others for which the alumina shell has crystallized previously to the DSC measurements. A possible application of the crystallization study is the evaluation of the fraction of particles for which the alumina layer has changed in structure during the oxidation. 3.4. Toward an Oxidation Scenario. The oxidation of aluminum powder is largely documented in literature. We propose a simple model, whose aim is to estimate the specific area of the various oxidized samples and to compare it to experimental values. Considering the heating rates used in this study, only diffusion based mechanism21,22 are considered. The following scenarios which are all based on simple geometrical consideration are tested: (i) A constrained core−shell model in which the total particle size is constant. The oxidation takes place by growing the alumina shell without growing the particle size. The reaction interface is at the aluminum/oxide interface and takes place via the diffusion of oxygen through the alumina shell. The aluminum core is constrained. (ii) A core−shell model with constant Al density in which the oxidation takes place by growing the alumina shell and the total size of the particles. In this model, it is not necessary to make a hypothesis on the location of the reaction interface. The final state of this model is equivalent to the ones given by Zhang and Dreizin21 when the alumina layer is considered as ductile, whatever the diffusion species. As proposed in the literature and summarized in the introduction, in this model, both diffusion of aluminum and oxygen can occur depending on the alumina crystallinity and defects. (iii) A hollow sphere model in which the inner diameter of alumina shell is constant. The oxidation proceeds via the diffusion of aluminum through the oxide shell. In this model it is assumed that aluminum wets alumina which results in the formation of a hole at the center of the particle. The reaction interface is at the oxide/surface interface and takes place via the diffusion of aluminum through the alumina shell. The changes in the density linked to the phase transformation of alumina (amorphous → crystallized) are taken into account in each model. For the modeling, we consider a unique mean particle having an initial equivalent diameter D0

that corresponds to the initial specific area of Nano-L-dep sample (D0 = 220 nm) The initial mass of metallic aluminum m0Al can then be written: 0 mAl =

⎞3 4 ⎛ D0 0 π⎜ − tox ⎟ ρAl 3 ⎝ 2 ⎠

(1)

and the initial mass of alumina is 0 mox

3 ⎡ ⎛ D0 ⎞3 ⎤ 4 ⎢⎛ D0 ⎞ 0 ⎥ = π ⎜ ⎟ −⎜ − tox ⎟ ρox_am 3 ⎢⎣⎝ 2 ⎠ ⎝ 2 ⎠ ⎥⎦

(2)

with beingthe initial alumina thickness, and ρAl and ρox_am being the density of aluminum and amorphous aluminum oxide, respectively. Defining the total initial mass (of one particle) as m0 = m0Al + 0 mox for a given weight gain w, the total mass can be written: t0ox

⎛ w ⎞⎟ 0 m m = ⎜1 + ⎝ 100 ⎠

(4)

The residual mass of aluminum is ⎞3 4 ⎛ D0 9 wm0 0 mAl = π ⎜ − tox ⎟ ρAl − 3 ⎝ 2 8 100 ⎠

(5)

(i) Core−shell model with constrained aluminum, that is, constant particle size: In that case, the specific surface area is directly given by a = πD°2 /m

(6)

(ii) Core−shell model: The diameter of aluminum core is DAl = 2 3

3mAl 4πρAl

(7)

The diameter of the particle is then equal to D = 23

⎛ D ⎞3 3mox + ⎜ Al ⎟ ⎝ 2 ⎠ 4πρox

(8)

with ρox being the density of crystalline alumina. (iii) Hollow sphere model with aluminum diffusion: The constraint is that now the inner diameter of alumina shell is constant and equal to D° − 2t0ox. The particle diameter can now be expressed by the following equation: D = 23

⎞3 ⎛ D0 3mox 0 +⎜ − tox ⎟ 4πρox ⎠ ⎝ 2

(9)

For the last two models, the specific surface area is calculated as a = πD 2 / m

(10)

The calculated values for specific surface area are shown in Figure 8 (filled symbol), together with the experimental values (open symbols). The evolution and the order of magnitude of the experimental and calculated values are in reasonable agreement. At the initial oxidation stages, the differences between experimental and calculated values can be attributed to changes in the surface polarity that directly influence the surface area obtained using nitrogen adsorption measurements. Jelinek and F


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This sequence mostly agrees with the one proposed by Rai et al.,20 except that in the present work hollow shells were observed before oxidation at 800 °C. Those results confirm the hypothesis drawn by Firmansyah et al.33 about a possible diffuse migration of Al atoms through the alumina shell around 660 °C. The observation of an Al-rich shell at the surface agrees also well with recent reactive molecular dynamics simulations suggesting that, when the Al core melts, the surface of the nanoparticles becomes Al rich due do the migration of ions.22

4. CONCLUSION Aluminum nanopowders were oxidized under air at different temperatures up to 900 °C. By combining specific surface area measurements and energy filtered TEM imaging with a geometrical model, a new sequence is proposed to explain the stepwise oxidation mechanism. At low temperature [20− 660 °C], the oxidation takes places through the oxide layer owing to the diffusion of oxygen or aluminum species. This mechanism is valid up to first pseudoplateau of TGA recording were a slow-down of oxidation is observed. At that point, the Al core has not melted and the onset of stress in the aluminum core has been evidenced. The oxidation rate restarts above 660 °C due to the melting of alumina that allows aluminum species diffusing faster than oxygen. The outward diffusion of aluminum proceeds through the oxide shell and the reaction interface is located at the interface alumina/surface. This result, expected by molecular dynamics simulation,22 is confirmed by EFTEM images that reveal a thin aluminum layer surrounding the external surface of the alumina shell. Some insights on the nucleation process during the crystallization of liquid aluminum were also proposed which are related to the particle microstructure: heterogeneous nucleation is proposed to govern the crystallization of liquid aluminum and to give a signature of the alumina layer structural state.

Figure 8. Variation of the specific area as a function of the oxidized aluminum mass fraction. The open (black) squares are the experimental values obtained from adsorption isotherm measurements. Filled symbols represent the values calculated using various oxidation models. Dotted lined are guidelines for the eyes.

Kovats32 have shown that these polarity changes can lead to variations as high as 20% of the BET surface area in respect to the actual surface area. Nevertheless, because the samples in this study stay rather polar as indicated by the C parameter value of the BET model (see Table 1) which is always higher than 90, an error bar of 5% is given in the graphs of Figure 8, in agreement with Jelinek and Kovats32 measurements on calcined and noncalcined samples. The core−shell models (constrained and unconstrained) give rise to a continuous decrease of the specific surface area and the decrease is even more pronounced for the constrained model. The hollow sphere model gives only slight variations of the specific area over the whole oxidation range. Both specific surface area measurements and TEM observation (even without statistical analyses) are consistent with a particle size increase. This allows the elimination of the constrained core−shell model. In contrast, EFTEM images (Figure 2e′) show the necessity of considering the diffusion of Al through the alumina shell and thus to consider the hollow sphere model. Taking into account that diffusion of Al species could be enhanced in the liquid phase, we may consider that the core−shell model is valid up to the melting of Al at which the hollow shell model starts to predominate. To this aim, we calculated a particle diameter for an oxidation state at 40% (i.e., at the melting point) using the core−shell model and used it to calculate the initial state for the hollow shell model. The values are represented by triangle symbols in Figure 8. It is shown that the hollow sphere model induces an increase of the calculated specific surface area as observed experimentally The following sequence can be proposed for the oxidation mechanism at low heating rates in Al nanopowders: (1) For low oxidation temperatures [20−660 °C], a core− shell oxidation takes place through the oxide shell by inward diffusion of oxygen or outward diffusion of aluminum. This geometrical model does not permit to conclude on the dominant diffusing species but is consistent with Jeurgens’4 work suggesting that oxygen diffusion may predominate once the alumina shell is crystallized. (2) At the melting of aluminum and high temperatures [660−900 °C], outward diffusion of aluminum ions predominates. The aluminum diffuses through the oxide shell and the reaction interface is located at the external surface of the particles.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Tel: +33413551809. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The “Groupe SNPE” is acknowledged for financial support and for providing the aluminum nanopowders. The authors are also indebted and grateful to the CRG-D1B at Institut Laue Langevin.

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REFERENCES

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Oxidation Mechanism of Aluminum Nanopowders - The Journal of

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