Controlling the Nanocontact Nature and the Mechanical Properties of

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Controlling the Nanocontact Nature and the Mechanical Properties of a Silica Nanoparticles Assembly Jeremy Avice, Christophe Boscher, Gwenaelle Vaudel, Guillaume Brotons, Vincent Juve, Mathieu Edely, Christophe Méthivier, Vitalyi E Gusev, Philippe Belleville, Herve Piombini, and Pascal Ruello J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/acs.jpcc.7b08404 • Publication Date (Web): 29 Sep 2017 Downloaded from http://pubs.acs.org on October 3, 2017

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The Journal of Physical Chemistry C is published by the American Chemical Society. 1155 Sixteenth Street N.W., Washington, DC 20036 Published by American Chemical Society. Copyright © American Chemical Society. However, no copyright claim is made to original U.S. Government works, or works produced by employees of any Commonwealth realm Crown government in the course of their duties.

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Controlling the Nanocontact Nature and the Mechanical Properties of a Silica Nanoparticles Assembly J. Avice,†,‡ C. Boscher,‡ G. Vaudel,† G. Brotons,† V. Juvé,† M. Edely,† C. Méthivier,¶ V. E. Gusev,§ P. Belleville,‡ H. Piombini,‡ and P. Ruello∗,† †Institut des Molécules et Matériaux du Mans, UMR CNRS 6283, Le Mans Université, 72085 Le Mans, France. ‡Commissariat ł’Energie Atomique et aux Energies Alternatives, Centre du Ripault, Monts, France. ¶Laboratoire de Réactivité de Surface, UMR CNRS 7609, LRC-CEA/UPMC/CNRS no 1, 4 Pl. Jussieu, Université Pierre et Marie Curie, 75252 Paris, France. §Laboratoire d’Acoustique, UMR CNRS 6613, Le Mans Université, 72085 Le Mans, France. E-mail: [email protected] Phone: +33 (0)2 43 83 32 68

Abstract Elaborating advanced nanomaterials based on the assembly of nanoparticles (NPs) is a versatile route for targeting and tuning a wide variety of properties like optical, magnetic, electrical properties or sensing. This route usually employs a so-called soft chemistry which has the advantage to be quite cheap and transferable to an industrial level. However getting quantitative informations on the quality of the mechanical consolidation of the nanoparticle assembly (ordered or disordered), in a non destructive

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manner is not often achieved although it is crucial for applications and integration of materials in devices. Therefore, we present in this article a complete method where we evaluate the elasticity of weakly (Van der Waals nanocontacts) to strongly (covalenthydrogen nanocontacts) interacting nanoparticles assemblies. This complete work is realized on a disordered silica nanoparticles network obtained by sol-gel method. A precise control of the chemical and physical properties of the nanoparticles surface molecular landscape is achieved thanks to infra-red, visible and ultra-violet spectroscopies as well as surface tension measurements and atomic force microscopy, while the nanoparticles assembly elastic stiffness is evaluated by ultrafast nanoacoustics based on a optical pump-probe method.

Introduction Nanomaterials engineering based on the assembly of nanoparticles (colloidal films) permits to tailor many properties of advanced functionalized film and smart coatings. 1–5 By adjusting the nanoparticles (NPs) volume fraction and the nature of the nanoparticles interconnection, the physical and chemical properties of colloidal films can be designed for specific applications in optics, plasmonics and photovoltaics. 3 These properties have also been studied as a function of the geometry of the packing and the role of the nanoparticle/nanocrystals shape has been previously deeply studied. 3,6,7 While, continuous efforts have been made to optimize the growth and the control of the properties of these nanoparticle assemblies, 3–5 including the recent programmable CND-coated assemblies, 8,9 it is admitted that nanoparticle assemblies suffer from a poor mechanical reliability and durability. Despite this situation, the control of the physical and chemical nature of the nanocontacts on the collective mechanical integrity has been poorly evaluated 3 although crucial for future integration in any devices and smart coatings. The nature of the chemical bonds involved in the nanocontact is well identified with traditional spectroscopic measurements 3–5 but the strength of the nanocontacts and the overall elasticity of the NPs network are not evaluated usually. 3–5 Beside the importance of 2


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the mechanical integrity of the NPs assemblies, 10,11 the quality of the nanocontact governs many other properties such as electronic 12,13 or thermal transport. 14 The nanocontact elastic stiffness can be studied with spectroscopic methods such as Brillouin 15 or Raman spectroscopyies. 16,17 While the spectroscopic methods are based on the analysis of thermally excited vibrations in the solids, the ultrafast optic methods offer the possibility to excite and detect with light in an efficient way some mechanical eigenmodes of nanomaterials and nanostructures. This so-called ultrafast nanoacoustics is known indeed as a powerful method to evaluate the elasticity at nanoscale and has been already used to probe nanostructured materials 11,18,19 and colloidal nanostructures. 11,20–23 However, to get a full understanding of the relation between the macroscopic elastic modulus of thin colloidal films and the nanoparticles surface chemical landscape, simultaneous investigations are needed including spectroscopic, thermodynamic and nanomechanical methods. In this article, we unravel the complex mechanisms of consolidation of a silica NPs network. By using a combination of techniques, we quantitatively evaluate the elastic modulus of the NPs network and we show how it evolves as a function of the colloidal film surface tension (contact angle) in line with the evolution of the nature of the interparticles chemical bonds (IR spectroscopy). Finally, we also evaluate the role of defects and we evidence a crossover in the nanomechanical properties from a regime where nanocontacts control the network elasticity to a situation where submicrometre cracks significantly influence the colloidal network elastic stiffness. The association of the several experimental methods demonstrates that it is possible to fully characterize and to control the chemical, physical and mechanical properties of silica colloids based thin films. This method could be extended to any colloidal film.

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Methods Sample Preparation The silica colloids have been prepared following first a sol-gel process as described in previous studies 24 and the thin films are obtained with a dip-coating method. 25 The NPs silica suspension we have used is a sol of amorphous silica prepared by the base catalyzed (NH3 ) hydrolysis of distilled tetraethylorthosilicate in pure ethanol according to the well known Stober method. 26,27 The basic medium favors the nucleation of hydrolyzed species and the silica sol consists of monodispersed roughly spherical particules of 10 nm diameter. The pH of the final silica is 6 after distillation of NH3 , its content is 4% in weight and the sol viscosity is 1.2 cP. The NPs assembly films are obtained by dip coating. The NPs packing has been realized either with Van der Waals nanocontacts or with Covalent-hydrogen bonds. This chemical/physical modification of the nanoparticles surface was achieved thanks to a postprocessing by an ammonia catalysor based treatment allowing a modification of nanocontact bonds from Van der Waals to hydrogen/covalent bonds. 26,27 This process is called the hardening process in the following. The chemical route and the colloidal film formation are sketched in Figure 5. The assembly of NPs was prepared on various substrates depending on the techniques of characterization used. For IR vibrational spectroscopy, the assembly was realized on a transparent silicon substrate, for near UV-visible-near IR optical transmission interferometry we used a silica substrate while for the picosecond acoustic methods (see Picosecond acoustic methods for more details), the assembly of NPs was obtained on a thin chromium film (20 nm) that is deposited on a silicon substrate. This thin chromium film is used as a local optoacoustic nanotransducer required for realizing ultrafast nanoacoustics experiments. 11,19,28

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Film thickness measurements The colloidal film thickness was determined from a spectral measurement in transmission (T) of a silica substrate coated by dip-coating (both silica surfaces were coated with the same colloidal film thickness, i.e. side 1 and side 2). The transmission measurements silica were made by a Perkin 900 spectrophotometer over [200-1500 nm]. Following the standard approach, 29 a full numerical fitting of the spectrum was performed to extract both the thickness and the refractive index values with a Cauchy model for the wavelength dependence of the refractive index. The fitting algorithm is given in details in the Supplementary Informations file with some examples of fits shown in Figures S1, S2 and S3 and Table S1.

Picosecond acoustics methods The ultrafast nanoacoustic method is based on the generation and detection of short acoustic pulse having acoustic wavelength of tens of nanometers suitable for probing nanomaterials. This is based on a 80 MHz repetition rate Ti:sapphire femtosecond laser. The beam is split with a polarizing beam splitter into a pump and a probe beams. The pump beam excites the system (generation of acoustic pulses) and the probe beam can follow in time how the acoustic pulse propagates within the nanostructure. It is somehow a nano-echography. All the information is contained in the modification of the optical reflectivity (∆R(t)) induced by the initial pump excitation. The time-resolved measurements signals (∆R(t)/R) are obtained thanks to a mechanical delay stage (delay line) which enables a controlled arrival time of the probe pulse regarding to the arrival of the pump pulse. The experiments were conducted with incident pump and probe beams perpendicular to the surface of the nanoparticles film as shown in Figure 3a with a typical pump and probe spot diameter of 10 µm. The pump beam is modulated in intensity which permits to extract small transient optical signal (∆R/R ∼ 10−5 /10−4 ) with a lock-in amplifier scheme. Since the silica nanoparticles network is transparent to both the pump (830 nm) and the probe (580 nm) beams, the femtosecond 5



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laser does not directly excite the NPs assembly film. For this reason the colloidal film was deposited on a thin chromium layer acting as a thermoelastic transducer to create acoustic nanowaves (this multilayer is sketched in Figure 3a). The use of such thermoelastic nanotransducer is commonly employed to probe the elasticity of transparent nanomaterials. 11,30 A typical signal of transient optical reflectivity is shown in Figure 3b. This signal is made of a sharp increase and decrease of the optical reflectivity in the timescale of 1 ps which corresponds to the optical excitation of the electrons of the chromium film, followed by a fast decay characteristic of the relaxation of the hot electrons. Then oscillatory signals evidenced over more than 1 ns are the signature of the nanoacoustic waves in the thin NPs assembly. In our particular case, we generate compression/dilatation acoustic waves in the chromium layer (thermoelastic optoacoustic transducer) that are transferred within the NPs assembly film (Figure 3a). After propagating and reflecting at the different interfaces (air/colloid and colloidal film/chromium layer), these acoustic waves make the NPs assembly film ringing. The analysis of the mechanical resonances provides direct information on the NPs assembly elastic properties as detailed latter on.

Results and discussion The modification of the nature of the chemical bonds on the nanoparticle surface during the VdW-CV/H crossover, has been first evaluated by IR spectroscopy as shown in Figure 2a. Figure 2a shows comparison between infrared absorption response of standard and NH3 posttreated silica layers revealing obvious chemical changes due to the ammonia treatment. Due to the hydrolysis-condensation process during the ammonia treatment, the methyl bonds (CH3 ) disappear while, on the opposite, the hydrogen bonds (H-OH and Si-OH) are created and the Si-O-Si bonds evolve due to the appearance of a covalent nanocontact. All IR-bands can be assigned either to Si-O, Si-OH or OH bond vibrations or to remaining ethoxy groups (bands appearing in the 1400 cm−1 region) in agreement with the literature. 31 After treat-

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ment, a strong modification occurred in the relative intensity of the Si-O-Si stretching broad band. The increase of the Si-O-Si LO asymmetric stretching vibration band at 1220 cm−1 and that of the Si-OH vibration band (950 cm−1 ) is interpreted to result from a strengthening of silicate gel network through cross-linking. As already mentioned, the ammonia-treatment effect on colloidal silica layers is consistent with the base-catalyzed condensation mechanism proposed by Iler. 32 With ammonia-curing, the particle-to-particle linking is enhanced via H-bonding of neighbor particles through vicinal silanols and condensation reactions via siloxane bridging. The first main important information from IR vibrational spectroscopy is that around 30 minutes of post-treatment is enough to make the methyl bonds disappear. After typically 30 minutes, we do not see indeed the IR transmission spectra evolution indicating, from this spectroscopic point of view, that the chemical transformation is complete (saturation effect). We have integrated the intensity of some selected IR bands and showed that saturation appears within less than one hour of duration of the hardening process (not shown). Whatever the NPs assembly film thickness (70 to 210 nm), the transformation of the IR spectrum is rapid.

The evolution of the surface energy of nanoparticles during sample hardening process (transition from VdW to CV-H bonds) has been also assessed by the contact angle measurements (Young-Dupré angle) and ultrafast nanoacoustics as described below. First of all, some typical contact angle profiles are shown in Figure 2b where we can see clearly the change of the surface thermodynamics. Figure 2c shows how this thermodynamic parameter evolves as a function of the hardening time. A drastic evolution from an hydrophobic (θ > 90◦ ) to hydrophilic (θ < 90◦ ) behavior (Figure 2c) is evidenced within the first 30 minutes of hardening process. The contact angle measurement was performed for two films (thickness of H=210 nm and H=70 nm) and both films exhibit similar evolution of the surface energy (see blue and red circles respectively). The rapid evolution of the contact angle versus hardening time (i.e. versus the amount of H-CV/VdW bonds ratio) is in

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correspondence with the chemical modification of the nanoparticle surface probed with IR spectroscopy. All these contact angles were measured around 1h hour after the sample was prepared. It is worth to underline that long hardening time leads to a continuous increase of the contact angle showing the hydrophilicity decreases. Similar phenomena was checked on bare silica surface submitted to NH3 hardening process (black circles in Figure 2c) while stable hydrophilic contact angle is preserved for silica in air or wet silica (see light blue circles in Figure 2c). This phenomenon is likely due to a nitrogen compounds adsorption coming from the NH3 catalyst evidenced by X-ray Photoemission Spectroscopy (XPS) as shown in Figure S4. However, as we will show in the following, this does not impact the nanoparticle elastic stiffness network. Indeed, as ammonia molecules adsorption increases onto the surface with hardening time, the hydrophilicity of the film decreases (contact angle increase above 50Âř) because of the covering of OH surface groups with NH3 molecules having a quite low dipole moment value compared to absorbed H2 O (1.47 D versus 1.86 D). Furthermore, while modifying the nature of the silica nanoparticles nanocontact, we observe (Figure 2d) a saturation behavior of the geometrical shrinkage of the film (∆H/H < 0) for long time of hardening process (the example is shown for H=210 nm but similar behaviors are observed whatever the sample). The shrinkage is due to the consolidation of the nanoparticle assemblies whose interparticles interactions are strengthened by the hydrogen and covalent bonds. Some solvents (alcohol, water) are also removed during the hardening post-treatment of the NPs assembly. The saturation of the thickness evolution demonstrates that the film reaches a stable packing and, compared to the contact angle, no continuous drift is observed when increasing the hardening time. This thickness measurement was realized thanks to a spectrophotometry measurement (Figure 2d) as detailed in Methods and SI. The consequence of this transformation of the nanocontact on the elastic stiffness of the 3D assembly is assessed by ultrafast acoustic methods. The laser induced acoustic waves signals are shown for various films in Figure 3c after the baseline substraction (electronic decay and thermal relaxation). The signals clearly evidence the existence of two characteristic

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modes f0 and f1 . The Fourier Transforms shown in Figure 3d indicate there is a ratio of 3 between these two frequencies indicating these modes are eigenmodes with f0 =V /4H and f1 =3V /4H, where V and H are the sound velocity and the thickness of the colloidal film respectively. 11,18 With a known thickness H determined by spectrophotometry, the sound velocity is straightforwardly deduced with the above eigenmodes frequency equation. Because the optical interferometry also provides an evaluation of the optical refractive index n, then by applying a mixing law that is relevant for large porosity p medium such as our NPs assembly, we can extract the mass fraction (1 − p) with the relation between ǫ = pǫair + (1 − p)ǫSiO2 √ with n = ǫ the refractive index of the NPs assembly. We thus deduce the mass density ρfilm = (1 − p)ρSiO2 . Finally, the elastic modulus is determined following M = ρf ilm V 2 and its values are given in Figure 4a versus the hardening time and the NPs assembly film thickness. As expected the nanocontact hardening leads to an increase of the elastic modulus by a factor of 5-6. The larger elastic modulus remains around 10 times smaller than the bare silica materials (∼ 60 − 90 GPa). Consistently with the shrinkage measurements (Figure 2d), an asymptotic behavior is also observed which confirms that after a given time, the NPs assembly is entirely consolidated. Beside the consolidation mechanism we observe some fluctuations of the values of the elastic modulus of so-called thick films (> 210 nm) that may come from sub-micrometer cracks that we will discuss in the following.

It is important to remind that the thickness and the sound velocity are physical properties related to the bulk of the NP assembly and both of them show an asymptotic behavior indicating that the hardening process efficiency (i.e. transformation of VdW into H-CV bonds) saturates in the NPs network after a critical time. The analysis of the IR-vibrational bands also confirms this saturation with even faster saturation effect as already mentioned. This is in contrast with the contact angle that continuously evolves with the hardening time after the transformation from a VdW to a H-CV bonds. The long exposure to H2 O-NH3 vapors appears to modify the film surface energy due very likely to NH3 adsorption as suspected by 9



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XPS measurements (Figure S4). Such presence of NH3 on silica surface, already observed previously on silica based materials, 33 might explain why the contact angle evolves with the hardening time (i.e. upon the NH3 exposure) while the nanocontact stiffness (hydrogencovalent bonds) are unaffected as demonstrated by nanoacoustics experiments. This first discussion is a clear evidence of the necessity to combine experiments to really disentangle the surface energy of NPs from the relevant nanocontacts participating in the effective elastic stiffness of the assembly of these NPs. In this particular case, we show that contact angle alone cannot provide relevant information to the mechanical network established between nanoparticles because it does not provide information on the nature and robustness of NPs nanocontacts. Moreover, it is remarkable to notice that the contact angle evolves in the same way for thin (H=70 nm) and thick (H=210 nm) films (see Figure 2c), indicating the thermodynamic properties of the surface does not depend on the film thickness. On the opposite, with the same set of materials, the ultrafast nanoacoustic measurements reveal that some variations of the elastic modulus appear when comparing the different samples (Figure 4a). We observe that systematic largest elastic modulus is achieved for the thinner samples (70 nm) whatever the hardening time. We have plotted, for a given hardening time, the thickness dependence of the elastic modulus as shown in Figure 4b. There is a tendency of elastic modulus softening with increasing thickness that we possibly attribute to the apparition of some submicrometer defects that we have revealed by different methods as shown below. First of all, while scrutinizing the near-UV optical reflectivity signals (Figure 4c), we have observed that for thick colloidal films, an increase of the UV light scattering clearly appears and is evidenced by the development of a tail in the optical reflectivity spectrum in the range 250-450 nm. Dark-field imaging also allows the observation of the diffuse light scattering induced by surface defects (Figure 4d). A proper model that is fully described in an up-coming paper 34 permits to extract the diffuse scattering intensity (ID ) which is ascribed to a volume of the defects. As shown in Figure 4d, these defects volume clearly increases with the thickness whatever the nature of the nanocontacts but with a variable threshold effect. This threshold

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is indeed dependent of the nature of the nanocontacts: the defects appear above a critical thickness which is larger when the nanocontacts are soft (VdW bonds) compared to the case of a network consolidated with CV-H bonds. This clearly indicates that while strengthening the network, the hardening process leads to larger defects volume and that hardening makes somehow the lattice more brittle. This optical measurements are consistent with Figure 4b where, within the thickness range of 70-300nm, only elastic modulus of CV-H colloidal films is affected by defects while the elastic modulus of VdW films does not evolve with thickness in this range. The morphology of these defects has been well evidenced by Atomic Force Microscopy (AFM) images shown in Figure 4e where characteristic cracks appear on the surface with a regular spatial distribution. The in-plane extension of the cracks is around 300 nm (Figure 4f). There are craters with typical width and depth of around 50 nm. We remind that ultrafast nanoacoustic experiments were conducted with laser focusing area of around πr2 with r ∼ 5 µm. This means that the extracted modulus is an average measurement that includes the contribution of tens of cracks. The cracks dimension is of the order of the optical wavelengths (250-450 nm) where near-UV optical scattering has been identified (Figure 4c). The effect of cracks on elastic properties of macroscopic solids is a long standing research and different effective models exist to describe it. 35 For non interacting cracks and with anisotropic distribution, an effective Young modulus E was proposed has a function of the cracks volume fraction pcracks : 35

E = E0 (1 + π × pcracks )−1

(1)

The definition of pcracks is not straightforward since it can depend on the shape of the P 3 defects. We will assume here the simplest situation where pcracks = V1 αφ , where V is a P 3 representative volume of the material within which a volume of defects Vcracks = αφ is present (φ is the characteristic radius of the defect). α is a parameter which depends on the

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geometry of the crack. If the cracks volume fraction pcracks is small, then the effective elastic modulus becomes E ≈ E0 (1 − π × pcracks ), revealing a linear decrease of the elastic modulus with defects concentration. This behavior is in line with our observations even if more data need to be accumulated to verify this model. Moreover, to establish a more definitive relationship between the measured longitudinal elastic modulus (M) and the above model that uses the Young modulus E, the Poisson coefficient ν of the silica NPs assembly is needed 1−ν since M = E ((1−2ν)(1+ν) 2 ) , but is unknown up to now.

Conclusions This complete study shows that it is possible to follow and to control the transformation of the nanocontacts physical and chemical properties in a disordered assembly of silica nanoparticles from weak coupling (Van der Waals bonds) to a strong coupling regime (covalent-hydrogen bonds). Moreover, we have established a correlation between the nanocontact properties (chemical nature of the molecular landscape, surface energy of the NPs) and the collective mechanical response with ultrafast nanoacoustic method which is a non-destructive/noncontact method. This nanomechanical analysis provides some new insight on the long-range mechanical interaction within the NPs network, and in particular has permitted to reveal the role of the cracks appearing when the NPs assembly film thickness increases or when the elasticity is strengthened from a VdW to a CV-H bonded nanocontact. We have discussed the effective elastic modulus of this NPs assembly in presence of a given volume fraction of cracks resulting from the strengthening of the nanocontact (CV-H bonds). This work shows that it is possible to quantify the mechanical quality of a NPs assembly and this approach could be extended to any kind of NPs network a priori.

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Supporting Information Available Algorithm giving the calculation the refractive index and the thickness of the colloidal layers. XPS measurements on colloidal layers before and after the amonia curing (hardening process). This information is available free of charge via the Internet at http://pubs.acs.org.

Acknowledgement J. A. Thanks the CEA and Region Pays de la Loire for his PhD grant.

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Controlling the Nanocontact Nature and the Mechanical Properties of

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