A tandem polyol process and ATRP used to design new processable hybrid exchange-biased Co_xFe_3−xO₄@CoO@PMMA nanoparticles

Abstract

We report herein the synthesis and characterization of exchange-biased Co_xFe_3−xO₄@CoO@PMMA nanohybrids in which the inorganic core comprised a highly crystallized polyol-made ferrimagnetic cobalt iron spinel oxide nanoparticles perfectly epitaxied to an very thin antiferromagnetic polycrystalline cobalt monoxide CoO shell, whereas the organic coating comprised a layer of dense poly(methyl methacrylate) (PMMA) chains. PMMA brushes were grown by atom transfer radical polymerization (ATRP). The polymerization was monitored by combining infrared absorption (FTIR) and X-ray photoelectron (XPS) spectroscopy. The resulting nanohybrids are considered as potential building blocks for flexible magnetic recording devices. Their magnetic properties were then investigated by low temperature field cooling (FC) and zero field cooling (ZFC) magnetometry experiments. The obtained results agree fairly with a direct dependency of the coercivity and the remanence with the polymer chain length and then with the interparticle distance, emphasizing the role of material processing in the design of tailored flexible polymer based hybrid materials.

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Introduction

Size tailoring magnetic particles at the nanometer scale have attracted increasing interest among the scientific community during the last few decades because of their original magnetic properties compared to those of their bulk counterparts, leading to the emergence of various research topics in different fields. Nowadays, they are indifferently and seriously considered as biomedical probes,¹ magnetic sensors² and magnetic recording media.³ One of the main challenges for the synthesis of such granular materials is the control of their aggregation state to control their collective magnetic response.⁴

Focusing on magnetic recording applications, high storage memory applications require dense but well-separated magnetic nanoparticles (NPs). Thus, in an optimized geometry, in the hypothesis of particles of a few nanometer in size, each one storing one bit of data, densities of 150 Gbit cm⁻² in range can be reached. The ability of such NPs to react with any magnetic stimulation is of course dependent on their intrinsic properties (composition, size, and shape) but also on their spatial arrangement.^5,6 For such a purpose, two issues have to be addressed. First, the intrinsic magnetic anisotropy of NPs, which should be high enough to assure the thermal stability of the magnetization at the operating temperatures, and second, the mutual influence between NPs that should be low enough to decrease the dipolar interaction (DI) effects.

In this context, several studies have been achieved to build exchange-biased NPs composed of ferro or ferrimagnetic (F) cores coated by nanocrystalline antiferromagnetic (AF) layers, leading to coupled F–AF interfaces. Thanks to the exchange-bias (EB), these nanostructures exhibit an enhanced magnetic anisotropy, an improved thermal magnetization stability and a significantly higher blocking temperature (TB).^7–11 Surprisingly, despite an impressive number of studies on this topic,^12–18 few of them focused on the interplay between the exchange bias (EB) and dipolar interactions (DI), while DI are known to be able to greatly affect the magnetic responses of any magnetic NPs in any type of assembly, particularly 2D⁸ and 1D¹⁹ ones.

Moreover, while some research groups have worked towards tuning the NP–NP distance to vary and control the dipolar interactions (DI), functionalizing NPs with dendrimers or fatty acids of different lengths,^20,21 or coating them by a more or less thin silica shell,^22,23 they have mainly focused on iron oxide particles, which are not the best magnetic systems for magnetic recording applications. They have shown that the NP–NP distance necessary to observe non-interacting particles is dependent on the nature of the spacer for similar NP sizes and have also shown that magnetic NP surface modification is a great challenge regarding the strong tendency of these nanomagnets to agglomerate. In the case of polymer spacers, it was established that optimization of the hybrid interface is required before building any efficient magnetic assembly. The hybrid interface, also called the interphase, is the area in which the chemical bonding between both the polymer chain and the magnetic particle is established. Therefore, despite polymer coatings having attracted increasing interest,²⁴ it has been necessary to wait for the development of fundamental processes of living radical polymerization (e.g. atom transfer radical polymerization ATRP, reversible addition–fragmentation chain transfer RAFT or nitroxide-mediated polymerization)²⁵ before producing robust polymer-particle links and then achieve polymer brushes of controlled molecular weight and polydispersity index, within a satisfying grafting density.

Clearly, for magnetic recording applications (as for others), prior to NP organization, (i) NPs have to be well defined in composition and morphology to obtain the best properties and (ii) the NPs must be judiciously functionalized to ensure their compatibility with the polymer matrix and (iii) the process used to assemble the resulting hybrids has to be optimized to fit the properties to the desired application.^6,26–28

To satisfy these requirements, we decided to (i) produce exchange-biased, almost isotropic in shape, about 10 nm sized and perfectly epitaxied Co_xFe_3−xO₄@CoO core@shell NPs using the polyol process,²⁹ (ii) grow, directly at their surface, poly(methyl methacrylate) (PMMA) chains using the ATRP grafting from technique and (iii) study their magnetic properties by comparison to those of bare NPs with a special emphasis on the interparticle distance effects.

Experimental section

Synthesis and characterization of the exchange-biased core–shell nanoparticles

Spherical oxide-based core@shell nanoparticles were synthesized via seed mediated growth in polyol.²⁹ In practice, 10 nm sized Fe_3−xO₄ iron oxide NPs were produced by forced hydrolysis of anhydrous iron(II) acetate salt (4.0 g) in diethylene glycol (250 mL) mixed with deionized water (1 mL). The as-produced NPs were then separated and used as seeds to grow a shell of about 1 nm thickness of cobalt oxide in a fresh diethylene glycol solution of cobalt(II) acetate (0.05 mol L⁻¹) in the presence of the required amount of deionized water (0.45 mL). The structural and microstructural analysis of the produced powder using mainly X-ray diffraction (XRD), ⁵⁷Fe Mössbauer spectrometry, SQUID-magnetometry and transmission electron microscopy (TEM) confirmed the production of the desired hetero-nanostructures and the starting iron oxide phase evolving to the Co_xFe_3−xO₄ solid solution during the CoO shell growth. This chemical evolution is attributed to the diffusion of free Co²⁺ ions from the reaction solution into the spinel lattice of the seeds and their localization in the octahedral iron cation vacancies.³⁰ Obviously, they exhibit an EB feature.

Synthesis and characterization of the thermal initiator for ATRP

The synthesis of 2-phosphonooxy-2-bromo-2-methylpropanoate was realized following the procedure described by Babu et al.³¹ Typically, 5.2 mL of pure ethyl ester were dissolved in 90 mL of anhydrous tetrahydrofuran (THF) and purged with argon. 3.5 mL of POCl₃ was added dropwise and the resulting mixture was maintained at 0 °C for 1 hour. It was further stirred for 3 hours at room temperature (RT). Then, it was diluted with 60 mL of water and extracted three times with 100 mL of chloroform. The organic extract was dried over anhydrous sodium sulfate, filtered and dried until a viscous yellow liquid was recovered. Routine spectroscopies confirmed the formation of the desired molecule:

FTIR (ν in cm⁻¹, KBr technique): 2933 (C–H), 2309 (P–O–H), 1733 (C=O), 1272 (P=O), 1160 (C–O of carbonyl), 1074 and 990 (C–O and P–O bond, respectively).

¹H NMR (400 MHz, δ in ppm, CDCl₃): 6.20 (2H), 4.24 (2H), 3.81 (2H), 1.88 (6H).

³¹P NMR (400 MHz, δ in ppm, CDCl₃): 0.67 ppm.

Polymer brushes grafting by ATRP

First, 280 mg of core@shell NPs and 700 mg of initiator (ratio: 2.5/1 in mass) were introduced to 28.3 mL of THF and sonicated at RT for 24 hours. The resulting nanohybrids, quoted herein as NP-Br, were recovered by several cycles of centrifugation and washing with ethanol, to be subsequently dispersed (80 mg) in 80 mL of acetonitrile before adding 22.7 μL of N,N,N′,N′,N′′-pentamethyldiethylenetriamine (PMDTA) (1 × 10⁻⁴ mmol) and 5 mL of methyl methacrylate (47 mmol). The mixture was degassed with argon and mechanically stirred for 1 hour. 16 mg of CuBr (11 mmol) was added and the solution was heating to around 30 °C and sonicated under an inert atmosphere. The polymerization time was tuned from 1 hour to 3 hours to influence the length of the polymer chains. The reaction was quenched by opening the system and diluting the solution in THF and hexane. Hairy hybrid NPs were recovered by centrifugation and washing with THF. They were designated in the further sections as NP-bPMMAh, h corresponding to the polymerization time, namely, 1, 2 or 3 hours.

Characterization of the hairy hybrid NPs

The different reaction steps were monitored by ATR-FTIR on a Thermo Nicolet 8700 spectrometer equipped with a diamond crystal (50 scans, 4 cm⁻¹ resolution). Thermogravimetric analyses (TGA) were performed in air on a Labsys-Evo at a heating rate of 10 °C min⁻¹. X-ray photoelectron spectroscopy (XPS) measurements were performed on the as-prepared and functionalized NPs using a Thermo VG ESCALAB 250 instrument equipped with a micro-focused, monochromatic Al Kα X-ray source (1486.6 eV) and a magnetic lens. The X-ray spot size was 500 μm (15 kV, 150 W). The spectra were obtained in the constant analyzer energy mode with pass energy of 150 and 40 eV for the general survey and the narrow scans, respectively. The samples were fixed on sample holders and out gassed in the fast entry airlock (2 × 10⁻⁷ mbar). “Avantage” software was used for data acquisition and processing. The C 1s line of 285 eV was used as a reference to correct the binding energies for charge correction. The molecular weight and polydispersity of the macromolecules were determined by size exclusion chromatography (SEC). Dionex Ultimate 3000 SEC apparatus comprised one Polymer Laboratories MesoPore 3 μm column (300 × 7.5 mm) located in an isothermal oven (35 °C) in series with a UV-visible detector (wavelength of detection λ = 254 nm). The flow rate of eluent (HPLC grade THF) was set at 1.0 mL min⁻¹. Calibration of the column was performed using a set of near-monodisperse polystyrene standards (EasiCal PS-2 Polymer Laboratories, M_p = 580–20 650 g mol⁻¹ range). For SEC analysis, each polymer solution (1 mg dissolved in 2 mL THF) was introduced via a manual injector loop (injected volume = 20 μL). Data acquisition and analysis were executed using Chromeleon V 6.80 software supplied by Dionex. Finally, magnetization measurements were performed at low temperature, 5 K, varying the magnetic field between −70 kOe and +70 kOe, using a Quantum Design MPMS 3 magnetometer in field cooling (FC) conditions. Typically, the hybrid powders were compacted in a sampling diamagnetic plastic tube and cooled from RT to 5 K under an applied magnetic field of 70 kOe before collecting the experimental data.

Results and discussion

Growth of the PMMA chains at the surface of the core@shell NPs

The general approach that was developed herein to produce the hybrid polymer decorated NPs is a two-step pathway. First, 2-phosphonooxy-2-bromo-2-methylpropanoate (1) was grafted on the surface of the CoO shells as an ATRP initiator (Scheme 1) through the phosphate acid functional group. Then, the PMMA chains were grown by ATRP at RT, in the presence of a CuBr/PMDETA complex, without the addition of a sacrificial initiator (Scheme 1).

Scheme 1 General approach to coupling the polyol process and ATRP to design new processable exchange-biased nanoparticles for magnetic recording applications.

The efficiency of the grafting process by the ATRP initiator has been checked by analysis of the surface state using XPS. The survey spectrum of the brominated NPs (NP-Br) was found to be similar to that of the as-prepared core–shell NPs (Fig. 1a and b), exhibiting all the peaks associated to O 1s (529.7 eV), Fe 2p_3/2 (711.4 eV) and Co 2p_3/2 (780.9 eV) and new peaks assigned to P 2p in a phosphonate compound (133.2 eV) and Br at 182.7 eV (3p_3/2) and at ∼69 eV (3d_5/2) as a shoulder of the peak assigned to Co 3p_3/2 (60.5 eV). The peaks at 781.3 eV and 796.9 eV correspond to the characteristic doublet related to Co and are assigned to Co 2p_3/2 and Co 2p_1/2; two shake-up satellite peaks at 786.6 eV and at 803.1 eV were observed, corresponding to Co 2p_3/2 and Co 2p_1/2. The presence of these two peaks and highly intense satellites is consistent with the presence of Co²⁺ cations in the high-spin state. The Fe 2p high-resolution spectrum exhibited a main peak at 711.5 eV and a satellite at 723.8 eV, which are characteristic of the presence of a majority of Fe³⁺ cations. Thus, we can conclude that the successful grafting of 2-phosphonooxy-2-bromo-2-methylpropanoate on the core–shell nanoparticles was achieved whose crystallographic structure was preserved during the functionalization process. The chemical composition of as-produced NP and NP-Br are gathered in Table 1.

Fig. 1 XPS survey spectra of (a) bare Co_xFe_3−xO₄–CoO NPs (NP), (b) NP-Br, (c) NP-bPMMA1, and (d) NP-bPMMA3.

Table 1 Elemental atomic composition of the as-prepared Co_xFe_3−xO₄–CoO nanoparticles (NP) and their relative hybrids

Sample	Elemental atomic composition (%)
	Fe	Co	O	C	P	Br
NP (CoxFe3−xO4–CoO)	14.5	24.3	40.2	21.0	—	—
NP-Br	6.6	15.5	46.2	29.9	1.4	0.4
NP-bPMMA1	1.3	3.8	30.2	64.1	0.6	—
NP-bPMMA3	0.8	2.9	28.6	67.3	0.4	—

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The C 1s peak before the functionalization process demonstrated the presence of organic matter (polyol and acetate molecules) on the core–shell NP surface. This is a consequence of the polyol process, during which the solvent molecules also play the role in ligand complexation on the NP surface to control the growth step.

The XPS spectrum of NP-bPMMA1 is depicted in Fig. 1c. The Co 2p and Fe 2p peaks were strongly attenuated after 1 hour of polymerization, revealing the presence of a non-negligible polymer coating (Fig. 2c). After 3 hours of reaction (NP-bPMMA3), the Co 2p and Fe 2p peaks have almost disappeared, suggesting a higher content of organic matter (Fig. 2d).

Fig. 2 The high-resolution Fe 2p (left) and Co 2p (right) scan of (a) bare Co_xFe_3−xO₄–CoO NPs (NP), (b) NP-Br, (c) NP-bPMMA1, and (d) NP-bPMMA3.

The decomposition of the C 1s peak specific to PMMA was difficult to achieve because many sources of organic matter (PMMA, initiator, little residual polyol and acetates) are present on the surface of NPs and that of O 1s for PMMA was difficult too because of the presence of oxides constituting the core–shell NPs.

The cobalt atomic percentage was found to be higher than that of iron for all the nanoparticles, either bare or functionalized. This trend can be explained by considering the mean free paths in both oxides, according to eqn (1) and (2) given by Dench and Seah for oxide materials:³²

λ(k) = 0.55 × E_k^0.5 × a_k^1.5 (1)

Considering E_k being the kinetic energy of ejected core electrons (eV) and a_k the dimension of atom k; E_k is tabulated while a_k is calculated using eqn (2):

With A the molecular weight of the oxide (g mol⁻¹), ρ its density (kg m⁻³), n is the number of atoms per molecular unit and N_A, Avogadro number.

The mean free paths were found to be 1.5 nm for iron in CoO and 1.4 nm for cobalt in CoO. However, as the thickness of the CoO shell is about 1 nm, electrons can penetrate the whole Co-rich shell but pass through only 0.5 nm of the Fe-rich core. The difference in depth analysis explains that both the core and shell are not characterised in a similar way because XPS provides only surface information. Moreover, one can notice that there is not as much phosphorus as bromine. Indeed, the C–Br bond is relatively sensitive to X-ray and tends to be altered, whereas the P–O bond is quite stable; this leads to a higher P atomic percentage than that of Br.

As a first approximation, the PMMA thickness could be estimated considering the C 1s and Co 2p peaks and their relative intensities (see ESI†) assuming the carbon and cobalt signals were mainly due to the CoO shell (the cobalt content into the spinel lattice is weak³⁰) and the PMMA brushes (the molecular weight of the phosphonate initiator being negligible front to the polymer brushes), respectively. It was found to be 1.3 ± 0.1 nm for NP-bPMMA1 and 2.7 ± 0.2 nm for NP-bPMMA3. However, one has to remember that under ultra-vacuum conditions (like in the XPS analysis chamber), polymer chains have a strong propensity to collapse; therefore, the calculated thickness have to be considered for a polymer in a random coil conformation, meaning that XPS measurements do not allow the determination of the exact chain length. However, we can conclude that increasing the polymerization time leads to the formation of a greater amount of polymer on the NP surface.

The XPS data were found to be in good agreement with the qualitative FTIR spectroscopy results (see ESI†) because the spectrum of NP-Br exhibits ν_P−O (1009 cm⁻¹) and ν_P=O bands (1278 cm⁻¹) corresponding to phosphonate and an intense peak at 1733 cm⁻¹ corresponding to ν_C=O in the bromoester. After ATRP, the C–C skeleton at 750 cm⁻¹, C–O–C bands (1149 and 1192 cm⁻¹) and very intense ν_C=O at 1727 cm⁻¹ (characteristics of PMMA) can be observed, whereas all peaks previously assigned to the initiator disappeared, which can be explained by the formation of a layer of PMMA of significant thickness.

To determine the molecular weight M_w and polydispersity index I_p = M_w/M_n of the grafted PMMA thus formed, the hybrid NPs were subjected to degrafting using concentrated hydrochloric acid, in the presence of chloroform, followed by precipitation in methanol and drying. The degrafted polymer chains were then analysed by SEC and the obtained results were correlated to those inferred from the TGA (see ESI†), assuming the observed total weight losses were due to polymer brushes decomposition.

In practice, we deduced from the weight loss observed at 150 °C, during the TG analysis of NP-Br, the grafting density of the initiator molecules as 2.8 molecules nm⁻². For that we assumed that the mass loss at this temperature is mainly attributed to the degrafting of the initiator molecules and we determined from BET measurements the specific area of the as-prepared core@shell NPs as 119 m² g⁻¹. Note that the obtained grafting density value agrees with that reported in previous study.³¹

Second, we supposed that the polymer chains on NP-bPMMA1 and NP-bPMMA3 are almost homogeneous in length. This hypothesis was supported by the fact that the TGA of the produced hybrids showed a unique and drastic weight loss at 300 °C: 33 and 43 wt% for NP-bPMMA1 and NP-bPMMA3, respectively. Assuming that all “Br” initiator ends can initiate a polymerization, we could estimate the mean degree of polymerization to be about 6 and 8 units, respectively. All the results allowed us to calculate not only the chain length (M_w(calc)), but also the monomer conversion ratio and the grafting efficiency (see ESI†). The results are presented in Table 2.

Table 2 The monomer conversion and calculated and corrected PMMA molecular weight determined from TGA

Sample	Weight loss (%)	Conversion (%)	Mw(calc) (g mol^−1)	Mw(corr) (g mol^−1)
NP-bPMMA1	33	0.62	674	1241
NP-bPMMA3	43	0.78	833	1561

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From Table 2, we can conclude that the conversion is very low in comparison with previous studies performed with a quite similar initiator.³¹ We believe that this discrepancy between our results and the former ones was due to the dilution of the NPs and monomers in the reaction medium. Indeed, to avoid NP aggregation during the functionalization process and to improve the control the radical polymerization with a polydispersity index close to 1 during the chain grow process, we worked with very low concentrated starting systems.

Finally, we deduced from SEC experiments the polymer chain density (eqn (3)) and the grafting efficiency after the chains had been degrafted:

with W₂ the weight loss observed for polymer–oxide hybrids and W₁ the weight loss of initiators (28% and 16%, respectively). M_w, the PMMA molecular weight, was measured by SEC and found to be equal to 816 g mol⁻¹. The specific area is 119 g m⁻² and N_A is Avogadro constant. The resulting chain density was found to be 1.4 chain per nm².

Then, the initiator efficiency could then be calculated as follows (eqn (4)):

where D_polym and D_init represent the densities of the polymer chain and initiator molecules, respectively. The initiator efficiency was found to about 53%, which is in good agreement with previous studies on similar materials (50%).³¹ This can be explained by the length of the carbon chain of the initiator. Green et al. obtained a higher polymer chain density using initiators with either very long (15C) or short (3C) carbon chains.³³ Short molecules exhibit less conformations and the brominated ester is easily accessible for the metallic complex and the monomer. In the case of longer initiators, the number of conformations increases and only few of them can initiate a polymerization reaction. For very long initiators, the number of conformations increases more and more, promoting rather more access for the monomer and Cu complex.

From the SEC measurements, we know that only half of the initiator molecules are active. On the basis of this information we propose a methodology to calculate the molecular weight (M_w) of the grown polymer chains from the TGA measurements (Table 2). Indeed, the degrafting step is quite hard to realize and requires more matter than TGA for doing the analysis. Therefore, we use eqn (14) in the ESI,† considering only half the quantity of initiators (2.3 × 10⁻⁵ mol instead of 4.3 × 10⁻⁵ mol). The results are compiled in Table 2 (M_w(corr)).

Magnetic properties of the produced PMMA-based nanohybrids

5 K-magnetization versus magnetic field variation was recorded on all the samples in FC conditions (Fig. 3). A hysteresis loop with significant coercivity and remanence is systematically observed. A horizontal shift along the filed axis of the loops is also observed, which is in agreement with the EB feature.

Fig. 3 5 K-magnetization versus magnetic field variation for the bare core–shell NP and their related hybrids, NP-bPMMA1 and NP-bPMMA3.

Then, for all the samples, we could calculate the coercive field, H_C, following eqn (5):³⁴

H_C (FC) = (−H_C− + H_C+)/2 (5)

where H_C+ and H_C− correspond to the positive and negative values of the field for a zero magnetization. The results are grouped in Table 3.

Table 3 The magnetic properties of the Co_xFe_3−xO₄–CoO NPs and their relative hybrids

	HC+ (Oe)	HC− (Oe)	HC (Oe)	Ms (emu g^−1)
NP (CoxFe3−xO4–CoO)	9743	−12 985	9648	65
NP-Br	7387	−10 010	9897	60
NP-bPMMA1 (1 h)	11 396	−13 699	11 503	46
NP-bPMMA3 (3 h)	10 964	−13 361	11 086	33

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One can notice the decrease in the magnetization as the amount of diamagnetic polymer content increases. Designing the saturation magnetization of free polymers (assumed to be zero), bare core@shell NPs and their related PMMA-hybrids as M_sat (polymer), M_sat (NP) and M_sat (NP–polymer), the PMMA weight content, x, can be deduced by eqn (6), neglecting in a first approximation the initiator weight content:

x × M_sat(polymer) + (1 − x) × M_sat(NP) = M_sat(NP–polymer) (6)

We found x = 31 and 45 wt% for NP-bPMMA1 and NP-bPMMA3, respectively, close to those determined by TGA and shows once again that the percentage weight grafted from the NP surface is significant.

Focusing on the coercivity of the produced hybrids, it is found that the coercive field measured at low temperature on bare NP, NP-Br and NP-bPMMA1 increases progressively in agreement with the inter-magnetic core distance with a slight decrease in the exchange field H_E defined by eqn (7):³⁴

H_E (FC) = (+H_C− + H_C+)/2 (7)

Interestingly, this tendency is not respected between NP-bPMMA1 and NP-bPMMA3 whereas more polymers are present on the surface of the NPs in NP-bPMMA3, its coercive field appears to be smaller than that of the less polymer rich NP-bPMMA1 sample. One may suspect that the polymer chains collapse around the inorganic cores when the polymerization degree increases, leading to a shorter NP–NP distance (Scheme 2).

Scheme 2 The polymer chain conformation on the surface of the magnetic NPs.

Conclusions

We successfully achieved the synthesis of magnetic exchange-biased core–shell Fe₃O₄@CoO nanoparticles via a polyol process and their further functionalization by PMMA brushes using ATRP, avoiding any degradation of the chemical composition of the nanoparticles and keeping the exchange bias property. Increasing the polymerization time induced the presence of more organic matter on the NP surface and longer NP–NP distances than in a sample of the as-prepared NPs. Then, it appeared that increasing the average distance between the NPs increased their low temperature coercive field (H_C) too.

These multifunctional magnetic nanoparticles, exhibiting exchange bias for more thermal stability, as well as a better miscibility in PMMA, are interesting for fundamental studies on the collective magnetic properties of extended arrays of nanoparticles.³⁵ They also would allow us to prepare flexible functional hybrid PMMA-NPs films with enhanced properties.³⁶

Acknowledgements

The authors would like to thank Dr Philippe Decorse and Dr Claire Mangeney for XPS measurements and fruitful discussions on the XPS results, Alexandre Chevillot for technical assistance for the TGA analyses. ANR (Agence Nationale de la Recherche) and CGI (Commissariat à l'Investissement d'Avenir) are gratefully acknowledged for their financial support of this study through Labex SEAM (Science and Engineering for Advanced Materials and devices) ANR 11 LABX 086, ANR 11 IDEX 05 02.

Notes and references

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Footnote

† Electronic supplementary information (ESI) available. See DOI: 10.1039/c6ra06963k

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