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Progress in Materials Science 67 (2015) 1–94

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Progress in Materials Science

journa l homepage : www.e lsev ie r . com/ loca te /pmatsc i

Nanoindentation in polymer nanocomposites

http://dx.doi.org/10.1016/j.pmatsci.2014.06.0020079-6425/� 2014 Elsevier Ltd. All rights reserved.

⇑ Corresponding authors.E-mail addresses: [email protected] (A.M. Díez-Pascual), [email protected] (A. Flores).

Ana M. Díez-Pascual a,⇑, Marián A. Gómez-Fatou a, Fernando Ania b,Araceli Flores b,*

a Department of Polymer Physics, Elastomers and Energy Applications, Institute of Polymer Science and Technology (ICTP-CSIC),Juan de la Cierva 3, 28006 Madrid, Spainb Department of Macromolecular Physics, Institute for Structure of Matter (IEM-CSIC), Serrano 119, 28006 Madrid, Spain

a r t i c l e i n f o

Article history:Received 12 June 2014Accepted 14 June 2014Available online 17 July 2014

Keywords:NanoindentationPolymer nanocompositesModulusHardnessFiller reinforcement

a b s t r a c t

This article reviews recent literature on polymer nanocompositesusing advanced indentation techniques to evaluate the surfacemechanical properties down to the nanoscale level. Special empha-sis is placed on nanocomposites incorporating carbon-based(nanotubes, graphene, nanodiamond) or inorganic (nanoclays,spherical nanoparticles) nanofillers. The current literature oninstrumented indentation provides apparently conflicting informa-tion on the synergistic effect of polymer nanocomposites onmechanical properties. An effort has been done to gather informa-tion from different sources to offer a clear picture of the state-of-the-art in the field. Nanoindentation is a most valuable tool forthe evaluation of the modulus, hardness and creep enhancementsupon incorporation of the filler. It is shown that thermoset, glassyand semicrystalline matrices can exhibit distinct reinforcing mech-anisms. The improvement of mechanical properties is found tomainly depend on the nature of the filler and the dispersion andinteraction with the matrix. Other factors such as shape, dimen-sions and degree of orientation of the nanofiller, as well as matrixmorphology are discussed. A comparison between nanoindenta-tion results and macroscopic properties is offered. Finally, indenta-tion size effects are also critically examined. Challenges and futureperspectives in the application of depth-sensing instrumentationto characterize mechanical properties of polymer nanocompositematerials are suggested.

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Abbreviations

AAER acrylic electrophoretic resinAFM atomic force microscopyAPTES aminopropyltriethoxysilaneCB carbon blackCCNT coiled carbon nanotubeCEAR cathodic electrophoretic acrylic resinCF carbon fibreCNC cellulose nanocrystalsCNF carbon nanofibreCOF coefficient of frictionCSH calcium silicate hydrateCSM continuous stiffness measurementCVD chemical vapour depositiondeox-ND decarboxylated nanodiamondDMA dynamic mechanical analysisDSC differential scanning calorimetryDSI depth-sensing indentationEG exfoliated graphiteEGO exfoliated graphene oxideEGS exfoliated graphene sheetsEOGCNF exfoliated oxidized graphitized carbon nanofibreEPDM ethylene-propylene-diene rubberFP-POSS fluoropropyl polyhedral oligomeric silsequioxaneFG few-layer grapheneFGS functionalized graphene sheetsGBL gamma-butyrolactoneGF glass fibreGlymo glycidoxy-propyltrimethoxysilaneGNP graphene nanoplateletsGO graphene oxideGP graphite plateletsGPTS glycidoxypropyltrimethoxysilaneGS graphene sheetsHA hydroxyapatiteHDDA 1,6-hexanediol diacrylateHDPE high density polyethyleneHEMA 2-hydroxyethyl methacrylateIBMA isobornyl methacrylateIF inorganic fullerene-likeISE indentation size effectKF kenaf fibreLbL layer-by-layerMBZ methylbenzoateMDMA monomeric dimethacrylatesMEH-PPV poly[2-methoxy-5-2(2-ethylhexyloxy-p-phenylenevinylene)]MEMO (3-(trimethoxysilyl) propyl methacrylateMMT montmorilloniteMPTES methacryloxypropyltriethoxysilaneMPTMS methacryloxypropyltrimethoxysilaneMWCNT multi-walled carbon nanotubesNA non availableNa-MMT sodium montmorillonite

2 A.M. Díez-Pascual et al. / Progress in Materials Science 67 (2015) 1–94

ND nanodiamondND-NH2 aminated nanodiamondOC-POSS octa[(epoxycyclohexylethyl) dimethylsilyloxy]silsesquioxaneODA octadecylamineOH-POSS octa[(epoxyhexyl) dimethylsilyloxy]silsesquioxaneOib-POSS octaisobutyl POSSOM optical microscopyOMMT organically modified montmorilloniteOP Oliver and PharrOTES octyltriethoxysilaneox- oxidizedPA polyamidePAA polyacrylic acidPAH polyallylamine hydrochloridePAM polyacryl amidePAN polyacrylonitrilePC polycarbonatePDDA poly(dimethyldiallylammonium chloride)PDMS polydimethylsiloxanePE polyethylenePEDOT-PSS poly(3,4-ethylenedioxythiophene)-poly(styrene sulphonate)PEEK poly(ether ether ketone)PEES poly(ether ether sulphone)PEI polyethyleneiminePEN polyethylene naphthalatePEO polyethylene oxidePET polyethylene terephthalatePHB poly(3-hydroxybutyrate)PHBV poly(3-hydroxybutyrate-co-3-hydroxyvalerate)PHEA poly(2-hydroxyethyl acrylate)PHO poly(3-hydroxyoctanoate)PHPS perhydropolysilazanePI polyimidePLA polylactic acidPLLA poly(L-lactic acid)PLS polymer/layered silicatePMA polymethacrylatePMMA poly(methyl methacrylate)PNIPAm poly(N-isopropylacrilamide)POSS polyhedral oligomeric silsequioxanePOSS-OH allyl alcohol terminated POSSPP polypropylenePPC poly(propylene carbonate)PP-g-MA polypropylene grafted-maleic anhydridePPS polyphenylene sulphidePPy polypyrrolePS polystyrenePSS polystyrene sulphonatePSBA poly(styrene-co-butyl acrylate)PU polyurethanePUF poly(urea–formaldehyde)PVA polyvinyl alcoholPVC polyvinyl chloridePVDF polyvinylidene fluoride

A.M. Díez-Pascual et al. / Progress in Materials Science 67 (2015) 1–94 3

PVK poly(9-vinyl carbazole)QD quantum dotrGO reduced graphene oxideRT room temperatureS sonicationSAXS small angle X-ray scatteringSB sonication and ball millingSC super critical solventSCCO2 supercritical carbon dioxideSCF short carbon fibreSEM scanning electron microscopySEBS styrene–ethylene/butylene–styrene(SiOEt)3 triethoxysilylSSE single-screw extrusionSWCNT single-walled carbon nanotubesTEM transmission electron microscopyTEOS tetraethyl orthosilicate (or tetraetoxysilane)TEG thermally expanded graphiteTGA thermogravimetric analysisTMOS tetramethyl orthosilicateTSE twin-screw extrusionTsp-POSS trisilanolphenyl POSSUHMWPE ultra-high molecular weight polyethyleneUPE unsaturated polyester resinUV ultra violet radiationVGCNFs vapour-grown carbon nanofibresWAXS wide angle X-ray scattering


Contents

1. Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 52. Recent advances on instrumented indentation applied to polymer materials: methods and

developments . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7
2.1. Depth-sensing indentation . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7
2.1.1. Determination of elastic modulus and hardness. Standard methods of analysis applied topolymer materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7

2.1.2. DSI time-dependent properties . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 102.1.3. Indentation size effect . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 13

2.2. Correlation mechanical properties-nanostructure. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 152.3. Comparison of indentation studies with bulk techniques . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 17

3. General aspects on polymer nanocomposites. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 184. Polymer nanocomposites incorporating carbon nanofillers. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 18

4.1. Carbon nanotubes . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 18
4.1.1. Synthesis methods and types of CNTs. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 184.1.2. Preparation of polymer/CNT nanocomposites. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 384.1.3. Modulus changes in polymer/CNT nanocomposites: comparison of thermoplastic and
thermoset matrices . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 394.1.4. Hardness . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 454.1.5. Indentation size effect . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 454.1.6. Creep . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 47

4.2. Graphene . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 48
4.2.1. Modulus enhancements in graphene/polymer nanocomposites . . . . . . . . . . . . . . . . . . . 49


4.2.2. Comparison with macroscopic mechanical properties . . . . . . . . . . . . . . . . . . . . . . . . . . . 504.2.3. Indentation size effect . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 524.2.4. Creep . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 52

4.3. Other organic nanofillers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 52
4.3.1. Modulus enhancements in organic/polymer nanocomposites . . . . . . . . . . . . . . . . . . . . . 534.3.2. Comparison with macroscopic mechanical properties . . . . . . . . . . . . . . . . . . . . . . . . . . . 564.3.3. Indentation size effect . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 574.3.4. Creep . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 57
4.4. Main features on organic fillers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 57
4.4.1. Carbon nanotubes. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 574.4.2. Graphene. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 584.4.3. Other organic fillers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 58
5. Polymer nanocomposites incorporating inorganic nanofillers. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 58
5.1. Layered silicates . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 59
5.1.1. Characteristics, preparation methods and types of nanocomposites. . . . . . . . . . . . . . . . 595.1.2. Modulus and hardness changes in polymer/layered silicate nanocomposites . . . . . . . . 605.1.3. Other layered silicate/polymer systems . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 645.1.4. Comparison with macroscopic properties. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 655.1.5. Indentation size effect . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 665.1.6. Creep properties . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 66

5.2. Spherical inorganic nanoparticles. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 66
5.2.1. Characteristics, preparation methods and types of nanocomposites. . . . . . . . . . . . . . . . 665.2.2. Modulus changes in nanocomposites containing spherical inorganic nanoparticles . . . 685.2.3. Indentation size effect . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 745.2.4. Creep properties . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 755.2.5. Comparison with macroscopic mechanical properties . . . . . . . . . . . . . . . . . . . . . . . . . . . 75
5.3. Other inorganic nanofillers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 765.4. Main features on inorganic nanofillers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 78

5.4.1. Layered silicates . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 785.4.2. Spherical inorganic nanoparticles . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 78

6. Comparison of the reinforcement effect of different nanofillers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 78
6.1. Effect of the type of nanofiller . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 786.2. Influence of the geometry of the filler and orientation effects. . . . . . . . . . . . . . . . . . . . . . . . . . . . 816.3. Comparison of properties at the macro and the nanoscale . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 82
7. Conclusions and future perspectives. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 83Acknowledgments . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 85References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 85

1. Introduction

Depth-sensing indentation (DSI) represents nowadays one of the principal techniques for the mechan-ical characterization of materials. The method monitors the penetration of an indenter into the materialsurface during the application and release of a load [1]. At the present time, the most advanced DSIinstruments can produce indentations with depths of only a few tens of nanometers, most of them alsooffering the possibility of approaching an upper limit in the micron regime [2]. The terminology ‘‘DSI’’defines the principle of measurement to attain indentation data. Usually, this technique is also referredto as ‘‘nanoindentation’’, even when penetration depths of a few microns are involved.

A typical DSI test includes a loading-hold-unloading cycle (see Fig. 1). Elastic contact consider-ations are usually adopted to analyse load-depth curves. Most commonly, hardness, H, and quasi-sta-tic elastic modulus, E, values are derived assuming linear elastic behaviour at the onset of unloading[3,4]. Because the new generation of DSI testers involves the most advanced engineering approachesto accurately record penetration depths in the nanoscale range, special care should be paid to calibra-tion procedures. These are usually friendly integrated in the software developed for collecting theindentation data. In addition, as DSI devices have sophisticated over the years to widen the range ofapplications, important considerations related to instrumental artefacts and to data analysis are mat-

Displacement [nm]

0 100 200 300 400 500

Load

[mN

]

0.0

0.2

0.4

0.6

0.8

1.0

s

hmaxhp

Pmax

Fig. 1. Load as a function of indentation depth for a semicrystalline PEEK sample (degree of crystallinity of �40%). Unpublishedresults.


ters of active discussion within the indentation community. This is reflected in the number of papersfully or partially dedicated to the problems encountered in obtaining meaningful DSI data [5–7].The application of DSI to polymers raises additional concerns [6,8–13]. On the one hand, linear elas-ticity cannot be straightforward assumed and particular strategies need to be taken into account. Onthe other hand, the calibration procedures should be carefully examined because some of themrequire pure elastic–plastic response. Finally, a comprehensive characterization of polymer materialsby means of DSI requires specific approaches to extract information on their time dependence.A step forward in this direction has been produced with the introduction of dynamic DSI, or continu-ous stiffness measurements (CSM) [14]. CSM superimposes an oscillatory force to the quasi-staticloading opening up the possibility of using DSI devices as micro or nanoscale dynamic mechanicalanalysers.

DSI data are most frequently acquired using purpose-designed instrumentation. Alternatively,Atomic Force Microscopy (AFM) can be used as a depth-sensing instrument provided that careful cal-ibrations to achieve meaningful load and depth data are carried out [15,16]. The advantage of usingatomic force microscopy for indentation purposes is the higher displacement sensitivity and superiorimaging facilities. In contrast, compared to instrumented indentation, the technique is more liable toinstrumental errors and miscalibrations, especially in force modulation experiments [16].

DSI is being progressively incorporated as a routine technique for mechanical characterization ofpolymer materials and yet a number of decisions concerning the device calibration, the test procedureand the test parameters are of great importance and should be critically examined by an experiencedresearcher. The attractiveness of instrumented indentation relies on the ability to extract mechanicalproperties from a small local deformation. This is extremely valuable for systems that are only avail-able in small amounts, or those with limited dimensionality (thin films and coatings). In addition, DSIallows spatially resolving the mechanical properties and this is of great importance for heterogeneousmaterials such as polymer nanocomposites.

Polymer-based composites incorporating nanoscaled fillers have attracted much attention overrecent years owing to their unique mechanical, thermal and electrical properties that are difficult toachieve using conventional fillers [17,18]. This superior performance combined with their low densitymake them suitable for a broad range of technological sectors such as telecommunications, electronicsand transport industries, especially for aeronautic and aerospace applications where the reduction ofweight is crucial to reduce fuel consumption [19]. The field of polymer nanocomposites has evolvedvery quickly since the last decade being layered silicates, carbon nanotubes (CNTs) and very recently


graphene the most widely used and highly successful examples of nanofillers incorporated into poly-mer matrices. However, nanofiller aggregation has been found to hamper mechanical propertyimprovements. Consequently, a great effort has been devoted to establish the most suitable conditionsfor the transfer of mechanical load from the matrix to the nano-reinforcement. A prerequisite for suchan endeavour is the homogeneous distribution of the nanofillers and the establishment of a strongaffinity with the surrounding polymer matrix.

The application of DSI to polymer nanocomposites has gained increasing interest in recent years.The technique has been proved to be sensitive to filler content, filler dispersion, as well as to the inter-facial nanofiller-matrix adhesion [20–22]. Information on heterogeneities of the composite material,either across the thickness or along the surface arising as a consequence of changes in the matrix mor-phology or uneven distribution of the filler, can be readily detected by means of DSI [23,24].

The present article reviews the most relevant contributions concerning the application of DSI topolymer-based nanocomposites, with special emphasis in those containing carbon-based (e.g. nano-tubes, graphene, nanodiamond) or inorganic (e.g. nanoclays, spherical nanoparticles) nanofillers.The first part of the review introduces the DSI method and summarizes the most importantapproaches developed to derive mechanical data in case of materials with time dependent propertiessuch as polymers. A number of specific recommendations for extracting meaningful mechanical datawill be offered to the reader. The central part of the review presents the state-of-the-art in the appli-cation of DSI to polymer nanocomposites. This technique has been frequently used to evaluate thereinforcing effect of different kinds of nanofillers. Results from different laboratories in similar rein-forced polymers often seem to be at variance. An effort has been done to present a comprehensiveunderstanding of the influence of multiple factors to the mechanical enhancement. Finally, challengesand future perspectives in the application of DSI to polymer nanocomposites will be addressed.

2. Recent advances on instrumented indentation applied to polymer materials: methods anddevelopments

2.1. Depth-sensing indentation

The usual aim of DSI experiments is to determine the material mechanical properties from load-indentation depth data. A characteristic DSI test includes a load-hold-unload sequence as that shownin Fig. 1. However, other methods including partial unloading, reloading, etc. are sometimesemployed. In order to optimize the data collection, certain test parameters such as the approach dis-tance, the allowable drift rate or the criteria to decide whether the indenter has contacted the surfaceshould be chosen. Pyramid indenters are most commonly adopted for DSI among the broad variety ofindenter tips because of their geometrical self-similarity and improved spatial resolution. Berkovichindenters are preferred to the Vickers geometry because the latter is prone to undesired ‘roof’ tipimperfections. Spheres and flat-ended cylindrical punches are particularly useful when the viscoelas-tic character of the material needs to be emphasized (see Section 2.1.2 on time-dependent properties).The use of flat punches offers additional advantages because this geometry eliminates uncertainties inthe calculation of the contact area.

2.1.1. Determination of elastic modulus and hardness. Standard methods of analysis applied to polymermaterials2.1.1.1. The compliance method. The methods commonly adopted to analyse load-depth data from DSIare based on elastic contact theories, first developed for spherical indenters by Hertz, extended toother indenter geometries by Sneddon and finally applied to the analysis of DSI curves by Doernerand Nix and later improved by Oliver and Pharr [2–4]. Linear elastic response is a major requirementto allow using any of these fundamental equations or protocols. For shallow indents, meaningful mod-ulus data have been achieved assuming linear elasticity during loading [15]. However, the proceduremost widely employed relies on the analysis of the initial portion of the unloading curve and is com-monly known as the ‘compliance method’ [3,4]. In this procedure, the following relationship betweenthe contact stiffness S, dP/dh, the reduced modulus Er, and the projected contact area A, is proposed:


Er ¼ffiffiffiffipp

2bSffiffiffiAp ð1Þ

where

1Er¼ ð1� m2Þ

Eþ

1� m2i

� �Ei

ð2Þ

E and m are the elastic modulus and Poisson’s ratio of the material respectively and Ei and mi those ofthe indenter. In Doerner and Nix’ analysis, linear unloading is assumed and S is calculated from the fitof a tangent to the unloading curve at maximum load Pmax (see Fig. 1). The extrapolation to zero load istaken as the plastic depth, hp, i.e. the depth of the indenter in contact with the sample at the onset ofunloading. In case of a Berkovich indenter, A is related to hp through: A ¼ 24:5 h2

p . Oliver and Pharrextended this procedure to account for the curvature in the unloading data. A power law was usedto fit the unloading curve and the derivative at maximum load was used to achieve S. The projectedcontact area A was obtained using A ¼ 24:5 h2

c , with:

hc ¼ hmax � €Pmax=S ð3Þ

where hmax is the maximum displacement (see Fig. 1). The constant €is often taken as € = 0.75 for aBerkovich indenter [25] although a rigorous study would require selecting the €value as a functionof the power law exponent deduced from the analysis of the unloading curve [26]. The b parameterin Eq. (1) is a correction factor. A thorough discussion on the preferred b values that should be adoptedfor each type of the indenter is given in Refs. [7,25]. In case of a Berkovich indenter, any value in therange 1.0226–1.085 can be taken, probably being b = 1.05 the best choice. Once the contact area isdetermined, the hardness can be derived using:

H ¼ Pmax

Að4Þ

Oliver and Pharr’s method (OP) is nowadays extensively used for the analysis of DSI data and mostcommercial instruments have incorporated the procedure into the software. In elastic–plastic materi-als with limited pile-up, the method has been shown to be most successful provided the appropriatecalibrations are carried out (mainly machine compliance, tip area function and thermal drift).

2.1.1.2. Calibrations procedures. The most widespread method for calibration of machine complianceand tip area function was initially proposed and further implemented by Oliver and Pharr [4,25].The methodology uses the indentation data of a reference sample that does not pile up (for example,fused silica) and relies on the assumption that H and E are independent of the penetration depth.Alternatively, the indenter shape area function can be achieved using material-independent methodssuch as AFM imaging [6]. Because the material-dependent procedure for tip calibration is usuallyfriendly integrated in the DSI software, this methodology is usually preferred to the material-indepen-dent ones. Some authors have suggested that the tip calibration constructed upon indentation on ahard material is not appropriate for soft materials such as polymers. Briscoe et al. [13] suggested thatHertzian-like deformations of the indenter tip when contacting hard materials may not resemble thecontact situation with softer surfaces and proposed that the indenter tip area shape for polymer mate-rials should be calibrated against the properties of a polymer reference sample. The concept that someinteraction between the tip and the indenter and/or an inadequate detection of the polymer surfaceproduce an apparent increase of E and H values lies behind the work of Loubet et al. [8] who developedan alternative approach to calibrate the indenter-polymer contact area at shallow penetrations byintroducing an ‘apparent’ tip-defect. Flores and Baltá-Calleja [10] adopted a similar assumption of con-stant E, H values as a function of indentation depth for a glassy polymer and used an alternativeapproach to correct for the ‘tip defect’ by extrapolating the shape of the load-depth curves at largepenetrations (>2 lm) to shallow penetrations. The consequence of this correction is clearly appreci-ated in Fig. 2 where the hardness values of glassy poly(ethylene terephthalate) (PET), before and afterthe ‘tip defect’ correction, are represented as a function of indentation depth. In this work, the OPmethod was used on the derivation of H. Nowadays, a common procedure adopted when polymer

h [μm]

0 2 4 6 8 10

150

180

210

H [M

Pa]

Fig. 2. Hardness values for glassy PET as a function of indenter displacement. The OP method was used to determine H valueswith ( ) and without ( ) application of the tip defect correction. Adapted from Ref. [10].


materials are tested at shallow depths is to calibrate for the tip defect using Oliver and Pharr’sapproach [4,25] on a polymer standard material such as polycarbonate (PC) or poly(methyl methac-rylate) (PMMA) [27,28].

The thermal expansion of both, the material and the equipment, produces an increase in displace-ment during the course of the indentation test that overlaps with the genuine response of the material.To correct for this thermal drift, a new hold period nearly at the end of the unloading (typically 10% ofPmax) that monitors the changes in depth is usually introduced in the indentation test [29]. This holdperiod is illustrated in the load-depth plot of Fig. 1 for a poly(ether ether ketone) (PEEK) sample(P � 0.08 mN). Conversely to elastic–plastic materials, the indentation depth significantly decreaseswhile maintaining the small load nearly at the end of the test. This is due to the characteristic visco-elastic nature of polymer materials. It is clear that the contribution of thermal drift to the measureddisplacement cannot be easily evaluated for materials with time-dependent properties. A criterion toachieve negligible thermal drift effects on the calculate modulus based on the duration of the unload-ing period in relation to the unloading rate, stiffness and contact depth has been proposed [30]. Inpractice, the procedure most commonly adopted is to minimize thermal drift effects by choosing shortindentation cycles and allowing thermal equilibrium before the experiment.

2.1.1.3. Application of the OP method to polymers. The application of the OP method to polymer mate-rials encounters certain difficulties as was soon realized by Oliver and Pharr themselves who stated intheir original paper that ‘‘displacements recovered during first unloading may not be entirely elastic,and because of this, the use of the first unloading curves in the analysis of elastic properties can some-times lead to inaccuracies’’. Indeed, this is nowadays widely accepted in the indentation community[6,8–13]. However, it is frequently considered that the OP method can be applied if there is no increasein displacement upon load release in the form of a ‘nose’. This ‘nose’ is usually avoided by introducinga hold period at peak load in combination with a rapid unloading rate. However, whether creep influ-ences the condition of pure elastic unloading even in the case that long holding times and high unload-ing rates are used is still a matter of concern [10,11]. In addition, it has been shown that the criterionadopted for the analysis of the unloading data (the portion of the unloading curve that is used for thecurve fitting) can also influence the stiffness values [6]. As an example, Fig. 3 shows the variation ofthe elastic modulus values of a glassy polymer sample as a function of the loading/unloading timefor two different holding periods. It is clearly seen that the E values increase with increasing unloadingtime. This effect is due to creep contributing to the total displacement, especially at long loading/

tload [min]

0 2 4 62.0

2.5

3.0

3.5

E [G

Pa]

thold

= 6 s

thold

= 200 s

Fig. 3. Variation of the elastic modulus values of a glassy PET sample as a function of the loading time to Pmax for two differentholding periods: ( ) 6 s, ( ) 200 s. Adapted from Ref. [10].


unloading times and short holding periods. To minimize the effect of creep on the calculation of theunloading stiffness, Loubet et al. [8] proposed that elastic unloading can be assumed if the strain rateat the beginning of unloading is 10 times greater than the strain rate measured at the end of the hold-ing time. A more elaborated approach was developed by Feng and Ngan [30] who proposed a proce-dure to correct the apparent contact stiffness data for creep and thermal drift based on themeasurement of the total displacement rate at the end of the hold period.

2.1.2. DSI time-dependent properties2.1.2.1. Creep analysis. Worthy efforts have been done to characterize the indentation response ofpolymers with time using creep or load relaxation studies [2,31–38]. Creep compliance (or relaxationmodulus) has been frequently employed as a quantitative measure of the capacity of the material toflow. Elastic–viscoelastic correspondence has been commonly adopted to generate solutions for thecreep compliance using the governing equations of a rigid body into an elastic medium [31–38]. Someof these studies describe the creep indentation behaviour in terms of a spring-dashpot mechanicalmodel with the ultimate goal of fitting the model to the experimental data and relate the fitted param-eters to the material properties (typically two elastic moduli and a viscosity coefficient) [31,34–38].Because the above approaches assume linear viscoelasticity, maintaining the indentation responsewithin this regime has been a major concern [32–34]. In this respect, spherical or cylindrical tip geom-etries are preferred to conical or pyramidal ones [32,39]. A generalized model of viscoelastic–plasticbehaviour has been proposed by introducing an additional plastic element in the dashpot-springmodel of viscoelasticity [40,41].

Creep by DSI has also been computed using a steady-state approach equivalent to that employed inuniaxial testing describing the strain rate as a power law of the stress. The method makes use of:

_e / Hn ð5Þ

Here, _e is the strain rate defined as the instantaneous displacement rate divided by the instantaneousdisplacement and n is the creep exponent. Strain rate and hardness values are usually recorded eitherin the hold period at maximum load or by performing a series of loading ramps at a constant strainrate [42,43]. The two methods differ in an important question: in the former procedure a wealth ofdata is computed during the hold period where the hardness and the strain rate are continuouslychanging as the test proceeds, while in the constant _e ramping, one H value is obtained in steady-stateconditions for each test. In turn, n values can be achieved in a log–log plot of _e vs. H. Constant


strain-rate ramping yields higher n values than those achieved from the hold experiments, being theformer in close agreement with those derived from uniaxial testing [42].

Alternatively, a simple way to compare the time-dependent deformation of different polymermaterials is to directly fit the variation of depth as a function of time during the hold period to a simpleequation of the form [44,45]:

hhð0Þ ¼

Ahð0Þ lnðBt þ 1Þ ð6Þ

The h(0) value defines the initial depth at the beginning of the hold period. The dimensionless param-eter A/h(0) is known as the creep strain rate sensitivity and defines the ‘‘extent’’ of the time-dependentdeformation while the B parameter in Eq. (6) is interpreted as the ‘‘rate’’ term.

2.1.2.2. Dynamic indentation testing. Dynamic indentation testing, also known as continuous stiffnessmeasurements (CSM) [4,14], represents an alternative route to study the time dependent behaviour ofpolymers. In addition, CSM represents an improved method for surface detection, which is fairly crit-ical for polymer materials. In dynamic DSI, a sinusoidal force is superimposed to the quasi-static load-ing allowing for a direct measurement of the harmonic contact stiffness as a continuous function ofdepth. This raises two important consequences. Firstly, a profile of the mechanical properties as afunction of indentation displacement in a single loading cycle can be achieved. Secondly, dynamicDSI can operate as a nanodynamic mechanical analyser (nanoDMA) yielding values of the storagemodulus E0 and loss modulus E00 as a function of the frequency of the harmonic force. It is well knownthat for a viscoelastic material the complex modulus:

E� ¼ E0 þ iE00 ð7Þ

where

E0 ¼ E cos d ð8Þ

and

E00 ¼ E sin d ð9Þ

being d the phase shift between the complex stress and the strain and E the usual quasi-staticmodulus.

It is important to note that often, in the DSI literature, no clear distinction is made between quasi-static indentation modulus E and dynamic indentation modulus E0, being both represented by E. In thisreview, the symbol EI will describe a general indentation modulus, either static or dynamic, while Eand E0 will keep their exact meaning.

In order to extract contact stiffness measurements from dynamic DSI, the instrument-sample sys-tem is usually represented as a mechanical model of springs and dashpots [2,46–51]. Appropriateelements are introduced in the mechanical model depending on the characteristics of the indenta-tion instrument and the tip-sample interaction [2]. The latter is usually modelled as a simple Voigtsolid of a spring and dashpot in parallel [46–49]. In this case, the materials dynamic stiffness K andcontact damping Ds can be directly related to the storage modulus and loss modulus valuesfollowing:

E0r ¼ffiffiffiffipp

2bKffiffiffiAp ð10Þ

E00r ¼ffiffiffiffipp

2bxDsffiffiffi

Ap ð11Þ

where x is the angular excitation frequency. K and Ds can be calculated from experimental parametersobtained from dynamic indentation testing, after subtracting the contribution of the instrument to thetotal measured response [2,47,49]. The projected area of contact A for a Berkovich indenter is usuallydetermined in a similar manner as that for the quasi-static experiments, i.e., using the hc valuesdefined in Eq. (3) where S is now the dynamic contact stiffness K. Cylindrical punches are

E´

[GP

a]

0

1

2

3

4

5

6

Displacement, h [nm]

0 500 1000 1500 2000

H [M

Pa]

0

100

200

300

400

Fig. 4. E0 and H values as a function of displacement into surface for a semicrystalline PEEK sample (degree of crystallinity of�40%). Unpublished results.


recommended in this type of experiments because having a constant contact area value at all depths(A = pa2 where a is the radius of the punch), the lack of accuracy introduced in the determination of Ais excluded. The area of contact can also be used to derive hardness values following Eq. (4). Fig. 4shows, as an example, the plot of elastic modulus (E0) and hardness values obtained following the pro-cedure above described as a function of displacement into surface for a PEEK sample using a Berkovichindenter.

Valuable work has been done to develop more sophisticated models that incorporate the materialas a standard linear solid [50,51]. An additional spring is introduced into the Voigt model to captureinstantaneous elastic behaviour. The solutions are more complex than for the above conventionalanalysis; however, once the elements of the standard linear solid system are determined, the samplestorage and loss moduli can be predicted as a function of frequency. In addition, more accurate valuesof the sample stiffnesses are claimed after careful subtraction of the instrument contribution than canbe relevant in the case of stiff materials [51].

A recent paper has reported a comparison between storage and loss modulus values obtainedby means of dynamic DSI and indentation creep using sharp indenters [36]. Interestingly, the E0

values were in close agreement with each other for small applied forces and were also found to bein the range of those determined by means of DMA, E0DMA. However, the E00 values obtained by meansof the different methods exhibited clear disparities [36]. The authors addressed an importantobservation, i.e. dynamic DSI measures recoverable viscoelastic deformation whereas indentation


creep experiments relies on the measurement of the total displacement that generally includes plas-ticity, viscoplasticity, as well as elasticity and viscoelasticity. Hence, in the latter case it is especiallydifficult to maintain the conditions of testing within the linear viscoelastic regime and the use ofgeometries of the indenter that can maintain the elastic contact over a wide force range is especiallyrecommended.

2.1.3. Indentation size effect2.1.3.1. Bulk materials. Early research in metals and ceramics using geometrically self-similar pyramidindenters in conventional hardness testers reported an increase in hardness with decreasing penetra-tion depth [52]. This observation is commonly known as indentation size effect (ISE) and has beenextensively investigated since. The ISE has been attributed to diverse sources, the most common onesrelated to the indenter-sample contact (frictional forces at the interface) [52] and to mechanisms asso-ciated to fundamental material processes [53]. With the advent of DSI testers, the study of the inden-tation size effect has been extended to shallower penetration depths. Because this moderninstrumentation requires a careful calibration and the data can be influenced by many factors, specialcare should be taken to separate experimental artefacts or effects arising from the testing conditionsfrom an ISE associated to the genuine material properties [53].

Among the number of artefacts that could introduce an apparent increase of hardness withdecreasing penetration depth, an inappropriate area function calibration is the most recognizedone. In this case, not only H but also EI values rise at shallow penetrations. If the rounding of the inden-ter tip is underestimated, the calculated contact area is smaller than the real one and an ISE will beapparent [53]. Tip blunting can produce an additional depth-dependence due to the transition fromthe deformation mode of a sphere to that of a pyramid [7,53]. In addition, an incorrect determinationof the point of initial contact towards higher displacement values can also lead to an ISE [53]. In gen-eral, for a correct evaluation of the mechanical properties at shallow depths, a critical examination ofall the possible effects that can introduce inaccuracies in the analysis of EI and H is needed. An exampleis the case of surface roughness that becomes especially significant in polymer materials due to thedifficulty in producing a surface finish with low roughness [27].

A number of models based on strain gradient plasticity have been proposed to explain the hardnessincrease with decreasing indentation size [54,55]. Nanoindentation data in two polymer glasses werefound to be well described by a strain gradient plasticity model based on a molecular kinking mech-anism [54]. Gao et al. [55] extended the classical expanding cavity model to account for strain gradienteffects. The model incorporates strain hardening characteristics of the material and arrived to anindentation size dependence of hardness that was validated against data for a soft metal.

In spite of the well-known hardness dependence on strain rate for polymer materials [13], this isfrequently overlooked in the ISE interpretation. It is of great importance to highlight that indentationexperiments using a constant loading rate (trapezoidal load function) will normally result in an ISEbecause the indentation strain rate during the loading segment continuously decreases with increas-ing displacement [8,42,56]. The study of the near-surface mechanical properties in polymer materialscan only be appropriately approached if the same constant indentation strain rate is applied at all pen-etration depths. For pyramidal indenters, this can be achieved by controlling that the instantaneousloading rate, P0 = dP/dt, divided by the instantaneous load, P, remains constant [56].

Most of commercial advanced DSI devices have incorporated the possibility of applying a constantstrain rate during the loading cycle. The combination of this option with the application of a smalloscillating force provides a rational route to obtain H (and E0) values at all penetration depths. Finally,it is noteworthy that a recent comprehensive study of the depth dependence of indentation data usinga dynamic method (CSM) has revealed that this methodology can induce a significant error in H and E0

at small depths [5] in such a way that it has even not been recommended for the characterization ofthe ISE [53]. In the case of a soft metal such as copper (Er/H � 225), misleading hardness and modulusvalues can be found for depths below �300 nm [5]. Reducing the amplitude of the displacement oscil-lation shifted the limit for reliable data to shallower depths. Although this source of error is minimizedfor materials with lower Er/H ratio (polymers typically exhibit Er/H � 10), the study of ISE in polymersusing a dynamic method should be carried out with some caution, using small amplitudes of oscilla-tion and allowing them to vary in order to discard possible artefacts.


All the above considerations help to understand why detecting a genuine gradient of mechanicalproperties across the sample thickness of a polymer or identifying a size-dependent mechanism ofdeformation is quite a difficult task. Indeed, most indentation studies do not elucidate whether ISEis related to instrumental conditions or to the distinct properties of the surface. In spite of all theseproblems, there are some useful hints on the indentation size effect in polymers that can provide valu-able information on the variation of the mechanical properties across the polymer sample thicknessrelated to morphological or structural changes in the material. For example, localized oxidation, age-ing or modification of the near-to-surface layers during fabrication have been suggested as mecha-nisms yielding an increase in H and EI at shallow depths [13]. Higher degrees of crystallinity at thenear-surface region have been also proposed to account for the enhanced hardness and modulus val-ues for low displacements found for semicrystalline polymer samples [11]. In the case of compositematerials, constant H and EI values found with increasing indentation depth have been interpretedin terms of a good dispersion of the filler in the matrix [57]. In injection-moulded parts, an initial hard-ness decrease with increasing indentation depth followed by a hardness increase above �1 lm hasbeen discussed on the basis of uncertainties at the very surface layer and an inhomogeneous distribu-tion of the crystalline morphology, respectively [58].

2.1.3.2. Thin films. In the particular case of thin films, it should be borne in mind the possible influenceof the substrate on depth sensing indentation results. This influence often takes the form of an ISE(normal o reverse) with increasing penetration depth. To minimize the problem, a rule of thumb com-ing from the early days of microindentation and without any sound physical basis, proposes a limit of10% of the film thickness for the penetration depth. From this point of view, the use of blunt indenterssuch as flat-ended cylindrical punches is in principle recommended. However, as film thicknessbecomes increasingly thin and/or for a large elastic modulus mismatch, substrate effects cannot beavoided and a number of models have been developed to extract the intrinsic elastic modulus of thinfilms from the composite film/substrate modulus value [3,59–65].

Doerner and Nix in 1986 [3] were the first to empirically model the influence of the substrate onthe elastic measurements assuming a rigid flat-ended cylindrical punch that indents a homogeneoushalf space. They proposed an equation for the compliance (reciprocal of S in Eq. (1)), adding two expo-nential terms depending on the relative indentation depth (depth divided by film thickness) whichinclude a constant to be determined empirically for each particular case. Soon afterwards, numericalcalculations performed by King were able to give values for the constant for different film-substratesystems and flat-ended indenter geometries [59]. The method of King was further extended by Sahaand Nix to a pyramidal indenter, by assuming that a flat punch is located at the tip of a Berkovichindenter [60]. Gao et al. also modelled the unloading process of a rigid indenter penetrating a thinfilm/substrate by a rigid cylindrical punch in frictionless contact with a layered elastic half space[61]. However, their analysis differs from previous studies in conceiving a first order rigorous mod-uli-perturbation method to derive a closed-form solution for the contact compliance of the compositesystem. The effective compliance is expressed as a function of film and substrate Poisson’s ratios andshear moduli. The method was proved to be valid for a limited range of shear moduli ratios betweenfilm and substrate and Poisson’s ratios. The Song–Pharr’s model can be considered as an extension ofthe work of Gao et al. modified to increase the applicable range of modulus mismatch between thefilm and substrate [62]. By using a reduced film thickness, subtracting the contact depth to the filmthickness, a good agreement between the model and experimental and numerical data was found.Mencík et al. also proposed some modifications to the Gao’s approach and compared it with other dif-ferent approximation functions [63]. Bec et al. proposed an analytical model based on the indentation,by a rigid cylindrical punch, of a homogeneous film deposited onto a substrate considered as a semi-infinite half space [64]. The composite system was simply modelled by two springs connected in ser-ies. The film reduced modulus can be calculated from the apparent modulus if the reduced Young’smodulus of the substrate and the film thickness are known.

According to Hay and Crawford [65], some of the previous models work reasonably well for com-pliant films on stiff substrates but are not so adequate if the film modulus is more than twice the sub-strate modulus. They put forward an elastic film/substrate model where the film is allowed to act inseries and in parallel with the substrate. The model has been found to work well for both compliant


films on stiff substrates and vice versa and yields accurate modulus values for indentation depths upto 25% of film thickness.

From a practical point of view, an alternate approach to the above mentioned models is describedin ISO standard 14577-4 [66]. It specifies a nanoindentation testing method which should be suitablefor thin coatings of all kinds of materials. The substrate independent elastic modulus should be deter-mined by measuring the elastic modulus according to the OP method as a function of penetrationdepth and then linearly extrapolating the data to zero displacement, where the influence of the sub-strate should be negligible.

Compared to the elastic moduli, hardness data of a thin film are less affected by the substrate. Thisis to be expected because the elastic field under the indenter is a long-range field that extends wellbeyond the plastic one [60]. The Han–Yu–Vlassak method [67] has been recently reported as animproved method for determining the hardness of thin film/substrate systems based on measure-ments of the contact stiffness as a function of indentation depth, provided that the elastic propertiesof the film and substrate are known. The method overcomes the limitations, at large indentationdepths, of the Saha and Nix model arising from the flat punch assumption. To calculate the film mod-ulus, the OP analysis, followed by extrapolating down to zero indentation depth by means of a poly-nomial fit, is used. This procedure, in addition to eliminating substrate effects, helps to minimizeproblems arising from surface roughness and/or pile-up effects.

2.2. Correlation mechanical properties-nanostructure

It is well known that the mechanical properties of a polymer material can be tuned by changing theinternal nanostructure. Establishing correlations between indentation magnitudes (typically H and EI)and characteristic nanostructural parameters of a polymer material (such as, for example, crystallinityand crystal thickness) is of great importance to control the material properties. DSI studies on poly-mers have been mainly focused on the method of analysis of load-depth data and there is limited lit-erature available on the mechanical properties-nanostructure correlations. However, there are well-established empirical models to explain the hardness dependence on the polymer nanostructure;the hardness numbers traditionally arise from conventional indentation testing based on the opticalmeasurement of the residual indentation. These models have been topic of research for many decades.The additivity law describing the hardness of a semicrystalline polymer material in terms of that of thecrystalline and of the amorphous regions, Hc and Ha respectively, has been extensively validated and isexpressed as [68,69]:

H ¼ Hcvc þ Hað1� vcÞ ð12Þ

where vc represents the volume fraction of crystalline material. An analogous additivity law for themodulus values extracted from DSI has not been sufficiently contrasted and only some indications thatEI could follow a parallel model with vc, as H does, have been proposed [70]. In addition, the Hc valueshave been also shown to be influenced by the crystal thickness, lc, through:

Hc ¼H1c

1þ blc

ð13Þ

Here, H1c is the hardness of an infinitely thick crystal and b is a parameter related to the ratio betweenthe surface free energy of the crystals, re, and the energy required to plastically deform them througha number of shearing planes, Dh (b = 2re/Dh). H1c is mainly related to the chain packing density withinthe crystals and represents the upper limit to the hardness of a polymer material.

As an example, Fig. 5 shows the hardness dependence with vc and lc for PEEK [71]. The glassy mate-rial displays the lowest H value (135 MPa) while the semicrystalline samples exhibit H values thattend to increase with increasing degree of crystallinity. However, the H changes cannot be only attrib-uted to variations in vc because lc is concurrently changing. A colour code has been used in Fig. 5 toidentify each group of data associated to a range of lc values. The straight lines follow the additivemodel of Eq. (12) with Ha = 135 MPa. Increasing crystal thickness values yield straight lines withincreasing slope (and hence, higher right-hand y-axis intercepts, Hc). One can now clearly perceive

vc

0.0 0.2 0.4 0.6 0.8 1.0

H [

MP

a]

100

200

300

400

Glass

l c = 2 nml c = 3-4

nm

l c=

5-6

nm

Fig. 5. Hardness dependence with vc for PEEK. The straight lines follow the hardness additivity law (Eq. (12)) withHa = 135 MPa. Each colour defines a group of data with similar crystal thickness values, lc. Note that increasing the crystalthickness yields higher Hc values (right-hand y-axis intercepts). Data taken from Ref. [71], copyright 1991, with permission fromSpringer Science and Business Media.


the separate influence of vc and lc on the H values. Moving along a straight line defines the hardnessincrease due to changes in the degree of crystallinity while ‘‘jumping’’ from one straight line toanother one of higher slope explains the hardness increase due to crystal thickening.

The hardness additivity law of Eq. (12) has been proved to be successful in describing the resistanceto plastic deformation of a polymer material in terms of the separate contribution of stiff and soft com-pliant elements. At the heart of this success is the connectivity between the crystalline and amorphousregions through tie molecules, interconnecting chains, etc.

Application of this concept to multicomponent materials such as polymer composites woulddepend to a large extent on the connectivity between the different phases. In reinforced polymermatrices, the filler-matrix interface that is responsible of the adequate load transfer across the mate-rial is a critical issue that will certainly influence the validity of any model predicting the properties ofthe composite in terms of those of the constituents.

Over one hundred years ago, there was a raising need of estimating the effective elastic moduli ofheterogeneous materials. The Voigt–Reuss rule of mixtures established upper and lower bounds forthe modulus of a composite material; pure elastic behaviour was assumed at constant strain (Voigt)or stress (Reuss) under uniaxial loading. In case of a two-phase composite material, both approachescan be expressed according to:

Ec ¼ vmEm þ v f Ef Voigt : upper bound; parallel model ð14ÞE�1

c ¼ vmE�1m þ v f E�1

f Reuss : lower bound; normal model ð15Þ

where v represents volume fractions and the subscripts c, m and f correspond to composite, matrix andfiller, respectively.

Using the elastic–viscoelastic correspondence principle, the expressions (14) and (15) could alsohold for viscoelastic materials by changing the modulus values by their complex counterparts,E�c ; E

�m and E�f [72]. Other micromechanics additivity models are: the Hashin–Shtrikman formulae

[73] and the Halpin–Tsai model [74]. The former were derived by applying variational principles oflinear elasticity and offer tighter bounds for the effective modulus values of isotropic composites thanthose of the Voigt–Reuss limits. The Halpin–Tsai equations are based on a rigorous self-consistentmethod which models the composite as a single fibre, surrounded by a cylinder of matrix and the


ensemble embedded in an unbounded homogeneous medium. These equations account for a widevariety of reinforcement geometries through an empirically adjustable parameter. The above models[73,74] have been used in a number of occasions to describe the modulus values of polymer-basedcomposites obtained by means of indentation techniques [75–78].

2.3. Comparison of indentation studies with bulk techniques

Comparing modulus data obtained from DSI with those achieved by conventional macroscopictechniques is quite a challenging task. In addition to the sources of error that can be introducedin DSI experiments concerning instrument calibration, other factors such as strain hardening, inter-action forces between the indenter and the sample, ISE or material heterogeneities could contributeto the discrepancies found between nanoindentation and traditional mechanical methods. Alsoimportant are the differences in the principle of measurement and the testing geometry, someof them are: (i) in DSI, the volume of deformation is continuously changing, as it is the inhomo-geneous distribution of stresses and strains generated beneath the indenter; this clearly differenti-ates DSI from other macroscopic techniques such as tensile testing. Indeed, the concept of an‘effective’ indentation strain rate enables application of well-established uniaxial tensile analysessuch as power-law creep [42]. (ii) The volume loaded is significantly smaller in DSI. In addition,for dynamic testing, the different amplitudes of oscillation between DMA and DSI (a few microm-eters in DMA vs. a few nm in DSI) has been suggested to give rise to different strain levels thatcould activate distinct mechanisms of deformation [36]; and (iii) the load direction is believedto evolve radially from the point of first contact in case of indentation testing while for macro-scopic tensile testing it is unidirectional. In the latter technique, mainly tensile stresses are appliedwhile in DSI a combination of compressive, tensile and shear forces are exerted on the material. Fornanofiller-reinforced systems, the disparity between results from the different techniques isexpected to be even larger because in uniaxial testing and in contrast to DSI, the longer fillerdimension would preferentially orient along the loading direction enhancing the overall materialreinforcement.

In spite of the above described limitations, it is rather frequent to attempt a comparison betweenDSI and conventional macroscopic data. However, one should bear in mind the phenomenologicalcharacter of such correspondence. In the early stages of application of dynamic DSI (nanoDMA)to polymers, DMA data were frequently taken as a reference to validate dynamic indentation mea-surements because the latter method assumes certain models and analogies (such as linear visco-elasticity and elastic–viscoelastic correspondence) that are not straightforward for a polymermaterial under standard testing conditions. A number of reports suggest that the modulus valuesE0 obtained from depth-sensing instrumentation are in fair agreement with those obtained byDMA or uniaxial compression with harmonic oscillation [47,48,79,80]. However, a significant dis-parity between E0 and E0DMA data has also been reported [20,22,76,81]. In many cases, the modulusvalues measured by dynamic nanoindentation have been found to be higher than those obtainedfrom DMA. The origin of the discrepancy between dynamic indentation and conventional bulk mea-surements with harmonic oscillation is still open to debate. In addition to the differences in the testgeometry and operating mode mentioned above, it seems that comparison of data from the differ-ent techniques is only valid if indentation data remain in the linear viscoelastic regime. However,this is often overlooked and still indentation measurements have been found to be in reasonableagreement with DMA. Interestingly, VanLandingham et al. [33] investigated a number of polymermaterials by means of traditional dynamic mechanical analysis and dynamic nanoindentation andfound an excellent agreement for the stiffer polymers while divergent results were found for themore compliant polymer samples. The authors argued that the current DSI calculation methodsmight be appropriate for materials where energy storage dominates the dynamic response butmight not be able to capture the response of materials with a pronounced viscous or rubbery flowcharacter.

In summary, although there is a technical similarity between dynamic DSI and DMA, neither thestresses applied and strain levels achieved, nor the volume and modes of deformation, amongst otherfactors, are identical and hence a perfect match of results cannot be expected.


3. General aspects on polymer nanocomposites

The optimization of the properties of polymer nanocomposites depends to a great extent on theinteraction between the matrix and the filler [18]. The state of dispersion of the filler and the natureof the interface/interphase with the host matrix are the main two factors accounting for the interac-tion between filler and matrix, and in turn, compromising the performance of the nanocomposite.Other important parameters such as the aspect ratio and orientation of the filler or the changes trig-gered in the matrix morphology upon addition of the reinforcement can significantly influence thecomposite properties. Microscopic techniques such as scanning and transmission electron micros-copies (SEM and TEM) can provide some direct information about the degree of orientation and dis-tribution of the nanofillers, or the matrix-filler interactions. Wide angle X-ray scattering (WAXS)and differential scanning calorimetry (DSC) experiments are frequently performed to assess the crys-talline structure, degree of crystallinity and crystal orientation of the matrix in the composite.

The following sections collect and summarize published data on nanofiller-reinforced polymernanocomposites obtained via nanoindentation, in order to extract some general conclusions aboutthe influence of the abovementioned parameters on the mechanical performance of the differentmaterials. Tables 1–6 present a comprehensive collection of room temperature (RT) DSI resultsreported to date in this field. The absolute values of EI and H and the percentage of change comparedto those of the matrix, DEI ¼ EI � E0

I

� �=E0

I � 100 and DH = (H � H0)/H0 � 100 being E0I and H0 the val-

ues of the neat material, have been included in the tables. Indentation modulus values are specifiedwhether obtained using a quasi-static (E) or a dynamic method (E0). Berkovich indenter tips are alwaysused unless otherwise stated. The nanofiller type, characteristics, concentration and state of dispersionas well as the composite processing method and the degree of crystallinity of the matrix are includedin the cases that this information is available. Furthermore, modulus data obtained by conventionalbulk techniques such as tensile testing, Y, or DMA are also listed in order to establish a correlationbetween nanoindentation results and macroscopic properties. Because the selection of the methodof analysis is relevant to achieve meaningful indentation data as discussed in Section 2, this informa-tion is offered to the reader in the different Tables. Finally, other aspects such as the variation ofmechanical properties with the indentation depth (ISE) are also reported.

Those indentation modulus data possibly affected by a large inaccuracy, due to the experimentalprocedure followed, are commented in the footnotes of the corresponding Tables. In addition, redsymbols are used for such data when appearing in the Figures. For example, it is well known thatat the beginning of an indentation test, the detection of the point of initial contact encounters impor-tant difficulties if the material exhibits low modulus [78,82]. This is especially critical under two cir-cumstances: (i) when sharp indenters such as Berkovich are used because the contact area betweenthe sample and the indenter is small; (ii) when the dynamic option enjoying improved resolutionto detect changes in stiffness is not used. Hence, from the above considerations, it is questionablethe significance of modulus values and increments on very low-modulus materials (E < 0.15 GPa)derived from quasi-static measurements using a sharp indenter. Other sources of error are the exis-tence of a clear ‘nose’ effect at the onset of the unloading curves and the inadequate correction ofthe load-depth data by thermal drift (Section 2.1.1).

4. Polymer nanocomposites incorporating carbon nanofillers

4.1. Carbon nanotubes

4.1.1. Synthesis methods and types of CNTsOne of the most efficient nanofillers to reinforce polymer matrices are carbon nanotubes (CNTs),

one dimensional carbon-based nanomaterials discovered by Iijima in 1991 [83] that possess extraor-dinary high Young’s modulus (up to 1.2 TPa) and tensile strength (ca. 50–200 GPa), combined withvery large aspect ratio (>1000), high flexibility and low density (�1.8 g/cm3) [84,85]. There are twomain types of CNTs: those consisting of a single graphite sheet wrapped into a cylindrical tube witha diameter in the range of 0.7–3 nm, the single-walled carbon nanotubes (SWCNTs), and those

POPOV

Highlight

Table 1DSI modulus (either quasi-static E or dynamic E0) and hardness data for CNT-reinforced polymer nanocomposites. Type of CNT, composite preparation method, morphological information onthe state of aggregation of the filler and/or extent of the filler-matrix interaction and crystallinity values are also included. Values of storage modulus from DMA, E0DMA , and Young’s modulusdata from tensile experiments, Y, are incorporated in the table when appropriate. The figures in parenthesis indicate the percentage of property increase in comparison to that of the neatmatrix. The method of analysis of the DSI data is included in the table with the following abbreviations: ‘‘From unloading (OP)’’ indicates that a single reading of E and H is determined from theonset of unloading using the OP method described in Eqs. (1)–(4); ‘‘Dynamic’’ is based on the derivation of the harmonic contact stiffness where E0 , E00 and H are calculated following Eqs. (10),(11) and (4) respectively. In the dynamic method, the contact depth values are determined using the OP procedure by substituting the quasi-static contact stiffness for the harmonic one (Eq.(3)). Usually, this option is accompanied by a continuous measurement of the material properties. The variation of mechanical properties with indentation depth is also included in the tablewhen this information available. Berkovich indenter tips are used unless otherwise stated. ‘‘NA’’ indicates that the information is ‘‘non-available’’.

Matrix Nanofiller type Processingmethod

Nanofillercontent(wt%)

Quasi-staticDSI E(GPa)(DE, %)

DynamicDSIE0

(GPa)(DE0 , %)

HardnessH (GPa)(DH, %)

Morphology(Dispersion,interaction)

Crystallinity(%)(% variation)

E0DMA

(GPa)(DE0DMA, %)

Y (GPa)(DY, %)

Methodof analysis

Depthvariation(range)

Ref.

Epoxy-acrylate

Raw MWCNTs Sonication/mixingphotocuring

0a 1.81 0.075 SEM (aggregatesat 0.75 wt%)

NA ISE (0.15–5 lm) [108]

0.25a 2.20(22) 0.098(30)0.75a 2.60(42) 0.133(77)0b 2.20 0.1140.25b 2.50(14) 0.137(20)0.75b 2.05(�7) 0.094(�9)

Epoxy CVD rawMWCNTs

Sonication/mixing Curing inoven

0 4.22 0.325 SEM(agglomerates)

2.00 From unloading (OP) [91,95]0.1 4.70(11) 0.388(19) 2.20(10)

Epoxy Raw CVDMWCNTs

Sonication/mixing Curing inoven

0 4.53 0.183 SEM(agglomerates)

From unloading (OP) [109]3.0 4.75(5) 0.187(2)

Epoxy Raw MWCNTs Sonication/mixing Curing inoven

0 3.16 0.138 SEM (gooddispersion up to0.5 wt%)

3.18 From unloading (OP)o [110]0.1 3.19(1) 0.141(2) 3.24(2)0.5 3.35(6) 0.150(9) 3.50(10)1.0 3.44(9) 0.157(14) 3.68(16)

Epoxy Raw CVDMWCNTs

Stirring/mixingCuring in oven

0 3.70 SEM, TEM (gooddispersion)

Dynamic ISE (0.2–1 lm) [111]5.0c 7.40(100)5.0d 4.62(25)

Epoxy Raw CVDMWCNTs

Stirring/mixingCuring in hot-press

0 2.26 0.125 2.14 From unloading (OP) ISE (0.1–0.4 lm) [112]1.0 3.10(37) 0.187(50) 2.50(17)2.0 3.64(61) 0.293(85) 3.09(44)

Epoxy Raw SWCNTs Stirring/castingCuring in oven

0 3.0 SEM, AFM (a lot ofaggregates)

From unloading (OP) ISE (0.15–0.3 lm) [113]1.0 3.0(0)

Epoxy Acid-functionalizedCVD MWCNTs

Stirring/mixingCuring in oven

0 3.70 SEM, TEM (gooddispersion)

Dynamic ISE (0.2–1 lm) [111]5.0c 7.55(104)5.0d 4.25(15)

Epoxy Fluorinated-SWCNTs

Stirring/castingCuring in oven

0 3.0 SEM, AFM (manyaggregates)


Epoxy Silanefunctionalized-SWCNTs


0 3.0 SEM, AFM (fewaggregates)


(continued on next page)

A.M

.Díez-Pascual

etal./Progress

inM

aterialsScience

67(2015)

1–94

19

Tab

le1

(con

tinu

ed)

Mat

rix

Nan

ofill

erty

pePr

oces

sin

gm

eth

odN

anofi

ller

con

ten

t(w

t%)

Qu

asi-

stat

icD

SIE

(GPa

)( D

E,%

)

Dyn

amic

DSI

E0 (GPa

)( D

E0,%

)

Har

dnes

sH

(GPa

)( D

H,%

)

Mor

phol

ogy

(Dis

pers

ion

,in

tera

ctio

n)

Cry

stal

lin

ity

(%)

(%va

riat

ion

)

E0D

MA

(GPa

)( D

E0D

MA,%

)

Y(G

Pa)

( DY,

%)

Met

hod

ofan

alys

isD

epth

vari

atio

n(r

ange

)

Ref

.

Epox

yPl

asm

afu

nct

ion

aliz

edC

VD

MW

CN

Tse

Son

icat

ion

/m

ixin

gC

uri

ng

inov

en

04.

220.

235

SEM

(goo

ddi

sper

sion

)2.

00Fr

omu

nlo

adin

g(O

P)[9

1,95

]0.

15.

05(1

9)0.

456(

40)

2.40

(20)

Epox

yM

BZ

fun

ctio

nal

ized

Arc

-dis

char

geSW

CN

Ts

Solu

tion

mix

ing

Cu

rin

gin

oven

04.

04.

00.

34SE

M(u

nif

orm

dist

ribu

tion

up

to1

wt%

)

From

un

load

ing

(OP)

and

Dyn

amic

Han

dE

con

stan

t(0

.2–1

lm

)[5

7]1.

04.

4(10

)5.

0(24

)0.

37(9

)3.

05.

6(40

)5.

9(48

)0.

41(2

0)5.

07.

0(75

)6.

7(67

)0.

44(3

0)Ep

oxy

Puri

fied

lase

r-gr

own

SWC

NTs

Son

icat

ion

/m

ixin

gC

uri

ng

inov

en

03.

900.

245

SEM

(poo

rdi

sper

sion

)5.

46D

ynam

icH

and

E0co

nst

ant

(2–4

.5l

m)

[24,

114]

0.1

3.78

(�3)

0.25

4(4)

4.78

(�12

)0.

53.

98(2

)0.

262(

7)4.

96(�

9)1.

04.

06(4

)0.

265(

8)6.

01(1

0)3.

04.

25(9

)–

–5.

04.

56(1

7)–

–Ep

oxy

Ali

gned

CV

DM

WC

NT

fore

sts

Mix

ing/

curi

ng

03.

60.

157

SEM

,TEM

(qu

ite

aggl

omer

ated

)N

A[1

15]

0.5

4.0(

11)

0.17

7(13

)

Epox

yA

lign

edC

VD

MW

CN

Tfo

rest

s

Insi

tuw

etti

ng/

Rea

lign

/cu

rin

g0

3.6

SEM

(den

seC

NT

arra

y)Fr

omu

nlo

adin

g(O

P)o

[116

]N

A6.

4(78

)N

A9.

6(16

6)Ep

oxy

Ali

gned

CV

DM

WC

NT

fore

sts

Mec

han

ical

den

sifi

cati

on/

Cap

illa

ry-

indu

ced

wet

tin

g/cu

rin

g

04.

8O

M,S

EM(d

ensi

fica

tion

)Fr

omu

nlo

adin

g(O

P)[1

17]

0.75

4.9(

2)5.

56.

0(25

)14

.28.

8(85

)

Epox

yC

VD

MW

CN

Tfo

rest

sSh

ear

mix

ing/

curi

ng

04.

8O

M,S

EM(g

ood

disp

ersi

on)

From

un

load

ing

(OP)

[117

]0.

754.

6(�

4)3.

05.

1(6)

Epox

yC

oile

dC

VD

CN

TsSo

nic

atio

n/

mix

ing

Cu

rin

gin

oven

04.

900.

240

SEM

(dis

enta

ngl

edan

dev

enly

spac

ed)

From

un

load

ing

(OP)

Han

dE

con

stan

t(1

.01–

1.03

lm

)[1

18]

1.0

5.60

(14)

f0.

271(

13)f

3.0

6.00

(23)

f0.

315(

31)f

5.0

7.00

(42)

f0.

350(

46)f

Epox

yC

VD

CN

Tsgr

afte

dsi

lica

fibr

e

Cu

rin

g0

2.23

0.13

Ram

an,A

FMFr

omu

nlo

adin

g(O

P)N

A[1

19]

NA

2.77

(24)

0.17

(28)

Ch

itos

anA

cid-

fun

ctio

nal

ized

CV

DM

WC

NTs

Solu

tion

-cas

tin

g0

0.11

9SE

M,T

EM(g

ood

disp

ersi

onu

pto

2.0

wt%

)

1.08

Dyn

amic

ISE

(<0.

3l

m)

Rev

erse

ISE

(1–5

lm

)

[99]

0.2

0.12

8(8)

1.33

(23)

0.4

0.16

3(37

)1.

92(7

7)0.

80.

150(

26)

2.08

(93)

2.0

2.15

(99)

Ch

itos

anPE

DO

T-PS

Sfu

nct

ion

aliz

edM

WC

NTs

Solu

tion

mix

ing/

cast

ing

02.

300.

130

SEM

(goo

ddi

sper

sion

)3.

05D

ynam

icIS

E(<

0.2

lm

)[9

8]0.

12.

60(1

3)0.

146(

12)

3.54

(16)

0.2

3.15

(37)

0.15

5(19

)3.

75(2

3)0.

53.

50(5

2)0.

223(

31)

4.09

(34)


Tab

le1

(con

tinu

ed)

Mat

rix

Nan

ofill

erty

pePr

oces

sin

gm

eth

odN

anofi

ller

con

ten

t(w

t%)

Qu

asi-

stat

icD

SIE

(GPa

)( D

E,%

)

Dyn

amic

DSI

E0 (GPa

)( D

E0,%

)

Har

dnes

sH

(GPa

)( D

H,%

)

Mor

phol

ogy

(Dis

pers

ion

,in

tera

ctio

n)

Cry

stal

lin

ity

(%)

(%va

riat

ion

)

E0D

MA

(GPa

)( D

E0D

MA,%

)

Y(G

Pa)

( DY,

%)

Met

hod

ofan

alys

isD

epth

vari

atio

n(r

ange

)

Ref

.

PHB

Raw

SWC

NTs

Solu

tion

-cas

tin

g0

5.66

0.31

OM

(poo

rdi

sper

sion

;la

rge

aggr

egat

esat

10w

t%)

Incr

ease

incr

ysta

llin

ity,

decr

ease

incr

ysta

llit

esi

ze

From

un

load

ing

(OP)

[97]

1.0

7.62

(35)

0.33

(6)

10.0

11.7

4(10

7)0.

35(1

2)

PHO

Raw

SWC

NTs

Solu

tion

-cas

tin

g0

0.12

g0.

0056

From

un

load

ing

(OP)

[97]

1.0

0.15

(25)

0.00

77(3

7)10

.00.

53(3

41)

0.01

37(1

45)

PPR

awM

WC

NTs

Insi

tupo

lym

eriz

atio

nC

ompr

essi

onm

ould

ing

01.

450.

065h

From

un

load

ing

(OP)

o[1

06]

0.4

2.22

(53)

0.09

2(41

)h

1.0

2.23

(54)

0.10

0(54

)h

1.3

1.58

(9)

0.08

1(24

)h

1.6

2.64

(82)

0.12

0(85

)h

2.2

2.02

(39)

0.09

3(42

)h

2.4

2.00

(38)

0.09

1(40

)h

PA-6

Aci

d-pu

rifi

edC

VD

MW

CN

TsM

elt-

blen

din

g/C

ompr

essi

on-

mou

ldin

g

01.

180.

06SE

M(u

nif

orm

disp

ersi

on)

CN

Tsac

tas

a-n

ucl

eati

ng

agen

ts

2.08

0.40

Dyn

amic

ISE

(<0.

2l

m)

[20]

0.2

1.33

(13)

0.08

(33)

2.55

(23)

0.68

(72)

0.5

1.60

(36)

0.10

(66)

2.57

(24)

0.77

(93)

1.0

2.02

(71)

0.11

(83)

3.14

(51)

0.85

(113

)PA

-6i

Aci

d-fu

nct

ion

aliz

edM

WC

NTs

j

Solu

tion

mix

ing/

elec

tros

pin

nin

g0

0.95

0.24

0SE

M,T

EM(g

ood

disp

ersi

on)

330.

901.

20Fr

omu

nlo

adin

g(O

P)p

[100

]1.

01.

56(6

4)0.

252(

5)36

(9)

–2.

00(7

0)2.

52.

14(1

25)

0.27

0(13

)N

A1.

50(6

7)2.

40(1

00)

5.0

2.38

(150

)0.

274(

14)

NA

–2.

60(1

20)

7.5

2.66

(180

)0.

280(

17)

NA

3.30

(267

)3.

00(1

50)

PLLA

Aci

d-pu

rifi

edA

rc-d

isch

arge

MW

CN

Ts

Son

icat

ion

/So

luti

on-c

asti

ng

03.

6SE

M(g

ood

disp

ersi

on)

From

un

load

ing

(OP)

[101

]1.

254.

4(22

)2.

55.

8(61

)3.

756.

6(83

)5.

07.

5(10

8)6.

257.

8(11

7)U

HM

WPE

Raw

MW

CN

TsM

illi

ng/

Elec

tros

tati

csp

rayi

ng

02.

020.

104

SEM

(qu

ite

aggl

omer

ated

)55

0.60

From

un

load

ing

(OP)

[105

]5.

02.

23(1

0)0.

116(

12)

43( �

22)

1.28

(115

)

UH

MW

PEPl

asm

afu

nct

ion

aliz

edA

rc-d

isch

arge

SWC

NTs

Solu

tion

mix

ing/

cast

ing

02.

750.

062

SEM

(goo

ddi

sper

sion

)X

PS(s

tron

gin

terf

acia

lad

hes

ion

)

Dyn

amic

[96]

0.05

3.1(

12)

0.82

0(32

)0.

14.

26(5

5)0.

100(

62)

0.2

4.35

(58)

0.10

3(66

)U

HM

WPE

Aci

d-fu

nct

ion

aliz

edM

WC

NTs

Bal

lm

illi

ng/

Com

pres

sion

-m

ould

ing

00.

617

0.26

1SE

M(g

ood

disp

ersi

on)

Ram

an(s

tron

gin

terf

acia

lad

hes

ion

)

520.

027l

–Fr

omu

nlo

adin

g(O

P)k

[104

]0.

50.

728(

18)

0.28

2(8)

55(5

)0.

034(

26)l

(13)

1.0

0.80

8(31

)0.

381(

46)

58(1

0)0.

23(7

63)l

(26)

1.5

0.92

5(50

)0.

412(

58)

61(1

5)0.

3(10

00)l

(38)

2.0

1.67

(170

)0.

457

(75)

64(2

3)0.

4(13

80)l

(44)

2.5

1.75

(183

)0.

483

(85)

66(2

7)0.

5(17

50)l

(40)

5.0

1.89

(206

)0.

564(

116)

71(3

7)0.

6(21

20)l

(24)

(con

tinu

edon

next

page

)


Tabl

e1

(con

tinu

ed)

Mat

rix

Nan

ofill

erty

pePr

oces

sin

gm

eth

odN

anofi

ller

con

ten

t(w

t%)

Qu

asi-

stat

icD

SIE

(GPa

)( D

E,%

)

Dyn

amic

DSI

E0 (GPa

)( D

E0,%

)

Har

dnes

sH

(GPa

)( D

H,%

)

Mor

phol

ogy

(Dis

pers

ion

,in

tera

ctio

n)

Cry

stal

lin

ity

(%)

(%va

riat

ion

)

E0D

MA

(GPa

)( D

E0D

MA,%

)

Y(G

Pa)

( DY,

%)

Met

hod

ofan

alys

isD

epth

vari

atio

n(r

ange

)

Ref

.

PVD

FiR

awM

WC

NTs

jN

ear-

fiel

del

ectr

ospi

nn

ing

00.

26SE

M(u

nif

orm

disp

ersi

on)

–N

A[1

20]

0.03

0.31

(19)

(25)

m

PVA

Raw

arc-

disc

har

geM

WC

NTs

Solu

tion

mix

ing

07.

00.

30TE

M(s

tron

gin

terf

acia

lbo

ndi

ng)

14N

A[1

03]

0.1

8.4(

20)

0.30

(0)

24(7

1)0.

259.

4(35

)0.

40(3

0)25

(75)

0.5

10.4

(48)

0.42

(38)

26(8

0)1.

012

.5(7

9)0.

47(5

7)27

(93)

PVA

Aci

d-fu

nct

ion

aliz

eA

rc-d

isch

arge

SWC

NTs

Solu

tion

cast

ing

00.

660.

038

42Fr

omu

nlo

adin

g(O

P)[1

21]

0.2

6.93

(950

)0.

278(

630)

52(2

4)0.

47.

30(1

000)

0.28

0(63

7)55

(30)

0.6

7.80

(118

0)0.

290(

663)

64(5

1)PE

EK/G

FA

rc-p

uri

fied

SWC

NTs

Mel

t-bl

endi

ng/

Hot

-com

pres

sion

03.

90.

198

SEM

(few

aggl

omer

ates

)39

16.7

15.6

Dyn

amic

Han

dE0

con

stan

t(0

.1–0

.4l

m)6

4(51

)

[122

]1.

04.

25(9

)0.

218(

10)

38(�

2)16

(3)

PEEK

/GF

PEES

-wra

pped

Arc

-pu

rifi

edSW

CN

Ts

Mel

t-bl

endi

ng/

Hot

-com

pres

sion

03.

90.

198

SEM

(goo

ddi

sper

sion

)39

16.7

15.6

Dyn

amic

Han

dE0

con

stan

t(0

.1–0

.4l

m)

[122

]1.

04.

60(1

8)0.

238(

20)

37(�

5)17

.8(1

4)

PDM

SR

awM

WC

NTs

Ult

raso

nic

atio

n/

curi

ng

00.

0067

g0.

0022

SEM

(agg

lom

erat

es)

0.00

150.

165

From

un

load

ing

(OP)

[124

]4.

00.

0089

(32)

0.00

27(2

3)0.

0028

(88)

0.23

4(42

)PC

CV

Dra

wM

WC

NTs

Mel

t-ex

tru

sion

04.

00.

38SE

M(h

omog

eneo

us

disp

ersi

on)

From

un

load

ing

(NA

)nR

ever

seIS

E(<

1l

m)

[125

]2.

04.

2(5)

0.40

(4)

4.0

4.4

(10)

0.41

(8)

6.0

4.5(

12)

0.43

(13)

15.0

5.2(

30)

0.50

(32)

PIR

awA

rc-

disc

har

geSW

CN

Ts

Son

icat

ion

/spi

n-

coat

ing

05.

230.

43A

FM,S

EM,O

M(s

ome

aggl

omer

ates

,go

odin

terf

acia

lbo

ndi

ng)

Dyn

amic

ISE

(<1

lm

)[1

26]

0.05

8.41

(59)

0.72

(67)

PVK

Raw

arc-

disc

har

geM

WC

NTs

Solu

tion

mix

ing

02.

00.

30N

A[1

03]

0.5

2.24

(12)

0.33

(10)

1.0

3.10

(55)

0.44

(48)

2.0

2.56

(28)

0.50

(68)

4.0

2.96

(48)

0.63

(110

)8.

06.

0(20

0)0.

60(1

00)

PMM

AA

mid

e-fu

nct

ion

aliz

edC

VD

MW

CN

Ts

Solu

tion

mix

ing/

spin

-coa

tin

g0

3.6

0.05

SEM

(non

un

ifor

mdi

sper

sion

)Fr

omu

nlo

adin

g(O

P)E

con

stan

tIS

Ein

H(0

.1–0

.5l

m)

[28,

102]

1.0

3.38

(�6)

0.04

8(�

4)3.

03.

67(2

)0.

051(

2)5.

03.

53(�

2)0.

05(0

)


Table 1 (continued)

Matrix Nanofiller type Processingmethod



DynamicDSIE0

(GPa)(DE0 , %)



Crystallinity(%)(%variation)

0DMA

GPa)DE0DMA,)

Y(GPa)(DY,%)

Methodof analysis


Ref.

PMMA Silica-coated CVDMWCNTs

Solutionmixing/spin-coating

0 3.6 0.05 SEM (aggregatesincreasing with CNTcontent)

From unloading (OP) E constant ISE inH (0.1–0.5 lm)

[28,102]1.0 3.24(�10) 0.05(0)2.0 4.93(37) 0.085(70)3.0 6.66(85) 0.10(100)4.0 8.28(130) 0.11(120)

P(MMA-co-MPTMS)

APTES functionalizedacid-treated MWCNTs

Sol–gel/spin-coating/curing

0 5.0 0.107 OM (homogeneousdispersion)

From unloading (OP) Reverse ISE(0.17–0.95 lm)

[127]2.0 6.3(26) 0.209(95)

P(MMA-co-MPTMS)

TEOS functionalizedPAH-dispersedMWCNTs

Sol–gel/spin-coating/curing

0 5.0 0.107 OM (aggregates) Reverse ISE(0.14–0.8 lm)

[127]2.0 4.9(�2) 0.189(76)

(See abbreviations in the appendix).a Cured for 12 h.b Cured for 24 h.c MWCNTs predispersed in acetone.d MWCNTs predispersed in gamma-butyrolactone (GBL).e Amine functionalized CNTs grafted to the matrix.f Results showing ‘‘nose’’ effect at the beginning of unloading.g Materials with matrix modulus <0.15 GPa, possibly affected by a large uncertainty in the detection of the point of ini l contact.h Martens hardness [2].i Composite fibre.j Aligned along the fibre axis.k Inadequate thermal drift correction.l Data from rheological measurements.

m Crystallinity of b phase.n Experiments carried out using cycling loading.o ISO 14577 test.p Spherical indenter.

A.M

.Díez-Pascual

etal./Progress

inM

aterialsScience

67(2015)

1–94

23

E((%

tia

Table 2Nanomechanical and macroscopic properties for graphene-reinforced polymer nanocomposites. Columns as indicated in Table 1.

Matrix Nanofiller type (thickness) Processingmethod



DynamicDSI E0

(GPa)(DE0 , %)




E0DMA

(GPa)(DE0DMA ,%)

Y (GPa)(D, Y %)

Method of analysis Depthvariation(range)

Ref.

Epoxy GP (350–550 nm) Sonication/mixing/curing

0 2.5 2.8 AFM (random andhomogeneousdispersion)

2.8 2.70 From unloading(OP) and Dynamic

ISE (0.5–1.5 lm)

[166,167]1.0 3.2(28)a 3.5(25) 3.5(25) 3.41(26)5.8 3.3(32)a 3.8(35) 3.8(36) –

Epoxy GNP (6–8 nm) Mixing/curing

0 3.7 0.130 From unloading(OP)i

[168]0.05 4.0(8) 0.165(26)0.1 4.1(11) 0.175(35)0.25 4.3(17) 0.175(35)0.5 4.6(25) 0.175(35)

Epoxy GNP (7 nm) Mixing/curing inoven

0 3.61 0.26 SEM (homogeneousand randomdispersion of theblocks)

2.72 Dynamic ISE (<0.2 lm)H and Econstant(>0.2 lm)

[169]1.0 3.68(2) 0.26(0) 2.82(4)2.0 3.90(8) 0.25(�2) 2.92(8)3.0 3.93(9) 0.25(�2) 3.04(12)4.0 3.97(10) 0.24(�4) 3.15(16)5.0 4.30(19) 0.26(0) 3.26(20)6.0 4.70(30) 0.26(0) 3.36(30)

Epoxy GNP Mixing/stirringcuring

0.1 (0)b NA AFM (grapheneclusters)

From unloading(NA)

[170]

Epoxy EG Sonication/photocuring

0 0.32 SEM (micrometeragglomerates)

3.0 From unloading(NA)

[171]1.0 0.34(6) 1.8

(�40)Epoxy FGS (stacks up to 7 sheets) Sonication/

photocuring0 0.32 SEM (stacked groups

of sheets)3.0 From unloading

(NA)[171]

1.0 0.34(6) 3.0 (0)Epoxy TEG Functionalized nitric

and sulphuric acid (Blocksof graphene nanosheets)

Mixing/curing

0 3.27 0.396 SEM (homogeneousdispersion of theblocks)

From unloading(OP)d

[172]1.0 3.82(17)b 0.386(�2)2.0 4.00(22)b 0.400(1)

Chitosan Arc-discharge N-doped GS(Few layered graphenesheets)

Solutioncasting

0 3.2 0.30 Raman/TEM(homogeneousdispersion of thesheets)

From unloading(NA)

[156]0.1 6.7(110)b 0.42(40)0.2 7.3(128)b 0.42(40)0.3 7.8(144)b 0.39(30)0.6 6.6(106)b 0.40(35)2.3 6.7(110)b 0.47(57)

PET CVD GSc (1–2 layers) Solutionmixing/coating

0 3.7 0.15 From unloading(OP)

[157]NA 4.3(16) 0.30(100)

PP GO (monolayer, bilayer (2–3 nm), multilayer)

Melt-blending/hot-pressing

0 0.459 0.00661 SEM (agglomerates) From unloading(OP)d

[154]0.1 0.539(15) 0.00715(8)0.5 0.592(28) 0.0080(21)1.0 0.666(45) 0.0087(30)

PP OTES-modified GO(monolayer, bilayer (2–3 nm), multilayer)

Melt-blending/hot-pressing

0 0.459 0.00661 SEM (homogeneousdispersion)

From unloading(OP)d

[154]0.1 0.613(33) 0.0218(228)0.5 0.651(42) 0.0254(283)1.0 0.743(62) 0.0256(286)

24A

.M.D

íez-Pascualet

al./Progressin

Materials

Science67

(2015)1–

94

Table 2 (continued)

Matrix Nanofiller type (thickness) Processingmethod



DynamicDSI E0

(GPa)(DE0 , %)




E0DMA

(GPa)(DE0DMA ,%)

Y (GPa)(D, Y %)

Method of analysis Depthvariation(range)

Ref.

PVA Acid-functionalized FG(few layer)

Solutioncasting

0 0.657 0.038 Raman (goodinteraction GS-matrix)

42 From unloading(OP)

[21]0.2 0.688(5) 0.039(3) 42.8(2)0.4 0.734(12) 0.0441(15) 46.2(10)0.6 0.885(35) 0.0557(47) 47.5(13)

PVA GNP (EOGCNF) Partiallyoxidized (monolayer)

Solutioncasting

0 0.230 0.0129 – From unloading(OP)d

[160]0.1 0.257(4) 0.0135(5) (10)0.5 0.388(20) 0.0153(19) (18)1.0 0.349(52) 0.0144(12) (�2)2.0 0.651(80) 0.0183(42) (�4)4.0 0.854(160) 0.0211(62) (�5)

PVA GOf,g Layer-by-layer

0 8.88 0.34 AFM (homogeneousdispersion)

From unloading(OP)

[173]– 17.64(99) 1.15(240)

PMMA Amide-functionalized FG(few layer)

Solutioncasting

0 2.12 0.140 From unloading(OP)

[21]0.2 3.36(58) 0.1449(4)0.4 3.61(70) 0.1456(4)0.6 3.65(72) 0.1531(9)

PMMA EG Melt-blending

0 4.3 0.42 SEM (regulardispersion,agglomerates at1.0 wt%)

From unloading(OP)

[155]0.25 4.5(5)b 0.45(8)0.5 4.7(10)b 0.51(23)1.0 4.7(10)b 0.49(16)

PAA PEI-modified GO Layer-by-layer

0 1.3 0.160 SEM, AFM, IR(increase inroughness)

Dynamic (NA)j ISE (<0.2 lm) [174]NA 1.9(46) 0.240(50)

PU Highly aligned rGOg,h

(1 nm)Solutioncasting

0 0.4 0.025 SEM (homogeneousdispersion)

0.3 From unloading(NA)

[158]1.0 1.0(150) 0.035(40) 1.25(320)2.0 3.1(675) 0.050(100) 3.00(920)5.0 7.4(1750) 0.820(228) 6.30(2000)

PU GOh (<50 nm) Solutioncasting

0 0.022e 0.061 SEM (homogeneousdispersion)

0.011 NA [159]1.0 0.027(23) 0.078(27) 0.031(180)4.0 0.061(175) 0.115(88) 0.092(735)

a Data for an average platelet length of 1 lm.b Results showing ‘‘nose’’ effect (too short holding time [155,170]; no holding time [156,172]).c Graphene overlayer onto PET substrate.d AFM nanoindentation with a silicon tip.e Materials with matrix modulus <0.15 GPa, possibly affected by a large uncertainty in the detection of the point of initial contact.f Planar orientation.g Partially reduced graphene oxide.h Nanofiller covalently bonded to the polymer matrix.i ISO 14577 test.j Lack of experimental details, indenter geometry not reported.

A.M

.Díez-Pascual

etal./Progress

inM

aterialsScience

67(2015)

1–94

25

Table 3Nanomechanical and macroscopic properties for polymer nanocomposites reinforced with different types of organic nanomaterials: carbon nanofibres (CNFs), nanodiamond (ND), carbonblack (CB), fullerenes, cellulose nanocrystals (CNC), and hybrids with two carbon-based nanofillers. Columns as indicated in Table 1.

Matrix Nanofillertype

Processingmethod



DynamicDSI E0

(GPa)(DE0 , %)


Morphology (Dispersion,interaction)

Crystallinity[%](%variation)

E0DMA

(GPa)ðDE0DMA ,%)

Y (GPa)(DY, %)

Method ofanalysis


Ref.

Epoxy VGCNF Calendaring/curing

0 4.50 0.375 TEM (Agglomeratesincreasing with nanofibrecontent. Good fibrewetting)

2.18 From unloading(OP)

[197]1.5 4.05(�10) 0.310(�17) 2.23(2)3.0 4.50(0) 0.370(�2) 2.36(8)5.0 4.70(27) 0.432(15) 2.82(30)

Epoxy AA-coatedCNF

Mixing/curing 0 3.60 0.21 SEM (agglomerates, fibrepull-out)

3.57 Dynamic ISE(<0.3 lm)

[198]0.5 3.59(0) 0.21(0) 3.58(0)

Epoxy ND�NH2a Stirring/curing 0b 1.90 TEM (some agglomerates,

particularly at highloadings)

Dynamicj [199]0c 2.500d 2.300e 1.910f 2.229.1g 1.84(0)10.0d 2.41(5)14.5e 2.43(30)16.7c 3.08(20)40.0f,i 3.94(78)53.3h,i 12.1(426)

Epoxy ox-ND Stirring/curing 0b 1.90 TEM (agglomerates) Dynamicj [199]0d 2.308.8b 2.22(15)8.8d 2.50(9)

Epoxy CB Ultrasonicmixing/curing

0 6.0 (<10) SEM (uniform dispersion) From unloading(OP)

Reverse ISE(0.2–0.25 lm)

[193]1.4 6.24(4)2.8 6.72(12)3.4 7.38(23)4.6 8.4(40)5.8 6.48(8)

Epoxy Mixedfullerenes

Mixing/curing 0 3.6 0.157 NA [115]10.0 1.8(�50) 0.058(�63)

PVA (90)/PAA(10)

CNC Sonication/Solution casting

0 2.20 AFM (some aggregates) 1.50 Single readingk

(Herzt)l[200]

10 3.10(40) 2.10(40)20 2.10(�5) 2.05(37)

PVA CNC Sonication/Solution casting

0 2.0 AFM (heterogeneousdispersion)

1.20 Single readingk

(Herzt)l[200]

15 3.0(50) 1.40(17)PVAm CNCn Mixing/

electrospinning0 2.10n Single reading

(OP)j[201]

1.0 2.31(10)n

3.0 2.72(30)n

5.0 3.16(50)n

10.0 4.26(105)n

15.0 5.88(180)n

20.0 7.58(262)n

26A

.M.D

íez-Pascualet

al./Progressin

Materials

Science67

(2015)1–

94

Tab

le3

(con

tinu

ed)

Mat

rix

Nan

ofill

erty

pePr

oces

sin

gm

eth

odN

anofi

ller

con

ten

t(w

t%)

Qu

asi-

stat

icD

SIE

(GPa

)( D

E,%

)

Dyn

amic

DSI

E0

(GPa

)( D

E0,%

)

Har

dnes

sH

(GPa

)(D

H,%

)

Mor

phol

ogy

(Dis

pers

ion

,in

tera

ctio

n)

Cry

stal

lin

ity

[%]

(% vari

atio

n)

E0 DM

A

(GPa

)ðD

E0 DM

A,

%)

Y(G

Pa)

( DY,

%)

Met

hod

ofan

alys

isD

epth

vari

atio

n(r

ange

)

Ref

.

PVA

ox-N

DSo

nic

atio

n/

solu

tion

cast

ing/

00.

670.

0383

TEM

(goo

ddi

sper

sion

)42

From

un

load

ing

(OP)

[121

,202

]0.

20.

87(3

0)0.

0437

(14)

52.6

(25)

0.4

0.96

(43)

0.05

28(3

8)55

.0(3

1)0.

61.

33(9

9)0.

0684

(80)

56.6

(35)

PVA

cC

NC

Solu

tion

cast

ing/

com

pres

sion

03.

000.

060

AFM

(agg

rega

tes

atco

nte

nts

>1w

t%)

From

un

load

ing

(OP)

[203

]1.

04.

20(4

0)0.

102(

70)

2.0

4.30

(43)

0.09

3(55

)3.

05.

40(8

0)0.

156(

160)

PLLA

as rece

ived

-N

D

Solu

tion

cast

ing

02.

60.

05TE

M(p

oor

disp

ersi

on)

Dyn

amic

(Her

zt)j,l

[187

]1.

02.

8(9)

0.12

(140

)3.

02.

7(5)

0.11

(120

)PL

LAO

DA

-m

odifi

edN

D

Solu

tion

cast

ing

02.

60.

05TE

M(g

ood

disp

ersi

on)

NM

R:

good

inte

ract

ion

wit

hm

atri

x

Incr

ease

incr

ysta

llin

ity

Dyn

amic

(Her

zt)j,l

[187

]1.

05.

3(10

7)0.

21(3

20)

3.0

5.5(

113)

0.25

(400

)5.

05.

9(12

9)0.

26(4

20)

7.0

6.8(

165)

0.31

(520

)10

.07.

9(20

6)04

6(82

0)PA

-11m

ox-N

DSo

nic

atio

n/

elec

tros

pin

nin

g0

2.0

0.07

5SE

M(i

ncr

ease

insu

rfac

ero

ugh

nes

sw

ith

incr

easi

ng

load

ing)

Dyn

amic

j[2

04]

2.5

3.5(

75)

0.08

2(9)

10.0

3.5(

75)

0.09

4(25

)20

.08.

0(30

0)0.

14(8

6)PA

-11

ox-N

DM

illi

ng/

ther

mal

spra

yde

posi

tion

01.

80n

0.15

0SE

M(p

oor

dist

ribu

tion

)D

ynam

ic[2

05]

7.0

1.62

( �10

)n0.

146(�

2)13

.01.

70( �

5)n

0.14

3(�

4)25

.01.

70( �

5)n

0.14

0(�

6)PS

-bas

edo

CB

Ult

raso

nic

atio

n/

curi

ng

01.

20.

045

SEM

(un

ifor

mdi

sper

sion

)0.

240

From

un

load

ing

(OP)

[206

]1.

01.

50(2

5)0.

054(

20)

0.26

5(10

)2.

02.

16(8

0)0.

061(

38)

0.30

0(25

)3.

53.

2(1

70)

0.10

1(12

5)0.

336(

40)

5.5

5.4(

350)

0.20

5(35

0)0.

430(

82)

PSPS

-coa

ted

CN

FM

ixin

g/cu

rin

g0

3.60

SEM

(un

ifor

mdi

sper

sion

)2.

10D

ynam

icIS

E(<

0.3

lm

)R

ever

seIS

E(0

.3–4

lm

)

[198

]1.

03.

70(3

)2.

15(2

)2.

03.

73(4

)2.

10(0

)3.

03.

90(9

)2.

30(1

0)5.

03.

80(6

)2.

27(8

)PM

MA

mC

NC

nSo

luti

onm

ixin

g/el

ectr

ospi

nin

g

05.

30SE

M(g

ood

disp

ersi

onat

low

load

ings

)D

ynam

icj

E0co

nst

ant

(0.1

–0.

25l

m)

[207

]5

5.94

(12)

96.

15(1

6)17

6.25

(18)

236.

15(1

6)33

6.90

(30)

416.

25(1

8)

(con

tinu

edon

next

page

)


Table 3 (continued)

Matrix Nanofillertype

Processingmethod



DynamicDSI E0

(GPa)(DE0 , %)



Crystallinity[%](%variation)

E0DMA

(GPa)ðDE0DMA ,%)

Y (GPa)(DY, %)

Method ofanalysis


Ref.

PAN ox-ND Solution casting 0 5.8 TEM (poor dispersion) 2.0 NA [208]5.0 6.32(9) 2.32(16)10.0 7.08(22) 2.50(25)20.0 9.0(55) 2.50(25)30.0 13.0(125) 2.45(22)

PAN ox-ND (de-aggregated)

Bead milling/Solution casting

0 5.8 TEM (good dispersion) 2.0 NA [208]5.0 6.13(10) 2.30(15)10.0 6.96(20) 2.58(28)20.0 7.37(27) 3.00(50)30.0 11.14(92) 3.60(80)

PAN deox-ND Solution casting 0 5.8 TEM (poor dispersion) NA [208]5.0 5.51(�5)10.0 4.52(�22)20.0 6.96(20)30.0 7.83(35)

PVA ox-SWCNTs/ox-FG

Solution casting 0 0.66 0.038 42.0 From unloading(OP)

[121]0.4/0.2 9.3(1300) 0.366(865) 56.5(35)0.2/0.4 8.6(1200) 0.337(787) 57.5(37)

PVA ox-SWCNTs/ox-ND


[121]0.4/0.2 7.5(1036) 0.314(726) 55.1(31)0.2/0.4 9.3(1300) 0.353(829) 57.2(37)

PVA ox-FG/ox-ND


[121]0.4/0.2 1.6(150) 0.066(75) 55.1(31)0.2/0.4 1.3(97) 0.061(61) 54.8(30)

a Nanofiller grafted to the polymer matrix.b r (overall amine groups/epoxide groups including curing agent and ND-NH2 when appropriate) = 0.39.c r (overall amine groups/epoxide groups including curing agent and ND-NH2 when appropriate) = 0.57.d r (overall amine groups/epoxide groups including curing agent and ND-NH2 when appropriate) = 1.e r (overall amine groups/epoxide groups including curing agent and ND-NH2 when appropriate) = 1.54.f r (overall amine groups/epoxide groups including curing agent and ND-NH2 when appropriate) = 0.64.g r (overall amine groups/epoxide groups including curing agent and ND-NH2 when appropriate) = 0.34.h r (overall amine groups/epoxide groups including curing agent and ND-NH2 when appropriate) = 1.18.i Without curing agent.j spherical indenter [199,187], spheroconical [204,205], conical [207] and cube-corner [201].k AFM nanoindentation.l Hertz model used ([200] applies Hertz to loading and unloading curves, values in the table corresponding to the load fitting).

m Composite fibre; naligned along the fibre axis.n Results showing ‘‘nose’’ effect at the beginning of unloading.o Shape memory polymer (PS: 30 wt%, styrene: 55 wt%, vinyl neodecanoate: 10 wt%).

28A

.M.D

íez-Pascualet

al./Progressin

Materials

Science67

(2015)1–

94

Table 4Nanomechanical and macroscopic properties for polymer layered silicates nanocomposites. Columns as indicated in Table 1.

Matrix Nanofiller type Processing method Nanofillercontent(wt%)

Quasi-staticDSIE(GPa)(DE, %)

DynamicDSI E0

(GPa)(DE0 ,%)

HardnessH(GPa)(DH, %)


Crystallinity(%) (%variation)

E0DMA (GPa)(DE0DMA, %)

Y (GPa)(DY, %)

Method ofanalysis


Ref.

PA-6 Organicallymodified clay(I30TC)

Melt blending/injection moulding

0 1.06 0.054 TEM (agglomeration at 10 wt%)SAXS (Exfol 65 wt% Inter/exf > 5 wt%)

15 1.50 Fromunloading(OP)

[223]2.5 1.84(74) 0.100(85) 13.75(�8) 2.25(50)5 1.97(86) 0.118(119) 12(�20)7.5 2.3(117) 0.138(156) 10.5(�30) 2.66(77)10 2.42(128) 0.141(161) 10(�33)

PA-6 CoAl-LHD(organo-modifiedlayered doublehydroxides)

In situpolymerization/compressionmoulding

0 1.13 0.063 TEM, X-ray (homogeneousdispersion. Exfoliation)

1.0 Dynamic ISE(<1 lm)

[224]0.5 1.46(29) 0.077(22) 1.6(60)1 1.49(32) 0.079(25) 2.0(100)2 1.50(35) 0.085(35) 2.1(110)

PA-6 Organicallymodified clay(OMMT-L, D,DD)

Solution blending/compressionmoulding

0 3.352 3.482 0.121 X-ray (intercalation) 27.6 Dynamicand fromunloading(OP)

[225]L-3 4.431(32) 4.392(26) 0.194(60) 26.4(�4)L-6 4.748(42) 4.855(39) 0.200(65) 26.2(�5)L-9 5.460(63) 5.445(56) 0.170(41) 23.1(�16)D-9 4.766(42) 4.392(26) 0.168(39)DD-9 4.417(32) 0.172(42) 25.5(�8)

PA-6 Surface modifiedMMT (1.34TCN)


0 1.621 0.179 X-ray (clay platelets uniformlydispersed. Exfoliation)

1.376 Fromunloading(OP)

[226]4 2.265(40) 0.198(11) 1.469(7)6 3.376(108) 2.136(55)

PA-6,6 Surface modifiedMMT (1.34TCN)




[226]4 2.917(24) 0.204(6) 1.800(8)6 2.909(24) 0.290(50) 1.819(9)

PA6/PA6,6(30/70)

Surface modifiedMMT (1.34TCN)




[226]4 2.875(21) 0.233(13) 1.648(0.7)6 2.763(17) 0.222(13) 1.705(4)

PA6/PA6,6(50/50)



0 1.938 0.172 X-ray (clay platelets uniformlydispersed Exfoliation)


[226]2 1.746(�10) 0.163(�5) 1.566(�2)4 3.004(55) 0.365(112) 1.686(5)6 3.149(63) 0.352(105) 1.704(7)

PA6/PA6,6(70/30)





[226]4 2.881(32) 0.283(41) 1.713(6)6 3.274(50) 0.341(70) 1.742(8)

PA-6,6 Surface modifiedMMT (1.34TCN)


0 2.30a 0.098a TEM, X-ray (homogeneousdispersion. Exfoliation)

52c 1.89 3.04 Dynamic ISE(<0.4 lm)

[227–229]0 2.27b 0.099b – – –

1 2.39(4)a 0.105(7)a 42(�19)c 2.22(18) 3.19(5)1 2.34(3)b 0.098(�1)b –2 2.57(12)a 0.116(18)a 41.5(�20)c 2.24(19) 3.41(12)2 2.42(7)b 0.106(7)b – – –5 2.72(18)a 0.125(28)a 38(�27)c 2.36(25) 3.91(29)5 2.44(8)b 0.106(7)b – – –

PA-6,6 OrganicallymodifiedHectorite


0 1.6d,e 0.0875d Fractography 2.6d NA [230]1 2(25)d,e 0.103(18)d 3.5(35)d

5 2.5(56)d,e 0.119(36)d 4(54)d


A.M

.Díez-Pascual

etal./Progress

inM

aterialsScience

67(2015)

1–94

29

Tabl

e4

(con

tinu

ed)

Mat

rix

Nan

ofill

erty

pePr

oces

sin

gm

eth

odN

anofi

ller

con

ten

t(w

t%)

Qu

asi-

stat

icD

SIE(

GPa

)( D

E,%

)

Dyn

amic

DSI

E0

(GPa

)( D

E0,%

)

Har

dnes

sH

(GPa

)( D

H,%

)

Mor

phol

ogy

(Dis

pers

ion

,in

tera

ctio

n)

Cry

stal

lin

ity

(%)

(%va

riat

ion

)

E0D

MA

(GPa

)( D

E0D

MA,%

)Y

(GPa

)(D

Y,%

)M

eth

odof

anal

ysis

Dep

thva

riat

ion

(ran

ge)

Ref

.

PA-6

,6/S

EBS-

g-M

AO

rgan

ical

lym

odifi

edM

MT

(Cl

30B

)

Mel

tbl

endi

ng/

inje

ctio

nm

ould

ing

01.

350.

110

TEM

,AFM

(goo

ddi

sper

sion

.Ex

foli

atio

n)

2.95

From

un

load

ing

(OP)

[231

]5

1.82

(35)

f0.

140(

27)f

2.65

(�10

)f

PA-1

1Su

rfac

em

odifi

edM

MT

(1.3

4TC

N)

Mel

tbl

endi

ng/

com

pres

sion

mou

ldin

g

01.

350.

086

1.08

0.61

Dyn

amic

ISE

(<1.

0l

m)

[232

]1

1.58

(17)

0.10

0(16

)1.

31(2

1)0.

62(2

)2

1.65

(22)

0.10

2(19

)1.

41(3

1)0.

75(2

3)5

1.78

(32)

0.10

9(27

)1.

84(7

0)0.

86(4

1)PA

-12

Surf

ace

mod

ified

MM

T(1

.34T

CN

)M

elt

blen

din

g/in

ject

ion

mou

ldin

g0

1.03

0.06

TEM

(hom

ogen

eou

sdi

sper

sion

.Ex

foli

atio

n)

1.40

1.03

Dyn

amic

ISE

(<0.

3l

m)

[233

]1

1.39

(35)

0.09

(50)

1.42

(1)

1.10

(7)

21.

44(4

0)0.

09(5

0)1.

45(4

)1.

12(9

)5

1.56

(52)

0.10

(67)

1.49

(6)

1.49

(45)

PA-1

2La

yere

dsi

lica

ten

anoc

lay

(LK

-PA

-C

R1)

Mel

tbl

endi

ng/

inje

ctio

nm

ould

ing

02.

212g

0.15

6SE

M,T

EM(I

nte

rcal

atio

n/

Exfo

liat

ion

)Fr

omu

nlo

adin

g(O

P)

[234

]1

2.34

0(6)

g0.

161(

3)3

2.90

0(31

)g0.

202(

30)

53.

835(

73)g

0.33

8(11

7)PL

AO

rgan

ical

lym

odifi

edM

MT

(Cl

30B

)

Mel

tbl

endi

ng/

com

pres

sion

mou

ldin

g

04.

360

0.21

9TE

M(c

lay

fin

ely

dist

ribu

ted

Inte

rcal

atio

n/E

xfol

iati

on)

403.

401

Dyn

amic

[235

]1

4.38

0(0.

5)0.

223(

2)40

(0)

3.91

4(15

)3

4.97

0(14

)0.

288(

32)

40(0

)4.

901(

44)

54.

990(

15)

0.29

8(36

)40

(0)

5.57

7(64

)PL

A-K

Ffi

bre

Org

anic

ally

mod

ified

MM

TM

elt

blen

din

g/in

ject

ion

mou

ldin

g0

6.25

0.25

From

un

load

ing

(OP)

[236

]3

6.21

(�1)

0.26

(4)

PHB

VO

rgan

ical

lym

odifi

edM

MT

(Cl

30A

)

Solu

tion

inte

rcal

atio

n/S

olve

nt

cast

ing

00.

761

0.04

6H

omog

eneo

us

disp

ersi

on0.

633

Dyn

amic

ISE

(<2.

0l

m)

[237

]1

0.88

1(16

)0.

051(

11)

1.04

3(65

)2.

51.

270(

67)

0.07

7(67

)1.

311(

107)

51.

546(

103)

1.67

7(16

5)PP

/PP-

g-M

A(9

0/10

)O

rgan

ical

lym

odifi

edM

MT

(Cl

15A

)

Sin

gle-

scre

wex

tru

sion

wit

hin

-lin

eSC

CO

2(S

SE)

Twin

-sc

rew

extr

usi

on(T

SE)

02.

18TE

M(d

iffe

ren

tde

gree

ofdi

sper

sion

for

TSE

and

SSE

Inte

rcal

atio

n/E

xfol

iati

on)

Dyn

amic

[218

]SS

E-3

2.29

(5)

TSE-

32.

33(6

)SS

E-5

2.22

(2)

TSE-

52.

50(1

3)PP

-g-M

AO

rgan

ical

lym

odifi

edM

MT

(I31

PS)

Mel

tbl

endi

ng/

inje

ctio

nm

ould

ing

0TE

MM

ixtu

reof

hig

hly

inte

rcal

ated

and

wel

lex

foli

ated

Dyn

amic

ISE

(<0.

2l

m)

[219

]2.

5(6

8)(1

2)5

(64)

(20)

7.5

(76)

(24)

10(9

8)(3

2)H

DPE

/HD

PE-

g-M

A(9

8/2)

Org

anic

ally

mod

ified

MM

T(C

l15

A,N

1.44

P)

Mel

tbl

endi

ng/

com

pres

sion

mou

ldin

g

01.

070.

050

SEM

Agg

lom

erat

ion

/wel

ldi

sper

sed

depe

ndi

ng

onty

pean

dco

mpo

siti

onIn

terc

alat

ion

/Ex

foli

atio

n

67.1

0.79

30.

663

From

un

load

ing

(OP)

[238

]15

A-1

1.11

(4)

0.05

0(0)

63.6

(�5)

0.73

8(�

7)0.

617(�

7)15

A-2

.51.

19(1

1)0.

065(

30)

63.7

(�5)

0.75

6(�

5)0.

670(

1)15

A-5

1.18

(10)

0.07

1(42

)60

.4(�

10)

0.86

4(8)

0.58

1(�

12)

N1.

44P-

11.

07(0

)0.

048(�

4)62

.4(�

7)0.

461(�

42)

0.56

9(�

14)

N1.

44P-

2.5

1.06

(�1)

0.05

7(14

)62

.9(�

6)0.

912(

15)

0.62

1(�

6)

N1.

44P-

51.

20(1

2)0.

061(

22)

63.9

(�5)

0.75

0(�

5)0.

664(

0.2)


Tab

le4

(con

tinu

ed)

Mat

rix

Nan

ofill

erty

pePr

oces

sin

gm

eth

odN

anofi

ller

con

ten

t(w

t%)

Qu

asi-

stat

icD

SIE(

GPa

)( D

E,%

)

Dyn

amic

DSI

E0

(GPa

)( D

E0,%

)

Har

dnes

sH

(GPa

)( D

H,%

)

Mor

phol

ogy

(Dis

pers

ion

,in

tera

ctio

n)

Cry

stal

lin

ity

(%)

(%va

riat

ion

)

E0D

MA

(GPa

)( D

E0D

MA,%

)Y

(GPa

)( D

Y,%

)M

eth

odof

anal

ysis

Dep

thva

riat

ion

(ran

ge)

Ref

.

PEN

Org

anic

ally

mod

ified

MM

TM

elt

blen

din

g/co

mpr

essi

onm

ould

ing

05.

02h

TEM

(in

terc

alat

ion

)29

.1h

2.83

hD

ynam

ic[2

39]

25.

39(7

)h28

.8(�

1)h

3.53

(25)

h

PEO

Na+

MM

T(G

105)

Org

ano-

clay

(I28

)So

luti

on(S

)M

elt

blen

din

g(M

)S-

00.

621

0.03

0X

-ray

(in

terc

alat

ion

)Fr

omu

nlo

adin

g(O

P)

ISE

(1–

5.5

lm

)[2

40]

S-G

105-

50.

421(�

32)

0.02

5(�

17)

S-G

105-

201.

694(

173)

0.05

4(80

)S-

G10

5-50

2.27

9(26

7)0.

088(

193)

S-G

105-

801.

619(

161)

0.10

4(24

7)S-

I28-

50.

751(

21)

S-I2

8-20

1.16

2(87

)0.

058(

93)

M-0

0.82

10.

043

M-G

105-

50.

966(

18)

00.3

6(�

16)

M-G

105-

201.

544(

88)

0.05

4(26

)

PPy

Org

anic

ally

mod

ified

MM

TLa

yer

stru

ctu

re

Elec

tro-

depo

siti

on0

2.3

0.12

5SE

M(h

omog

eneo

us

disp

ersi

on)

NA

[241

]0.

013.

3(44

)0.

17(3

6)0.

052.

7(17

)0.

18(4

4)0.

12.

5(9)

0.12

5(0)

0.2

1.3(�

44)

0.03

(�76

)0.

51.

2(�

48)

0.05

(�60

)PP

CO

rgan

ical

lym

odifi

edM

MT

(Cl

20B

)

Solu

tion

inte

rcal

atio

n0

1.41

40.

105

TEM

(agg

rega

tion

/wel

ldi

sper

sed

Inte

rcal

atio

n/E

xfol

iati

on)

Dyn

amic

ISE

(<0.

2l

m)

[242

]In

terc

1.95

0(38

)0.

114(

9)Ex

f2.

340(

66)

0.12

1(15

)PS

Org

anic

ally

mod

ified

MM

T(C

l15

A)

Mel

tbl

endi

ng/

com

pres

sion

mou

ldin

g

05.

230.

361

Inte

rcal

atio

nD

ynam

icN

A[2

43]

0.5

5.51

(5)

0.38

1(6)

PSSu

rfac

em

odifi

edM

MT

(Cl

20A

)M

elt

blen

din

g/sp

inco

ated

03.

60.

290

Dyn

amic

ISE

(<0.

4l

m)

[244

]10

5.1(

42)

0.39

0(35

)PD

DA

Na+

MM

TM

ult

ilay

erfi

lmLa

yer-

by-l

ayer

depo

siti

onpr

oces

s50

laye

r–

0.42

AFM

(sm

ooth

surf

ace

Inte

rcal

atio

nM

MT

laye

rs)

Dyn

amic

Rev

erse

ISE

(0.2

–2

lm

)

[245

]10

0la

yer

9.5

0.32

Ch

itos

an/H

A(6

0/40

)N

a+M

MT

Solu

tion

inte

rcal

atio

nC

hi-

05.

290.

16A

FM(w

ell

dist

ribu

ted

nan

opar

ticl

es)

FTIR

(str

ong

inte

rfac

ial

inte

ract

ion

.In

terc

alat

ion

)

NA

[246

]C

hi-

108.

28(5

7)0.

28(7

5)C

hi/

HA

P-0

7.02

(33)

0.22

(38)

Ch

i/H

AP-

109.

15(7

3)0.

28(7

5)

Epox

yO

rgan

ical

lym

odifi

edM

MT

(Cl

93A

)

Stir

rin

g+

son

icat

ion

/ca

stin

gC

uri

ng

inov

en

02.

870.

209

TEM

,X-r

ay(a

gglo

mer

atio

n/w

ell

disp

erse

dde

pen

din

gon

com

posi

tion

.Exf

olia

tion

<2.5

wt%

Inte

rcal

atio

n>2

.5;w

t%)

2.37

From

un

load

ing

(OP)

ISE

(<2.

0l

m)

[247

]1

3.08

(7)

0.21

2(1)

2.68

(13)

2.5

3.22

(12)

0.21

6(3)

2.73

(15)

53.

30(1

5)0.

218(

4)2.

79(1

8)7.

53.

45(2

0)0.

220(

5)2.

85(2

0)

(con

tinu

edon

next

page

)


Table 4 (continued)


Quasi-staticDSIE(GPa)(DE, %)

DynamicDSI E0

(GPa)(DE0 ,%)

HardnessH(GPa)(DH, %)



E0DMA (GPa)(DE0DMA, %)

Y (GPa)(DY, %)

Method ofanalysis


Ref.

Epoxy Na+MMTSilylated MMT(MMT-A1100,MMT-A1120)

Sonication (S)Sonication + ballmilling (SB)/castingCuring in oven

0 4.01 0.145 SEM (clay aggregates.Intercalation)


[248]Na-S-3 4.10(2) 0.162(12) 1.79(7)A1100-S-3 4.30(7) 0.195(35) 1.9(16)A1120-S-3 4.34(8) 0.220(52) 1.94(16)Na-SB-3 3.70(�8) 1.83(10)A1100-SB-3

4.20(5) 0.178(23) 1.84(10)

A1120-SB-3

4.29(7) 0.157(8) 1.87(12)

CEAR Na+ MMT(NMMT)

Stirring/electrodepositionCuring in oven

0 3.6 0.172 TEM (clay randomly dispersed.Intercalation/Exfoliation)

1.1 NA [249]1 4.0(11) 0.175(2) 1.18(7)2 4.1(14) 0.175(2) 1.2(9)4 4.2(17) 0.188(9) 1.2(9)5 4.25(18) 0.188(9) 1.4(27)

AAER Na+ MMT(NMMT)HydrothermalmethodLaminates(HMMT)

Stirring/electrodepositionCuring in oven

NMMTfilmHMMTfilm (8%polymer)

2.9 X-ray (intercalation) Dynamic [250]5.0 (72)

UPE Organicallymodified layeredsilicate


0 5.39 0.301 X-ray (intercalation) Fromunloading(OP)

[251]1 6.19(15) 0.387(29)3 6.07(13) 0.372(24)5 6.65(23) 0.343(14)

PU Organicallymodified MMT(C20, C30)

In situpolymerization/casting Curing inoven

0 0.014i 0.0017 X-ray (intercalation/Exfoliation) Fromunloading(OP)

[252]C20-1 0.017(21) 0.0018(6)C20-3 0.021(50) 0.0020(18)C20-5 0.023(64) 0.0024(41)C30-1 0.021(50) 0.0022(29)C30-3 0.025(79) 0.0026(53)C30-5 0.029(107) 0.0028(65)

a Polished samples.b Unpolished samples.c Measured by WAXS.d Dried as moulded samples (DAM [230]).e Results showing ‘‘nose’’ effect at the beginning of unloading.f PA-6,6 reinforced with OMMT first and then blended with SEBS-g-MA (N3 in [231]).g Reduced modulus.h Annealed samples.i Material with matrix modulus < 0.15 GPa, where it is difficult to determine the initial point of contact.

32A

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Science67

(2015)1–

94

Table 5Nanomechanical and macroscopic properties for polymer nanocomposites incorporating inorganic spherical nanofillers. Columns as indicated in Table 1.


Quasi-static DSIE (GPa)(DE, %)

DynamicDSI E0

(GPa)(DE0 , %)

Hardness H(GPa)(DH, %)



E0DMA (GPa)(DE0DMA , %)

Y (GPa)(DY, %)

Method ofanalysis


Ref.

Epoxy Nanosilica 25 nm ø Mechanical mixingpreheatingultrasonication

4(40 �C) 2.3–5.5a SEM (temp. dependentcluster size)

Fromunloading(OP)

[262]4(80 �C) 7.4a

4(100 �C) 8.8a

Epoxy Fumed silica 20 nmø forming 200–300 nm ø clusters

Ultrasonicationmechanical mixingcuring RT + 90 �C

0 3.39 0.318 SEM, Optical microscope(clusters)


[263]0.5 3.33(�2) 0.299(�6) 3.25(�2)1 3.00(�12) 0.295(�7) 2.90(�12)2 3.05(�10) 0.273(�14) 2.96(�10)3 2.55(�25) 0.256(�19) 2.48(�25)

Epoxy Nanosilica 10 nm ø Ultrasonicationmechanical mixingcuring RT

0 3.30a 0.143 TEM (good dispersion) 1.46c Fromunloading(OP)

[264]1 3.58(8)a 0.160(12) 1.77(21)c

3 3.87(17)a 0.177(24) 2.20(51)c

5 4.00(21)a 0.180(26) 2.86(96)c

Epoxy Surface-modifiednanosilica 20 nm ø

Sol–gel 0 2.51 0.190 3.27 Fromunloading(OP)d

[76,267]5.3 2.66(6) 0.210(11) 3.60(10)10.3 2.72(8) 0.215(13) 3.71(13)13.3 2.96(18) 0.222(17) –16.7 3.07(22) 0.225(18) 3.87(18)24.2 3.51(40) 0.250(32) –

Epoxy Surface-modifiednanosilica 25 nm ø

Sol–gel 0 5.88e 2.90e Dynamic ISE (<200 lN) [75,267]1.8 6.08(5)e 3.13(8)e

5.3 6.69(14)e 3.38(17)e

10.3 6.74(15)e 3.55(22)e

16.7 6.94(18)e 3.71(28)e

22.7 8.59(46)e 4.07(40)e

Epoxy PMMA + fumednanosilica 14 nm ø

Blending PMMA + SiO2

mechanical mixingcuring 50 �C + 150 �C

0 3.80a 0.143f SEM, AFM (good dispersion ofPMMA domains and SiO2

particles in the domains)

Fromunloading

[268]0.1 3.74(�2)a 0.136(�5)f

0.2 3.75(�1)a 0.138(�3)f

0.3 3.96(4)a 0.147(3)f

0.4 3.96(4)a 0.146(2)f

0.5 3.96(4)a 0.146(2)f

Poly-phenyleneSiLK™I

Nanosilica 8 nm ø Spin coating thermalcuring

0 3.84 0.283 TEM (local aggregations) Dynamic [77]5.2 4.02(5) 0.285(1)12.7 4.50(17) 0.315(11)21.4 4.32(13) 0.320(13)31.9 4.66(21) 0.330(17)42 5.77(50)

Poly-phenyleneSiLK™D

Nanosilica 8 nm ø Spin coating thermalcuring

0 3.26 0.240 TEM (local aggregations) Dynamic [77]10.5 3.45(6) 0.250(4)16.2 3.74(15) 0.251(5)21.4 4.01(23) 0.268(12)26.7 4.81(47) 0.308(28)

MDMA Nanosilica 10–15 nm ø

Ultrasonication UVcuring

0 1.91 0.078 0.650 Dynamic ISE (<0.2 lm) [210]1 3.09(62) 0.131(68) 1.20(85)2 3.62(89) 0.203(160) 1.45(123)


A.M

.Díez-Pascual

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aterialsScience

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33

Tabl

e5

(con

tinu

ed)

Mat

rix

Nan

ofill

erty

pePr

oces

sin

gm

eth

odN

anofi

ller

con

ten

t(w

t%)

Qu

asi-

stat

icD

SIE

(GPa

)( D

E,%

)

Dyn

amic

DSI

E0

(GPa

)( D

E0,%

)

Har

dnes

sH

(GPa

)( D

H,%

)

Mor

phol

ogy

(Dis

pers

ion

,in

tera

ctio

n)

Cry

stal

lin

ity

(%)

(%va

riat

ion

)

E0 DM

A(G

Pa)

(DE0 D

MA

,%)

Y(G

Pa)

( DY,

%)

Met

hod

ofan

alys

isD

epth

vari

atio

n(r

ange

)

Ref

.

33.

78(9

8)0.

234(

200)

1.44

(123

)5

4.11

(115

)0.

250(

221)

–PM

AN

anos

ilic

a5–

15n

mø

sila

niz

edgr

afte

d

UV

curi

ng

01.

850.

057

TEM

(goo

ddi

sper

sion

belo

w35

wt%

)Fr

omu

nlo

adin

g(O

P)

ISE

[270

]19

.63.

64(9

7)0.

262(

360)

35.5

4.41

(138

)0.

360(

532)

48.5

5.39

(191

)0.

492(

763)

PVA

Nan

osil

ica

fun

ctio

nal

ized

80n

mø

Solu

tion

mix

ing

00.

3SE

M(s

ome

aggr

egat

ion

prev

ente

dby

HA

coat

ing)

Gra

dual

decr

ease

wit

hfi

ller

con

ten

t

0.3

0.3

From

un

load

ing

(OP)

Eco

nst

ant

(0.2

–2.0

lm

)[7

8]28

.13.

4(10

00)

3.22

(970

)3.

0(90

0)37

.54.

8(15

00)

4.87

(150

0)5.

0(16

00)

44.0

6.8(

2200

)6.

65(2

100)

6.5(

2100

)54

.68.

1(26

00)

7.20

(230

0)8.

0(26

00)

54.6

+H

A11

.8(3

800)

11.5

(370

0)13

.6(4

400)

PEEK

Surf

ace-

trea

ted

fum

edsi

lica

13n

mø

+25

wt%

CF

Mel

tm

ixin

g40

0�C

0g10

.7a

0.63

1SE

M,A

FM(a

ggre

gate

s)Fr

omu

nlo

adin

g(O

P)

[272

]1

11.2

(5)a

0.65

7(4)

1.5

14.6

(36)

a0.

727(

15)

215

.4(4

4)a

0.76

8(22

)PM

MA

Nan

osil

ica

3n

mø

form

ing

aggr

egat

esSi

lan

izat

ion

ofSi

O2

son

icat

ion

insi

tupo

lym

eriz

atio

nth

erm

alcu

rin

g

03.

840.

243

SEM

(goo

dde

aggl

omer

atio

nby

SC)

1.62

7Fr

omu

nlo

adin

g(O

P)

Rev

erse

ISE

(<10

mN

)[2

73]

33.

88(1

)0.

258(

6)1.

885(

16)

3(S

Cge

l)3.

88(1

)0.

276(

14)

2.38

4(47

)3

(SC

sol)

4.56

(19)

0.32

6(34

)2.

502(

54)

PSSi

lica

nan

odom

ain

s35

–45

nm

øB

lock

copo

lym

eriz

atio

nm

icro

phas

ese

para

tion

calc

inat

ion

05.

1a0.

28TE

M(v

inyl

phen

olfa

vou

rsm

icro

phas

ese

para

tion

)N

AIS

E(<

0.2

lm

)R

ever

seIS

E(>

0.2

lm

)

[274

]17

.27.

2(47

)a0.

40(4

3)20

.85.

5(8)

a0.

34(2

1)PS

Sili

can

anod

omai

ns

35–4

5n

mø

Blo

ckco

poly

mer

izat

ion

mic

roph

ase

sepa

rati

onca

lcin

atio

n

07.

4a,g

0.25

gTE

M(v

inyl

phen

olfa

vou

rsm

icro

phas

ese

para

tion

)N

AIS

E(<

0.2

lm

)R

ever

seIS

E(>

0.2

lm

)

[274

]17

.213

.5(8

2)a,

g0.

34(3

6)g

20.8

10.8

(46)

a,g

0.37

(48)

g

Poly

-sil

oxan

eN

anos

ilic

afu

nct

ion

aliz

ed15

nm

øgr

afte

d

Spin

coat

ing

curi

ng

100

�C0

3.0

TEM

(som

eag

greg

atio

n)

From

un

load

ing

(OP)

[276

]5

vol%

3.6(

20)

10vo

l%4.

3(43

)15

vol%

5.1(

70)

Poly

-sil

oxan

eN

anos

ilic

afu

nct

ion

aliz

ed35

nm

øgr

afte

d

Spin

coat

ing

curi

ng

100

�C0

3.0

TEM

(som

eag

greg

atio

n)

From

un

load

ing

(OP)

[276

]5

vol%

3.6(

20)

10vo

l%4.

0(33

)15

vol%

4.3(

43)

PMM

AA

l 2O

339

nm

øFe

3O

490

nm

øV

orte

xm

ixin

gin

solu

tion

03.

3SE

M(c

lust

ers)

1.8

2.3

Dyn

amic

[278

]1.

5A

l 2O

31.

7(�

48)

0.9(�

50)

2.1(�

9)1.

2Fe

3O

41.

3(�

61)

1.3(�

28)

2.0(�

13)

PSA

l 2O

339

nm

øFe

3O

490

nm

øV

orte

xm

ixin

gin

solu

tion

04.

3SE

M(c

lust

ers)

2.1

2.8

Dyn

amic

[278

]1.

4A

l 2O

30.

2(�

95)

2.0(�

5)2.

7(�

4)1.

1Fe

3O

43.

8(�

12)

1.8(�

14)

2.5(�

11)

PDM

Sc-

Al

fun

ctio

nal

ized

graf

ted

Son

icat

ion

Hyd

rosy

lila

tion

0.21

0.00

43b

SEM

From

un

load

ing

(OP)

[279

]0.

360.

0029

b

0.56

0.00

21b

0.96

0.00

22b

1.11

0.00

24b


Table 5 (continued)


Quasi-static DSIE (GPa)(DE, %)

DynamicDSI E0

(GPa)(DE0 , %)

Hardness H(GPa)(DH, %)



E0DMA (GPa)(DE0DMA , %)

Y (GPa)(DY, %)

Method ofanalysis


Ref.

iPP IF-WS2 80 nm ø Melt-blending/Hot-compression

0 2.70 0.153(0) 0.604 1.5 1.27 Dynamic E0 and Hconstant(P1.5 lm)

[22]0.05 2.41(�11) 0.147(�4) 0.582(�4) – 1.29(2)0.1 2.70(0) 0.152(�1) 0.598(�1) 1.6(7) 1.38(9)0.25 3.00(11) 0.175(14) 0.626(4) – 1.53(21)0.5 2.95(9) 0.172(12) 0.619(2) – 1.57(24)1.0 2.93(9) 0.165(8) 0.634(5) – 1.75(38)2.0 3.10(15) 0.182(19) 0.649(7) – 1.71(35)4.0 3.13(16) 0.175(14) 0.641(6) 1.8(20) 1.81(43)8.0 3.19(18) 0.180(18) 0.657(9) 2.0(33) 1.86(46)

PPS IF-WS2 80 nm ø Melt-blending/Hot-compression

0 3.70 0.242 0.519 1.8 2.18 Dynamic E0 andHconstant(P1.5 lm)

[22]0.05 3.70(0) 0.251(4) 0.473(�9) 2.3(28) 2.15(�1)0.1 4.00(7) 0.285(18) 0.521(0) 2.1(17) 2.35(8)0.5 4.70(26) 0.359(48) 0.596(15) 3.25(81) 2.77(27)1.0 4.50(20) 0.336(39) 0.582(12) 3.20(78) 2.84(30)2.0 4.95(34) 0.382(58) 0.607(17) 2.75(53) 2.58(18)4.0 4.80(29) 0.350(45) 0.603(16) 2.70(50) 2.86(32)8.0 5.00(35) 0.390(61) 0.619(19) 3.25(81) 2.97(36)

PEEK IF-WS2 80–220 nmø

Aerosol-assisteddeposition

0 6.7 0.33 SEM (some aggregation) Possibleincrease dueto filler

NA [282]1 7.5(12) 0.32(�3)2.5 7.2(8) 0.35(6)5 7.9(18) 0.43(30)10 10.1(51) 0.48(45)20 10.7(60) 0.55(67)

Collagen NanoHA 100–150 nm ø

Electrostatic co-spinning

0 0.2 0.026 SEM (aggregations at higherload)

0.0062 Dynamic [286]10 0.5(150) 0.060(130) 0.23(3600)20 0.6(200) 0.100(290) –

Collagencrosslinked

NanoHA 100–150 nm ø

Electrostatic co-spinning

0 0.52 0.068 SEM (aggregations at higherload)

Dynamic [286]10 0.96(85) 0.100(47)20 1.10(110) 0.118(74)

MEH-PPV Cd Se Solution mixingultrasonication spincasting annealing120 �C

0 9.6h 0.23h TEM (better dispersion athigher loads)

Fromunloading(OP)

[287]75 24.3(150)h 0.70(200)h

95 38.4(300)h 0.83(260)h

MEH-PPV Cd Se Solution mixingultrasonication spincasting annealing120 �C

0 11.7i 0.23i TEM (better dispersion athigher loads)

Fromunloading(OP)

ISE (<40–80 nm)Reverse ISE(>80 nm)

[287]75 27.5(140)i 0.43(87)i

95 40.2(240)i 0.79(240)i

a Results showing ‘‘nose’’ effect at the beginning of unloading.b Materials with modulus <0.15 GPa, where it is difficult to determine the initial point of contact.c Measured at 0 �C.d Conical indenter.e Dynamic 10 Hz.f Martens hardness.g Composites on glass.h Constant load 10 lN/s.i Constant load 100 lN/s.

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Table 6Nanomechanical and macroscopic properties for miscellaneous inorganic nanoparticle-reinforced polymer composites. Columns as indicated in Table 1.


Quasi-staticDSI E (GPa),(DE, %)

DynamicDSI E0 (GPa),(DE0 , %)



Crystallinity[%] (%variation)

E0DMA

(GPa)(DE0DMA,%)

Y (GPa)(DY, %)

Method ofanalysis


Ref.

Epoxy IF-WS2 nanotubes60 nm ø

Solution mixing/curing

0 3.47a 0.192 SEM/TEM (good dispersion) Fromunloading(OP)b

[289]3.0 4.44(28)a 0.190(�1)

PA-6 Oib-POSS Melt-blending 0 0.137a,c,d 2.2 0.024 AFM/TEM (large irregularaggregates)

24.4 1.34 Fromunloading(OP) anddynamic

ISE(<0.3 lm)

[292]5.0 0.776(466)a,c 4.0(82) 0.075(212) 29.2(20) 1.77(32)10.0 2.23(1527)a,c 4.8(118) 0.166(590) 35.3(45) 1.81(35)

PA-6 Tsp-POSS Melt-blending 0 0.137a,c,d 2.2 0.024 AFM/TEM (small regularlydispersed POSS domains)

24.4 1.34 Fromunloading(OP) anddynamic

ISE(<0.3 lm)

[292]5.0 0.987(620)a 4.3(95) 0.066(175) 27.4(12) 1.91(42) –10.0 2.380(1637)a 5.0(127) 0.157(554) 29.4(21) 1.96(46)

PVDF FP-POSS Solution casting 0 2.61e 0.174 OM (POSS aggregates) gradualincrease withfiller content

1.80 1.71 Fromunloading(OP).

E,Hconstant(0.5–2 lm)

[293]3.0 2.77(6)e 0.177(2) 2.28(27) 2.00(17)5.0 2.33(�11)e 0.158(�9) 2.13(18) 1.59(�7)8.0 2.21(�15)e 0.130(�25) 1.86(3) 1.17(�32)

PU GPTS-modifiedAlOOH nanorodf 10–20 nm ø

Sol–gel/Spincoating/thermalcuring

0 8.98 – SEM (homogeneousdispersion)

NA [290]40.0 8.86(�1) 0.83

PU GPTS-modifiedAlOOH nanorodf 10–20 nm ø

Solution mixing/casting/UV curing

0 8.98 SEM (homogeneousdispersion)

NA [290]5.0 9.22(3)

PU GPTS-modifiedAlOOH nanoparticle

Sol–gel/Spincoating/thermalcuring

0 8.98 – SEM (isotropic dispersion) NA [290]40.0 9.88(10) 0.98

PU GPTS-modifiedAlOOH nanoparticle

Solution mixing/casting/UV curing

0 8.98 SEM (isotropic dispersion) NA [290]5.0 9.43(5)

PU Silica networkgeneratedin situfromTEOSg

Sol–gel/UV curing 0 2.75 0.13 OM (homogenousdispersion)

Fromunloading(OP)

[298]20 3.28(19) 0.16(23)50 3.73(36) 0.18(38)

PMMA Eu:Gd2O3

nanoparticles 40 nm øSolution casting 0 4.28a 0.263 SEM (agglomerates) 2.15 From

unloading(OP)

[300]3.0 4.82(13)a 0.268(1) 2.60(21)

PMMA APTES-modifiedEu:Gd2O3

Solution casting 0 4.28a 0.263 SEM (good dispersion) 2.15 Fromunloading(OP)

[300]3.0 5.60(31)a 0.293(11) 3.05(42)

PMMA Trioctylamine-functionalizedgraphene-like BN

Solution mixing/annealing

0 2.20 0.137 Fromunloading(OP)

[288]3.0h 4.81(118) 0.285(108)3.0i 5.11(132) 0.310(126)

PMMA-co-MPTES

Silica networkgenerated in situ fromTEOSg

In situpolymerization/Sol–gel/Spin coating/curing

0 4.10 0.25 SEM (transparent, no phaseseparation)

Fromunloading(OP)

ReverseISE

[296]22 7.60(85) 0.50(100)47 6.60(61) 0.54(116)72 9.50(132) 0.85(240)

36A

.M.D

íez-Pascualet

al./Progressin

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Science67

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94

Table 6 (continued)


Quasi-staticDSI E (GPa),(DE, %)

DynamicDSI E0 (GPa),(DE0 , %)



Crystallinity[%] (%variation)

E0DMA

(GPa)(DE0DMA,%)

Y (GPa)(DY, %)

Method ofanalysis


Ref.

PMMA Silica networkgenerated in situ fromTEOS

In situpolymerization Sol–gel/Spin coating/curing

0 6.90 0.32 SEM (transparent, no phaseseparation)

Fromunloading(OP)

ReverseISE

[297]17 6.30(�9) 0.45(41)43 8.20(19) 0.56(75)

PDDAj CSH nanoparticles Stirring/precipitation

– 4.44(�83)k 0.11(�80)k TEM (high packing density) Fromunloading(OP)

E,Hconstant(1–512mN)

[291]100 26.8 1.09

PI POSS-OH Solution casting/thermal imidization

0 2.9 0.22 SEM (homogenousdispersion. Aggregates at10 wt%; strong H-bonding)

NA [294]3.0 3.0(3) 0.22(0)5.0 3.2(10) 0.23(5)10.0 3.5(17) 0.24(9)

PI OC-POSS Solution casting 0 4.80 0.48 3.30 Fromunloading(OP)

[295]7.8l 4.70(�2) 0.48(0) 3.10(�6)15.4l 4.70(�2) 0.44(�8) 3.00(�10)30.2l 4.80(0) 0.40(�17) 3.30(0)44.5l 5.20(8) 0.40(�17) 3.00(�10)

PI OH-POSS Solution casting 0 4.80 0.48 3.30 Fromunloading(OP)

[295]7.0l 4.75(�1) 0.45(�6) 3.40(3)13.9l 4.60(�4) 0.40(�16) 3.70(12)27.9l 4.65(�3) 0.38(�21) 3.20(�3)41.8l 4.20(�12) 0.34(�29) 3.00(�10)

a Reduced modulus.b Spherical indenter.c Results showing ‘‘nose’’ effect at the beginning of unloading due to short holding time.d Materials with matrix modulus <0.15 GPa, where it is difficult to determine the initial point of contact.e Results affected by thermal drift correction.f Aligned along the surface direction.g Covalently grafted to the matrix.h DMF solvent.i Chloroform solvent.j Intercalated between the nanoparticle lamellae.

k Percentage of variation compared to the pure nanoparticles (cement).l Calculated theoretically.

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.Díez-Pascual

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37


composed of more than two coaxial cylinders, each rolled out of single sheets, with diameters between2 and 40 nm, the multi-walled carbon nanotubes (MWCNTs). Both types of CNTs can be synthesizedusing different routes [86], including high-temperature evaporation using arc-discharge or laser abla-tion, and gas-phase catalytic growth from carbon monoxide or chemical vapour deposition (CVD) fromhydrocarbons. CVD materials generally contain residual catalyst particles, while the main contami-nants in the products of the high-temperature reactions are carbonaceous impurities. The laser pro-cess generates CNTs of the highest quality (few defects, high crystallinity, extremely high aspectratio) compared to the other synthesis methods [87], while nanotubes produced by the CVD techniquepresent a larger number of defects as a result of the lower growth temperature. The purity, quality,aspect ratio and nature of impurities, hence the source of the CNTs, can have strong influence onthe final properties of CNT-reinforced composites.

In order to fabricate polymer/CNT composites with enhanced mechanical performance, two mainchallenges need to be addressed. Firstly, it is very difficult to attain homogenous CNT dispersion sincethese nanofillers have a strong tendency to gather and form bundles owing to strong van der Waalsinteractions. Moreover, their large aspect ratio causes extremely high viscosity in the polymer-CNTblend, which in turn affects their uniform distribution. Secondly, the inert nature of the CNT surfaceresults in poor interfacial adhesion with the host matrix, which prevents efficient stress transfer dur-ing loading [88]. To solve these issues, several methods have been reported such as mechanical disper-sion (i.e. ultrasonication [89], ball-milling [90]), plasma treatment [91] and chemical modificationinvolving either non-covalent or covalent bonding between nanotubes and polymers [92]. The non-covalent approach consists in the physical adsorption and/or wrapping of polymers to the CNT surface.The graphitic CNT sidewalls can interact with conjugated polymers via p � p stacking as well as withthose containing heteroatoms with free electron pairs. This route preserves the nanotube integrity andproperties, since it does not destroy the conjugated system of the CNT sidewalls. The covalent methodinvolves the chemical bonding (grafting) of polymer chains to CNTs, and can be performed via ‘‘graft-ing to’’ or ‘‘grafting from’’ strategies. The first approach is based on the synthesis of a polymer deriv-ative that can react with the surface of pristine, oxidized or functionalized CNTs [93]. However, theoxidation/functionalization treatments performed in acid media generally shorten the CNTs and canbring significant damage such as sidewall opening or tube breakage [94], introducing defects in thetubular framework that can adversely impact their mechanical properties. Another disadvantage ofthis method is that the grafted polymer content is restricted due to the low reactivity and high sterichindrance of the polymer chains. In the ‘‘grafting from’’ strategy the polymer is grown from the CNTsurface via in situ polymerization of monomers initiated by chemical species immobilized on the CNTsidewalls and edges. The high reactivity of monomers makes the grafting from process efficient andcontrollable, enabling the preparation of nanocomposites with high degree of grafting. However, thismethod requires strict control of the amounts of each reactant and the polymerization conditions.Alternatively, CNT functionalization can be performed by plasma treatment [91,95,96], a time savingand environmentally friendly technique for modifying CNTs by directly introducing a high density offunctional groups. The CNT surfaces can readily change from hydrophobic to hydrophilic with plasmaprocessing, thus facilitating the dispersion within the polymer matrix. Furthermore, this treatmentdoes not alter the intrinsic mechanical properties of the tubes.

4.1.2. Preparation of polymer/CNT nanocompositesSeveral methods have been reported for the preparation of polymer/CNT nanocomposites, includ-

ing solution mixing [28,97–103], milling [104,105], melt-compounding [20] and in situ polymerization[106,107]. In the solution processing route, the CNT powder is typically dispersed in a liquid mediumby stirring and/or sonication and then mixed with the polymer solution, followed by evapouration ofthe solvent. Ball milling is a method employed to reduce particle size and, in turn, alter surface prop-erties of materials by modifying surface area, porosity and shape. This technique can be used to pre-pare a fine polymer/CNT power that is subsequently melt-mixed with a twin-roll or extruder. Melt-processing is the most widespread technique for the fabrication of CNT-reinforced thermoplastic com-posites, and is also suitable for polymers with low solubility in common solvents. It consists in theblending of the polymer melt with the CNT material by application of intense shear forces, and is oftencombined with a shape-forming step like injection moulding, spinning or hot-compression. The in situ


polymerization of monomers in the presence of CNTs has been extensively explored for the prepara-tion of polymer–grafted nanotube composites. This technique is particularly important for insolubleand thermally unstable matrices, and produces homogenous composites with high CNT weightfraction.

4.1.3. Modulus changes in polymer/CNT nanocomposites: comparison of thermoplastic and thermosetmatrices

Table 1 gathers nanoindentation results collected from the literature on polymer/CNT nanocom-posites [20,24,91,95–106,108–127]. The percentage of change in the elastic modulus DEI (eitherquasi-static E or dynamic E0) of thermoplastic and thermoset matrices as a function of the nanotubeloading is represented in Fig. 6a and b, respectively. These figures show that in general, CNTs have

0

100

200

300

1000

1200 Chitosan/MWCNT [98] PHB/SWCNT [97] PP/MWCNT [106] PA-6/MWCNT [20] PA-6/MWCNT [100] PLLA/MWCNT [101] UHMWPE/SWCNT [96] UHMWPE/MWCNT [105]UHMWPE/MWCNT [104] PVDF/MWCNT [120] PVA/MWCNT [103] PVA/SWCNT [121]PHO/SWCNT [97] PVK/MWCNT [103] PMMA/MWCNT [28] PMMA/MWCNT [28] PMMA/MWCNT [127] PMMA/MWCNT [127] PI/SWCNT [126] PDMS/MWCNT [124]

CNT [wt%]

Δ EΙ [

%]

0 2 4 6 8 10

0 1 2 3 4 5 6

0

20

40

60

80

100

120

CNT [wt%]

Δ EΙ

[%]

[115] [91] [91] [57] [24] [109] [110] [111] [111] [117] [117] [112] [113] [113] [108]

(a)

(b)

Fig. 6. Percentage of change in the elastic modulus (DEI) of carbon nanotube-reinforced polymer nanocomposites as a functionof the CNT loading: (a) thermoplastic and (b) thermoset matrices. Solid and open symbols correspond to raw and functionalizedCNTs, respectively; s and h in Fig. 6a represent TEOS and APTES functionalized MWCNTs, respectively; ⁄ indicates CNTs graftedto the polymer matrix, + and j denote aligned MWCNTs and � designates silica-coated CNTs. Symbols in red highlight the datathat should be taken with caution (see Table 1 for details).


a positive effect on EI, reinforcing the polymer matrix already at very low loadings. The maximum CNTcontent incorporated is 10 wt% due to the nanotube tendency to agglomerate into bundles. It is note-worthy that the maximum filler concentration is significantly lower for the thermoset matrices(�5 wt%) compared to the thermoplastic ones (�10 wt%). As will be seen below, this is a direct con-sequence of the commonly encountered difficulties in producing epoxy-based composites with lowagglomeration and high degree of polymer bonding. In the case of thermoplastic matrices, significantDEI are achieved at high loadings, reaching 200% increase for filler loadings of 8 wt% (see Fig. 6a), whilemuch more limited EI improvements are found for the thermoset matrices.

4.1.3.1. Thermoplastic matrices. A number of papers dealing with thermoplastic-based composites[20,96–98,100,101,103,104,121] have reported a systematic increase of EIwith increasing CNT contentup to around 1–2 wt% loading. Higher filler loadings are frequently not considered because agglomer-ation of the nanotubes into bundles starts to play a significant role and mechanical enhancementslevel-off or even decrease. Fig. 6a shows that for low nanotube contents (<2 wt%), the stiffness incre-ments for composites incorporating raw [106,120] or functionalized CNTs [20,96,98,100,101,104] liewithin the same range; the diversity of methodologies and indentation testing conditions employedin the analysis of DSI values can induce some data scattering that hides any significant difference. Sim-ilarly, it is not clear that composites incorporating SWCNTs exhibit larger modulus improvementsthan those using MWCNTs in spite of the superior modulus of the former [128]. The most noticeableEI increments for low CNT contents (see Fig. 6a) have been explained as summarized below[96,98,103,106,121,126].

Satyanarayana et al. [126] found a noticeable increase in the dynamic modulus of glassy polyimide(PI) (�59%) upon addition of only 0.05 wt% raw arc-discharge SWCNTs albeit the existence of smallagglomerates was inferred from AFM analysis. The strong nanotube-matrix interfacial bondingattained via p � p stacking interactions between the aromatic rings of PI and the sp2 hexagonal net-works of the nanotubes was discussed by the authors of this work to be responsible for the notewor-thy improvement at such low loadings. Moreover, a partial alignment of the CNTs that could have beenattained during the spin-coating process is also suggested to influence the measured data.

Interestingly, Cadek et al. [103] found an 80% increment in the modulus of polyvinyl alcohol (PVA)with the incorporation of 1.0 wt% raw arc-discharge MWCNTs, value above those reported for similaramounts of acid-functionalized MWCNTs embedded in polyamide 6 (PA-6) [20,100] or poly(L-Lactide)(PLLA) [101] (Fig. 6a). Such significant improvement is ascribed to the drastic increase in the crystal-linity of PVA (�93%) upon addition of the MWCNTs. The benefits of such outstanding crystallinityincrease are twofold: on the one hand, the mechanical properties of the matrix are enhanced; onthe other hand, the nucleation of PVA crystals resulted in an extremely strong CNT-matrix interfacialinteraction, as demonstrated by TEM observations, which showed that the matrix did not fail at thepolymer–nanotube interface. Surprisingly, Prasad et al. [121] reported a DE of 1180% for PVA uponaddition of only 0.6 wt% acid-functionalized SWCNTs, attributed to the strong bonding between thefunctionalized CNTs and the matrix combined with the very high SWCNT aspect ratio. Nevertheless,this result could be inaccurate since the matrix modulus is quite low (E = 0.7 GPa) and hence proneto a difficult determination of the point of initial contact, as discussed in Section 3. In addition, the roleof water content on the indentation property improvement of this type of hydrophilic matrices has notbeen conveniently explored and could certainly shed some light on the mechanisms for such extraor-dinary mechanical enhancement.

A very efficient non-covalent functionalization approach was reported by Wu et al. [98] for chito-san based nanocomposites. MWCNTs were wrapped by poly(3,4-ethylenedioxythiophene)-poly(sty-rene sulphonate) (PEDOT-PSS), a surfactant that improved the nanofiller-matrix compatibility. Thesulphonic acid groups of PSS enabled ionic linkage with positively charged amine groups of chitosan,while PEDOT was bound to the MWCNT surface through p � p and hydrophobic interactions. As aresult, a 52% improvement in modulus was attained at 0.5 wt% loading.

On the other hand, Samad and Sinha [96] reported a D E0 of 58% for ultra-high molecular weightpolyethylene (UHMWPE) coatings upon addition of only 0.2 wt% plasma-treated SWCNTs (seeFig. 6a). The functionalization via plasma treatment resulted in a large number of oxygen containing


groups on the nanotube surface that are envisaged to promote bundle disentanglement and improvethe filler interaction with the surrounding matrix, leading to a very uniform dispersion.

Focusing on nanotube concentrations P2 wt%, semicrystalline matrices incorporating acid-func-tionalized MWCNTs [100,101,104] are found to exhibit enhanced stiffness compared to those rein-forced with pristine CNTs (Fig. 6a). As mentioned earlier, the oxidation in acid medium leads to ashortening of the tubes and frequently induces some damage on their sidewalls, drawbacks that seemto be overweighed by the more homogenous nanotube distribution and stronger CNT-matrix interfa-cial adhesion attained through this functionalization process, which is reflected in larger modulusincrements. As an example, Baji et al. [100] reported remarkable E enhancements (�180%) for an arrayof randomly electrospun PA-6 fibres containing 7.5 wt% of acid-functionalized MWCNTs aligned alongthe fibre axis. The confinement effect of the nanotubes on the mobility of PA-6 chains, together withthe increase in crystallinity and the strong interactions between the carboxylic groups of the MWCNTsand the amide moieties of the matrix, were suggested to give rise to the noticeable stiffening effect.However, the work does not address the influence of orientation effects on the mechanical properties.Sreekanth and Kanagaraj [104] described one of the largest DE values in thermoplastic matrices. About206% increase in the modulus was reported for UHMWPE reinforced with 5.0 wt% oxidized MWCNTs(Fig. 6a). A very homogenous dispersion, combined with a large increase in crystallinity of almost 40%,are believed to be the reasons for such strong mechanical enhancement. These increments however,should be contemplated with care because they can be affected by the low modulus of the neat matrix,in addition to the inadequate correction of the load-depth data for thermal drift (see Section 2.1.1 for adetailed discussion). Contrary to the aforementioned results, very small E increase (�10%) was foundupon addition of 5.0 wt% raw MWCNTs to UHMWPE [105], ascribed to the poor nanofiller-matrixinterface, the presence of voids and the nanotube waviness that limits the efficiency of thereinforcement.

Other functionalization approaches like amidation reactions have been employed to improve thedispersion of large amounts of MWCNTs (>2 wt%) within thermoplastic polymers such as amorphouspoly(methyl methacrylate) (PMMA) [28,102]. However, SEM images revealed the presences of someaggregates, and hardly increase in modulus was attained (see Fig. 6a). It was suggested that the flex-ibility of CNTs, their curved morphology and inferior bending properties may contribute to thereduced reinforcement action. In order to overcome these drawbacks, MWCNTs were coated with asilica shell encasing the nanotubes. The strategy was found to be most successful yielding DE valuesof 130% for 4.0 wt% loading (see Fig. 6a). Mammeri et al. [127] described two novel strategies to func-tionalize MWCNTs in order to incorporate them in poly(methacrylic acid-co-methacryloxypropyltri-methoxysilane) P(MMA-co-MPTMS): the covalent grafting of hydrolyzable triethoxysilyl (SiOEt)3

groups on oxidized MWCNTs and the non-covalent adsorption of a polycation, polyallylamine hydro-chloride (PAH), on pristine CNTs. While the first approach resulted in very homogenous compositeswith moderate improvement in the matrix modulus (�26% at 2.0 wt% loading), the second procedurewas insufficient to disaggregate the nanotubes, and no increase in modulus was attained.

4.1.3.2. Thermoset matrices. With regard to nanocomposites based on thermoset matrices, as men-tioned above, DEI values are in general lower than those found for the thermoplastic counterparts,by up to �100% in the concentration range of 0.1–5.5 wt% (Fig. 6b) [57,91,95,108–113,115,118], andeven some decrements in stiffness are reported [24,114,117]. This discrepancy is due to the poornanofiller dispersion in epoxy composites (Table 1), which could be related to diffusion and re-aggre-gation of CNTs during the curing process [129]. Furthermore, contrary to the tendency observed forthermoplastic/CNT samples, no significant stiffness difference is found for composites reinforced withpristine or functionalized CNTs in all the range of concentrations tested. Results hint at an importantmodulus increment for plasma functionalized MWCNTs grafted to the epoxy matrix yielding aDE = 20% for a loading of only 0.1 wt% [91]. The plasma method enabled incorporating a high densityof primary amine groups onto the nanotube surface that allowed direct covalent bond to the epoxidegroups, thus resulting in improved filler-matrix adhesion and nanotube dispersion, as revealed by SEMmicrographs (Fig. 7). Surprisingly, similar moderate EI increments, with a tendency to increase athigher filler loadings, have been reported for composites incorporating raw MWCNTs [108,112],despite their poor dispersion.

Fig. 7. (a) Representative DSI load–displacement curves on neat epoxy and 0.1 wt% unfunctionalized and plasma functionalizedMWCNT-reinforced nanocomposites. (b) and (c) SEM images of the fractured surfaces of the nanocomposites withunfunctionalized and functionalized CNTs, respectively, showing the difference in the state of nanotube dispersion. Adaptedfrom Ref. [91], copyright 2013, with permission from Elsevier.



Clear examples on the insignificant moduli differences between epoxy composites filled with eitherraw or functionalized CNTs were provided by Lagoudas et al. [113] and Mionic et al. [111]. Interest-ingly, Mionic demonstrated that the improvements in the dynamic modulus were dependent on thesolvent used for the CNT pre-dispersion and not on the functionalization approach, indicating the lackof formation of chemical bonds between the constituents. The largest stiffness increase (about 104% at5.0 wt% acid-treated MWCNTs) was found for acetone, a highly volatile solvent that would evaporaterapidly enabling a fast polymerization of the matrix. On the other hand, the smallest increment (onaverage 20% for the same concentration) was observed for gamma-butyrolactone (GBL), which copo-lymerizes with the epoxy and acts as a plasticizer, weakening the mechanical properties of the mate-rial. Li et al. [57] also found relevant DEI of 75% with the incorporation of 5.0 wt% methylbenzoatefunctionalized arc-discharge SWCNTs (see Fig. 6b). The functionalized tubes were easily dissolvedin polar solvents and integrated in the matrix via solution processing, resulting in relatively good dis-persion for all the samples. Intercalation of the polymer into the bundles was proposed as the key rein-forcing mechanism for these nanocomposites.

In contrast, only small dynamic modulus improvements (around 17% at 5.0 wt% loading) werereported for epoxy nanocomposites incorporating purified laser-grown SWCNTs [24,114] (Fig. 6b),regardless of the high quality and aspect ratio of these nanofillers, attributed to their curved morphol-ogy, the presence of a large number of aggregates, the sliding between individual nanotubes within abundle and the composite microvoids. Even more striking is the limited DE found when CNT ‘‘forests’’are incorporated to the epoxy matrix. CNT ‘‘forests’’ are dense, millimeter-tall and vertically alignedstructures grown by water-assisted CVD process and incorporated into the epoxy matrices [115–117]. These nanocomposites are expected to exhibit enhanced mechanical performance when the loadis applied along the vertical direction of the aligned CNTs. However, only modest EI improvementshave been reported [115,117] and, also surprisingly, no significant differences were observed com-pared to samples filled with the random counterparts [117]. CNT waviness was suggested to be themajor factor accounting for the deviations of the modulus values of the composites from the rule ofmixtures (Eq. (14)). The authors conclude that a fine control of this parameter would be highly desir-able to reach a significant stiffness contribution from the CNTs.

Finally, it is worth mentioning that coiled carbon nanotubes (CCNTs) represent an interesting alter-native to traditional straight nanotubes for reinforcement of polymer matrices [130]. They exhibit aYoung’s modulus of �0.7 TPa, and their coiled configuration enables an improved bonding to the sur-rounding matrix, resulting in an efficient reinforcement. Li et al. [118] investigated a series of CCNT-reinforced epoxy composites with different weight percentage of nanotubes by means of DSI. How-ever, no conclusion can be drawn from this work because modulus values derived from the onset ofunloading present the well-known ‘‘nose effect’’, a fact that critically compromises the validity ofthe stiffness data (see Section 2.1.1).

4.1.3.3. Other CNT-reinforced polymer systems. Recently, the layer-by-layer (LbL) assembly of oppo-sitely charged polyelectrolytes on a charged surface has emerged as an alternative method for inte-grating functionalized CNTs into polymers [131–135]. It is an environmentally sound and cost-effective approach that can result in nanocomposites with high nanofiller weight fraction and con-trolled internal structure, having the potential to reach the desired mechanical properties through tai-lored design, and with eventual applications in the development of a new generation of scaffolds. Forinstance, in poly(dimethyldiallylammonium chloride) (PDDA)/poly(acrylic acid) (PAA) multilayersprepared via LbL, the addition of 4.7 wt% acid-treated SWCNTs to PAA led to 120% increment in thein-plane modulus of the polyelectrolyte film [131,132], attributed to a homogenous nanotube distri-bution combined with their alignment parallel to the film plane. Contrary to the aforementionedresults, Firkowska et al. [133] and Pavoor et al. [135] reported that the presence of acid-treatedMWCNTs in polyethyleneimine (PEI) or PAH films was detrimental to the mechanical performanceof LbL multilayers of polystyrene sulphonate (PSS)/PEI or PAH/PAA, respectively. The striking factabout these results is that the modulus values were significantly lower than those of the LbL filmsmade solely of polyelectrolytes. This effect was explained in terms of the lubricant role of the CNTs,easily rolling or sliding between the different layers of the assembly. It is noteworthy that CNT


bending, waviness, curved morphology or sliding ability have been posed to explain the limited rein-forcing effect of the nanotubes in epoxy matrices [24,114,117].

On the other hand, about a 40-fold enhancement in modulus has been reported for isocyanatemicrocapsules, with poly(urea–formaldehyde) (PUF) as a shell, upon incorporation of only 1.1 wt%oxygen plasma functionalized MWCNTs [107]. These novel materials were prepared by in situ poly-merization, which enabled a very uniform dispersion of the nanotubes within the microcapsule shell.In addition, the authors argued that oxygen groups from the nanotube surface could strongly interactwith the amide groups of PUF, resulting in extraordinary improved mechanical properties. The inden-tation data analysis is, however, a matter of concern because modulus values are too low (54 MPa) toallow for a precise determination of the point of initial contact.

CNTs have also been incorporated into conventional fibre-reinforced polymers to fabricate multi-scale (also named hierarchical) composites with improved mechanical performance [136]. Two alter-native strategies for the preparation of CNT-based hierarchical composites have been reported: thedispersion of CNTs into the matrix [122] or their direct anchoring onto the fibre surface [119]. Nano-indentation in these systems is usually directed to produce a map of local mechanical properties aim-ing at distinguishing between the polymer matrix, the fibres and the interface/interphase. Forexample, Qian et al. [119] combined nanoindentation studies with AFM and Raman spectroscopy toprobe the local mechanical properties of epoxy composites incorporating MWCNTs grafted onto silicafibres via CVD process. The authors demonstrated that the grafting of MWCNTs to the fibres increasedthe modulus of the epoxy matrix in the fibre surroundings. More recently, the local mechanical prop-erties of SWCNT-reinforced PEEK/glass fibre (GF) laminates have been investigated using nanoinden-tation [122]. About 9% improvement in the PEEK matrix modulus was attained upon addition of1.0 wt% purified arc-SWCNTs, increment that was doubled with the incorporation of the same amountof SWCNTs wrapped in poly(ether ether sulphone) (PEES). More important, the paper demonstratesthat the polymer–fibre interface can be easily detected using the dynamic technique, and providesexperimental evidence that the extent of the interphase increases if the matrix is reinforced withSWCNTs, particularly those wrapped in the compatibilizing agent.

4.1.3.4. Comparison with macroscopic properties. The data of Table 1 show that the increments in inden-tation quasi-static modulus (calculated from OP analysis) increase with increasing filler content fol-lowing a similar trend than those of tensile or flexural moduli [91,100,104,110,112], although thecomparison of the increment values exhibit some disparity [104,110,112]. Y increments have beenfound to be higher [105,110], lower [104,112] or in the range [100] of those for E. The differencesbetween these numbers sometimes approach one order of magnitude, in one or the other direction[104,105]. The larger DE values obtained by means of indentation have been occasionally attributedto an indentation size effect due to agglomeration of CNT bundles at the surface [112]. Most surprisingis the fact that modulus increments for indentation and tensile testing were quite close only in thecase of electrospun mats [100]. In this work, molecular chains were aligned along the loading directionfor tensile testing and perpendicularly for indentation.

Although it seems that comparison between E and E0DMA does not have a reasonable ground, some-times it yields an excellent match. For example, Chen et al. highlighted that the large E increase foundupon addition of functionalized MWCNT to an epoxy resin was in agreement with DMA results (seeTable 1) [91]. In contrast, Wu et al. reported substantial differences between the modulus incrementsobtained by dynamic DSI and tensile testing (see Table 1) [98]. The authors of this work highlightedthe consistency of both techniques in revealing a rising trend of modulus with increasing MWCNTswithout further discussion on the relative increments. This limited analysis seems quite judicious inview of the large differences between both mechanical indentation tests.

Table 1 only offers one example in which the E0 and E0DMA can be contrasted [20]. It is found that themodulus increments exhibit the same trend but different values. This discrepancy was discussed by theauthors in terms of the different frequencies employed in DMA and DSI (1 Hz vs. 45 Hz, respectively) andthe different loading directions with respect to the plane of the compression moulding films.

In summary, both indentation analysis and traditional macroscopic mechanical techniques revealsimilar rising trends of the modulus with increasing CNT content. However, the increment values donot exhibit a clear correspondence.


4.1.4. HardnessA limited amount of H data, compared to EI, is available for CNT/polymer nanocomposites. This is

especially apparent for CNT-reinforced epoxies (see Table 1). The table shows that the H behaviourwith filler content is qualitatively similar to that discussed previously for the modulus. In the follow-ing sections, it will be found that this is usually also the case for other nanofiller-reinforced polymersystems. Hence, in order to avoid unnecessary repetitions, the main discussion will be focused on themodulus data and only some hints of the differences between the elastic and the plastic responses aregiven below and also remarked in other sections when appropriate.

As an example, Fig. 8 illustrates the percentage of change in indentation hardness of thermoplasticmatrices as a function of the nanotube loading. The figure shows that the H increments are in generallower than those of EI (see Fig. 6a), suggesting that CNTs are more effective in enhancing the stiffnessthan the hardness of thermoplastic matrices. Mammeri et al. [127] observed that the elastic modulusis more sensitive than the hardness to the type of surface modification of the CNTs. While similar Henhancements were recorded by incorporating 2.0 wt% of MWCNT–PAH–SiO2 or MWCNT–COOH–APTES to P(MMA-co-MPTMS), a higher elastic modulus increment was determined when using thesecond type of nanofiller, which exhibited improved dispersion. Results suggest that the interfacesand the degree of dispersion of CNTs influence in a different manner the elastic and plastic deforma-tion processes. These findings are in agreement with the general trend found for larger EI enhance-ments compared to those of H. Fig. 8a also offers some examples of hardness enhancementsmodulated by changes in the degree of crystallinity and crystal size [97,100,103,104], analogous tothose discussed in Section 4.1.3.1 for the modulus. In short, CNTs act as nucleating agents at low load-ings generally enhancing the matrix crystallinity and decreasing the crystal size. Such structuralchanges contribute to the H variations with filler content in terms of the models described in Eqs.(12) and (13). At large CNT contents, agglomeration has a dual detrimental effect: on the one hand,the CNT-matrix interfacial area decreases; on the other hand, the reduced efficiency of the CNTs asnucleating agents limits the amount of crystallinity in the polymer matrix.

4.1.5. Indentation size effectA wide number of studies on CNT/polymer composites have reported a dependence of EI or H on

indentation depth [20,28,98,99,108,111–113,125–127]. The difficulty in discerning whether an ISEis attributed to genuine material properties or is it due to instrumental effects was pointed out in Sec-tion 2.1.3. This is especially critical at penetrations below a few hundred nanometers.

0 2 4 6 8 10

0

50

100

150

650 Chitosan/MWCNT [98] Chitosan/MWCNT [99] PHB/SWCNT [97] PP/MWCNT [106] PA-6/MWCNT [20] PA-6/MWCNT [100] UHMWPE/SWCNT [96] UHMWPE/MWCNT [105]UHMWPE/MWCNT [104] PVA/MWCNT [103] PVA/SWCNT [121]PHO/SWCNT [97] PVK/MWCNT [103] PMMA/MWCNT [28] PMMA/MWCNT [28] PMMA/MWCNT [127] PMMA/MWCNT [127] PI/SWCNT [126] PDMS/MWCNT [124]

CNT [wt%]

ΔH [

%]

Fig. 8. Percentage of change in the indentation hardness (DH) of carbon nanotube-reinforced thermoplastic polymernanocomposites as a function of the CNT loading. Symbols as indicated in Fig. 6.


As an example, Fig. 9a illustrates the variation of hardness as a function of indentation depth for achitosan sample and a MWCNT-reinforced one [99]. A conspicuous increase of H below penetrationdepths of �200 nm is clearly seen for both materials. Analogous EI and H enhancements have alsobeing observed in other CNT-reinforced chitosan composites [98], polyamide/CNT materials [20]and in the 100/0 PI/SWCNT composite [126]. All this research makes use of the dynamic method inthe determination of E0 and H as a function of indentation depth [20,98,99,126]. A recent paper sug-gests that such a method may produce an apparent reverse ISE at shallow depths [5]; a hint of sucheffect is given in Ref. [126]. Depth dependences at small penetrations have also been reported forepoxy/CNT systems although significantly lower E or H enhancements were found [112,113]. Inden-tation size effects at shallow depths are most commonly attributed to inaccuracies in the determina-tion of the point of contact and of the tip area function.

In contrast to the above studies, Olek et al. [28] reported constant E values in the range 75–500 nm(in contact depth) for PMMA and CNT-reinforced PMMA composites. It should be noticed that theauthors used a soft polymer material for calibration of the tip area function. Indeed, this could beat the heart of the E tendency to remain constant at such small penetration depths. In contrast, the

Fig. 9. (a) Variation of H with displacement into surface for neat chitosan (solid symbols) and chitosan/MWCNT (0.4 wt%)nanocomposite (open symbols). Adapted from Ref. [99], copyright 2005, with permission from the American Chemical Society.(b) Dynamic modulus E0 as a function of indentation depth for epoxy composites reinforced with nonfunctionalized (blue) andfunctionalized (red) MWCNTs prepared in acetone as solvent. Adapted from Ref. [111], copyright 2010, with permission fromJohn Wiley & Sons, Inc.


H values were found to diminish in the same depth interval. This was explained in terms of surfaceroughness, time-dependent plastic deformation or other nanostructural features in the spin-coatedfilms. However, it is not clear why these phenomena would influence the depth dependence of Hand not that of E.

Most interesting are the studies on CNT-reinforced polymer composites reporting a variation of EI

and H as a function of indenter displacement that approaches the micrometer scale [108,111,126]. Inthis range, the sources of error influence the indentation data to a lesser extent and the ISE is mostlikely associated to the genuine properties of the material. As an example, Fig. 9b shows the dynamicmodulus E0 of epoxy-based samples with pristine and functionalized MWCNTs as a function of theindenter displacement. The continuous decrease of modulus values above h � 200 nm has been dis-cussed as a consequence of a morphology gradient across the sample thickness arising from solventmolecules trapped in the reinforced matrix [111]. In other CNT-reinforced polymer composites,because appropriate information on the experimental procedure and data analysis are not given, itis difficult to evaluate whether the ISE is associated to a gradient of material properties [108,126].For example, it seems that Dos Santos et al. have used a trapezoidal load function for the loading/unloading cycles [108]. As indicated in Section 2.1.3, because the strain rate is continuously changingas the test proceeds, this will normally result in an ISE. Other authors have found that EI and H valuesremain constant in a wide range of indentation depths [57,114] and have interpreted this result asindicative of the good level of CNT dispersion throughout the matrix [114].

Fig. 9a also shows that an H increase is found with increasing displacement for h > 500 nm [99].This behaviour represents a clear example of reverse ISE, attributed in this case to substrate effects.The film properties are expected to be represented by the values shown at h � 500 nm-1000 nm, thatis, in the small interval between the normal ISE and the reverse one. A detailed analysis of the reverseISE due to substrate effects is given by Mammeri et al. [127] on the basis of indentation studies onCNT-PMMA based hybrid coatings. It was clearly shown that the ISE influences the elastic modulusvalues to a larger extent than the hardness ones because the elastic field of deformation encompassesthat associated to plasticity and hence, the former would ‘‘feel’’ the substrate at an earlier stage (seeSection 2.1.3.1).

4.1.6. CreepThe incorporation of CNTs into polymer matrices is found to have a beneficial effect on the creep

behaviour [24,109]. The time-dependent deformation of epoxy/SWCNT nanocomposites with differentfiller content has been analysed in terms of the power law relationship between the strain rate and thehardness through the creep exponent n (see Eq. (5)) [24]. Experimental data during the hold period atpeak load were employed. The double logarithmic plot of H and _e was found to exhibit a linear trend inalmost all the range of strain rates (Fig. 10). However, two observations are worth commenting. Thefirst one is that, in contrast to the findings of Lucas and Oliver [42], no transient period at the begin-ning of the holding time is observed. The second observation is that the power law fit does not hold forlow H and _e values. The n values are presumably obtained in the linear part of the log H–log _e plot.Eventually, the work reports increasing values of n(reduced creep deformation) with increasingamount of CNT incorporated into the epoxy matrix.

Tehrani et al. [109] analysed the creep behaviour of epoxy/MWCNT nanocomposites with 3 wt% offiller using two different approaches. The first one described creep experiments in terms of the creepstrain rate sensitivity parameter defined in Eq. (6). The second method monitored the time-dependentdeformation from the measurement of the creep compliance. A general tendency for the strain ratesensitivity to decrease when CNTs are incorporated into the epoxy matrix was found. This wasexplained by the authors in terms of several factors, the most important one possibly being CNTs act-ing as blocking sites hindering the displacement of the amorphous epoxy chains subjected to the fieldof deformation. In addition, the work also analysed the increased strain rate sensitivity found atincreasing temperatures above RT for both, the neat polymer and the reinforced composites. Thisresult was discussed in terms of increasing free volume as the temperature rose towards the glasstransition. The authors concluded that strain rate sensitivity was a more useful parameter in the studyof creep deformation than creep compliance studies.

Fig. 10. Variation of indentation strain rate with hardness for epoxy composites with different SWCNT content. Reprinted fromRef. [24], copyright 2004, with permission from Cambridge University Press. Values of the g parameter included in bracketsrefer to the creep exponent n defined in Eq. (5).


4.2. Graphene

Graphene is one of the most exciting nanomaterials being currently investigated both for funda-mental studies and potential applications. It consists in a flat, atomically thick, two-dimensional sheetcomposed of sp2 carbon atoms arranged in a honeycomb structure, and it is the basic building blockfor graphitic materials of all other dimensionalities. Graphene sheets offer extraordinary electronic,thermal and mechanical properties [137–142], superior electron mobility (15,000 cm2/V s), excellentthermal conductivity (�5000 W/m K) [140], very large surface area (�2630 m2/g) and the highestelectrical conductivity known at RT (6000 S/cm) [140]. Moreover, graphene is one of the strongestmaterials on earth, the elastic modulus of a suspended monolayer measured by nanoindentationwas found to be �1 TPa [142]. In addition, an ultimate strength of 130 GPa and a breaking strength200 times greater than that of steel were reported [143]. These unique properties make graphenean excellent additive for reinforcing polymer matrices, and the resulting nanocomposites have thepotential to rival or even surpass the performance of their carbon nanotube-based counterparts pro-vided that cheap, large-scale production and processing methods for graphene become available[141,144–147]. Graphene has already been used in a variety of applications such as sensors, batteries,supercapacitors, hydrogen storage devices and so forth. The first isolation of free-standing single-layergraphene from micromechanical cleavage of highly oriented pyrolytic graphite was reported by Nov-oselov et al. [138] in 2004. To date, several methods have been reported for the preparation of graph-ene sheets such as epitaxial growth on SiC wafers [148], chemical vapour deposition of hydrocarbonson metal surfaces [149], thermal exfoliation of graphite oxide [150], chemical reduction of graphiteoxide [151] or electrochemical intercalation [152]. Among them, the thermal and chemical exfoliationof graphite oxide are effective ways to prepare graphene in relatively large quantities at a low cost.Graphite oxide is a layered material produced by the treatment of graphite using strong mineral acidsand oxidizing agents. Each layer consists of covalently attached oxygen-containing groups such ashydroxyl, epoxy and carboxylic acid. The reduction of highly oxidized graphene oxide (GO) sheetsfrom the exfoliated graphite oxide leads to reduced GO (rGO) or functionalized graphene sheets (FGS).

Three main approaches have been reported for the preparation of polymer/graphene nanocompos-ites: melt-blending, solution intercalation and in situ polymerization of monomers [153]. In the firstapproach, the matrix is mechanically mixed with graphene at elevated temperature via extrusion orinjection moulding. However, because of thermal instability of most chemically modified graphene,the use of melt blending has been limited to a few studies [141,154,155]. The solution techniqueinvolves the dispersion of graphene in an appropriate solvent, the adsorption of the polymer onto


delaminated graphene sheets in solution and the elimination of the solvent, resulting in sandwich-likenanocomposites [21,156–160]. In the in situ polymerization strategy, graphene is first swollen withinthe liquid monomer, the initiator is then added and subsequently the polymerization is started eitherby heat or radiation [161]. To attain stable dispersions of graphene and control the microstructure ofthe resulting nanocomposites, polymer-covalent functionalization of graphene is highly desired [153].The direct modification of graphene flakes with polymer matrices by using the oxygenated groups (i.e.hydroxyl, epoxy and carboxyl) in graphene oxide (GO) represents an interesting strategy. In particular,thermoplastic polymers such as PVA [162], polyvinyl chloride (PVC) [163] or PMMA [161] and thermo-sets like polyurethane (PU) [158,159] have been covalently bonded to the GO surface. Moreover, clickchemistry reactions have been recently successfully used to modify graphene for its incorporation inpolymer nanocomposites [164,165]. An alternative approach is the non-covalent modification, whichenables the attachment of molecules through p � p stacking or hydrophobic (van der Waals) interac-tions, preserving the intrinsic electronic properties of graphene [154].

4.2.1. Modulus enhancements in graphene/polymer nanocompositesNanoindentation has been successfully applied to study the reinforcing effect of graphene-like fill-

ers in polymer matrices. Table 2 collects literature data published to date on this topic [21,154–160,166–174]. Similarly to Tables 1 and 2 includes relevant information for the discussion, e.g. samplemorphology, state of aggregation of the filler or methodology employed for analysis of the DSI data.

Fig. 11 illustrates DEI values for different polymer matrices as a function of graphene content. It isworth remarking that the term ‘‘graphene’’ encompasses a wide range of graphene-like structuresincluding graphite platelets (GP), graphene nanoplatelets (GNP) and different types of functionalizedgraphene sheets such as GO, exfoliated graphite (EG), thermally expanded graphite (TEG) and rGO. Inaddition, the graphene-like stacks can adopt different average sizes. Fig. 11 shows that, similarly toCNT-reinforced materials, most of the data apply to low filler contents (<1 wt%) and there are onlya few examples at higher loadings. A gradual reinforcement with increasing filler concentration isobserved for different glassy and semicrystalline thermoplastic matrices such as PMMA [21], PP[154] and PVA [160]. In the case of PVA, the trend surpasses the low concentration range approachinggraphene levels of 4 wt%. This behaviour contrasts with that of epoxy matrices where the modulushardly increases with the addition of graphene [166–169], especially at high filler loadings[166,167,169]. It is noteworthy that in all cases graphene sheets have been introduced in epoxy

0 1 2 3 4 5 6

0

50

100

150

1000

2000

ΔEΙ [%

]

Epoxy/GP [166]Epoxy/GNP [168]Epoxy/GNP [169]Epoxy/GNP [170]Epoxy/TEG [172]Chitosan/GS [156]

PP/GO [154] PP/GO [154] PVA/FG [21] PVA/GNP [160] PMMA/FG [21] PMMA/EG [157]

PU/rGO [158]PU/GO [159]

Graphene [wt%]

Fig. 11. Percentage of increase in the elastic modulus (DEI) for different graphene-reinforced polymer nanocomposites as afunction of the nanofiller content. Symbols in red indicate data that should be taken with caution (see Table 2 for specificdetails), and those in blue correspond to thermoset matrices. Solid and open symbols designate raw and functionalizednanofillers, respectively; M and s correspond to GO and OTES-modified GO, respectively; ⁄ indicates GO covalently bonded tothe polymer matrix and + denotes aligned rGO anchored to the matrix.


matrices without functionalization. This route seems to give rise to micrometer or sub-micrometerblocks that are homogenously distributed in the matrix [166,169]. Hence, the limited EI (and H, seeTable 2) increments detected in case of epoxy matrices must be related to the aggregation of grapheneinto stacks and the poor matrix-filler interaction.

Fig. 11 clearly shows that the largest DEI (�1800%) has been attained for amorphous PU compositesincorporating 5 wt% of well aligned, ultralarge-size rGO layers being the direction of the load perpen-dicular to the graphene sheets [158]. Although this DE value was calculated relative to E = 0.4 GPa ofneat PU and hence can be affected by some inaccuracies, it is reasonable to suggest that the orientationand size of the graphene sheets would play an important role. These rGO sheets exhibit mean lateralsize of more than 10 lm that minimizes the occurrence of matrix-filler interfaces. Most important,graphene sheets are self-aligned in the polymer matrix with their basal plane perpendicular to thefilm thickness. Indentation studies on suspended graphene or reduced graphene oxide monolayershave yielded modulus values of 1 TPa [142] and 0.25 TPa [175] respectively, while only 36.5 GPahas been found for the modulus of exfoliation of graphene sheets [169]. Hence, orientation shouldbe a major contribution to the outstanding reinforcement of PU/graphene composites.

Concerning thermoplastic matrices, the largest EI improvements are found for PMMA-based nano-composites approaching a modulus increase of 70% with the incorporation of only 0.6 wt% amide-functionalized few-layer graphene (FG) [21]. This functionalization route seems to be most effectivein providing a good load transfer across the polymer–matrix interface. The same authors reportedlower increments (DE = 35%) but still significant when incorporating the same amount of FG function-alized with –COOH and –OH groups to PVA [21]. Fig. 12a shows the plot of the DSI load-depth data forthis series of PVA/FG nanocomposites incorporating different percentages of filler. It can be clearlyobserved that the stiffness (slope at the onset of unloading) rises with increasing graphene loading.The modulus increments were partially attributed to an increase in the degree of crystallinity of thematrix, as illustrated in Fig. 12b. Crystalline phases are suggested to nucleate at the interphasebetween the filler and the matrix resulting in a strong polymer–matrix interaction. Other authors havereported similar modulus enhancements for PVA with the incorporation of partially oxidized graphene(see Fig. 11 and Table 2) [160]. Interestingly, the linear increment in modulus found with increasinggraphene loading was persistent up to the high loading of 4 wt% [160]. Enhanced crystallinity valueswere suggested to contribute to the modulus improvement for low filler loadings, in agreement withthe work of Das et al. [21]. However, for filler contents above 1 wt%, the crystallinity decreased due tophysical hindrance of graphene to the crystal growth, and the modulus increments were attributedsolely to the reinforcing effect of the nanofiller, its homogenous dispersion and the strong interfacialadhesion between the two nanocomposite phases. Comparable increments to those of PVA have beenreported for semicrystalline PP based nanocomposites loaded with 1.0 wt% GO [154]. Surface modifi-cation of GO with a coupling agent (octyltriethoxysilane, OTES) was found to further enhance E of PPnanocomposites (see Fig. 11) due to increased reactive sites on the GO surface that enhanced the inter-facial interaction with the matrix and hence promoted mechanical interlocking [154].

Planar orientation of the graphene sheets is also expected to play a fundamental role in achievingenhanced mechanical properties in films of graphene/polymer bilayer and multilayer assemblies. Infact, a recent study on ultrathin PVA/GO multilayer films fabricated by the LbL technique has attrib-uted the significant enhancement found for Er and H, 99% and 240% respectively, to the high degree ofplanar orientation of the graphene-like nanosheets [173]. The LbL technique has also been used tosequentially assemble a number of layers to produce polyelectrolyte/GO membranes. Incorporationof GO was found to be beneficial for the indentation mechanical properties [174]. The deposition ofa top graphene layer onto a flexible polymer substrate is possibly the simplest layer assembly in whichgraphene has been found to enhance the indentation mechanical properties of the film [157]. A mod-erate DEI of 16% was reported while a 97% H improvement was attained [157]. The authors discussedthe limited increment found for EI as a consequence of the far-field nature of the elastic stress fieldcompared to the plastic deformation zone.

4.2.2. Comparison with macroscopic mechanical propertiesIn spite of the a priori difficulties in comparing nanoindentation and mechanical data from

conventional macroscopic techniques due to technical dissimilarities, different scale of the volume

Fig. 12. (a) Representative load–displacement curves for PVA with different contents of acid functionalized FG. (b) Change incrsytallinity as a function of FG content. Adapted from Ref. [21], copyright 2009, with permission from IOP publishing.


of deformation, diverse load directionality and distribution of stresses, amongst other important fac-tors discussed in detailed in preceding sections, Table 2 shows an outstanding correspondencebetween the increments in modulus derived from indentation and macroscopic techniques. Compar-ison of the Y and EI values [158,169] shows a reasonable agreement, the latter values being slightlyhigher. Most interesting is the excellent agreement found in the only example where modulus valuesfrom dynamic indentation and DMA are compared, an epoxy matrix reinforced with GP [166,167]. Inthis work, special care was paid to carry out a comprehensive comparison between nano and macrotechniques. In the first place, three-point bending DMA and indentation measurements were com-pared at the same applied frequency. In addition, it was demonstrated that, for 1 lm-size GP, the vol-ume of deformation in DSI encompassed a platelet distribution that was representative of that of thebulk material. In fact, increasing the size of the graphite platelets to 15 lm produced divergent resultsbetween DSI and DMA because the local nanoindentation measurements were found to be no morerepresentative of the macroscopic properties.


4.2.3. Indentation size effectBecause the dimensions of the graphene-like platelets typically approach a few micrometers in lat-

eral size, the depth dependence of modulus and hardness in graphene/polymer composites is espe-cially sensitive to the distribution of platelets throughout the sample. Dynamic DSI studies onepoxy-based composites with 5.8 wt% GP (1 lm of mean lateral size) reported different E0 profilesdepending on the location of the initial contact that finally converged to the same modulus value atlarge penetration depths (h � 1500 nm) [166]. In contrast, similar experiments using 15 lm size GPyielded E0 profiles vs. indentation depth that exhibited a number of undulations, which were associ-ated with the occurrence of new platelets in the immediacies of the indenter tip [166]. In this case,the depth dependence did not exhibit a final plateau because the volume of deformation is not a goodrepresentation of the composite composition at any stage of the indentation test. In contrast to thesestudies, King et al. [169] reported constant E0 and H values with penetration depth above h = 200 nmfor epoxy matrices reinforced with 5 wt% of graphene nanoplatelets with average particle diameter of15 lm. Results seem to be at variance with the E0 depth dependence discussed above for similar epoxyreinforced matrices [166]. The apparent disagreement could be due to the size of the graphite blocksemployed. For shallow penetration depths (h < 200 nm), a conspicuous increase of E0 and H withdecreasing depth has been observed in a number of studies, possibly associated to an incorrect tip cal-ibration or uncertainties in the determination of the point of initial contact [169,174].

4.2.4. CreepThe incorporation of graphene-like nanofillers is generally found to enhance the creep resistance of

polymer matrices [21,157,169,173,176]. This is immediately detected in a standard load-hold-unloadcycle. During the hold period, depth increases at a lower rate when the nanofiller is present[21,157,173]; an illustration of such effect can be seen in the load-depth curves of Fig. 12 on PVA/FG composites. A thorough analysis of the creep behaviour has only been recently offered[145,169,176].

Chen et al. [157] analysed in detail the creep behaviour of a bilayer film of graphene on top of PET.The creep strain, defined as the indenter displacement at any point during the hold period divided bythe displacement at the beginning of the hold period, was found to be substantially smaller for thePET/graphene nanocomposite compared to that of the neat polymer. Analysis of the creep compliancealso revealed that the graphene overlayer significantly improved the creep resistance of PET. In thecase of epoxy matrices, the incorporation of 0.1 wt% GNP significantly decreased the creep displace-ment at peak load [176]. It was found that the creep displacement vs. creep time curves obtained witha cylindrical punch are well described by a power law, the exponent values being smaller for the nano-composite compared to the neat polymer. In contrast, King et al. [169] only found a slightly lowercreep compliance for GNP/epoxy composites than that of neat epoxy for relatively low loadings(2 mN) while the same creep compliance was found for both filled and pristine material for appliedloads in the range 10–45 mN.

4.3. Other organic nanofillers

In addition to CNTs and graphene, other carbon nanomaterials such as carbon nanofibres (CNFs),nanodiamond (ND), fullerenes, carbon black (CB) and so forth have been successfully used for thedevelopment of polymer nanocomposites. CNFs are intimately related to CNTs regarding their struc-ture and properties. They exhibit a tubular hollow configuration, with length in the range of 50 to1000 lm and outer diameter of 40–150 nm, smaller than that of conventional CFs (5–10 lm) but con-siderably larger than that of CNTs (1–40 nm). CNFs possess very high thermal conductivity (�1900 W/m K [177]), electrical conductivity close to that of graphite (�2 � 104 S/cm [178]) and superior Young’smodulus (in the range of 300–600 GPa [179]) and tensile strength (�8.7 GPa [180]). They are generallysynthesized by CVD process [181] in which the fibre grows around particles of a transition metal cat-alyst. This procedure involves several stages including gas decomposition, carbon deposition, fibregrowth, fibre thickening, graphitization, and purification. CNFs produced by this method are knownas vapour-grown carbon nanofibres (VGCNFs).


Among diamond-based nanomaterials, detonation nanodiamond (ND), also named ultradisperseddiamond or ultrananocrystalline diamond, is most attractive for nanotechnological applications. Det-onation nanodiamond can be easily suspended in water, has moderate cost and is non-cytotoxic andbiocompatible, hence, highly suitable for the biomedical field. It is produced by detonation of carbo-naceous explosives in a closed chamber either in a gaseous atmosphere, e.g. CO2 (dry method), or inwater (wet method) [182], and can also be synthesized from a suspension of graphite in organic liquidat atmospheric pressure and RT using ultrasonic cavitation [183]. ND is composed of crystalline nano-particles with typical diameter of 3–6 nm, specific surface area in the range of 300–500 m2/g, thermalconductivity of �1200 W/m K [184], hardness of �300 GPa [185] and modulus close to 500 GPa [186].When mixed with a polymer matrix, non-functionalized diamond nanoparticles exhibit poor disper-sion and tend to aggregate resulting in irrelevant property improvement. Therefore, a major issuein manufacturing ND-reinforced composites is to improve filler dispersion. To achieve this goal, var-ious approaches similar to those described previously for CNTs have been utilized, including highshear mixing, extrusion, in situ polymerization, wrapping the nanoparticles with surfactants, andcovalent chemical functionalization [187].

Fullerenes are hollow molecules in the form of spheres or ellipsoids composed entirely of sp2 car-bon atoms. They are similar in structure to graphite, but they also contain pentagonal and/or heptag-onal rings [188]. The simplest and most symmetric fullerene (named buckminsterfullerene C60), iscomposed of 20 hexagons and 12 pentagons, and was discovered in 1985 by Curl, Kroto and Smalley[189]. Unlike other allotropes of carbon, fullerenes are superconductors and exhibit low thermal con-ductivity (�0.4 W/m K [188]). This material is harder than graphite but softer than diamond, showinga hardness of around 170 GPa [185].

Carbon black (CB) is an amorphous form of carbon produced by the incomplete combustion ofheavy petroleum products. It is composed of roughly spherical particles with diameters between 10and 400 nm that are fused together into aggregates [190]. CB is extensively employed for polymerreinforcement due to its high electrical (�102 S/cm�1 [191]) and thermal conductivity (300–700 W/m K [192]), abundant availability, low density and low cost. Compared to the above mentioned nanof-illers, CB exhibits modest mechanical properties, modulus values approaching 15 GPa [193].

Recently, nanofillers based on natural polysaccharides like cellulose, starch or chitin have beenused for polymer reinforcement. Cellulose nanocrystals (CNC) can be obtained from native cellulosefibres by an acid hydrolysis [194], giving rise to highly crystalline (�63%) and rigid nanoparticles thatare approximately 100–300 nm long and 3–15 nm wide, referred to as nanowhiskers. CNC extractedfrom wood pulp display has good mechanical properties for use in material applications: elastic mod-ulus in the range of 100–150 GPa and tensile strength of �7.5 GPa [195], comparable with those ofKevlar and superior to those of GF. The most important restraint to the use of CNC as reinforcementin nanocomposites is the poor compatibility between the hydrophilic polysaccharides and the typi-cally hydrophobic polymer matrices. In order to prevent nanocrystals from aggregation and toimprove their adhesion with the matrix, surface modification is required via graft polymerization, sily-lation of the OH group or the use of surfactants [196].

Table 3 presents DSI data for polymer nanocomposites reinforced with different types of organicnanomaterials [115,121,187,193,197–208], together with other relevant information. The incrementsin modulus DEI are plotted as a function of the nanofiller content in Fig. 13.

4.3.1. Modulus enhancements in organic/polymer nanocompositesFig. 13 shows that limited DEI values are reported for epoxy matrices upon addition of a number of

fillers including carbon nanofibres, either vapour grown or coated, nanodiamond, fullerene and carbonblack. Remarkable increments are only found for very high filler loadings (�50 wt%) as commentedbelow. Analogous moderate improvements, or even decrements, are observed for the hardness values(see Table 3). This behaviour resembles that observed for CNT and graphene-reinforced epoxies (seeSections 4.1 and 4.2) and similarly to these latter studies, poor nanofiller dispersion and weak fil-ler-matrix bonding are envisaged to be the main factors limiting the mechanical enhancement of car-bon fillers/epoxy nanocomposites [197,198]. Interestingly, Sánchez et al. [197] suggested that thepresence of CNFs hinders the curing reaction of the epoxy matrix giving rise to reduced crosslinkingdensity which results in lower stiffness and hardness. This effect explains the high nanofiller content

0 10 20 30 40 50

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ΔEΙ [

%]

Epoxy/CNF [197]Epoxy/CNF [198] Epoxy/ND [199]Epoxy/ND [199]Epoxy/fuller. [115] Epoxy/CB [193] PLLA/ND [187] PLLA/ND [187]PA-11/ND [205] PA-11/ND [204] PVA/ND [202]PVA/CNC [201] PVA/CNC [200] PVAc/CNC [203] PS/CNF [198] PS/CB [206] PAN/ND [208] PMMA/CNC [207]

nanofiller [wt%]

Fig. 13. Percentage of change in the elastic modulus (DEI) for different polymer nanocomposites reinforced with carbonnanofibres (CNF), nanodiamond (ND), carbon black (CB), fullerenes (fuller.) or cellulose nanocrystals (CNC), as a function of thenanofiller content. The open symbols in Refs. [199,202,204,205,208] correspond to ox-ND. ⁄ indicates nanofiller grafted to thepolymer matrix, + and j denote aligned nanofiller. Red symbols refer to data possibly affected by inaccuracies (see Table 3).


(5 wt%) that is necessary to achieve an apparent improvement of mechanical properties. The authorsreported AFM studies that provided evidence of limited pile-up for indentations on the neat resin andthe CNF/epoxy composites with low nanofiller content while significant pile-up was detected forhigher filler loadings of 3 wt% presumably due to an incomplete curing process (see Fig. 14). Thepile-up effect was reverted for higher filler loadings of 5 wt% due to enhanced stiffness as a conse-quence of the incorporation of CNFs.

Most interesting is the work of Neitze et al. that reports a route to covalently bond nanodiamond tothe epoxy network [199,209]. Amino groups were incorporated into the surface of nanodiamond par-ticles and participated actively in the curing chemistry of the epoxy resin [199,209]. Modulus andhardness values of an aminated ND reinforced epoxy were found to be significantly higher (50%and 200% respectively) than those of a carboxylated ND/epoxy material with the same filler content[209]. It was found that amino groups of the nanodiamond particles compete with those of the curingagent in such a way that they can compromise the mechanical properties of the composite [199].Table 3 collects the modulus increments for some of these aminated ND/epoxy nanocomposites. Out-standing DE values are found at high loadings achieving�400% upon incorporation of�50 wt% of ami-nated nanodiamond (see Fig. 13). Such remarkable increment was explained as due to the formation ofa ND network crosslinked by epoxy. Other amine-ND reinforced polymers have been prepared usingoctadecylamine (ODA) covalently attached to ND [187]. A significant increase of mechanical propertiesin amine-ND/PLLA composites was achieved at filler concentrations up to 10 wt% (see Fig. 13). In con-trast, addition of the as-received ND nanofiller terminated with oxygen containing functional groupsproduced irrelevant changes in modulus and reduced increments in hardness (see Table 3). The goodaffinity between ND-ODA and the PLLA matrix together with the hydrophobicity of the filler providedthe framework for enhanced filler dispersion. TEM studies showed the occurrence of small agglomer-ates, tens of nanometers in size, at all filler concentrations that turned out to interconnect at highloadings. Uniform filler distribution, enhanced filler-matrix interaction and increased crystallinity val-ues account for the remarkable enhancement of mechanical properties.

The addition of ND to other polymer matrices has proved most successful in enhancing themechanical performance when significant functional groups at the surface of ND are promoted[121,202,204,208]. Indeed, the increased surface reactivity of oxidized ND (ox-ND) with respect todecarboxylated ND (deox-ND) has been shown to be critical for the improvement of the elastic mod-ulus values of polyacrylonitrile (PAN) [208]. Ox-ND has also been shown to produce marked modulus

Fig. 14. AFM image of indentations produced on: (a) neat epoxy and (b) epoxy/CNF (3.0 wt%). (c) Line profiles for neat epoxyand the nanocomposites. (d) Comparison of the elastic modulus of epoxy and CNF/epoxy nanocomposites measured by DMA(black symbols) and nanoindentation tests at two different maximum loads. Adapted from Ref. [197], copyright 2011, withpermission from Elsevier.


and hardness increments when incorporated to PVA at low filler loadings (�80–100% increase in E andH for only 0.6% ND, see Table 3 and Fig. 13) [121,202]. It was demonstrated that part of this improve-ment is due to the nucleating effect of ND producing significant matrix crystallinity enhancements andfavouring a strong matrix-filler interaction [121,202]. Most significant is the modulus increase of PA-11 (�300%) with addition of 20 % oxidized ND, although the authors do not describe the procedureemployed for the DSI data analysis and hence, results should be taken with caution [204]. Neverthe-less, this work proposes an interesting approach to achieve uniform dispersion of ND that consists inthe use of electrospinning as a vehicle to deposit nanodiamond with minimal agglomeration followedby heating the electrospun nanofibres to produce ND reinforced thermoplastic coatings.

Concerning CNC-reinforced polymers, the study of the nanomechanical properties reported so farreveals moderate improvements [200,203,207]. Composites in the form of fibres or films, as well asglassy and semicrystalline polymer matrices have been investigated. Loadings as high as 41 wt% havebeen added to electrospun PMMA fibres, and yet they produce mechanical enhancements of no morethan �30% as revealed by nanoDMA [207]. The authors argued that these limited increments couldarise from the anisotropy of the mechanical properties. CNC were aligned along the PMMA fibre axis,while indentations were produced in the perpendicular direction and hence, more significantincreases should be expected along the fibre direction. In CNC-thermoplastic films, nanofiller agglom-eration is found to be the main limitation for mechanical improvement. This is indeed behind the AFMnanoindentation studies on PVA/CNC and PVA/PAA/CNC (CNC wt% 620) [200] and polyvinyl acetate(PVAc)/CNC (CNC wt% 63) [203]. As an example, Fig. 15a shows the inhomogeneous distribution of

Fig. 15. (a) Topographic images of the PVA/CNC (15 wt%) nanocomposite showing the heterogeneous CNT dispersion. (b)Comparison of the elastic modulus obtained by tensile testing with those derived from the application of Hertz’s model to theloading and unloading indentation curves on (A) neat PVA; (B) PVA 90 wt%/PAA 10 wt%; (C) PVA 85 wt%/CNC 15 wt%; (D) PVA80 wt%/PAA 10 wt%/CNC 10 wt%; (E) PVA 70 wt%/PAA 10 wt%/CNC 20 wt%. Adapted from Ref. [200], copyright 2012, withpermission from IOP publishing.


CNC in a PVA/CNC 85/15 composite [200]. Interestingly, indentations produced at different locations,on a CNC agglomerate and on the matrix, revealed only small differences in the load-depth curves interms of maximum penetration depth and slope because the polymer underneath CNC contributes tothe recorded deformation behaviour. This work also highlights an important consideration for theanalysis of load-depth curves: E values derived from the application of Hertz’s model to the loadingand unloading seem to be at variance (see Fig. 15b) due to the deviation of the unloading behaviourfrom the condition of full elasticity.

Finally and most surprising is the remarkable modulus enhancement shown in Fig. 13 (up to 350%for 5 wt%) for a PS-based shape memory nanocomposites incorporating 1–5.5 wt% of carbon black par-ticles [206]. CB particles were incorporated as received in the crosslinked system, and nevertheless thereinforcement seems to be most efficient. This result contrasts with the modest increments attainedfor other thermoset materials incorporating similar amounts of CB (see Fig. 13). At the heart of thisinconsistency could be the fact that low storage modulus values have been determined by means ofDMA for all the PS-based shape memory nanocomposites (see Table 3). Indentation testing on compli-ant materials is known to display additional concerns in the detection of the initial contact that canintroduce some inaccuracy in the modulus data (see Section 3).

Recently, novel PVA-based hybrid nanocomposites incorporating two acid-functionalized carbonnanofillers with different dimensionalities have been successfully developed [121]. Results demon-strated extraordinary mechanical improvements due to synergistic effects (Table 3). In particular,DE values in the range 1000–1300% were found upon addition of binary combinations of SWCNTs/FG or SWCNTs/ND in small amounts (0.2 or 0.4 wt% each). Surprisingly, almost no change in thematrix crystallinity was detected, suggesting that an increase in this parameter did not contributeto the observed synergy. The reader should be here aware that, as repeatedly indicated in this review,the low E values reported for the PVA matrix (E = 0.7 GPa) are prone to some inaccuracy.

4.3.2. Comparison with macroscopic mechanical propertiesThe interest in correlating indentation modulus values with those determined from macroscopic

techniques is manifested in the significant number of papers that approach this issue[197,198,200,206,208].

The reinforcing effect of the filler is usually detected in a similar manner by DSI and conventionalmacroscopic techniques [197,206,208] although some discrepancies have also been reported[200,208]. For PS/CNF composites, Y values have been successfully used as a reference to correct forpile-up effects in the determination of the contact area necessary for DSI analysis [198]. However, it


is a most common finding that considerably higher modulus values are achieved by means of inden-tation testing [197,200,206,208]. Sánchez et al. [197] attributed these differences to indentation sizeeffects and the occurrence of local higher loadings. Pakzad et al. [200] argued in terms of less availablefree volume in DSI, occurrence of defects and microscopic cracks in macroscopic testing, differencesbetween compression (identified with DSI) and tensile testing and influence of dissimilar strain ratesin nano and macro techniques.

4.3.3. Indentation size effectFew studies have approached the depth-dependence of EI and H and only some examples on poly-

mer nanocomposites containing CNF, CB and CNC are available [193,198,207]. A conspicuous ISEbelow h � 300 nm and a minor reverse ISE above this penetration depth have been reported for PS/CNF and epoxy/CNF composites [198]. The marked E0 and H rise at small displacements was similarto that reported in the preceding sections for other carbon-based reinforced polymers and can be asso-ciated to an underestimation of the contact area and other sources of error at shallow penetrationdepths. In carbon black/epoxy composites, a gradual rise of modulus values at h > 200 nm has beenassociated to the influence of the stiffer silicon substrate [193]. NanoDMA storage modulus data havebeen reported for PMMA and CNC/PMMA electrospun fibres of 500 nm average diameter over a nar-row range of indentation depths covering 20–30% of the relative depth (indenter depth/fibre diameter)[207]. The constant E0 values encountered within this range of relative depths was interpreted by theauthors as a means of reassuring that issues arising from probe shape and sample depth were notinfluencing the genuine values of the mechanical properties.

4.3.4. CreepThe creep behaviour of CNF, ND, CB and CNC polymer nanocomposites by means of indentation

testing has been little explored so far, and yet a quick look to the data during the hold segment at max-imum load can provide worthy information on the creep properties. For example, the introduction ofCB and amine functionalized-ND in PS and PLLA respectively has been shown to significantly reducecreep [187,206]. Some analysis of the creep behaviour is only offered by Kaboorani et al. [203] usingthe relative change of the indentation depth at the peak load for PVAc composites incorporating CNC.A progressive reduction of creep was found upon addition of the nanofiller reaching prominent valuesof 55% decrease with only 3% of CNC.

4.4. Main features on organic fillers

4.4.1. Carbon nanotubesA wealth of data on CNT-reinforced polymers is available although with significant data scatter.

This is possibly due to the different experimental procedures and criteria adopted for the load-depthanalysis. However, and overall, some conclusions can be drawn that teaches us important aspects onthe reinforcement mechanisms.

Reinforcement of thermoset matrices with CNTs is quite a difficult task. Moderate increments ofmodulus or even some decrements are usually found. A poor dispersion of CNTs is possibly the mainfactor limiting the effective mechanical enhancement. Functionalization approaches do not seem to besuccessful; a few examples reveal that similar results are obtained for functionalized and raw CNTs.Perhaps the most successful route is the use of plasma functionalized MWCNTs grafted to the matrix,although more research is needed to extrapolate this behaviour to higher loadings.

In case of thermoplastic matrices, a fair dispersion of the filler, either raw or functionalized, is usu-ally achieved at low CNT loadings (<2%). Hence, this factor does not seem to be critical in obtainingmoderate EI and H increments that rise with increasing filler content. An optimum CNT-matrix inter-action seems to play here a relevant role. It has also been shown that the changes in crystallinity trig-gered by the CNT on the polymer matrix can significantly contribute to the EI and H enhancement.Results for high CNT loadings (>2%) suggest that in addition to the filler-matrix interface adhesion, fil-ler dispersion becomes now an important factor. In this respect, acid-functionalization of CNTs seemsto be the most efficient approach.


Interestingly, both in epoxy and thermoplastic matrices, some authors have explained the poormechanical enhancement found in certain composites in the light of the specific properties of CNTs:waviness, curved morphology, inferior bending properties or sliding ability.

Finally, creep studies on CNT-reinforced polymers are quite limited but results published suggestthat the incorporation of CNTs into polymer matrices can diminish the time-dependent deformation.

4.4.2. GrapheneNanoindentation studies on epoxy/graphene composites reveal irrelevant mechanical improve-

ments. It is noteworthy that in this case the filler has always been introduced without functionaliza-tion. In contrast, the incorporation of graphene-like fillers in thermoplastic matrices has been provedto be most successful always following a functionalization route. Occasionally, in semicrystallinematrices, the modulus increments have been partially attributed to the nucleating effect of graphenethat, on the one hand, increases the degree of crystallinity of the matrix and on the other, nucleatescrystalline phases at the interphase between the filler and the matrix, giving rise to a strong poly-mer–matrix interaction.

Overall, the factor that seems to produce the most significant enhancement of mechanical proper-ties is the orientation of the graphene sheets. Incorporation of well-oriented micrometer-sized graph-ene sheets into a polymer matrix produces a remarkable increment of indentation modulus in thedirection perpendicular to the plane of the graphene.

For this type of composites, indentation size effects can be expected because filler size is of thesame order as indentation displacement. It has been shown that the modulus of epoxy/graphene com-posites at large indentation depths is only characteristic of that of the bulk material when the volumeof deformation encompasses a representative number of platelets. In this latter case, three point bend-ing DMA studies and dynamic nanoindentation measurements, both carried out at the same fre-quency, were found to be in excellent agreement.

Finally, the incorporation of graphene-like nanofillers has been shown to enhance the creep resis-tance of thermoplastic matrices and to a lesser extent of epoxy resins. Moreover, a graphene overlayerhas been also found to reduce the creep strain of a flexible PET film. Detailed creep analysis has onlybeen offered recently and it is expected that the field progresses in the next years.

4.4.3. Other organic fillersSimilarly to CNT and graphene–reinforced epoxies, the incorporation of other carbon-based nanof-

illers (carbon nanofibres, nanodiamond, fullerenes and carbon black) to epoxy polymers produces lowor moderate modulus and hardness increments, probably due to the poor filler dispersion and weakmatrix–filler interaction. In addition, it has been suggested that the nanofiller can hinder the curingreaction of the epoxy giving rise to a matrix with reduced stiffness and hardness. Unprecedentedmechanical increments have been found in one example of aminated-ND/epoxy composite wherethe filler was covalently attached to the epoxy matrix (DE0 � 400% for a loading of �50 wt%).

Nanodiamond seems to be the most effective reinforcing filler for thermoplastic matrices, aminefunctionalization being the most successful route for optimum mechanical performance. Improved fil-ler dispersion, enhanced filler–matrix interaction and enhanced levels of crystallinity in case of semi-crystalline matrices account for the remarkable increase of mechanical properties.

Analysis of the indentation data at peak load reveals that the incorporation of CB, ND and CNC tothermoplastic matrices significantly enhances the creep resistance.

Recent indentation results on PVA reinforced with binary combinations of carbon nanofillers (CNT,graphene and ND) point towards significant synergistic effects when two of these nanofillers areemployed. Further research in this direction should be of great value.

5. Polymer nanocomposites incorporating inorganic nanofillers

Inorganic–organic nanocomposites generally refer to polymer composites composed by nanoscaleinorganic building blocks and a polymer matrix. These building blocks include: layered silicates (e.g.,montmorillonite, hectorite, saponite), metal nanoparticles (e.g., Au, Ag), oxides (e.g., SiO2, TiO2, Al2O3),


semiconductors (e.g., PbS, CdS) and so forth. They aim to combine the characteristic properties of poly-mers with those of an inorganic material. The small size of the nanofillers yields a very large interfacialarea, which may give rise to a significant relative amount of interphases with properties different fromthe bulk. This fact opens the possibility to have synergistic phenomena that produce effects greaterthan the sum of the individual components. On the contrary, the very high surface energy of thenanofillers favours the agglomeration into larger particles. The consequence is poor nanoparticle dis-persion within the nanocomposite that usually leads to degradation of the properties and, for this rea-son, the preparation process becomes critical [210,211].

As previously done in Section 4, dedicated to organic based nanofillers, this part of the review col-lects most of the published data obtained by DSI techniques on polymer nanocomposites with inor-ganic nanofillers. The main aim is, likewise, to understand the influence of the different parametersthat characterize the indentation response of these materials and to compare the results with thoseobtained by macroscale conventional techniques. The discussion on different types of inorganic nanof-illers will be subdivided in three sections: layered silicates or nanoclays (5.1), spherical nanoparticles(5.2) and other inorganic nanofillers (5.3).

5.1. Layered silicates

5.1.1. Characteristics, preparation methods and types of nanocompositesPolymer/layered silicate (PLS) nanocomposites have attracted great scientific and industrial inter-

est due to the extraordinary improvements in properties attained at very low filler contents. Enhancedstrength, modulus and toughness, increased tear, radiation and fire resistances and reduced thermalexpansion and permeability to gases have been reported in these nanocomposites making them idealmaterials for applications in food packaging, structural automotive components and electronicsamong others [212–217].

Montmorillonite (MMT) and hectorite are the layered silicates most commonly used for thepreparation of the nanocomposites. These clays exhibit a very high aspect ratio with layer thick-nesses of 1 nm and lateral dimensions that may vary from a few hundred nanometers to microns.Their structure include charge compensating counter-ions such as Na+ located in the inter-layerspace which can be replaced by organic ammonium and phosphonium cations with long alkylchains leading to organically modified clays having better compatibility with polymers. In fact,the ability of layered silicates to tune their surface chemistry through ion exchange reactions withorganic and inorganic cations, together with the possibility to disperse them into individual layers,are the two main characteristics that determine their adequate incorporation in polymer nanocom-posites [213,214].

Depending on the strength of the interfacial interactions between the polymer matrix and theclay, two main types of PLS nanocomposites can be obtained: intercalated and exfoliated struc-tures. The first type is formed when polymer chains are inserted between the layers of the clayand a separation of several nanometers between the platelets is observed, while the second oneis obtained when the layers of the clay are completely separated and dispersed throughout thepolymer matrix. The latter structure usually leads to better mechanical properties than the inter-calated one because it maximizes the polymer–clay interactions facilitating the stress transfer.However, complete exfoliation of the clays is difficult to achieve and most of the polymer nano-composites reported have mixed intercalated and exfoliated structures [212–214]. Moreover, forhighly non-polar polymers such as polyolefins, intercalation is thermodynamically unfavourableand it becomes necessary to incorporate a compatibilizer, usually a functionalized polymer orblock copolymer, in order to favour interactions with the clay and miscibility with the matrix[218–222].

The ability to promote organoclay exfoliation is strongly influenced by the preparation of the PLSnanocomposites. Four main preparation strategies have been developed: in situ template synthesis(sol–gel technology), in situ intercalative polymerization, solution intercalation and melt intercalation.The last method presents environmental advantages and is most commonly used in the industry. Adetailed description of the methods for producing PLS nanocomposites can be found in a number ofreviews [212–214].


5.1.2. Modulus and hardness changes in polymer/layered silicate nanocompositesRegarding the mechanical performance of PLS nanocomposites, significant improvements in

Young’s modulus and tensile strength have been observed for very low clay concentrations (2–5%)[212–215]. In addition to the preparation method, the nature of the matrix and the content and degreeof dispersion of the filler, other factors such as surface treatment of the clay, orientation of the dis-persed platelets or the influence of the nanofiller on the development of crystallinity in thermoplasticbased PLS nanocomposites need to be considered when analysing the mechanical properties of thistype of materials.

DSI results on PLS nanocomposites are collected in Table 4 and in Fig. 16 [218,219,223–252].Young’s modulus data determined by tensile testing and DMA storage modulus values are alsoincluded in the table. Similarly to preceding tables, other characteristics such as nanofiller type andcontent, composite processing method and morphological features related to the filler dispersionand filler-matrix interaction have also been included for the sake of the discussion.

5.1.2.1. Polymer matrix influence. Fig. 16 shows that the largest DEI values have been reported for nano-composites with semicrystalline polymer matrices, mostly polyamides and polyesters. A similar trendhas been observed for hardness data (Table 4). Enhancements of 130% in E and 160% in H have beenreported for a PA-6 nanocomposite containing 10 wt% of an organically modified clay [223], while val-ues of DE = 110% and DH = 75% were found with addition of 6 wt% of a surface modified MMT [226],both nanocomposites prepared by melt blending. In addition, there are many examples of polyamidenanocomposites (PA-6, PA-6,6, PA-11 and PA-12) with significant DEI and DH at low filler contents[224,226–234] regardless the method of preparation (see Table 4 and Fig. 16). In the case of polyesterssuch as poly (3-hydroxybutyrate-co-3-hydroxyvalerate) (PHBV), the addition of 5 wt% organoclay by asolution intercalation method resulted in E0 and H enhancements of 100% and 90%, respectively [237].Outstanding DE and DH values have been reported for poly(ethylene oxide) (PEO)/clay nanocompos-ites with very high loading (270% and 250% respectively at 50 wt% clay content), however at low con-centrations the improvements are not significant [240]. One of the reasons for the large improvementsdescribed in both mechanical properties for polyamide and polyester PLS nanocomposites is thestrong interaction between the matrix and the silicate layers that in some cases has been attributed

0 2 4 6 8 10

0

20

40

60

80

100

120

140

Δ EΙ [

%]

clay [wt%]

Epoxy/Cl 93A [247]Epoxy/A1100 [248]Epoxy/A1120 [248] PA-6/I30TC [223] PA-6/CoAl-LHD [224] PA-6/1.34TCN [226] PA-66/1.34TCN [226] PA-66/1.34TCN [227] PA-66/1.34TCN [227]PA-66/Hectorite [230] PA-11/1.34TCN [232] PA-12/1.34TCN [233] PA-12/layered silicate [234] PLA/Cl 30B [235] PHBV/Cl 30A [237] PP/PP-g-MA/Cl 15A [218] PP/PP-g-MA/Cl 15A [218] PP-g-MA/I31PS [219] HDPE/HDPE-g-MA/Cl 15A [238] HDPE/HDPE-g-MA/N1.44P [238] PS/Cl 15A [243] PS/Cl 20A [244] CEAR/NMMT [249] UPE/layered silicate [251]PU/C20 [252]PU/C30 [252]

Fig. 16. Changes in DSI modulus DEI as a function of nanoclay loading for different polymer/layered silicate nanocomposites.Organically-modified clays have been used in all cases except in Ref. [249]. Red symbols highlight data that should be takenwith caution (see Table 4 for details). Blue symbols have been employed for epoxy-based composites. Symbols h and Mcorrespond to values for polished and unpolished samples, respectively; � and � denote composites prepared by SSE and TSE,respectively.


to the formation of hydrogen bonds [213,226]. In contrast, PLS nanocomposites based on semicrystal-line non-polar polymers such as polyolefins, exhibit small DEI and DH values even with the incorpo-ration of compatibilizers due to poor clay-matrix interactions (see Fig. 16 and Table 4) [218,238].Indeed, the enhancements in modulus measured in PP/organically modified MMT (OMMT) nanocom-posites with maleic anhydride functionalized polypropylene (PP-g-MA) as a compatibilizer [218] andin high density polyethylene (HDPE)/OMMT nanocomposites compatibilized with polyethylene-grafted maleic anhydride (PE-g-MA) [238] are not higher than 13% for clay contents 65 wt%. Thesevalues can be raised to DE0 � 100% at 10 wt% of clay using PP-g-MA as the polymer matrix becausemore MA functional groups are incorporated [219].

Fig. 16 and Table 4 show moderate H and E0 increments for PLS nanocomposites based on amor-phous polymers [242–244]. On the other hand, DEI and DH for thermoset PLS nanocomposites areusually lower than those of thermoplastic based ones [247–249,251], as already remarked for organicfillers. Indeed, the largest DE reported for epoxy/OMMT nanocomposites is 20% for a clay content of7.5 wt% [247] and 23% for a nanoclay reinforced unsaturated polyester (UPE) with 5 wt% filler[251]. Analogously, the strongest rise in H found for epoxy/OMMT nanocomposites is 50% for a claycontent of 3 wt% [248]. One exception to this trend is found in PU/OMMT nanocomposites with exfo-liated structures [252], where very large improvements in modulus and hardness (107% and 65%respectively at 5 wt% filler content) are reported. The consistency of these results will be commentedbelow. Functionalized organic cations are commonly employed to enhance the thermoset/clay inter-action. However, it is worth mentioning that some organic modifiers, such as alkyl ammonium chains,may lower the matrix cross-linking density near the clay surfaces [247]. Therefore, the degree of cross-linking is affected by the degree of interaction and has a competitive effect on the clay reinforcement.In epoxy/OMMT nanocomposites the smaller increments in properties at higher loadings were attrib-uted to the reduced crosslinking density of the epoxy matrix due to the higher presence of surfactant,together with the increased population of intercalated clay clusters [247].

At this point, it becomes apparent that only some general trends regarding EI and H values of PLSnanocomposites in relation to the nature of the polymer matrix (semicrystalline, amorphous or ther-moset) can be drawn. To fully understand the differences observed in these properties other crucialfactors must be taken into account.

5.1.2.2. Clay dispersion and type of structure. One of the most important parameters that has to be con-sidered when analysing the mechanical performance of PLS nanocomposites is the degree of disper-sion of the clay and, in particular the formation of intercalated and/or exfoliated structures.Exfoliated structures have been reported to display higher EI and H values than the intercalated ones[218,242,252]. Yusoh and Song [252] found significant differences in the surface mechanical proper-ties of exfoliated and intercalated PU/organoclay nanocomposites thin films. However, these resultsshould be taken with caution: in the first place, the load-depth curves of the neat polymer exhibit aclear artefact at the beginning of unloading (most probably associated to the test exceeding the max-imum penetration allowed); in addition, E and H values are remarkably small and possibly affected byan incorrect determination of the point of initial contact. Higher E0 and H values for the exfoliated sys-tem compared with the intercalated one have also been reported for poly(propylene carbonate) (PPC)/organoclay composites prepared by solution intercalation [242]. Treece and Oberhauser [218]reported the influence of two different melt-blending strategies (conventional twin-screw extrusion(TSE) and single-screw extrusion (SSE) with in line supercritical carbon dioxide (SCCO2) feed) forthe preparation of compatibilized PP/PP-g-MA/OMMT nanocomposites. Different degrees of clay exfo-liation and dispersion were observed by TEM. The high shear of the TSE was very effective in exfoli-ating and dispersing the clay than the SSE method. DSI experiments were most worthy in evaluatingthe effect of processing, dispersion and exfoliation on the modulus. A significant mechanical enhance-ment was only found for nanocomposites prepared by the TSE process (see Table 4). In PA-6/claynanocomposites E and H increased with clay loading (1–10 wt%), as expected, but a slowdown ofthe growth was observed above 5 wt% clay content [223]. These results were related to the changein the clay morphology observed when loading exceeded 5 wt%, from exfoliation-dominated to inter-calation/exfoliation mixture as evidenced by TEM and X-ray diffraction. The same trend was observedin epoxy/organoclay nanocomposites [247], the mechanical properties exhibiting smaller increments


above 2.5 wt% clay content due to the formation of intercalated structures besides the exfoliated ones.Piscitelli et al. [248] investigated epoxy based nanocomposites with sodium montmorillonite (Na-MMT) untreated and modified with two different aminosilanes prepared by two different dispersionmethods: sonication (S) and combination of sonication and ball milling (SB). Hardness and modulusincreased with the addition of the clay but the enhancement was found to depend on the dispersionmethod and the functionalization of the layered silicate (see Table 4). Sonication allowed better nano-filler dispersion and better interface interaction between the clay tactoids and the epoxy matrix. How-ever, the combination of sonication and ball milling increased the aggregation of clay tactoidsreducing the interfacial interactions. Consequently, all the nanocomposites prepared by the combinedmethod showed lower E and H than their counterparts prepared by simple sonication. On the otherhand, those nanocomposites with silylated MMT showed higher values for both properties than thenanocomposite with Na-MMT. The authors concluded that the interfacial interactions between theamine moieties anchored to the clay surface and the epoxy matrix play a more significant role thandoes the clay morphology over the composite properties.

From the examples discussed above it can be concluded that dispersion of the clay nanoplatelets inthe polymer matrix is one of the most important factors controlling the EI and H improvement in PLSnanocomposites independently of the nature of the matrix. Exfoliated structures make the entire sur-face of clay layers available for the polymer and maximize polymer–clay interactions, boosting thereinforcement. Other factors playing a role in enhancing the nanocomposite mechanical propertiesinclude the organic modification of the clay and the addition of compatibilizers to the polymer matrix.

5.1.2.3. Crystallinity of the matrix. Another parameter that needs to be taken into account is the influ-ence of the nanofiller on the development of crystallinity in PLS nanocomposites based on semicrys-talline polymers and its impact on the nanoindentation measurements. It is well established that clayplatelets can affect the crystalline structure and crystallization behaviour of the polymer nanocompos-ites but controversial results have been reported [212]. A number of DSI studies have determined thatthe incorporation of clay nanofillers in systems with a semicrystalline matrix results in an increase inEI and H with clay content but simultaneously reduces the degree of crystallinity [223,225,227–229,238]. Between the two competing factors affecting the nanocomposites mechanical behaviourwith increasing clay loading, i.e., the reinforcing effect of the nanoclay and the decrease in crystallinityof the polymer matrix, the first one seems to dominate. Sikdar et al. clearly observed this behaviour inPA-6/OMMT using three organic modifiers with the same composition and different end functionalgroups [225]. Moduli and H were found to increase as the levels of crystallinities were reduced uponaddition of the OMMTs (Fig. 17). Results were explained invoking different clay-polymer interactions

Fig. 17. Effect of organic modifiers on the DSI modulus of PA6/OMMT nanocomposites. Reprinted from Ref. [225], copyright2007, with permission from John Wiley & Sons, Inc.


for each OMMT that give rise to distinct polymer chain entanglements and interphases. It is importantto point out that in this study the nanoclay filler also introduced a polymorphic change in the crystal-line structure of the matrix, from the a phase in pure PA-6 to the c phase in the nanocomposites. Asimilar H and EI rise concurrent to a crystallinity decrease was observed in HDPE nanocomposites withorganically modified clays and a compatibilizer (PE-g-MA) [238] and in PA-6,6/OMMT nanocompos-ites [227]. It has been suggested that the presence of the clay confines polymer chains and segmentalmobility hindering rearrangement during crystallization and restricting the formation of crystals.However, there are also examples in which the enhancements in mechanical properties were notaccompanied by any modification in the matrix crystallinity and thus, were attributed exclusivelyto the reinforcing effect of the stiff clay nanofiller such as PLA/OMMT nanocomposites prepared bymelt intercalation [235].

5.1.2.4. Processing method. Nanoindentation can also be very useful to explore mechanical anisotropyinduced by the processing method. Injection moulded composites of crystallisable polymer matricesrepresent a clear example in which valuable information on the local mechanical properties can beachieved by means of DSI. Shen et al. [223] studied E and H of injection moulded PA-6/clay nanocom-posites and found that both properties adopted higher values along the injection direction than per-pendicularly to the melt flow and increased from the surface to the core of the samples. The differencein experimental values across the samples was found to be as large as 20%. The results were explainedas due to a parallel rise of clay content and crystallinity from the outer to the inner region of the sam-ples and of clay orientation along the injection direction. Crystallinity changes were associated to thetemperature gradient effect induced by the injection moulding process. Fig. 18 illustrates the variationof the mechanical properties at different locations in the moulded bars, exposed by polishing, from thenear surface (A), through the intermediate zone (B), to finally reach the core position (C). Uneven dis-tribution of the clay has been also put forward to explain the enhanced E0 and H values found in theinner regions of PA-6,6/clay of the injection moulded bars [229].

Fig. 18. DSI (a) modulus and (b) hardness for three exposed surfaces at 0.5 cm (A), 3 cm (B) and 5.5 cm (C) from the outersurface of the injection-moulded bars for neat PA6 and its nanocomposites with 5 and 10 wt% clay. Reprinted from Ref. [223],copyright 2005, with permission from Elsevier.


Preferred orientation of the clay along the injection direction in poly(ethylene naphtalate) (PEN)/OMMT nanocomposites has been revealed by small angle X-ray scattering (SAXS) [239]. SAXS resultsalso showed that, upon annealing, this preferred orientation induced an anisotropic crystalline mor-phology in the PEN matrix with more secondary lamellae aligned parallel to the clay basal plane.The higher modulus values measured by DSI in the flow direction were then attributed to the addi-tional reinforcing effect of the anisotropic arrangement of the crystalline lamellae together with theclay orientation in this direction. Here, in annealed samples, the degree of crystallinity of the nano-composites was similar to that of the matrix.

Other effects, such as multiple reprocessing cycles, were investigated in clay/nanocomposites byDSI tests and correlated with changes in their nanostructure. In polystyrene (PS)/clay nanocomposites[243], X-ray diffraction experiments determined that the state of intercalation of the clay wasimproved after eight processing cycles using TSE. H values of the nanocomposites increased by 10%after reprocessing (from 380 to 420 MPa) and by 16% compared to that of neat PS (360 MPa). Resultswere attributed to better dispersion of the clay in the matrix. In contrast, H of the neat polymer wasslightly reduced to 350 MPa after 8 cycles due to a reduction in molecular weight caused by thermaldegradation. The same behaviour was observed for E0 which increased by 5% with clay incorporationand by an additional 3% after reprocessing the nanocomposites. However, E0 of neat PS decreased pro-gressively with the reprocessing cycles due to chain scission.

The effect of gamma irradiation and natural weathering (up to 130 days) on the degradation behav-iour of PLA/OMMT nanocomposites and its impact on the local mechanical properties were also stud-ied [253,254]. While neat PLA was strongly degraded by gamma irradiation the correspondingnanocomposites were less affected. The steric interactions of the clay layers were modified by irradi-ation which promoted a better dispersion of the clay within the matrix. As a consequence, E0 and Hvalues of the irradiated nanocomposites showed certain improvement [253]. A slight increase in theseproperties, depending on clay content and exposure time, was also observed in the nanocompositesunder natural weathering [254].

As a final point, the influence of strain rate on the mechanical properties of PLS nanocomposites hasbeen investigated by dynamic DSI [228]. In exfoliated PA-6,6/organoclay nanocomposites it wasobserved that E0 was not affected by the strain rate whilst H increased with increasing strain ratefor both neat PA-6,6 and the nanocomposite samples. It was suggested that the former elasticresponse is insensitive to the strain rate because the motion of the amorphous chains in semicrystal-line materials is frozen in the glassy state.

5.1.3. Other layered silicate/polymer systemsPLS nanocomposites using polymer blends as matrices have been studied by DSI. Jarrar et al.

reported significant enhancements in the reduced elastic modulus and hardness of PA-6/PA-6,6 blendswith the incorporation of an organically modified clay, the greatest augment being for the 50/50 com-position [226]. The enhancements in mechanical properties were supported by FTIR spectroscopyresults which showed the formation of hydrogen bonding and possible formation of ionic bondsbetween the polymers and the nanoclays.

Additionally, the influence of the combination of clays with other micro or nanofillers on the nano-mechanical properties of PLS nanocomposites has also been investigated [246,236]. Biopolymer basedchitosan/MMT/hydroxyapatite (HA) nanocomposite exhibited significant E and H enhancements ascompared to pure chitosan as well as to chitosan/MMT and chitosan/HA composites [246] (seeTable 4). The improvements were attributed to better nanoparticle dispersion as observed by AFMand X-ray diffraction and stronger interactions between the three components as evidenced by FTIRstudies. In PLA/Kenaf fibre (KF)/OMMT nanocomposites H was improved as compared with PLA/OMMT and PLA/KF due to good interfacial bonding [236].

Finally, several authors have investigated laminated or multilayer PLS nanocomposites by DSI[241,245,250]. E and H values of layer-structured polypyrrole (PPy)/MMT nanocomposite films pre-pared by electrodeposition increased at very low clay loadings, but above 0.01 wt% both propertiesdecreased due to MMT agglomerates [241]. Another organic–inorganic nanocomposite with laminatedstructure was prepared by a hydrothermal-electrophoretic method [250]. In this assembly an acrylicanodic electrophoretic resin (AAER) was intercalated into the interlayer space of Na+MMT (NMMT) by


the hydrothermal process (HMMT films). A remarkable DE0 value (70%) was observed for the HMMTfilms compared with neat NMMT considering that the polymer content in the former is very low(8 wt% as determined by TGA). This improvement was attributed to the laminated structure that playsa crucial role in absorbing energy during the elastic deformation of the nanocomposite. The last exam-ple of PLS layered assemblies is constituted by PDDA/MMMT thin films prepared by stepwise alternat-ing polyelectrolyte and clay deposition from solution [245]. The hardness H of the multilayer thin filmwas 0.42 GPa, which is the highest value reported for PLS nanocomposites (Table 4) and is comparableto that of a soft metal such as copper (0.46 GPa). The well-ordered anisotropic structure also yielded ahigh modulus value of E0 = 9.5 GPa.

5.1.4. Comparison with macroscopic propertiesMost of the DSI studies on PLS nanocomposites offer, as well, the Young’s modulus values mea-

sured by tensile testing (Table 4). Comparison with the corresponding DSI moduli, reveals that a sim-ilar trend with clay content and material modification is found although, in general, the absolutevalues derived from both techniques are different. EI values have been found to be higher [235], orlower [226,227] than Y. The discrepancies have been ascribed to scale factors and size effects, differ-ences in the loading direction and local variations of crystallinity and crosslinking density. Interest-ingly, injection moulded nanocomposites having a processing-induced mechanical anisotropyexhibited EI values parallel to the injection flow that were comparable with tensile data. In both cases,the loading direction was applied parallel to the flow direction [223]. A detailed study of local vs. glo-bal stiffness has been reported for PP-g-MA/MMT nanocomposites in which DSI was performed in theclay-polymer intercalated region, the boundary region between clay aggregate and matrix and thematrix alone [219]. Fig. 19 shows that the local stiffness enhancements determined for indentationsat a depth of 500 and 1000 nm are considerably higher than results from tension and compression.This is partially due to the increment in local filler content. In addition, other factors such as polymerchain confinement and topological constraints at the nanometer scale need to be considered.

The storage modulus obtained by DMA has also been compared with the storage modulus mea-sured by dynamic DSI (nanoDMA) and differences have been reported [227,232,233,238,239]. The ori-gin of the discrepancies was explained again based on size effects and differences in loading directions

Fig. 19. Relative increase in stiffness for PP-g-MA/MMT nanocomposites vs. clay content. Stiffness values were measured by DSIat two different indenter displacements (500 nm and 1000 nm) and by tensile and compressive testing. Reprinted from Ref.[219], copyright 2006, with permission from Elsevier.


and load frequency used. It is worth mentioning that in one DSI study [225] both, E and E0, were deter-mined. It was found that E0 followed the same trend as E as a function of clay content and of the dif-ferent organic modifiers of the clay. The absolute values from both methods were in this case veryclose to each other (see Table 4).

5.1.5. Indentation size effectA variation of EI and H as a function of indentation depth has been reported in a number of DSI

studies on PLS nanocomposites [219,224,227–229,232,233,240,242,247]. Most of them used thedynamic method [219,224,227–229,232,233,242] and found an increase of the mechanical propertieswith decreasing penetration depth. At shallow depths (300 nm), the behaviour was attributed to anumber of factors, as have been explained in preceding sections of this review (see Sections 2.1.3,4.1.5, 4.2.3 and 4.3.3). At higher indentation depths, in the range 1000–6000 nm, Beake et al. [240]investigated PEO/clay nanocomposites by means of a number of load-partial unload experiments.The authors observed a continuous decrease of E and H values with increasing penetration depth inthe whole interval studied, the variation being more pronounced in the solution processed samplesthan in the melt processed ones. Results were partially associated to a strain rate effect arising froma trapezoidal load function and not to a real gradient in mechanical properties (see Section 2.1.3 for adetailed explanation).

5.1.6. Creep propertiesThe creep behaviour of PLS nanocomposites has also been investigated by DSI

[223,224,238,240,244,247,255]. In general, a decrease in creep displacement and rate has beenobserved in PLS nanocomposites with increasing clay content showing that the incorporation of rigidclay nanofillers improves the creep resistance [223,238,244,255]. In injection moulded PA6/MMTnanocomposites the creep behaviour was investigated on the skin and core regions of the samplesand compared with creep measured by DMA cantilever-bending [255]. Both techniques showed adecrease in creep compliance with the addition of the clay but, only the cantilever measurementsrevealed an additional improvement in the time-dependent compliance behaviour. The authorsexplained this behaviour considering the higher capacity of the organoclay to restrict the molecularmobility in the bulk compared to the surface. In other systems, factors related to the nanostructureof the polymer matrix seem to have an influence on the creep behaviour. In PA66/MMT nanocompos-ites, an unexpected increase in creep displacement was observed with increasing clay content [227].This reduction of creep resistance was attributed to the decrease of crystal size and crystallinity in thenanocomposites as revealed by X-ray diffraction. In epoxy/OMMT nanocomposites the creep resis-tance increased when clay loading was less than 2.5 wt% and the opposite behaviour was found athigher loadings [247]. This result was discussed on the basis of a reduced crosslinking density nearthe clay surface.

5.2. Spherical inorganic nanoparticles

The high aspect ratio of the nanofillers in most of the aforementioned polymer nanocomposites(e.g., carbon nanotubes, carbon nanofibres, graphenes or nanoclays) often results in a large increaseof the nanocomposite melt viscosity as fillers resist shear. The melt viscosity rise can cause undesir-able slower production rates and higher processing costs. The use of spherical nanoparticles opens thepossibility of achieving both a viscosity reduction together with a reinforcement of the tensile mod-ulus. It has been shown that the size and shape of nanoparticles can be a key parameter to modify rhe-ological and mechanical properties of polymer nanocomposites [256].

5.2.1. Characteristics, preparation methods and types of nanocompositesThe most common inorganic nanofillers, generally adopting a spherical shape, can be classified in

three main groups: metallic nanoparticles, oxide nanoparticles and other miscellaneous sphericalfillers.

Currently, nano-sized metal particles are a focus of interest in biomedical sciences, engineering andmany other fields. These materials can be synthesized and modified with various chemical functional


groups which allow them to be conjugated with antibodies, ligands, and drugs, opening a wide rangeof potential applications in biotechnology, drug delivery, and diagnostic imaging [257]. Other signif-icant applications result from their optical and electromagnetic properties, especially those relatedto quantum-size effects such as photo and thermo-luminescence, dichroism, size-dependent ferro-magnetism, superparamagnetism, electromagnetic wave absorption, and interference shielding orsuper-catalytic activity [258]. The difficult handling of metallic nanoparticles makes embedding intopolymer matrices a valid solution for many applications because most of the mentioned propertiesremain to a certain extent in the nanocomposites. Concerning the mechanical properties of polymernanocomposites including metallic nanofillers, the literature search performed for the present surveydid not yield any DSI study, representing an area of great potential for near future investigations.

Inorganic oxide nanoparticles share many of the properties of the metallic nanofillers. They are alsoessential in modern technologies such as microelectronics (silica, high-k dielectrics), sensors (tinoxide), corrosion protection (alumina, zirconia), optoelectronics (transparent conductors), energy pro-duction and storage (titania, zeolites), heterogeneous and environmental catalysis (transition metaloxides), drug carriers or biocompatible materials (HA). Common methods for nanocomposite prepara-tion are solution blending techniques, graft polymerization, sol–gel methods and surface modificationof the nanoparticles [259].

Among the numerous inorganic–organic nanocomposites, polymer composites reinforced withnanosilica (SiO2) are the most commonly reported in the literature, having attracted substantial aca-demic and industrial interest and being employed in a variety of applications. Silica nanoparticlespresent several advantages, for instance ease of preparation at a relatively low cost, possibility of per-forming surface modifications with different functional groups and acceptable biocompability. Thesenanocomposites can be prepared following various synthetic routes, according to the way that eachphase is introduced. The organic polymer matrix can be introduced as a precursor (monomer or oli-gomer), as a preformed linear polymer (in molten, solution, or emulsion states), or as a polymer net-work, physically or chemically cross-linked. The nanofiller, in turn, can be introduced as pre-existingnanoparticles or as precursors such as tetraethyl orthosilicate (TEOS), tetramethyl orthosilicate(TMOS) or perhydropolysilazane (PHPS). This leads to three general methods for the preparation ofpolymer/silica nanocomposites according to the starting materials and processing techniques: blend-ing, sol–gel processes, and in situ polymerization (Fig. 20) [211]. Blending involves simple mixing ofthe polymer matrix and the silica fillers, and agglomeration of nanoparticles is a usual drawback.The second method is based on direct mixing of monomers with silica particles followed by a poly-merization process. In situ polymerization begins with a mixture of both silica precursors and polymermonomers, followed by the polymerization of precursors and monomers [78].

Because of the strong tendency towards aggregation among inorganic nanoparticles, which maydepress the overall physical–chemical properties, considerable efforts have been devoted to optimizethe interfacial interaction between the two nanocomposite phases, i.e. to enhance the compatibilitybetween the polymer (hydrophobic) and the nanosilica. A common procedure is to modify the surfaceof the silica nanoparticles (especially for the blending and in situtechniques), which simultaneously

Fig. 20. Scheme showing the three general approaches to prepare polymer/silica nanocomposites. Reprinted from Ref. [211],copyright 2008, with permission from the American Chemical Society.


can enhance the dispersion of nanosilica in the polymer matrix [260]. The surface modification ofnanosilica can be carried out by either chemical or physical methods. The former involve modificationeither with coupling agents or by grafting polymers while physical methods use surfactants or mac-romolecules adsorbed onto the surface of the silica nanoparticles. Composites with nanoparticles con-taining an inorganic core and an organic shell or surface modified nanoparticles can be considered aspecial type of nanocomposites and they are often designated as hybrid nanoparticles [261]. For fur-ther details, the reader is referred to the recent and comprehensive review by Zou et al. on polymer/silica nanocomposites [211].

5.2.2. Modulus changes in nanocomposites containing spherical inorganic nanoparticlesThe DSI results published for the nanocomposites containing spherical inorganic nanoparticles are

summarized in Fig. 21 and Table 5. The figure summarizes the percent variation, in relation to thevalue of the corresponding neat polymer matrix, of DSI elastic modulus data for different polymermatrices reinforced with spherical nanoparticles as a function of the nanofiller content in weight per-cent. A general look at the figure shows that the proportion of loading is usually much higher thanthose reported for other types of nanofillers, probably due to a comparatively lower cost, but alsobecause of the especial applications in which they are involved. In the following, a discussion onthe effect of the type of matrix on the mechanical properties of polymer/silica nanocomposites willbe presented, together with a description of nanocomposites including different oxide nanoparticlesand other miscellaneous spherical shaped fillers.

5.2.2.1. Polymer/silica nanocomposites. (a) Thermoset matricesSeveral studies have reported the characterization of epoxy/nanosilica nanocomposites by DSI

[75,76,262–264,268]. Lam and Lau [262] emphasized the importance of studying the localized elasticmodulus by means of DSI to understand the distribution of nanoparticle agglomerations in epoxy/nanosilica composites. The nanocomposites were prepared by preheating and ultrasonication at dif-ferent temperatures followed by in situ polymerization. DSI experiments on the cross-section of thesecomposites revealed that at lower sonication temperatures (40 �C) E decreased as the indenter movedfrom the bottom to the top sample surface. This result was attributed to the slow curing time whichfavours the gravitational effect of the nanoclay. At higher temperatures, however, local mechanicalproperties of the composites were evenly distributed throughout the entire sample. At a filler concen-tration of 4 wt%, EI was found to increase from 2.3 to 5.5 GPa across the cross-section of the sample for

Fig. 21. Changes in DSI modulus DEI as a function of nanofiller loading for different polymer nanocomposites incorporatinginorganic spherical nanofillers. See Table 5 for comments on the data in red colour. The symbols M, O, s and } correspond tofunctionalized nanoparticles. ⁄ indicates nanosilica grafted to the polymer matrix.


a sonication temperature of 40 �C, and up to a constant value of 8.8 GPa when the temperature waselevated to 100 �C. It must be noted that the remarkably high absolute values provided (see Table 5)should be taken with caution because of the notorious nose effect observed in the experimental load-depth curves. Very recently the work by Tzetsis et al. [263] illustrated the negative effects of clusteringon the mechanical properties of epoxy nanocomposites reinforced with fumed silica. Both, uniaxialtensile tests and DSI results, showed that Y, E and H of the nanocomposites steadily decreased withincreasing fumed silica content. The decrease in E was around 25% for a composite with 3 wt% SiO2,attributed to the strong tendency to cluster formation of fumed silica as also revealed by opticalmicroscopy (OM) and SEM analyses. In this case, the sonication procedure and the high mechanicalmixing followed by curing at RT and postcuring at 100 �C for 2 h were not sufficient to yield a gooddispersion. A discussion on different possible sources of error concerning E data was also given, beingone of them that the tip area function can be material-dependent (see Section 2.1.1 and [12]). Directincorporation of dehydrated nanosilica with 10 nm average diameter in an epoxy matrix, by simplymechanical mixing and ultrasound waves, was also used by Allahverdi et al. [264]. This study howeverfound significant increments in E and H obtained by DSI with the addition of the nanofiller to the poly-mer. A nanocomposite with 5 wt% of well dispersed nanosilica showed DE and DH values of 21% and26%, respectively, as compared with the neat epoxy. It is noteworthy that the DSI experimental pro-cedure used does not seem to contemplate a hold period at peak load which might cause inaccurateresults. The reinforcement effect of the silica nanoparticles was also found to improve the storagemodulus E0DMA of the composite in both glassy and rubbery regions. The surface modification of SiO2

has been shown to more effectively enhance the tensile elastic modulus and yield strength of an epoxyresin [265,266]. Following this route, a sol–gel technique was employed by Wang et al. [76] and Zhanget al. [75] to prepare epoxy based nanocomposites containing surface-modified nanosilica particleswith average size of 20 and 25 nm, respectively. This procedure resulted in homogeneously dispersednanoparticles even at high filler contents [267] (Fig. 22, left). In the former study [76], the authors per-formed micro and nanoscale indentation and scratch tests on the nanocomposites to determine theinfluence of the load content on the mechanical and tribological properties. Indentation resultsshowed that H and E of the composites increased monotonously with the particle content reachingincrements of 33% and 40%, respectively, by the addition of 24 wt% nanosilica. The experimental val-ues were well fitted using the Hashin–Shtrikman lower bound model [73], which is situated slightlyabove the Reuss limit. In addition, scratch tests also revealed an effective improvement in the tribo-logical properties of the epoxy with the addition of proper amounts of SiO2. Zhang et al. [75], in turn,reported the viscoelastic properties of similar nanocomposites by dynamic nanoindentation and

Fig. 22. (a) TEM picture of silica/epoxy nanocomposite with 5 wt% SiO2; (b) Dynamic storage modulus vs. indentation loadusing a frequency of 75 Hz for the different composites investigated (from top to bottom: 0, 2, 5, 10, 17, 23 wt% silica). Adaptedfrom Ref. [75], copyright 2009, with permission from John Wiley & Sons, Inc.


derived their dynamic storage modulus E0 and loss tangent. E0 was found to increase with the forcefrequency, denoting a loading rate effect and also with the nanosilica volume fraction (Fig. 22, right).Using the proper indentation loads, DE0 yielded a value of 46% with the addition of 23 wt% nanosilica.

Castrillo et al. [268] used PMMA microparticles filled with nanosilica (10 wt%) to modify an epoxymatrix. An addition of 5% microparticles resulted in a nanosilica content in the nanocomposite of only0.5%. This final small content of filler was not intended to achieve high mechanical reinforcement, butto enhance other properties such as absorption of ultraviolet light, electric or magnetic shielding, con-ductivity or dielectric behaviour. However, the good dispersion of the silica nanoparticles in thePMMA domains, as well as that of the silica-filled PMMA microparticles in the epoxy matrix, main-tained or even slightly increased Martens hardness (2%) and modulus (4%) as compared to those ofthe neat epoxy.

Apart from epoxy resins, other DSI studies on thermosets such as crosslinked polyphenylenes orpolyacrylates with nanosilica as reinforcement have been found in the literature [77,210,269,270].In the case of polyacrylate matrices, Fig. 21 shows that the enhancement in mechanical propertiesis much higher than the one found for epoxy thermosets. Lin et al. [77] reported the fabrication oflow dielectric constant (low-k) nanocomposite thin films to be used in small semiconductor devices.The films were prepared by spin coating and thermal curing of solution mixtures of low-k thermosetprepolymers (highly crosslinked polyphenylenes) and silica nanoparticles with an average diameter of8 nm. The mechanical properties of these nanocomposites were characterized by means of DSI using adynamic method at very small loads. It was found that the addition of the nanofiller enhanced both E0

and H (around 50% and 40%, respectively, for 40 wt% SiO2). DE0 was smaller than that predicted by theHalpin–Tsai equations [74]. This result was tentatively explained due to poor interfacial adhesion and/or aggregation of the hydrophilic silica nanoparticles in the hydrophobic thermoset matrices. Theaddition of SiO2, however, had strong negative effects on the insulating properties of the nanocompos-ites which were attributed to the presence of impurities in the silica nanoparticle solution. The DSIdynamic method was also used to test acrylate-based nanocomposites prepared by UV in situ poly-merization [210,271]. The results showed strong E0 and H increases, by more than 100% and 200%,respectively (Table 5) for a nanocomposite with 5 wt% SiO2. Devaprakasam et al. [269] aimed tounderstand the different tribo-mechanical performance of a micro and a nanosilica reinforced polymercomposite produced by visible light curing. In both cases the load content was 56 vol% in matrices ofmonomeric dimethacrylates. The nanocomposite included 40–70 nm diameter nanosilica and the pri-mary particles in the microcomposite presented a bimodal size distribution ranging from 200 to500 nm and 1 to 4 lm, respectively. DSI results showed that E and H of the nanocomposite werehomogeneous throughout the sample surface yielding extremely high absolute values (around 15and 0.7 GPa, respectively). On the other hand, results for the microcomposite were found to be highlyheterogeneous. The sharp microparticle edges increased friction and wear, giving also rise to a high-energy dissipation. Such behaviour was only locally found in the nanocomposite in the presence ofnanoparticle clusters and aggregates. Soloukhin et al. [270] reported the preparation of hybrid nano-sized silica particles (5–15 nm) by mixing a colloidal silica dispersion in a water-containing organicmedium with a silane-coupling agent, 3-(trimethoxysilyl) propyl methacrylate (MEMO) and subse-quently with different methacrylic monomers and a UV photoinitiator. The obtained cross-linked sil-ica-methacrylate hybrid coatings were deposited on PC substrates and the mechanical properties weredetermined using DSI. They found a modulus E increase of almost 200% for a silica content of 48 wt%.The mechanical properties attained were found to be dependent on the chemical nature of the mono-mer used that considerably influenced the degree of crosslinking density. The authors concluded that,in general, the properties of hybrid coatings cannot be simply discussed in the light of filler content,but other factors such as the mutual influence of filler content on cross-link density, coating thicknessand chemical composition have to be taken into account.

(b) Thermoplastic matricesBhattacharya and Chaudhry [78] tested silica-reinforced PVA biocompatible nanocomposites with

the final aim of mimicking the properties of human bone. A homogeneous silica dispersion in watercontaining PAH was added to a water solution of PVA and cast on glass, forming 40 lm thick films.The indentation modulus E increased from 300 MPa for neat PVA to 8.1 GPa for the nanocompositewith 55 wt% nanosilica (2600% increase), as shown in Fig. 21. The results from this work were reported


by the authors as one of the largest increases in modulus found in the literature, being certainly thelargest in this review. However, the reader should be aware of the very low modulus values foundfor the neat polymer which can affect the obtained D E values. Results were explained as a conse-quence of the attachment of the molecular chains of PVA to the surface of the functionalized silicananoparticles through hydrogen-bonding interactions between the hydroxyl groups of PVA and theamide groups of PAH. The good dispersion in the PVA matrix and the absence of particle agglomera-tion also contributed to the superior mechanical properties. A further factor that should be taken intoaccount for a semicrystalline matrix is the crystallinity variation in the presence of the filler. Here, as agradual decrease of polymer crystallinity with increasing filler content was observed, it would havecontributed to a reduction in the modulus of the matrix. Hence, the increase in mechanical propertiesdue to filler reinforcement was even more remarkable. A similar behaviour was observed for polymer/layered nanocomposites as discussed in Section 5.1. Moreover, by coating the surface of the silicananoparticles with calcium HA, the resulting nanocomposite (55 wt% of silica) further increased themodulus to 11 GPa, a value that is close to that of cortical bone.

As a new example of hierarchical polymer composites (see Section 4.1.3.3) that are reinforced bymore than one reinforcing agent, Molazemhosseini et al. [272] reported the fabrication by the melt-mixing process at 400 �C of PEEK based hybrid composites reinforced with short carbon fibres (SCFs)and nano-SiO2 particles. Nanoindentation and nanoscratch methods were used to evaluate the nano-mechanical phases, namely the bulk PEEK matrix, the SCFs and the interphase region. Results revealedan instantaneous elastic recovery in the SCFs while a time-dependent contribution was found in theneat matrix denoted by a nose in the unloading cycle. Such an effect could be favoured by the absenceof a holding time at the end of the loading cycle and as a consequence the accuracy of the experimen-tal results can be questioned. It was found that the incorporation of 25 wt% SCFs into neat PEEKyielded a remarkable DE of 140%. In addition, the joint presence of nano-SiO2 particles in the conven-tional composite was shown to effectively contribute to additional E and H improvements of 44% and22% for a nanosilica content of 2 wt%, respectively.

Stojanovic et al. [273] reported a different route for the surface modification of nanosilica, namelyusing SCCO2 as an environmentally friendly solvent to improve the mechanical properties of PMMA. AMEMO-type coupling agent was also used. A comparison of conventional and supercritical coatingmethods revealed enhanced properties for the nanocomposites obtained through the latter processesregarding particle size distribution, amount of coated silane, and homogeneous dispersion in thePMMA matrix. Thus, the nanocomposite containing 3 wt% of SiO2 obtained by supercritical processingof the nanosilica sol showed an increase in E and H of 19% and 34%, respectively.

Low temperature sol–gel methods, although frequently used to produce nanosilica composites,have two main drawbacks: silica synthesized contains many lattice defects and the composites areoften contaminated by catalysts. For this reason, Saito et al. [274] synthesized PS/silica nanocompos-ites from a silica precursor such as PHPS that could overcome both difficulties. A blended organic solu-tion of PHPS and PS derivatives with hydroxyl groups gave rise to grafted block copolymers containingPHPS branches. Solution casting and subsequent calcination at around 100 �C produced compositefilms with microphase separated domains of PHPS and PS. Clear microphase separation was observedfor the composites when the degree of grafting was close to 100% and in some cases silica spheres of�40 nm of diameter surrounded by the organic matrix were obtained. EI and H were found to increasewith the silica content, although one can notice a ‘nose’ effect at the beginning of unloading; no sig-nificant differences were observed among the different morphologies. Organic/silica nanocompositesprepared with PHPS could then be used as hard coatings, lighter than the pure silica counterparts. Thesame authors [275], prepared nanocomposites of PMMA-rich spheres in a silica-rich matrix followinga similar PHPS route to provide a convenient hard coating for neat PC films. They obtained a maximumsurface hardness for the composites on the substrate of 1.07 GPa.

Douce et al. [276] synthesized silica nanoparticles from TEOS, a frequent precursor. The fillers haddifferent sizes (from 15 to 60 nm diameter) and two types of surface modification in order to promoteor minimize interactions with the matrix. The obtained nanosilica were prepared as colloids and intro-duced at different concentrations in glycidoxy-propyltrimethoxysilane (Glymo) type coatings. Suchcoatings are meant to protect transparent polymer lenses which usually have a poor scratch resis-tance. Both, indentation and scratch experiments were used to characterize the mechanical properties


of the films. The moduli E gradually increased with the addition of fillers, being this effect larger (70%increase at 15 vol% SiO2) for the filler with smaller size (15 nm). On the other hand, the weakness ofmatrix-filler interactions had no significant effect on the modulus of the Glymo films, while core–shellfillers improved it slightly.

Polymer/silica nanocomposites based on poly(2-hydroxyethyl acrylate) (PHEA) were prepared byRodríguez Hernández et al. [277] by the simultaneous polymerization of the organic and the silicaphases in a sol–gel process using TEOS as silica precursor. The structure of this system was investi-gated using AFM in the tapping mode and DSI experiments. AFM nanoindentation explores the visco-elastic behaviour up to 30 nm depth from the surface, but the obtained stiffness also depends on thestructural characteristics of the surrounding material. The stiffness increase was quite significant forsilica contents lower than 10 wt% (from around 0.08 N/m for the PHEA homopolymer to 0.27 N/m forthe composite containing 10 wt% of silica, i.e. more than 200% increase in stiffness). However, therewas a further qualitative change in stiffness for silica fractions above 15 wt%, resulting in highly scat-tered stiffness values. The explanation offered was that for a silica content lower than 15 wt%, the sys-tem consisted of isolated silica aggregates dispersed in the organic matrix, while above thatconcentration the structure became co-continuous with that of the organic matrix. Hence, silica fillersbehaved as an inorganic scaffold improving the mechanical properties of the nanocomposite.

5.2.2.2. Other oxides. This subsection includes a few examples of other metal oxides (e.g., ZnO, Al2O3,Fe3O4) frequently used as nanofillers and investigated by means of DSI. Ciprari et al. reported the char-acterization of the mechanical properties of four polymer nanocomposite systems by DSI tests andDMA and investigated the role of the interphase structure on their behaviour [278]. Alumina(Al2O3) and magnetite (Fe3O4) nanoparticles were dispersed in PS and PMMA matrices, which werechosen on the basis of their differing reactivity with metal oxides. The structure of the interphasewas investigated and correlated with the mechanical properties of the composites. The results indi-cated that Al2O3 nanoparticles were more reactive with the polymer matrix than Fe3O4 counterparts,but neither showed strong interactions with the matrix as compared to other studies. The low inter-phase density around the high number of nanoparticles resulted in a decrease of the tensile modulusof the nanocomposites compared with those of the neat polymers.

Conducting polymer materials used for instance in photovoltaic devices are easily degraded by thecombined action of oxygen and moisture. Polymer nanocomposites are suitable materials for theencapsulation and protection of such devices. Gupta et al. [279] reported a technique based on graftinghydride-terminated PDMS to allyltrimethoxysilane functionalized c-alumina nanoparticles to makenanocomposites that can be used for the abovementioned purpose. The DSI technique was used toanalyse the mechanical characteristics of the composites. In all cases E values were extremely low(see Table 5) but, from the unloading curves, it was clear that a large elastic recovery took place afterthe removal of the load which is an important characteristic of an encapsulant. The work by Chakr-aborty et al. [280] aimed to compare the scratch hardness with the indentation hardness of PMMA/ZnO nanocomposites produced by solution blending and in situ spin coating. A good correlationwas found between both properties at low indentation depths, where a linear relationship betweenhardness and the ZnO content was observed. Ramezanzadeh and Attar [281] studied the corrosionresistance of polymer composites to be used as coatings to protect metal parts. For this purpose epoxynanocomposites containing different contents of nanozirconia particles were prepared by in situ poly-merization. The nanocomposites were exposed to 3.5 wt% NaCl solution up to 60 days, and theirmechanical properties (before and after exposure to NaCl solution) were studied by DMA and DSItechniques. Results showed that the blank sample was severely deteriorated after exposure to the cor-rosive electrolyte due to a significant decrease of cross-linking density with a concurrent diminutionof indentation hardness. It was shown that these properties improved with the addition of nanopar-ticles giving rise to nanocomposites with better corrosion resistance.

5.2.2.3. Miscellaneous spherical shaped fillers. A few studies [22,282] have been devoted to investigatethe nanomechanical properties of polymer composites reinforced with nanoparticles of layered metaldichalcogenides like WS2and MoS2. These non-carbon materials possess similar structure to fuller-enes, and are named as inorganic fullerene-like (IF) nanoparticles. They have exceptional properties


such as very high stiffness and strength, attributed to their small size (typically in the range of 40–180 nm); closed-cage layered structure and excellent solid lubricant behaviour. Moreover, becauseof their chemical inertness and low agglomerating tendency, they can be homogenously dispersedwithin polymer matrices via simple melt-blending technique [283]. The reinforcing effect of IF-WS2

nanoparticles on three different thermoplastic matrices, PP and polyphenylene sulphide (PPS) [22]and PEEK [282] has been recently investigated by DSI (Table 5). In the first study [22,284,285],dynamic DSI was used to obtain the mechanical properties of both polymer composites and they werefound to initially rise with increasing filler content and then level-off at loadings of �0.5 wt%. Themechanical improvement was discussed in terms of the influence of the nanoparticles on the nano-structure of the polymer matrices and the reinforcement due to intrinsic properties of the filler,including matrix-particle load transfer. The analysis of the H data suggested that the nanostructuralchanges induced in the polymer matrices can contribute up to DH � 10%, while the filler reinforce-ment can reach DH � 50%. Remarkable DE0 and DH values were found for PPS based nanocomposites(by up to 35% and 60%, respectively, at 8 wt% IF-WS2), while only moderate enhancements wereobserved for the PP counterparts (around 18% for the same concentration). This discrepancy wasattributed to the different nature of the amorphous phase for both nanocomposites.

In the second study [282], IF-WS2 nanoparticles were incorporated into PEEK coatings using an aer-osol-assisted deposition process with the aim of reducing the coefficient of friction (COF) and improv-ing the wear resistance. Higher proportions of IF-WS2than the abovementioned ones were added andthe results showed significant enhancements in the mechanical properties of the nanocomposite coat-ings. In particular, E and H continuously increased up to �60% at 20 wt% IF-WS2 concentration. Theinfluence of possible crystallinity changes was not discussed. On the other hand, the COF was reducedby 70% upon incorporation of 2.5 wt% of the nanofiller, but a further increase in the IF-WS2 content didnot produce any additional reduction in the COF of the coatings.

Nanosized spherical HA has also been used to improve the mechanical properties of biopolymers. Astudy by Thomas et al. reported for the first time the electrostatic cospinning of collagen (type I) andHA (Fig. 23a and b), representing a potential nanofibrous osteoconductive and bioactive nanobiocom-posite scaffolds [286]. The HA nanoparticles exhibited a mean diameter size of 100–150 nm (Fig. 23c).SEM analyses showed a well interconnected pore network structure with nanofibrous morphology ofrandomly oriented fibres in the diameter range of 500–700 nm, depending on the composition(Fig. 23a and b). The fibre diameter increased with increasing HA content. Tensile testing and dynamicDSI were used for the mechanical characterization (Fig. 23d). The neat collagen fibrous materialshowed a dynamic modulus E0 of 0.2 GPa, value that increased to 0.6 GPa (200% difference) with theaddition of 20 wt% HA. This strong rise was attributed to an increase in rigidity of the polymer dueto the addition of HA and the strong adhesion between the two composite components. It is worthmentioning that the small modulus values measured in this work bring up serious experimental prob-lems, especially in the detection of the point of initial contact (see Section 3), a fact that was mini-mized by using the dynamic DSI option. Nevertheless, the significant enhancement of mechanicalbehaviour achieved by the incorporation of HA can be clearly discerned from a quick inspection tothe load-depth curves of Fig. 23d. As a final point it was established that the chemical crosslinkingof collagen (x-collagen) further increased the absolute values of the mechanical properties of thenanobiocomposites.

Organic electronics based on conjugated polymers constitute a promising route for low cost andflexible optoelectronic applications. A method to increase device efficiency is to incorporate semicon-ducting nanostructures such as quantum dots (QDs) into the polymer. QDs are nanoscale particleswith tunable optical and electrical properties. McCumiskey et al. [287] reported the preparationand mechanical characterization of nanocomposite thin films consisting of CdSe QDs and the electro-luminescent polymer poly[2-methoxy-5-2(20-ethylhexyloxy-p-phenylenevinylene)] (MEH-PPV). Thecomposite films were prepared by solution blending and spin casting on glass substrates, using differ-ent concentrations of QDs. QD dispersion was assessed by means of TEM. The QDs were found to bemore homogeneously dispersed at 90 wt% than at 50 wt% (Fig. 24a and b). The former concentration ismore typically employed for QD polymer devices. The modulus values E of the polymer matrix linearlyincrease with increasing QD volume content (Fig. 24c and d). This is a common result among polymernanocomposites, however, it is quite exceptional to keep a linear increase up to very high nanofiller

Fig. 23. SEM micrographs of electrospun scaffolds: (a) collagen fibres and (b) collagen +10% nanoHA (uncrosslinked), (c) TEMimage of nanoHA powders and (d) Nanoindentation load–displacement curves of collagen/nanoHA biocomposites. As nanoHAcontent in the fibre increases, the maximum load also increases. Adapted from Ref. [286], copyright 2007, with permission fromElsevier.


contents (i.e. 95 vol%). In this case, a DEI of 240% was observed for a QDs content of 95% (from 11.7 to41.2 GPa).

5.2.3. Indentation size effectThe EI and/or H variation with penetration depth have been observed in a number of papers

[75,210,270,273,274,287]. For silica nanocomposites and using a dynamic DSI method, an ISE wasfound for penetration depths below 200 nm [210]. This range was considered of little value to deriveany material information. Some authors have recommended the application of indentation loads aslarge as possible (over 200 lN) to get constant and reliable properties [75].

Soloukhin et al. investigated the typical response of a coating-substrate system where the substrate(PC) is softer than the polymer methacrylate/silica nanocomposite [270]. Thus the strong decrease in Ewith increasing indentation depth, using loads as high as 1 N, was attributed to the influence of thesofter PC substrate. On the other hand, the absence of a constant E plateau at shallow indentationdepths where the influence of the substrate was negligible, was tentatively explained as due to a com-bination of sink-in and pile-up effects. The effect of a harder glass substrate, typically producing areverse ISE which often comes together with a normal ISE at smaller penetration depths, has been alsoreported [274,287]. The influence of a glass substrate on the H variation with indentation depth of PS/silica nanocomposites [274] and on the E dependence for MEH-PPV/CdSe nanocomposites was studied[287]. In both cases, a discussion on which penetration was appropriate for mechanical characteriza-tion of the material was carried out. In the first case, H values were taken at 20% displacement [274],while McCumiskey et al. determined the average modulus E and hardness values for each samplewithin an optimum range, in between the observed ISE and reverse ISE effect (Fig. 24c) [287]. Finallyit is worth mentioning that a reverse ISE was also observed for indentation loads >10 mN in PMMA/nanocomposites [273].

Fig. 24. Plan-view TEM images showing CdSe QD dispersion at (a) 50 wt% in MEH-PPV, and (b) 90 wt% in MEH-PPV, drop-castfrom pyridine–chloroform binary solvent mixtures. (c) Elastic modulus as a function of contact depth and (d) as a function of QDloading (wt% and vol%) in MEH-PPV (error bars represent values within one standard deviation of the mean). Adapted from Ref.[287] with permission from IOP Publishing.


5.2.4. Creep propertiesThe only study involving creep behaviour concerns MEH-PPV/CdSe films [287]. The indentation tip

displacement was monitored as a function of time during the 30 s hold period at maximum load. Itwas shown that the amount of viscoelastic creep decreased with increasing CdSe QD loading. Creepresults also determined that high QD contents suppressed the viscoelastic behaviour of the matrix.One could expect that it was a consequence of a granular character of the films due to insufficientpolymer–polymer bonding that would lead to a stiffness decrease. However, the actual stiffnessincrease had to be explained by a different bonding mechanism which could be envisaged as sintering.

5.2.5. Comparison with macroscopic mechanical propertiesDynamic DSI (nanoDMA) and DMA data, concerning nanocomposites incorporating spherical nano-

particles, have been compared in a number of studies [22,75,278]. In most cases, although the relativevariations DE0 and DE0DMA are quite similar, the absolute modulus values usually are higher for dynamicindentation, E0 > E0DMA. For epoxy/silica nanocomposites [75], E0 and E0DMA results were also comparedwith quasi-static E data [76], Young’s modulus values Y and with theoretical results using microme-chanical models (Voigt–Reuss and Hashin–Shtrikman [73]). It was shown that E0 � 2E0DMA using thesame force frequency for both techniques. In turn, E and Y were found to be more similar to E0DMA thanto E0. Flores et al. [22] discussed the higher E0 values obtained by DSI in PP/IF-WS2 and PPS/IF-WS2

nanocomposites as compared with those obtained by DMA in the light of the different directionalityof the stresses applied in both methods. Quasi-static DSI moduli E have also been found to be higher


than E0DMA in a number of cases [264,273]. On the other hand, two studies [210,286] reported higher E0

values than the corresponding tensile properties. In the second one, concerning porous electrospunmaterials [286], the difference was attributed to the fact that tensile properties, in bulk, mainlydepend on their porosity, total fibre content, and orientation of fibres whereas nanomechanical prop-erties should be closer to those of the individual electrospun fibres. The weakness of this argument isthat using penetration depths of 500 nm with a Berkovich indenter, and being the fibre diameter of thesame order, it is difficult to figure out a deformation field not including a good number of fibres, poresand different orientations. Tzetsis et al. [263], working on fumed silica/epoxy systems showed that Eand Y modulus values practically coincided when using a modified area function related to the neatepoxy polymer as a reference. It was considered that the typical calibration procedure which involvescalibration on a hard reference material such as fused silica may not resemble the contact situationwith softer surfaces. Finally, and in contrast with other studies, Bhattacharya et al. for PVA/silica nano-composites obtained modulus values derived from DMA and tensile tests which were comparablewith those of quasi-static DSI and seemed to agree well with the ones obtained using the Halpin–Tsaiequations [78].

5.3. Other inorganic nanofillers

In addition to layered silicates and spherical nanoparticles, other inorganic nanofillers such as two-dimensional graphene-like boron nitride (BN) [288], inorganic fullerene-like tungsten disulphide IF-WS2 nanotubes [289], AlOOH boehmite nanorods [290], calcium silicate hydrate (CSH) [291], polyhe-dral oligomeric silsequioxane (POSS) [292–295], to mention but a few, have been incorporated intopolymer matrices. Amongst them, POSS has received extensive research attention for the preparationof hybrid nanocomposites. It has a nanometer-sized cage-like structure composed of a cubic octamericmolecule with an inner inorganic silicon and oxygen framework that is externally surrounded byorganic functions. It can be depicted by the formula (RSiO1.5)n where n is an even number andR = H, Cl, or a variety of organic groups. POSS exhibits superior thermomechanical properties in termsof wearability, thermal stability, oxidation resistance and mechanical strength. The dispersion behav-iour of POSS molecules in a polymer matrix, and consequently the mechanical properties of the result-ing nanocomposites depend strongly on the nature of the R groups. On the other hand, organic–inorganic hybrid thin films based on methacrylates or thermosets reinforced with metal oxidesobtained from different precursors via sol–gel processes have been recently developed [296–298],and their mechanical properties have also been characterized by DSI. The results published on thenanomechanical properties of different inorganic hybrid nanocomposites are summarized in Table 6.

Very large quasi-static E improvements, up to 16-fold increase at 10 wt% loading, were reported forPA-6 nanocomposites incorporating homogenously dispersed polar trisilanolphenyl (Tsp)-POSS [292](Table 6). These results, however, should be contemplated with caution due to the extremely lowreduced modulus obtained for the neat PA-6 (0.137 GPa) as compared with typical values that arein the range 1–3 GPa (Table 4) [223,225,226]. Dynamic storage modulus data provided by the sameauthors were E0 = 1.34 GPa for PA-6 and DE0 = 130% for the nanocomposite with a 10 wt% Tsp-POSScontent. The increment was ascribed to the strong stiffening effect of the robust POSS cages, theirhomogenous dispersion within the matrix and the efficient role played as nucleating agents. Interest-ingly, Tsp-POSS yielded greater improvements in mechanical properties than non-polar octaisobutyl(Oib)-POSS, since the hydroxyl groups of the former nanofiller allowed hydrogen bonding interactionswith the amide linkages of PA6. On the other hand, only 30–40% values were observed for DE0DMA deter-mined using conventional DMA tests, probably due to the preferential location of the fillers on thenanocomposite surface. This morphology could also explain the observed decrease in E0 with increas-ing indentation depth below 300 nm. Further, this ISE effect could also be related to the different cool-ing rate between the nanocomposite bulk and surface during the melt-blending process. In fact, thefastest cooling at the surface should lead to restrained chain mobility and, hence, to a higher modulus.

Strong E improvements have also been observed in PMMA nanocomposites filled with 3.0 wt%graphene-like BN [288], showing a maximum increase of 130% for samples prepared using chloroformas solvent. Moreover, Mammeri et al. found noticeable E enhancements in hybrid PMMA-SiO2 thinfilms prepared by sol–gel processes and exhibiting different interfaces between the polymer and


the silica network [296,297]. In order to extract the mechanical properties, the authors estimated thecreep rate of the material that was used to correct the contact stiffness following Tang and Ngan equa-tion [299]. It was found that the viscoelastic effects influenced the modulus to an extent lower than10%. Hybrid materials, where the organic and inorganic components were linked through covalentchemical bonds (class II) [296], showed better mechanical properties than those where the two com-ponents exchange only weak interactions (class I) [297]. Class I and II materials were synthesized byin situ polymerization of TEOS with PMMA monomers, which were free of or functionalized with tri-ethoxysilane groups, respectively. Hybrid materials of class II were, in a final step, reinforced with SiO2

nanoparticles [297]. It was concluded that the mechanical response was mainly determined by thesize of the hybrid interface, since the lowest mechanical properties corresponded to materials formedfrom silica nanoparticles which exhibited a more defined interface.

On the other hand, moderate DE values (30%) have been reported for epoxy composites incorporat-ing 3.0 wt% IF-WS2 nanotubes [289], being superior to those found for equivalent SWCNT-reinforcednanocomposites. This result is highly interesting since the SWCNTs possess higher strength and elasticproperties than the IF-WS2, and was attributed to the improved dispersion of the inorganic fillerswithin the matrix. In addition, the IF-WS2 nanotubes (70 nm diameter) are expected to restrict themovement of advancing cracks during the deformation process, thus enhancing the compositestrength and toughness, while the SWCNTs (1–2 nm in diameter) are much smaller than the crack-opening displacement and, hence, they are unlikely to cause crack pinning. Chen et al. [290], dealingwith PU composite coatings reinforced with GPTS-modified AlOOH nanorods or nanoparticles onlyfound small modulus increments compared to that of the neat resin, despite the formation of a strongfiller-matrix interface in the presence of the compatibilizing agent. Further, E values for nanorod-filledcoatings were slightly lower than the nanoparticle-filled counterparts, ascribed to the different fillerorientation in the composites since nanorods were horizontally aligned and thus offered lower resis-tance to the normal force applied during the indentation tests, while the isotropically distributednanoparticles were more suitable to bear the load under compression. Interestingly, different defor-mation morphologies were observed in the indents made on these samples (Fig. 25). Large cracks,located along or at an angle to the indenter’s edges, were detected in the nanoparticle-reinforced coat-ing (Fig. 25a), while the nanorod filled counterpart exhibited no sign of chipping (Fig. 25)b, and onlysmall radial cracks were found between contact edges together with minuscule cracks perpendicularto the radial ones (see the arrows marked on the image). Such features could be due to the release ofpressure in the film as the indenter was removed, which allowed relaxation of the deformation causedby buckling in the coating layer under indentation and should lead to significantly higher fracturetoughness in the nanorod-filled coating. Slight stiffness improvements were also reported for otheramorphous polymers like PMMA upon incorporation of Eu:Gd2O3 nanoparticles (Table 6); the modi-fication of the fillers with a silane coupling agent enhanced their dispersion within the polymer,resulting in somewhat increased modulus, in good agreement with data derived from DMA tests [300].

Fig. 25. Indents made of PU-based coating samples filled with GPTS-modified AlOOH nanoparticles (a) and nanorods (b).Adapted from Ref. [290], copyright 2009, with permission from Elsevier.


Contrary to the aforementioned results, some authors have reported a diminution in stiffness uponaddition of inorganic fillers. For example, for PVDF/fluoropropyl (FP)-POSS composites, it was shownthat small contents of filler caused an incomplete nucleation of PVDF and generated big spherical par-ticles that slightly reinforced the matrix, while the incorporation of higher concentrations provoked acomplete nucleation and formed separated and low-crosslinked particles that behaved as soft fillers,reducing the matrix modulus [293]. DSI results confirmed that DE was slightly positive (6%) for a3 wt% loading while at higher loadings, 5.0 and 8.0 wt%, the modulus decrease was about 11 and15%, respectively [293]. Tensile data followed similar trend to that found from nanoindentation. Inparallel, however, DEDMA showed increases for all FP-POSS loadings although after a 27% increase ata filler content of 3%, there was a decreasing tendency. DMA results also revealed that the viscosityof the nanocomposites was much higher than that of the neat polymer and, together with theobserved increasing fracture toughness, made these materials to favour plastic flow and become moreefficient on energy dissipation with increasing addition of FP-POSS.

Analogously, Huang et al. [295] found a decrease in E and H of neat PI with the addition of epoxide-modified POSS, despite the significant increase in crosslinking density due to the reaction between ter-minal amine groups of the matrix and epoxy groups of POSS. This behaviour was attributed to the for-mation of a soft interphase around the POSS molecules in the nanocomposites due to the presence ofeight flexible aliphatic substituents attached to the OH cage. Pelisser et al. [291], in the search of Port-land cement-based materials, studied PDDA/CSH nanocomposites and reported a decrease in E and Hof around 80%, with respect to the CSH values, due to the intercalation of the polymer between theCSH lamellae that resulted in lower nanoparticle packing density.

5.4. Main features on inorganic nanofillers

5.4.1. Layered silicatesDSI experiments on polymer/layered silicates suggest that the largest mechanical enhancements

are associated to nanocomposites with semicrystalline thermoplastic matrices. In contrast to organicnanofillers, nanoclays hinder the development of crystallinity and in spite of this EI and H increasewith increasing filler load. Matrix crystallinity can also change due to temperature gradient effectsduring processing. In addition, local orientation of the nanofiller and/or the matrix can be generatedgiving rise to mechanical anisotropy effects that are readily detected by nanoindentation. Other fac-tors such as the modification of the clay and the interfacial interactions between the polymer matrixand the layered silicates produce a strong effect on the degree of dispersion of the filler and on thestructure of the final layered structures, being the exfoliated ones those exhibiting the best perfor-mance as measured by DSI.

5.4.2. Spherical inorganic nanoparticlesResearch efforts on DSI results concerning inorganic spherical nanoparticles are mainly devoted to

silica reinforced composites. In most cases a continuous increase of the mechanical properties withthe nanosilica content is observed, in spite of the typical high loadings used. Several authors havecompared the experimental results of the modulus with values obtained from different models[75–78]. The lower bound of the Hashin–Shtrikman model [73] and the Halpin–Tsai equations [74]seem to offer the best approximations. Polyacrylate matrices reinforced with silica seem to yieldthe most remarkable properties [210,270] (see Fig. 21), taking into account that the highest mechan-ical reinforcement found for PVA may be an issue for further assessment due to the low modulus of thematrix [78]. Collagen with HA nanofillers also yields remarkable mechanical properties (see Fig. 21).

6. Comparison of the reinforcement effect of different nanofillers

6.1. Effect of the type of nanofiller

Figs. 26–28 compare the reinforcing effect of different carbon and inorganic nanofillers on the DSImodulus of epoxy, glassy and semicrystalline matrices, respectively. To provide a clearer comparison,

0 2 4 6 8 10-20

0

20

40

60

80

100

Δ EΙ [

%]

MWCNT [117]MWCNT [112]MWCNT [111]MWCNT [91]MWCNT [91]MWCNT [108]SWCNT [24]SWCNT [57]GNP [168]GP [166]ND [199]

Cl 93 A [247] A1 100 [248] A1 200 [248]fumed SiO

2 [263]

SiO2 [76]

SiO2 [75]

nanofiller [wt%]

Fig. 26. Comparison of the reinforcing effect of different carbon and inorganic nanofillers on the DSI modulus of thermosetmatrices. Data for composites incorporating the same filler have been represented using identical colour. Open and solidsymbols refer to functionalized and raw fillers respectively. ⁄ indicates nanofiller grafted to the polymer matrix and j denotesaligned nanofiller.

0 2 4 6 8 10-20

0

20

40

60

80

100

120

140

nanofiller [wt%]

PMMA/MWCNT [28]PMMA/MWCNT [28]PMMA/FG [21]PS/CNF [186]PMMA/CNC [195] PS/Cl 15A [231] PS/Cl 20A [232] PMMA/SiO

2 [261]

Δ EΙ [

%]

Fig. 27. Comparison of DEI values for glassy matrices incorporating different nanofillers. Colour code as in Fig. 26. Open symbolsrefer to functionalized fillers. � refers to CNTs functionalized and additionally coated with silica; + denotes electrospun fibres.


Fig. 26 only includes the most representative data regarding epoxy/CNT composites, covering thewhole range of modulus increments reported. The figure shows that limited EI improvements areobserved at all nanofiller contents, most of them lying in the range 0–40%. Especially low are theDEI values at large filler loadings (above 5 wt%). Analogous moderate improvements, or even decre-ments, are observed for the hardness values (see Tables 1–6). Surprisingly, the alignment of nanofillerswith high aspect ratio such as CNTs in the direction of the loading [115] or the grafting of crystallinenanoparticles to the polymer matrix [199] do not result in higher modulus improvements comparedto their random or non-grafted filled epoxy counterparts respectively. These limited mechanicalenhancements have been frequently associated, as discussed in preceding sections, to a poor disper-sion of the nanofiller including aggregation into micro or submicro-domains and/or inhomogeneousdistribution of the blocks, together with weak matrix–filler interaction. These features are commonly

0.0 2.5 5.0 7.5 10.0

0

50

100

150

200Δ E

Ι [%

]PP/MWCNT [106] PA-6/MWCNT [100]PA-6/MWCNT [20]

PVA/MWCNT [103]PLLA/MWCNT [101]PP/GO [154]

PVA/GNP [160] PVA/FG [21]PVA/ND [202]PLLA/ND [187]

PP/PP-g-MA/Cl 15A [218] PP/PP-g-MA/Cl 15A [218] PA-6/I30TC [223] PA-6/1.34TCN [226] PLA/Cl 30 B [235] PP/IF-WS

2 [22]

nanofiller [wt%]

Fig. 28. Effect of carbon and inorganic nanofillers on EI of different semicrystalline thermoplastic matrices. Colour code as inFig. 26. Solid symbols designate raw fillers. The rest of symbols refer to functionalized fillers. � and � denote composites withOMMT prepared by SSE and TSE, respectively.


encountered in all types of filled epoxy materials and possibly arise as a consequence of an inadequatedispersion of the filler in the blend and the latter interfering in the curing process of the epoxy matrix.Some authors have suggested that the nanofiller influences the cross-linking density in the epoxy[197,270], especially in the surroundings of the matrix-filler interface [247], hence, affecting the stiff-ness and hardness of the matrix. From the above considerations it seems that the intrinsic propertiesof the filler are not fully exploited and this turns up in similar matrix reinforcements for fillers thatsignificantly differ in mechanical properties. For example, the addition of 2D graphene platelets[166,168] results in analogous modulus enhancements compared to the incorporation of similaramounts of layered silicates [247,248] (see Fig. 26). It can also be observed that the largest mechanicalincrements are reported for CNT-reinforced epoxies. However, it is not clear whether this observationis related to the development of improved strategies for filler dispersion and filler–matrix interaction.In fact, the largest mechanical increment in Fig. 26 (DEI = 100% for 5 wt%) has been explained as aris-ing from a minimum interference of the filler with the curing process of the epoxy through an ade-quate selection of the solvent employed for CNT pre-dispersion [111].

Fig. 27, in turn, gathers DEI values of glassy matrices incorporating different nanofillers. Repre-sented data are restricted to PMMA and PS because these two matrices are the only examples inthe literature offering indentation data on nanocomposites using more than one type of filler. In caseof PS, the addition of clays or carbon nanofibres produces small increments of the mechanical proper-ties possibly due to agglomeration of the filler [198,243,244]. Most interesting are the data on PMMAwhere a wider collection of nanofillers can be compared. As shown in Fig. 27, few-layer graphene andsilica-coated CNT seem to be most effective. In both cases, the nanofiller is functionalized to improvethe filler-matrix interaction. Compared to CNTs, graphene exhibits a large surface area that promotesthe interaction between the oxide groups of the acid-functionalized graphene sheets and those of theglassy matrix [21]. In fact, functionalized CNTs have been shown to produce quite limited reinforce-ment on PMMA (see Fig. 27). It has been suggested that the curved morphology and inferior bendingproperties of CNTs could account for the reduced reinforcement action [28]. Only when CNTs arecoated with silica shell a significant reinforcing effect comparable to that of graphene is detected. Nev-ertheless, it is noteworthy that lower amounts of FG are needed with respect to silica-coated CNT toproduce a similar reinforcement effect on PMMA.

Finally, Fig. 28 illustrates the effect of carbon and inorganic nanofillers on EI of different semicrys-talline thermoplastic matrices. Only those matrices using different types of filler are included in thefigure. It is noteworthy that functionalization of the filler is a common route for all nanocompositesrepresented. In the glassy matrices of Fig. 27, as discussed above, the intrinsic properties of the filler


and the distribution and interaction with the matrix have been demonstrated to be the key features toachieve enhanced properties. In the case of semicrystalline matrices, the nucleating effect of the filleris additionally expected to contribute to the mechanical enhancement through the changes triggeredin the matrix nanostructure and via the nucleation of crystals at the interface. To what extent eachaspect contributes to the overall mechanical enhancement is usually not approached in most inden-tation studies, however, some valid analyses have been done in this direction [22].

In spite of the difficulties in comparing indentation data from different sources, clear trends can befound in Fig. 28. First, one can see that DEI conspicuously increases with increasing filler content up to�1 wt% where the rate of increase diminishes or levels off. It is also clearly seen that carbon-based fill-ers such as CNTs, graphene and nanodiamond tend to lie in the upper part of the plot, especially forfiller contents above 1 wt%; hereafter, D EI still rises at a significant rate [20,100,101,160,187]. In con-trast, data for clays and inorganic particles are mainly located in the lower part of the plot and exhibita tendency to remain constant above 1 wt% [22,218,235] although there are also examples where EI

further increases at higher loadings [223,226].All carbon-based fillers seem to act as reinforcement to approximately the same extent within the

limitations inherent to the comparison of data from different sources. For instance, DEI values for PVAreinforced with CNTs, graphene-like fillers and ND lie within the same range [21,103,160,202]. More-over, PA6, PLLA and PP provide examples in which the reinforcing effect of carbon-based fillers andinorganic particles can be clearly discerned. In case of PA6, the reinforcing effect of MWCNT is signif-icantly higher at all filler loadings than that of clays (see Fig. 28 [20,100,223,226]). Moreover, incorpo-ration of MWCNT or ND to PLLA produces a marked reinforcement compared to the one originated inPLA with the addition of clay [101,187,235]. For PP, higher enhancements are achieved when grapheneoxide is used with respect to organically modified MMT and inorganic fullerene-like nanoparticles[22,154,218]. Data for PP/MWCNT composites exhibit too large scatter (see Fig. 28) and hence, havebeen not taken into account in the discussion. A detailed study on PP/IF-WS2 composites has shownthat enhancements associated to the filler itself contribute more significantly to the overall reinforce-ment than those related to changes induced in the PP morphology. The latter changes, however,account for the initial drop of mechanical properties at the lowest filler content [22].

6.2. Influence of the geometry of the filler and orientation effects

The aspect ratio of the nanofiller itself does not seem to introduce an additional value to themechanical properties of reinforced polymers by means of indentation. This appears to be quite rea-sonable for polymer matrices incorporating randomly distributed nanofillers taking into account thatthe applied stress during an indentation test evolves in a triaxial manner. Anisotropy of the fillers isnot exploited because the force is not applied along the direction of higher stiffness. Indeed, thereis no evidence that carbon-based fillers with different geometries such as ND, CNTs and grapheneexhibit different reinforcing effects that can be attributed to the geometry when they are randomlydistributed in a polymer matrix. On the contrary, in case of CNTs, it has been argued that the cylindri-cal shape can be detrimental for the indentation properties due to the curved morphology and inferiorbending properties of the nanotubes [28]. However, the anisotropic shape of the filler can be used toenhance the filler–matrix interaction. For example, the large surface area of graphene sheets can beused to create a large number of active sites to promote the interaction with the matrix [21].

In the case that the anisotropic filler exhibits some preferential orientation in the composite mate-rial, mechanical properties are expected to be dependent on the relative orientation of the compositewith respect to the indentation loading direction. The application of the load perpendicularly to thedirection of highest stiffness in whiskers of cellulose nanocrystals has been proposed as one of the pos-sible reasons for the minor mechanical enhancements encountered in CNC-reinforced PMMA fibres[207]. Most strikingly are the indentation results reported in epoxy composites including CNT ‘‘for-ests’’ where application of the load along the vertical direction of the aligned CNTs only revealed mod-est modulus improvements [115–117]. Again, CNT waviness was suggested as the major factoraccounting for the limited stiffness enhancements.

Graphene represents the best example where mechanical anisotropy seems to be fully exploited.The planar orientation of graphene sheets in a composite material has been proved to be a most


successful route in providing enhanced mechanical properties [158,173]. Indentation studies on com-posites of self-aligned ultralarge-size rGO sheets and PU, and graphene/polymer multilayer assemblieshave demonstrated that modulus and hardness values can significantly benefit from the planar orien-tation of graphene. However and as several times stated throughout this review, the reader should beaware that such achievements could be questionable due to the very low modulus values reported forthe neat PU matrix.

6.3. Comparison of properties at the macro and the nanoscale

Fig. 29 illustrates the plot of the modulus values obtained by means of indentation, EI (E or E0), as afunction of those determined by means of macroscopic techniques, Emacro (Y or E0DMA). A colour codehas been used to distinguish between dynamic DSI or quasi-static DSI data vs. macroscopic tensileor DMA measurements. Some correlation between EI and Emacro can be distinguished that allows estab-lishing clear limits to the EI/Emacro ratios most commonly encountered. This observation is purely phe-nomenological, as extensively discussed in Section 2.3. The figure shows that most data lie within thearea defined by two dotted lines, EI = 2Emacro and EI = Emacro/2, that are drawn as a guide to the eye inthe figure. Only two series of EI � Emacro data lie outside this area. One of the series is associated to PS-based shape memory polymers reinforced with carbon black [206], exhibiting very low E0DMA values, infact the lowest of all the Emacro data included in Fig. 29. As discussed in Section 4, this can produce rel-evant consequences in the determination of the point of initial contact during an indentation testintroducing important inaccuracies in the indentation data. On the other hand, it is difficult to ascer-tain the possible reason for the significant deviation of the series of PAN/ND data from the generaltrend because only limited information on the indentation experimental procedure and method ofanalysis is available [208].

It can also be observed from Fig. 29 that the correspondence between E from quasi-static indenta-tion and Y from tensile testing (blue symbols) is fairly good, most of the data lying in the line EI = Emacro

(also drawn in the figure). All the nanocomposites in this group of data include randomly distributednanofillers except the one constituted by self-aligned ultralarge-size rGO layers incorporated into PU.In the latter case, it is interesting to note that the modulus of indentation achieved with the graphenesheets perpendicular to the load direction seems to be in agreement with tensile testing data obtained

0 2 4 60

2

4

6

8

EΙ [

GP

a]

Emacro

[GPa]

Epoxy/MWCNT [91]Epoxy/MWCNT [110]

Epoxy/MWCNT [112]Epoxy/SWCNT [24]Chitosan/MWCNT [98]PA-6/MWCNT [100]

PA-6/MWCNT [20]UHMWPE/MWCNT [105]

Epoxy/GP [166]Epoxy/GNP [169]Pu/rGO [158]Epoxy/CNF [197]Epoxy/CNF [198]PS/CNF [198]PVA/CNC [200]PS/CB [206]PAN/ND [208]

PA-6/I30TC [223]PA-6/CoAl-LHD [224] PA-6/1.34TCN [226]PLA/Cl 30B [235]PHBV/Cl 30A [237]Epoxy/Cl 93A [247]Epoxy/fumed SiO

2[263]

Epoxy/SiO2 [75]

PMMA/SiO2 [273]

PP/IF-WS2 [22]

PPS/IF-WS2 [22]

Fig. 29. DSI moduli EI (E or E0) vs. values determined by means of macroscopic techniques Emacro (Y or E0DMA) for polymernanocomposites reinforced with different nanofillers. Open and solid symbols refer to functionalized and raw fillersrespectively. + as in Fig. 11. The dashed lines represent EI = 2Emacro, and EI = Emacro/2, while the dotted-dashed line corresponds toEI = Emacro. Symbols in blue, green, magenta and black correspond to E vs. Y, E vs. E0DMA , E0 vs. Y and E0 vs. E0DMA, respectively.


directing the load along the graphene plane [158]. The majority of the EI data in blue colour have beenanalysed using Oliver and Pharr’s method [100,110,112,223,226,247,263]. Only one paper uses a dif-ferent method of analysis consisting in applying Hertz to the loading curve [200]. It seems that thelatter procedure yields somehow higher EI/Emacro values approaching 2. According to the availabledata, it is then suggested that OP analysis offers modulus values in the range of those determined fromtensile testing. This contention is most probably mediated by the fact that all the polymer matricesincluded in this group of data (epoxies, PA6, PVA, PAN) exhibit E > 1 GPa and a limited viscous char-acter. Indeed, a recent paper examined the validity of Oliver and Pharr’s analysis to polymer materialsand suggested that this method can yield consistent results for polymers with glass transition temper-ature well above the temperature of measurement [301].

Fig. 29 additionally demonstrates that in spite of the large scatter detected in the plot of E0 vs. Y(magenta symbols), data are found to fluctuate around the line defined by EI = Emacro. This is quite sur-prising taking into account the technical dissimilarities between dynamic indentation and quasi-statictensile testing. The reason for this behaviour is probably grounded on a similar argument to that pre-viously discussed: the limited rubbery flow character at room temperature of most polymers matricesrepresented. Accordingly, plotting E0 vs. Y or E vs. Y should look quite similar.

At the present stage, no clear correspondence is found between modulus from indentation andthose of DMA. The large scatter of data in the representation of E0 vs. E0DMA (black symbols inFig. 29) does not allow drawing any significant tendency other than that the EI/Emacro ratio can takeany value between 0.5 and 2. On the other hand, modulus values from quasi-static DSI vs. those ofDMA (green symbols) tend to lie along the line EI = 2Emacro. However, the amount of data is so limitedthat this contention should be only taken as a hint for further research.

7. Conclusions and future perspectives

An effort has been done to gather indentation data on polymer nanocomposites. This is quite achallenging task. On the one hand, DSI devices are too often employed as routine instrumentationwithout a basic knowledge of the physics underlying. Moreover, the methods developed for the deter-mination of the mechanical properties from DSI curves rely on a number of assumptions that are oftenoverlooked. As a consequence, large errors in the determination of the modulus and hardness can beintroduced. This has been known for over more than two decades and yet, too many published papersreport values for mechanical properties that at first sight are recognized to be incorrect. On the otherhand, we have frequently observed that many authors offer little experimental information, in such away that an adequate evaluation of the relevance of the results is not feasible. It is noteworthy that themost outstanding reinforcing effects are commonly related to matrices with very low modulus values.This observation applies to quite different fillers and polymer matrices, e.g., carbon nanotubes andPHO (E = 0.12 GPa) [97] or UHMWPE (E = 0.6 GPa) [104]; graphene-like fillers and PU (E = 0.4 GPa)[158] or PVA (E = 0.2 GPa) [160]; a combination of different carbon fillers and PVA (E = 0.7 GPa)[121]; clays and PU (E = 0.01 GPa) [252]; spherical nanoparticles such as nanoHA and collagen(E0 = 0.20 GPa) [286] or silica and PVA (E = 0.30 GPa) [78]. Results on low modulus materials need a care-ful consideration, in view of the well-recognized technical difficulties. Some useful recommendationsare given in Refs. [302,303].

In spite of all the limitations outlined, we think that the following general conclusions can bedrawn from the present review.

Indentation data on epoxies clearly reveal quite limited modulus and hardness increments withrespect to thermoplastic materials. Functionalization does not seem to play here a relevant role. Ithas been suggested that the limited nanofiller dispersion in the blend, together with the interferenceof the filler with the curing chemistry are key parameters affecting the final properties. In thermoplas-tic materials, conspicuous mechanical property enhancements with increasing filler content followedby a levelling-off at loadings of 1–2% have been observed as a general performance for most nanofil-lers. The state of distribution and dispersion of the filler is usually the basis of this behaviour. At thispoint, a significant difference is found between raw and functionalized fillers, the latter usually exhib-iting larger enhancements and increasing their positive difference with the raw ones at high loadings.


In case of semicrystalline thermoplastic polymers, the influence of the filler on the nanostructure ofthe matrix has also been shown to be of great importance for a full understanding of the reinforce-ment. Creep results, that can be easily evaluated using depth-sensing instrumentation, provide anadditional source of information on the time-dependent mechanical behaviour of nanocomposites.Most indentation studies show an enhanced creep resistance with the incorporation of the filler.

Hitherto, the shape of the nanofiller does not seem to play a clear role on mechanical improvement,as measured by means of DSI. However, it should be noted that a limited amount of data are available.One should recall that the compressive stress field developed underneath the indenter is approxi-mately radial from the point of first contact. This setup does not favour orientation of anisotropicnanofillers in the direction that yields their highest stiffness, as could be the case in uniaxial tensiletesting, thus limiting the potential benefits of nanofiller anisotropy. Nevertheless, in some cases theanisotropic shape of the filler promotes mechanical reinforcement because an enlarged surface areacan accommodate a higher number of interaction sites with the matrix.

Most important, nanoindentation provides worthy information on the local mechanical propertiesat the submicroscale. Uneven distribution of the filler or the matrix morphology can be readilydetected by means of DSI [111,223,229,262]. At shallow depths (below 200–300 nm), a large numberof studies report remarkable indentation size effects that are commonly discussed in the light ofinstrumental issues such as: area calibration, tip blunting, frictional forces between the indenterand sample or determination of the point of initial contact [20,98,99,126,169,174]. For indentationdepths exceeding h � 200 nm, the continuous decrease of modulus and/or hardness values has beendiscussed as a consequence of a morphology gradient across the sample thickness [111,229]. On theother hand, the constancy of EI and H values over a broad range of indentation depths has been inter-preted as indicative of the good dispersion of the filler throughout the matrix [114]. For nanofillerssuch as graphene where at least one dimension approaches a few micrometers, DSI can clearly mon-itor the continuous change of the mechanical properties as the indenter penetrates into the surfaceand finally converge towards the mean value of the material at large penetration depths [166].

Finally, a phenomenological correspondence has been found between indentation modulus (eitherquasi-static or dynamic) and Young’s modulus values from tensile testing. Such a fair correlation issomewhat unexpected due to evident experimental differences (e.g., loading direction, scale of volumedeformed, occurrence of material heterogeneities). A larger data scattering is observed when compar-ing EI with DMA measurements. In this case further research is needed taking into account the limiteddata available.

What should we expect for the next decade? It is clear that the field of polymer nanocompositeswill continue growing with the incorporation of new nanofillers and the development of complexhybrid and hierarchical materials to be applied as thin films, coatings, multilayers or bulk. The DSItechnique, although still in its infancy, has been shown to be an effective tool to characterize themechanical properties of polymer nanocomposites, particularly when there are limitations in sizeor arrangement of the specimen nanocomposite, or when a local probe of the mechanical propertiesis required. These features will permit to address the new challenges presented by innovative mate-rials, especially from the point of view of understanding the nanocomposite heterogeneity, on the oneside, and the macroscopic mechanical properties on the other.

We have noticed that little research has been devoted to the correlation between nanoindentationresults and structural information that can be extracted by other techniques. Usually, DSI data of poly-mer nanocomposites are correlated to microscopic studies but no further structural analysis is offered.This is indeed an area that still needs to be developed, as it is the analysis of creep or, in general, thetime-dependent behaviour. The integration of nanoindentation with real-time techniques such aselectron imaging, X-ray scattering, electrical and thermal characterization among others, representsone of the most attractive challenges for the coming years in polymer nanocomposites and materialsscience [304]. The spread use of dynamic indentation, i.e. nanoDMA, in a wider range of frequenciesand temperatures should also be encouraged. We can anticipate that at short term, dynamic DSI willbecome a convenient technique for recording information on the relaxation modes of polymer spec-imens in the nanoscale range. Some research has already been done in this direction [305]. It is alsonoteworthy that few studies employ indentation testing to assess the mechanical properties of inter-phases. Improving the force and depth resolution of the current depth-sensing instrumentation will


indeed foster the research in this area. In parallel to the rising concern on environmental problems,bionanocomposites are likely to be increasingly used and the investigation of their mechanical perfor-mance in comparison with traditional non-biodegradable composites will certainly be of interest[306]. Taking a step forward, the application of DSI to biological samples, that are hierarchical nano-composites by nature, is also an area of great potential. Hard biological materials such as bone struc-tures and dentine are easy to measure by indentation techniques because modulus values lie in theGPa range [307–309]. On the contrary, indentation measurements on soft polymer biomaterials andtissues, with modulus values in the MPa or kPa range represent quite a challenge [310,311]. Nanoin-dentation on soft materials has been frequently carried out by means of AFM because of the ease ofuse of fluid cells for experiments with hydrated specimens, together with superior displacementand force sensitivity [312–314]. Nevertheless, the analysis of AFM data faces severe difficulties relatedto detection of initial contact, force-depth calibration and tip geometry [315]. With respect to AFM,DSI represents a robust method arriving from the submicron scale with well-established calibrationmethods that provide reliable mechanical properties. In addition, some DSI instruments have beenclaimed to account for the necessary stiffness sensitivity to properly measure the properties of mate-rials as soft as 10 kPa in modulus [316]. Other commercial nanoindenters incorporate a force feedbackcontrol that ensures an accurate measurement of the load (and not of the stiffness) as a means ofimproving the surface detection. Initial DSI studies on hydrated biological samples, with modulus val-ues in the kPa range, highlight the potential of the technique to the study of biological nanocomposites[316,317]. The precise knowledge of the viscoelastic properties of biological systems in a controlledtest environment reproducing in vivo conditions can help, for instance, to understand the mechanismsinvolved behind a pathological behaviour or to fabricate biological replacement parts [311]. To prop-erly extract the true elastic modulus of very soft materials, such as biological systems, data analysescapable of separating first the influence of external effects and then the non-linear viscous compo-nents from the linear ones should be employed [318].

An effective use of DSI for soft biological materials will require exploiting both the full potential ofthe new generation of nanoindenters with improved depth and force resolution and the mostadvanced analysis routes.

Acknowledgments

The authors wish to thank the MICINN (Ministerio de Ciencia e Innovación), Spain, for financialsupport under the Grants MAT2010-21070-C02-01 and FIS2010-18069. AD would like to thank theCSIC for a JAE postdoctoral contract cofinanced by the EU.

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