Depth-sensing indentation involves applying a specific force-time history on a rigid indenter while continuously monitoring the displacement of the indenter into the surface. Frequency specific depth-sensing indentation testing entails adding a small harmonic force on the indenter and measuring the harmonic response of the indenter at the excitation frequency. While often taken for granted, understanding the dynamics behind these frequency specific measurements is of vital importance in the determination of quantitative mechanical properties. This paper will focus on the dynamics of a variety of depth-sensing indentation systems and how these dynamics affect such parameters as detecting the point of surface contact, environmental sensitivity, dynamic frequency range, and the range over which contact stiffnesses and moduli can be accurately measured.

To determine the mechanical properties of materials, using spherical ultra-microindentation techniques, the geometry of the part of the indenter in contact with the surface must be well known. However, crystallographic anisotropy ensures that nominally spherical diamond indenters with very small tip radii will not be truly spherical. This paper examines the shape of typical indenters, with nominal tip radii of 5 μm and 10 μm. Atomic force microscopy techniques are used to determine the general form and then using a technique based on Hertzian contact mechanics in combination with load/penetration data obtained by indentation, to determine the effective radius of the part of the indenter in actual contact with the surface.

The mechanical properties of thin films can be measured by a variety of different techniques, with nanoindentation being one of the most recent developments in this growing field. By using a depth-sensing indentation method it is possible to obtain quantitative values for the hardness and modulus, and thus gain better insight into the response of a material to controlled deformation at such small scales. However, previous work [1] has shown that the effects of pile-up, particularly in soft films deposited on hard substrates, can produce significant overestimation of the hardness and modulus due to an underestimation of the true contact area by common nanoindentation analysis procedures. By measuring the topography of the residual indent using Scanning Force Microscopy (SFM) and combining this information with the indentation data, it is possible to gain a fuller understanding of the indentation method and its effects on the material being tested. In addition, the true contact area can be directly measured from the SFM images and subsequently used to calculate the hardness of the material more accurately. Experimental results are presented for a selection of soft films on hard substrates where SFM analysis of indentations at varying depths gives significant additional information concerning the true response of the system to instrumented indentation at a nanometric scale. Pile-up effects can be precisely monitored as a function of depth and correlated to hardness variations encountered across the coating/substrate interface.

In indentation literature, there is support for the proposition that the extents of the elastic and plastic fields scale with the contact radius. For example, in Hertzian contact, the locations of constant-shear-stress lines scale with the contact radius. This idea has significant consequences with respect to measuring substrate-independent properties of thin films. If it is the contact radius that determines the extent of the elastic and plastic fields, then any general rule expressed in terms of the indentation depth is only appropriate for one indenter geometry - the one for which it was determined. In this work, three different indenter geometries were used to measure the hardness of one thin film. For all geometries, hardness results are expressed in terms of both depth and contact radius

The spherical indentation process has been modeled by the finite element method to study the pile-up behavior of elastic-plastic materials with different degrees of strain hardening. A wide range of materials was examined characterized by different elastic moduli, yield stresses, and strain hardening exponents. The geometry of the contact impressions was examined in both the loaded and unloaded conditions. Results show that pile-up behavior in elastic-plastic materials cannot be related solely to the strain hardening exponent, as has often been done in the past. Relating the pile-up behavior strictly to the strain hardening exponent may lead to significant errors in the calculated contact area for materials with modulus to yield stress ratio lower than 1000. In addition, an assessment of the combined influences of material properties and stage of development of the contact impression on the pile-up behavior is described.

The Sneddon solution, as it is implemented in the Oliver-Pharr method, deviates from the indentation experimental data in a manner which depends on both the indenter angle and the Poisson ratio of the sample. These effects are demonstrated experimentally by performing indentations in tungsten and aluminum using a cube-cube comer indenter where the effects are exacerbated by the small indenter angle. The first objective was to experimentally support and validate an approximate analytical solution in conjunction with finite element simulations which illustrate the Poisson ratio and indenter angle effects. Second, a review of data analysis procedures is presented which leads to a better understanding of the systematic errors which percolate through in the measurement of Young's modulus and hardness.

Indentation at the nanometer scale using a Berkovich indenter is a widely used technique for measuring mechanical properties of materials. Other indenter geometries may find application in assessing elastic contact, the mechanism of fiber push-out in composites or fracture toughness. This study undertakes a comparison between the various indenter types, explores their suitability for acquiring elastic modulus, hardness and fracture toughness data, and discusses some limitations and cautions based on the data obtained with a Nanoindenter® and a Topometrix® scanning force microprobe.

We have begun to explore the use of continuous stiffness techniques for studying the mechanical properties of coated systems using nanoindentation and thus to critically assess the differing types of sample information available from simple load-displacement (P-δ) curves, P-δ2 analyses and P-S2 (i.e. load-vs-(contact stiffness)2) data. We have also examined whether the small superposed AC signal used to continuously measure contact stiffness has had any marked effect on the sample load-displacement response. A range of fully calibrated Nano Indenter II and Nano Indenter XP instruments have been used for data acquisition, while transmission electron microscopy (TEM) and high resolution scanning electron microscopy (HRSEM) have both been employed to characterise the resultant indentation appearances and deformation structures. Samples have spanned a number of monolithic and coated materials systems including fused silica, 6H-SiC, 3C-SiC on Si and TiN on both M2 and 304 steels. Plots of P-δ and P-δ2recorded with and without the additional AC signal were found to be indistinguishable. Further, apart from some small statistical variation in the loads at which crack initiation was observed in a SiC-on-Si coated system, TEM and HRSEM studies showed no detectable differences in deformation substructures between nominally identical indentations made with and without the AC signal.

While the parameter P/S2 is independent of detailed tip shape and is a constant with displacement for monolithic systems, it was found to display unexpected variations with displacement for the coated systems. In order to explore the origins of these variations, we have examined plots of system elastic modulus (Es) and system hardness (Hs) as a function of indenter displacement which have shown that the effective contact elastic modulus decreases more rapidly with increasing contact scale than does the effective system hardness P/S2 can also be related to the plasticity index (ψ) widely used to describe the balance between the elastic and plastic responses of materials subjected to contact damage. Thus, the observations of maxima in plots of P/S2 with displacement, for at least some coated systems, suggests that there may be an optimum contact scale for maximising the elastic contribution to the contact response of such systems

The early stages of plasticity have been studied by nanoindentation at room temperature, with a new apparatus, in GaAs and MgO single crystals. On both materials a very well defined pop-in has been observed, whose characteristics depends on the initial dislocation density in the indented zone. On (100) cleaved surfaces of MgO crystals, several special surface features, including the presence of rosette arms, have been observed by SFM. Starting from these observations a dislocation mechanism has been proposed.

Defect structure beneath and near Vickers indentations made with loads of 10, 25 and 50g in Si has been studied in detail by TEM. Both plan-view and cross-sectional observations have been made. Beneath the 10 and 25g indentations an amorphous phase is formed, but beneath the 50g indentation no amorphous phase is formed.

In order to study the plasticity in Cu-Ni multilayers deposited on single crystals of Cu and etched Si, we have investigated the mechanical properties of the multilayers both by nanoindentation measurements, and by the transmission of well characterized dislocations from the underlying substrate by tensile deformation of Cu single crystals.

Various multilayers were deposited by physical vapor deposition with layer thicknesses varying between 1000 and 20 Angstroms (for a total thickness between 0.8 and 1 μm). Two types of experiments were designed. The first one aimed at injecting, in a controlled way, some dislocations from the substrate into the multilayers; the second type of experiment concerned the structure of the multilayer surface after having plastically pushed the material away from a nanoindenter.

This communication reports the results from the nanoindentation measurements, as well as the observations of slip on the surface. We observed through the injection of dislocations by nanoindentation that the multilayers increase in strength with refinement of the layer structure but at thicknesses below 35Å exhibits a softening behaviour. Also observation of the upheaval around the nanoindent showed an evolution from slip lines to more spread plasticity with refinement of the layer structure.

The transition from purely elastic to plastic deformation during the initial stages of indentation is explored. The yield point is dependent on crystal orientation, and the extent of the transition from elastic to plastic deformation is compared to a model utilizing a superdislocation to accommodate deformation based on a change in shear stress before and after the yield point. The thickness of an oxide film is shown to impact, but not be solely responsible for, this phenomenon. High dislocation densities limit, and eventually eliminate, the ability to sustain elastic loading. Finally, the yield point is clearly both stress and time dependent; yielding occurs when a tip is held at a constant load over a period of time ranging from seconds to minutes.

The yield point phenomenon observed in nanoindentation of metallic single crystals has been attributed to dislocation nucleation. The event is preceded by elastic reversible indentation followed by a rapid increase of indenter tip displacement. A similar event has been seen in epitaxial Co films on GaAs substrates. As these films are on the order of 10 nm thick, the substrate plays a significant role in dislocation behavior as dislocations are emitted at the nucleation site. Epitaxial Fe films 23 and 100 Å in thickness have been grown on GaAs and GaAs with a Sc0.3Er0.7As interlayer. Nanoindentation experiments have been performed using the Hysitron Triboscope, an add-on device to an atomic force microscope. The influence of film thickness, epitaxial interlayer, and biaxial misfit stress on the yield point behavior in the overlayer will be considered.

A new method for hardness–depth profiling on the nanometer scale has been developed. In this method, the hardness–depth profile was obtained by a stepwise removal of a surface layer by ion sputtering and measurement of the hardness by nanoindentation at each step using the same maximum load (250 μN). During removal of each surface layer the composition was obtained using Auger electron spectroscopy. For comparison, the hardness–depth profile was also assessed by a stepwise increase of the maximum load applied to the original surface. The sample used in these experiments was a Ti–6Al–4V cast alloy, doubly implanted with carbon (120 kV/3 · 1017 cm−2 and 60 kV/4 · 1017 cm−2). The hardness–depth profiles were related to the composition–depth profile. It followed that only the constant load and removal method provides a direct relation between hardness and composition at a given depth.

In recent literature, the ratio of the residual depth of indentation to the total depth, hf/hmax, has been shown to be a useful parameter in predicting constraint factor and pile-up during indentation in bulk materials. In this work, hf/hmax is shown to have a simple one-to-one relationship to Sl/S, or the ratio of the slope of the load vs. displacement curve when loading to the elastic stiffness of the contact. Experimentally, the parameter Sl/S is attractive because unlike hf/hmax, it can be determined continuously during loading, using only data near the depth of interest. The elastic stiffness, S, is measured using a forced-oscillation technique, thus allowing Sl/S to be determined as a function of depth. Previous predictions for constraint factor and pileup are recast in terms of Sl/S.

Experimental equipment has been developed consisting of a compression machine interfaced with an atomic force microscope. This equipment is particularly suited to the analysis of fine slip line structures on a nanometer scale and to follow in situ plastic flow from the emergence of the first few dislocations from the bulk crystal to the stages of work hardening. The motivations, instrumentation and results on the γ phase of Ni based superalloy are described.

As part of a programme investigating whether slip (by dislocation motion at some critical resolved shear stress) or densification (by structural collapse at some critical hydrostatic pressure) dominates the plastic deformation response of open-crystal-structured ceramics having the diamond, zinc blende or simply-related structures, transmission electron microscopy (TEM) and high resolution scanning electron microscopy (HRSEM) are being used to characterise the deformation structures in, and around, nanoindentations made over a range of loads in single crystal samples of Si, Ge, SiC and various III-V and II-VI compound semiconductors. Since SiC is believed to lie close to the boundary between those materials which slip and those which densify as primary plasticity responses to contact with Vickers and Berkovich indenters, this part of the study has focused on establishing the deformation mechanisms of 6H-SiC during nanoindentation.

TEM of nanoindentations in (0001) 6H-SiC samples has established that dislocation slip is indeed the sole mechanism of plastic deformation from the nucleation of a few dislocation loops - at or near the theoretical strength - to extensive dislocation plasticity. Also, HRSEM observations have revealed slip steps of limited extent in 6H-SiC samples with more extensive slip steps arrays found in all the other compounds. By contrast, both Si and Ge show evidence of heavily deformed - and sometimes extruded - material believed to be a characteristic of structural collapse accommodating at least some part of the indentation strain before any dislocation slip occurs.

A dependence of elastic response on the stress-state of a thin-film has been demonstrated using the interfacial force microscope (IFM). Indentation response was measured as a function of the applied biaxial stress-state for 100 nm thick Au films. An increase in measured elastic modulus with applied compressive stress, and a decrease with applied tensile stress was observed. Measurements of elastic modulus before and after applying stress were identical indicating that the observed change in response is not due to a permanent change in film properties.

Nanoscale indentation experiments were performed on Au using the Interfacial Force Microscope (IFM) in an indentation mode. The indentation modulus and yield point were measured for three orientations of the Au surface: (111), (110), and (001). The indentation modulus for the (111) surface was found to be 36% greater than for the (001) surface and only 3% higher than the (110) surface. Additionally, the yield point was found to vary between orientations, but the shear stress resolved on {111} slip planes beneath the indenter on the axis of symmetry was found to be approximately 1.8 GPa for yield points of all three orientations.

Nanoindentation is a method of measuring the mechanical properties of materials that cannot be tested by conventional techniques. A method for determining yield stress from indentation has been developed by Huang and Gerberich based on Johnson's spherical cavity model. The method requires only a knowledge of the peak load and corresponding plastic zone radius to determine yield stress. This has previously been demonstrated for large scale indentations, but not for nanoindentation where assumptions such as isotropy and homogeneity may no longer be valid. Nanoindentation experiments have been performed on a series of cold rolled and annealed OFHC copper samples. Different annealing temperatures resulted in a range of initial yield stresses which were measured by conventional tensile testing. Electropolished samples were then tested using the Hysitron Triboscope (an add-on device to an atomic force microscope) which allowed for collection of both indentation and image information from which load and plastic zone radius could be measured. Comparison of yield stress measurements taken from nanoindentation and tensile tests shows an agreement to within 10% for initial yield stresses as low as 190 MPa and as high as 360 Mpa.