Department of Physics, Chemistry and Biology, IFM

Master’s Thesis

Nanoindentationin situ a

Transmission Electron Microscope

Lars JohnsonLITH-IFM-EX--07/1732--SE

Department of Physics, Chemistry and Biology, IFMLinköpings universitet

SE-581 83 Linköping, Sweden

Master’s ThesisLITH-IFM-EX--07/1732--SE

Nanoindentationin situ a

Transmission Electron Microscope

Lars Johnson

Supervisor: Finn GiulianiIFM, Linköpings universitet

Examiner: Lars HultmanIFM, Linköpings universitet

Linköping, 7 March, 2007

Avdelning, InstitutionDivision, Department

Thin Film Physics DivisionDepartment of Physics, Chemistry and Biology, IFMLinköpings universitetSE-581 83 Linköping, Sweden

DatumDate

2007-03-07

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� Svenska/Swedish

� Engelska/English

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�

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� Licentiatavhandling

� Examensarbete

� C-uppsats

� D-uppsats

� Övrig rapport

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URL för elektronisk versionhttp://urn.kb.se/resolve?urn=urn:nbn:se:liu:diva-8333

ISBN

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ISRN

LITH-IFM-EX--07/1732--SE

Serietitel och serienummerTitle of series, numbering

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TitelTitle Nanoindentation

in situ aTransmission Electron Microscope

FörfattareAuthor

Lars Johnson

SammanfattningAbstract

The technique of nanoindentation in situ a transmission electron microscope has been imple-mented on a Philips CM20 instrument. Indentations have been performed on Si and sapphire(α − Al2O3) cut from wafers; Cr/Sc multilayers and Ti3SiC2 thin films. Different sam-ple geometries and preparation methods have been evaluated. Both conventional ion andfocused ion beam milling were used, with different ways of protecting the sample duringmilling. Observations were made of bending and fracture of samples, dislocation nucleationand dislocation movement. Basal slip was observed upon unloading in sapphire. Dislocationmovement constricted along the basal planes were observed in Ti3SiC2. Post indentationelectron microscopy revealed kink formation in Ti3SiC2 and layer rotation and slip acrosslayers in Cr/Sc multilayer stacks. Limitations of the technique are presented and discussed.

NyckelordKeywords nanoindentation, electron microscopy, mechanincal deformation, thin films, MAX phases,

functional ceramics

AbstractThe technique of nanoindentation in situ a transmission electron microscope hasbeen implemented on a Philips CM20 instrument. Indentations have been per-formed on Si and sapphire (α − Al2O3) cut from wafers; Cr/Sc multilayers andTi3SiC2 thin films. Different sample geometries and preparation methods havebeen evaluated. Both conventional ion and focused ion beam milling were used,with different ways of protecting the sample during milling. Observations weremade of bending and fracture of samples, dislocation nucleation and dislocationmovement. Basal slip was observed upon unloading in sapphire. Dislocationmovement constricted along the basal planes were observed in Ti3SiC2. Postindentation electron microscopy revealed kink formation in Ti3SiC2 and layerrotation and slip across layers in Cr/Sc multilayer stacks. Limitations of the tech-nique are presented and discussed.

v

Acknowledgments

I would like to thank:

Finn Giuliani, my supervisor, for his patience and help.

Lars Hultman, my examiner, for his encouragement, help and for giving methe opportunity to do this work.

Thomas Lingefeldt for his help with “difficult” equipment.

Karl Brolin for some critical precision lathe work.

The members of the Thin Film and Plasma Physics groups; for their kind-ness and help, and in particular Fredrik Eriksson and Jens Emmerlich forproviding Cr/Sc multilayer and Ti3SiC2 samples, respectively.

vii

Contents

1 Introduction 11.1 Aims . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 21.2 Guide to the chapters . . . . . . . . . . . . . . . . . . . . . . . . . . . 2

2 Background 32.1 Crystal structure and defects . . . . . . . . . . . . . . . . . . . . . . 3

2.1.1 Kink Bands . . . . . . . . . . . . . . . . . . . . . . . . . . . . 52.2 Silicon . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 52.3 Chromium-Scandium multilayers . . . . . . . . . . . . . . . . . . . . 52.4 Sapphire . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 52.5 MAX phases . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 62.6 Transmission Electron Microscopy . . . . . . . . . . . . . . . . . . . 72.7 Scanning Electron Microscopy . . . . . . . . . . . . . . . . . . . . . . 82.8 Focused Ion Beam Microscopy . . . . . . . . . . . . . . . . . . . . . 82.9 Nanoindentation . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 8

3 Experimental Details 93.1 Nanoindentation in situ TEM . . . . . . . . . . . . . . . . . . . . . . 9

3.1.1 Instrument . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 93.2 Sample preparation . . . . . . . . . . . . . . . . . . . . . . . . . . . . 12

3.2.1 Ion milling . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 133.2.2 Focused Ion Beam milling . . . . . . . . . . . . . . . . . . . . 13

3.3 Samples . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 143.3.1 Silicon . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 143.3.2 Chromium-Scandium multilayers . . . . . . . . . . . . . . . 143.3.3 MAX-phase Ti3SiC2 . . . . . . . . . . . . . . . . . . . . . . . 153.3.4 Sapphire . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 15

3.4 Nanoindenter tips . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 15

4 Results 174.1 Silicon . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 17

4.1.1 PIPS-prepared samples . . . . . . . . . . . . . . . . . . . . . 174.1.2 FIB-prepared samples . . . . . . . . . . . . . . . . . . . . . . 174.1.3 Berkovich & chisel indenters . . . . . . . . . . . . . . . . . . 18

4.2 Chromium/scandium multilayers . . . . . . . . . . . . . . . . . . . 18

ix

4.3 Sapphire . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 194.4 MAX-phase titanium-silicon-carbide . . . . . . . . . . . . . . . . . . 194.5 Issues and limitations . . . . . . . . . . . . . . . . . . . . . . . . . . . 22

5 Discussion 275.1 Nanoindentation in situ TEM . . . . . . . . . . . . . . . . . . . . . . 27

5.1.1 Cr/Sc and Sapphire . . . . . . . . . . . . . . . . . . . . . . . 285.1.2 Ti3SiC2 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 28

5.2 Conclusions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 285.3 Future work . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 29

References 31

A Guidelines for using the Indenter-Holder 35

B Log data extractor 37

Chapter 1

Introduction

MECHANICAL properties of materials have always been of great interest toanyone who makes a tool or component for a demanding task, and the un-

derstanding of these properties allows not only for the proper choice of material,but for the design of new alloys with desirable properties.

A common way to improve the properties of an object is to coat it with athin film of some other material. This allows for the use of more expensive orexotic materials in small quantities but most important is that the nature of thedeposition process may create structures not producible in bulk samples.

An important tool in studying mechanical properties is nanoindentation. Thebasic premise is that of probing the specimen with an indenter, most often inthe form of an pyramid, and recording the displacement response to the appliedforce, and from this curve various properties such as hardness and elastic modu-lus can be calculated. Traditional nanoindentation has a major weakness in thatit is only possible to observe the indent after the indentation.

This problem was first overcome by Wall and Dahmen[1] when they devel-oped a nanoindentation device for use in a Transmission Electron Microscope(TEM). There have been some similar instruments reported in the literature [2, 3],but as the devices have been highly bespoke their use has been limited. Recentlya Swedish company called Nanofactory Instruments AB put a commercial nano-indentation device for in situ work on the market. This allows scientists whoare not experts in Micro Electrical Mechanical Systems (MEMS) to utilise in situnanoindentation.

The Thin Film Physics Group at Linköpings University has recently acquiredsuch an in situ indenter and the aim of the work presented in this thesis has beento investigate the possibilities offered by it and to develop a basic knowledge ofthe workings of the instrument.

1

2 Introduction

1.1 Aims

The aim of this thesis is to evaluate the Nanofactory Instruments Indenter-Holderfor application to thin film science and in particular for investigations of the de-formation of brittle materials.

1.2 Guide to the chapters

The thesis is laid out as follows:

Chapter 1 is a short introduction to the subject and information on the thesis (thischapter).

Chapter 2 gives background to the thesis, treating both simple materials scienceand experimental methods.

Chapter 3 contains the experimental details.

Chapter 4 presents the results of the experiments.

Chapter 5 discusses the results, with conclusions and recommendations for fu-ture work.

Chapter 2

Background

THIS is a short introduction to basic subjects that is needed to follow the laterchapters. It should be noted that for each of these subjects there exists a large

body of works and they all deserve a more in depth description than what can bepresented here. I urge the interested to consult the references given herein, andI would particularly recommend Introduction to Solid State Physics [4], PhaseTransformations in Metals and Alloys [5], Introduction to Materials Science ForEngineers [6] and Physical Methods for Materials Characterization [7].

2.1 Crystal structure and defects

Most materials in the solid state form a crystal structure, meaning the atoms —ideally — sit in a periodic infinite lattice. In reality this crystal is not infinite buthas limits and when a sample is made out of many crystals it is called polycrys-talline where the individual crystals are called grains.

A basic property of a crystal is its periodicity, or in other words, it can beconstructed from a repetition of a simple “building block” in all directions. Thissmallest block is called the primitive unit cell, a conventional unit cell which isa larger but easier to deal with is often used instead. Figure 2.1 shows the threesimplest unit cells.

Figure 2.1. Unit cells for a) Simple cubic, b) Body centered cubic and c) Face centered cubic

3

4 Background

For indexing planes and vectors in a crystal Miller indices are used. For cubicsystems, these are constructed for a plane by taking the reciprocals of the fractionsof the intersection to the corresponding length of the unit cell and then reducingthe fractions to the smallest integers. For example, a plane perpendicular to thex-axis and intersecting the x-axis two lattice positions away from the origin (2a)has the Miller index: ( a

2aa∞

a∞ ) or (100). For cubic systems the normal of a plane

is the same as its Miller Index.

(001) — all planes with the z-axis as normal.

{001} — all planes with a normal along a directional axis {(100), (010), (001)}.

[001] — the direction along the z-axis.

<001> — all equivalent directions {[100], [010], [001]}.

There are many ways to construct a crystal lattice, and most materials andall alloys may be stable in several different lattices, depending on composition,temperature or pressure. The different possible configurations for a material aretermed phases, and the properties of the material can vary considerably depend-ing on them.

In reality crystals are not perfect, even when they consist of a single grain (asingle crystal or monocrystal ). The simplest defects are point defects which aredefects arising from the addition or removal of an atom. Line defects, or dislo-cations, are more complex defects where atoms are removed or shifted along aline. Next there are surface defects such as grain boundaries and volume defects,precipitates (inclusions) of another phase or voids.

For this thesis dislocations will be the most important type of defect, requiringmore explanation. The two basic types of dislocation are edge and screw dislo-cations. An edge dislocation can be thought of as formed by introducing a halfplane between two atomic planes in a crystal, where the edge of the plane corre-sponds to the dislocation line. As for a screw it can be formed by cutting halfwaythrough a crystal and shifting the edges along the cut.

To characterize a dislocation the construction known as the Burgers vectoris used. The Burgers vector is constructed by making a loop, a Burgers circuit,around the line of the dislocation, stepping from atom to atom, and then per-forming the same number of steps in an undeformed lattice. The vector neededto close the loop in the perfect lattice is the Burgers vector. It is a measure of thedistortion caused by a dislocation.

Deformation and dislocations are intrinsically connected. When a materialis put under stress it will first deform elastically (shifting the atoms from theirequilibrium positions but in such a way that the shift is reversible). As more andmore stress is applied eventually a point is reached when it is more energeticallyfavourable for the material to shift the atoms by introducing dislocations. Thisrelieves some of the strain in the material but the change is not reversible, whenthe stress is removed the dislocations will still remain and the material will retainsome of the strain now arising from the dislocations.

2.2 Silicon 5

2.1.1 Kink Bands

Kink bands are a deformation mechanism that has been observed for materialswith a unit cell that is longer in one dimension than the others or that have adifferent symmetry (hexagonal lattices are the typical example). By convention,the longer side has the length c, and the shorter ones the length a. It is when theratio c/a grows large that kinking has been observed.

Kink bands change the direction of the basal planes with a sharply definedangle (a kink), and are always followed by a kink of opposite direction changingthe angle of the basal planes back to the previous one. The kinks consist of dislo-cations that line up in the kink and it is their density that defines the angle of thekink [8].

2.2 Silicon

Silicon and its properties are well known on account of its widespread use inthe semiconductor industry; this coupled with its availability makes it a perfectstarting point for testing and learning.

Silicon has a diamond crystal structure (Take an fcc lattice and for every pointadd another at ( 1

4 , 14 , 1

4 ) from the point.) and deforms by dislocations slippingalong the {111} planes, with a perfect dislocation having a Burgers vector of1/2<110> [9].

2.3 Chromium-Scandium multilayers

Given the correct period of the layers, Chromium-Scandium multilayers work asX-ray mirrors [10] with the potential for constructing optics for X-rays. Thesemultilayers are soft, with a hardness of 5 GPa [10], compared to the other materi-als in this work but with a hardness comparable to that of steel.

Multilayers have been reported to deform by rotation of the layers under theindenting tip, shear across layers and shear along column boundaries [11].

2.4 Sapphire

Sapphire or α−Al2O3 is a ceramic with a hexagonal crystal structure (i.e. havingplanes with a hexagonal arrangement of the atoms, see [12] for a more completediscussion) and a hardness of 24 GPa [13]. The interest for Al2O3 in this thesisis mainly as a substrate for Ti3SiC2, but sapphire itself has a variety of uses insemiconductors, as hard coatings or jewelry.

Deformation in sapphire is governed by slip along three systems: basal {0001},prismatic {1210} and pyramidal {1011}; in addition, twining has been observed[13, 14].

6 Background

2.5 MAX phases

MAX phases — or more properly MN+1AXN phases where M is an early tran-sition metal, A is an element from the A-group of the periodic system and X iseither carbon or nitrogen (see figure 2.2)— were first synthesized in the 1960’sby Nowotny [15] while researching carbides and nitrides. Though they possesssome interesting properties not usually found in ceramics they were mostly for-gotten until the mid 1990’s when Barsoum started to take an interest in them.They managed to synthesize bulk samples of Ti3SiC2 [16] and started an inves-tigation of the properties of MAX phases.

Figure 2.2. Elements making up the MAX phases. Image courtesy of P. Eklund, used withpermission.

The reason that MAX phases are so intriguing is that they combine proper-ties of both ceramics and metals while their nanolaminated structure — atoms ofthe different elements form layers in the crystal structure — also contributes tothe unique set of properties. They are very good electrical conductors and elasti-cally stiff and at the same time machinable, softer than expected of a ceramic andrelatively less brittle than other ceramics [17, 18].

There are currently three structures identified of MN+1AXN phases, for N ∈{1, 2, 3} which will be referenced as 211, 312 and 413 onwards and their unit cellscan be seen in figure 2.3. The largest subgroup is 211 followed by 312, as for 413only four alloys have been found to date [18, 19].

The main deformation mechanism in MAX phases has been shown to be for-mation of kink bands. These materials also exhibit a nonlinear elastic deformationregime, and this has been linked by Barsoum [20] to a proposed formation of kinkbands put forward by Frank and Stroh [8]. This theory invokes the creation andannihilation of proto-kink bands, termed incipient kink bands, to explain the re-versible nature of the deformation process. Drawing parallels to other materialsthat deform by kinking he has coined the expression Kinking Nonlinear Elastic

2.6 Transmission Electron Microscopy 7

Figure 2.3. Illustration of 211, 312 and 413 MAX phases, courtesy of M. Beckers, used withpermission.

(KNE) solids, where he counts MAX phases, graphite and sapphire among othersas belonging to the group [21, 22, 23].

2.6 Transmission Electron Microscopy

Transmission Electron Microscopy is — as its name implies — a technique wherean electron beam is shone through a thin sample and the image is formed byscattering of the electrons on the atomic lattice of the sample.

The basic components of a TEM are an electron gun, electromagnetic lensesand something to view the image on, often a fluorescent plate in combinationwith a (digital) camera. This is very similar to an ordinary microscope in de-sign, and so most of the concepts of optical microscopy carry over into electri-cal microscopy. The main difference is that optical lenses can be ground veryprecisely while electromagnetic ones are about as good as the bottom of a sodabottle. Fortunately this is not such a great limit as can be imagined so the deBroglie-wavelength of a highly energetic electron is small enough to still allowfor the resolution of individual atoms.

8 Background

2.7 Scanning Electron Microscopy

A Scanning Electron Microscope (SEM) is another electron microscopy technique.In many ways it much simpler than TEM and the images are more easily inter-preted. The instrument produces images of the surface of the sample by scanninga fine electron beam across it and recording the electrons that are ejected backfrom the sample. The image is then formed by assigning each pixel a grayscalevalue according to the recorded intensity of that spot. The whole image is builtup in such a way by rastering the beam across a section of the sample.

As electrons and not visible light is used the resolution of a SEM can be asgood as a nanometer under ideal circumstances, this being determined mostly bythe size of the interaction volume of the electron beam and the sample.

2.8 Focused Ion Beam Microscopy

The Focused Ion Beam (FIB) instrument is in many ways similar to a SEM, themain difference is that the scanning beam is composed of ions of a heavy element(e.g. Gallium). These ions not only allow for imaging but also for milling of thesample. Precise control of the beam allows for the milling of fine features andin some cases complex shapes. They are also commonly integrated with a SEMwhich allows for nondestructive observation of the sample as well.

Examples of FIB use are preparation of TEM samples and testing, repair ofsemiconductor chips and other sculpting and fabrication in the nanoscale. Anexample of a sample prepared in a FIB can be seen in figure 3.7.

2.9 Nanoindentation

Nanoindentation is a technique where one directly probes the mechanical re-sponse of a material by probing it with a more or less sharp object, an indenter,which is usually made of diamond for its high stiffness (high elastic modulus)and high hardness. The reason for the name nanoindentation is that both the tipand the indent are in the nanoscale (10−9m)

The most common indenter shape in normal indentation is the Vickers pyra-mid — a four sided pyramid with a slope of 68◦. For nanoindentation the mostcommon shape is the Berkovich indenter which is a three sided pyramid with aslope adapted to give the same projected area for a certain depth as the Vickers in-denter. The reason for not using the same shape is that when cutting a four-sidedpyramid the tip will not be a point but a line which destroys the symmetry of theindenter. In contrast a three-sided pyramid can be cut with much better precision[24]. Another shape that has gained in popularity is the cube corner, giving asharper tip than the Berkovich that therefore penetrates more easily, promotingplastic response at lower loads.

During the indentation the force necessary to achieve each depth is recorded,both in the loading and unloading phases. From these force-displacement curvesproperties such as hardness and elastic modulus can be calculated [24].

Chapter 3

Experimental Details

WHEREIN the details of the work are explained. The instrument is described,as is sample preparation and selection of materials for samples. The dif-

ferent indenter geometries used are illustrated.

3.1 Nanoindentation in situ TEM

The goal of this experiment is to plastically deform a sample while viewing it inthe TEM. This is achieved by indenting a thin foil edge-on while looking throughit in the microscope. The plastic deformation, which is, in reality, the introductionof dislocations in the material, will give rise to contrast in the image through theinteraction between the electron beam and the dislocations.

Performing nanoindentation in a TEM puts special demands on both sam-ple and instrument. The sample must be thin enough to be electron transparent(approx. 200 nm in a 200 kV TEM) and at the same time mechanically stableenough to allow meaningful deformation. On the other hand, the instrumentmust be small enough to fit inside the microscope and precise enough to allowfor controlled motion on the nanometer scale. These problems have begun to beovercome but a standard way of solving them has yet to develop. A detaileddescription of the use of the instrument can be found in Appendix A.

3.1.1 Instrument

The instrument used is the Nanofactory Indenter-Holder, which is a TEM speci-men holder fitted with a nanoindenter (figure 3.1). In this device the indenter tipand force sensor are fixed while the sample is mounted on piezoelectric crystalswhich move the sample. There are two separate possible ways to move the sam-ple, for gross motion inertial sliding is used where the sample hat’s legs slide ona metal ball (figures 3.2 and 3.3). For fine movement one controls the voltage tothe piezoelectric crystals giving sub-nanometer control.

9

10 Experimental Details

Figure 3.1. The Indenter-Holder

Figure 3.2. Close-up of the holder, showing the sensor with tip and the sample holderwithout a sample.

Figure 3.3. The sample holder with sample present.

3.1 Nanoindentation in situ TEM 11

The diamond tip is parallel with the sample and perpendicular to the electronbeam, giving a top down view of sample and tip in the microscope. Figure 3.4defines the geometry.

Figure 3.4. Geometry of indentation

The whole system is made up by the Indenter-Holder, a control electronicsbox and the computer that runs the control software.

The controlling computer is equipped with a video capture card, and by con-necting that to the computer running the camera software, videos that are corre-lated in time with the indentation data can be produced.

Measuring the force is achieved by measuring the capacitance between twoplates separated by a spring; knowing the area of the plate and the spring con-stant the force may be calculated. The capacitance expressed in spacing “d” is,where ε0 is the permittivity of vacuum and A is the area of the capacitor:

C =ε0A

d(3.1)

The force acting on a spring to compress it is, where deq is the equilibrium spacingand k is the spring constant:

F = k(deq − d) (3.2)

Equations 3.1 and 3.2 combined yield:

F = kε0A

(1

Ceq− 1

Cmeasured

)(3.3)

12 Experimental Details

The value of the equilibrium capacitance (Ceq) is affected by two things, parasiticcapacitance between cables, contacts etc. and charging due to exposure to theelectron beam in the microscope, and these are corrected separately. The first oneis measured without the beam turned on and is thereafter subtracted from allreadings, this is called “Tune Without Electron Beam” in the NFControl program,and the second is corrected automatically at the start of an indentation . Formanual indentation the second step is not performed automatically.

3.2 Sample preparation

The process of preparing a sample for use in a TEM is quite involved, and evenmore so when one edge of the sample must be available for indentation. Thegoal is to create a sample that is so thin that it is electron transparent and at thesame time large and sturdy enough to be mechanically deformed by an indenterin a meaningful way. The four basic steps in the process are: cutting, polishing,mounting and thinning.

For cutting out samples two methods were used. In the first the cut is madeby a wheel coated with finely ground diamond, and in the second, by a metalwire lubricated with water mixed with fine silicon carbide powder.

For brittle samples wire cutting is more appropriate (see image 3.5 for com-parison) as it will make a finer cut with less damage on the edge. To further limitthe influence of the cuts on the final specimen the samples were cut to a width of1.5 mm.

Figure 3.5. Silicon cut with diamond wheel (a) and wire saw (b)

If the final thinning step is to be ion milling the sample must be large enoughto withstand the etching. A solution to this is to glue two samples together usinga strong compound. In this case “M-Bond 610” from M-Line Accessories wasused.

With the sample mounted on its long edge with wax on an objective glass itis polished down from both sides to a thickness of about 50 µm. The grindingpapers used were diamond cloths with a roughness from 32 µm to 1 µm. As withthe cutting it is important not to be too rough on the sample; in particular, it isimportant to get the edges of the sample as smooth as possible as it is on the edgethe indentation will be performed.

3.2 Sample preparation 13

When the sample is thin enough (around 50 µm) it is removed from the holderand cleaned of the wax by bathing it in acetone. A holder, a “grid”, is then cutopen on one side and the sample is then glued onto the grid with the epoxycompound “M-Bond 610”(figure 3.6). This is left to cure for an hour; then the tabon the grid is broken off and the sample is ready for ion milling or FIB-ing.

Figure 3.6. Sample mounted on a TEM grid.

3.2.1 Ion milling

Ion milling is a technique where the sample is bombarded with an ionized gas,usually argon, incident at an oblique angle. The ions etch the sample and it ismilled selectively where the gas flow is focused. During the milling stage theions were accelerated with a voltage of 4.5 kV using a Gatan 691 Precision IonPolishing System (PIPS).

When the sample is polished down to a thickness proper for TEM (around ahundred nanometers) an interference effect will show up as fringes and this isthe telltale sign for stopping the milling. After this the sample is further polishedin the same way but with ions accelerated by 2.5 kV. This is to remove damageintroduced by the rougher milling.

3.2.2 Focused Ion Beam milling

Another way of making an electron transparent area is to cut out a window usinga FIB. This allows the creation of small windows which are more mechanicallystable than the large thin areas made by conventional ion milling. It is also pos-sible to select an area of interest (perhaps an indent) in making the sample whichis not possible with an Ion miller.

14 Experimental Details

These windows are made by milling trenches from both sides of the sam-ple, converging on a bit in the middle which has been covered with platinumto protect the sample from damage from the ions. The main drawback with theplatinum is that it needs to be removed or pierced by the tip when indenting.

The FIB used (a Zeiss 1540 EsB Crossbeam FIB) has a 30 keV Gallium ionsource and the ability to deposit Platinum for protection of the sample. The win-dow length was typically 15-20 µm with a a thickness of 150 nm. Figure 3.7 showsa window milled in silicon that also shows support pillars inserted in the windowto decrease the flexibility of the window.

Figure 3.7. Sample milled in FIB, showing platinum protective layer and supportive pillarsin the window.

3.3 Samples

Several different samples were chosen for testing, both bulk single crystals andthin films.

3.3.1 Silicon

Silicon was chosen for the initial tests for its availability and its widespread use inboth industry and academia. It is also a standard test specimen for transmissionelectron microscopy. The samples were produced using both the PIPS and theFIB from Si(001) wafers.

3.3.2 Chromium-Scandium multilayers

For testing on a softer sample a thin film of Chromium/Scandium multilayersgrown by Fredrik Eriksson et al. [25] was chosen. The period of the layers was 36

3.4 Nanoindenter tips 15

Å and with a total number of 100 periods the thickness was 360 nm, on a Silicon(001) substrate. Samples were prepared using the FIB method.

3.3.3 MAX-phase Ti3SiC2

MAX-phase Ti3SiC2 was chosen as a sample because of the interesting problemof its deformation, i.e. the formation of kink bands during the plastic deforma-tion. Samples consisting of Ti3SiC2 (0001) thin films deposited on sapphire byJens Emmerlich et al. [26] were prepared using a modified FIB method, in whichthe samples were first covered with a thin (roughly 50 Å) layer of gold by sput-tering before being put in the FIB.

3.3.4 Sapphire

As the substrate for the Ti3SiC2 was sapphire it was easily available and there-fore included for testing. Only FIB thinning was used. The orientation was (0001).

3.4 Nanoindenter tips

Two different geometries for tips were tested. The standard Berkovich tip and atip resembling a chisel in that it has a straight edge (see fig 3.8) instead of a point.

Figure 3.8. Indenter tip geometries, a) Berkovich and b) chisel.

16 Experimental Details

Chapter 4

Results

HERE the results are presented from the indentations of silicon, sapphire, chro-mium/scandium multilayers and MAX phase titanium silicon carbide with

limitations of the technique and the instrument discovered through usage.

4.1 Silicon

As silicon was the first material used it was tested with both the PIPS and FIBmethods of sample preparation and was indented with both Berkovich and chiselindenter tips.

4.1.1 PIPS-prepared samples

Samples prepared with the PIPS showed large thin areas. Unfortunately the sameproperty makes the samples mechanically weak which means that the samplewill bend or even break before any meaningful plastic deformation from inden-tation can be introduced.

4.1.2 FIB-prepared samples

The sample prepared by FIB proved much more resistant to bending and break-ing than the PIPS-produced ones, at the expense of a lower quality sample forTEM imaging. This is for to two reasons: firstly the electron thin area is muchsmaller on a FIB-prepared sample and secondly, the milling is rougher (more en-ergetic ions) in the FIB which restricts the achievable quality. There was also somedifficulty in removing the platinum deposited to protect from FIB damage. Thiswas solved later for the Ti3SiC2 samples (see section 3.3.3 and 4.4).

17

18 Results

4.1.3 Berkovich & chisel indenters

Using a Berkovich tip proved to induce slipping of the sample along the tip giv-ing rise to a complex strain including pronounced bending, which also greatlyincreases the risk of fracture. This is due to the blunt tip geometry and is evenmore pronounced on badly machined or used tips.

The use of a chisel tip provided one great advantage over the Berkovich inthat it was less prone to slipping. On the other hand, the geometry makes accu-rate positioning of the tip harder in that it is impossible to focus the TEM on themiddle of the tip; this in turn forces the user to gauge the position of the middleof the tip. In most cases failure to position the sample at the middle of the tip willresult in slipping.

4.2 Chromium/scandium multilayers

As the period of the layers in the sample was just 36 Å it was not possible toresolve them at the magnifications practical for the indentations (see section 4.5).The material, being softer than the other materials [10] studied, deformed easily.In addition some flaking off was observed, most likely due to the sidewise loadproduced by the cross-coupling of the piezo-crystal.

Post-indentation studies revealed that the layers had deformed in line withwhat is expected for a multilayer [11], with rotation of the layers under the tipand slip across the layers observed. In figure 4.1 the rotation of the layers can beobserved at 1, slip across layers is found in several places at 2 and bending andslip can be seen at 3. The flaking is confirmed by the area under the indent witha marked brighter contrast than the surrounding film.

Figure 4.1. Overview of indent in Cr/Sc multilayers. Shown are Si substrate (bright) Cr/Scmultilayers and Pt protective coat (black).

4.3 Sapphire 19

Figure 4.2. Closeup of indicated regions in figure 4.1.

4.3 Sapphire

Indentation of sapphire with a maximum load of 350 µN showed only elasticdeformation while indentation with a maximum load of 400 µN displayed heavydislocation activity. Even at the highest load the depth of the indentation is veryshallow and when the load is removed the indent would hardly be visible if itwas not for the large plastically deformed area under it. Basal slip was observedboth in the movie and in post indentation observation (see figures 4.3-4.5). In themovie strain is visible from frame b, bend contours are visible from frame c andplastic deformation begins at d. Frame e shows further dislocation movementduring unloading with basal planes appearing to the right side of the indenter (atarrow) and the last figure shows the deformed region after removal of the load.

4.4 MAX-phase titanium-silicon-carbide

The gold sputtered on the sample before the platinum protective strip made theremoval of the protective Pt coat easy and nondestructive to the thin film under-neath. The platinum was removed by breaking it off with the tip using up- anddownward movements. This produced the best and cleanest surface for indenta-tion of those covered in this thesis.

On loading the material exhibited large dislocation movements along the basalplanes. On loading up to the maximum load of the sensor there war no apparentslip on other planes than the basal. With the exception of a few movements themain dislocation activity started suddenly at a load of 350 µN . Figure 4.6 shows asequence from an indentation and 4.7 shows the force plot. The maximum pene-tration depth was 90 Å. In the sequence (fig. 4.6) the major dislocation movementstarts at frame c in the series; the movements can be seen as changes in contrastin the film above the tip where the dislocations flow along the basal planes awayfrom the indenter.

In frame b, the arrow points to the strain contrast which has the form of ahalf-circle, indicating that the load is distributed evenly. Frames c and d showthe onset of the dislocation nucleation and glide. Some bending is also observed.

20 Results

Figure 4.3. Extracted frames showing live indentation of sapphire in situ the TEM. Themovie is available at LiU Electronic Press [27].

4.4 MAX-phase titanium-silicon-carbide 21

Figure 4.4. Overview of indent in sapphire showing deformation phenomena: 1) Basalslips, 2) Bend contours and 3) region of high dislocation density.

Figure 4.5. Force-displacement and force-time curves for the indent in sapphire, displace-ments are corrected for the compliance of the sensor. Letters correspond to the frames infigure 4.3.

22 Results

Arrow no. 2 points to more distinct bending contrast curves; in the same framethe other arrow (no. 3) marks the clearest basal slip.

Post-indentation microscopy (figure 4.8) showed that the deformation wasmostly limited to the depth of the indentation, aside from buckling of the win-dow. Kinks were observed on each shoulder of the indent with clear changes inorientation, 30◦ for the left and 20◦ for the right kink. Examinations of indentsmade at low loads (less than 200 µm) showed only bending of the layers aroundthe indent and no kinking.

4.5 Issues and limitations

A problem with the Indenter-Holder is that the tip is not really fixed in placerigidly but is spring loaded in two places. Firstly, the whole sensor is held in placeby springs which also serve as connectors to the sensor, secondly the sensor itselfis constructed as a capacitor with a spring between the plates. This introduces acompliance in the system which is undesirable from both the view of the nano-indentation and from the TEM part as the tip will move considerable distancesunder moderate load often removing the tip from view in the TEM, putting aneffective limit on the magnification of the microscope. The springs holding thesensor package were found stiff enough not to be noticeable under the loads thesensor is rated for. Unfortunately it is impossible to do anything about the sensoritself, the compliance of the spring was 1000 N

m2 giving a movement of 0.1µm for aload of 100µN . The upper limit of the magnification was 50 000 times for loadingup to the maximum load.

The setup is also very sensitive to electrical charging, and the sensor will ceasefunctioning if exposed to a too large charge (around a force reading of 150 µN ).The only way to restart the sensor is to turn off the electron beam and wait forthe capacitor to discharge. For normal usage this is usually not a problem, as thecharging equilibrium is within the sensors tolerance. Inserting an objective aper-ture in the microscope leads to more charging and generally makes the sensorunusable. This is unfortunate as it precludes the use of the objective aperture,which effectively decreases image contrast (figure 4.10 gives an example).

Another limitation is the camera. The way the system is connected the videostream that is delivered to the control computer has a resolution of 570x480 pixelswhich is quite limited at around half of the full resolution of the camera. Like-wise the speed of the acquisition matters here; too slow and it will be difficultto see anything. On the other hand as the needed exposure time depends on theintensity of the beam it cannot be raised too high and a tradeoff is required.

The sensor is also limited to a maximum load of 500 µN . While this is lowcompared to normal nanoindentation [24] it has proven enough to plastically de-form the samples. The samples tested in this thesis are all very hard materialswith the exception of the metal multiayers.

Crossover is always present to some degree in the piezocrystals meaning thatthere will be some unwanted motion in other directions than the indented. Forour system, this is particularly expressed in a cross coupling between the forward

4.5 Issues and limitations 23

Figure 4.6. Frames taken from movie showing indentation of Ti3SiC2. The movie isavailable at LiU Electronic Press [28].

24 Results

Figure 4.7. Force-displacement and force-time curves for indentation of Ti3SiC2. Dis-placements are corrected for the compliance of the sensor. Letters correspond to the framesin figure 4.6.

4.5 Issues and limitations 25

Figure 4.8. Overview of the indent in the Ti3SiC2 sample. Arrows point to positions ofthe kinks.

Figure 4.9. HRTEM image of the kinks on left and right sides of the indent in the Ti3SiC2

sample. Angles are 25◦ and 20◦ respectively.

26 Results

Figure 4.10. Difference in contrast between not using (a) and using an objective aperture(b). Indents in Si by Berkovich indenter on a PIPS-prepared sample.

and side motions. This forces the tip sideways into the sample when indenting,complicating the load on the sample, which may produce unwanted friction be-tween sample and tip as well.

It is also important not to move the indenter too fast, as the image quality willgreatly degrade if the CCD cannot keep up with the motion. Best results wereachieved by manually depressing the “down” button while having the fine Z-motion slider focused giving an indentation speed of approximately 30 nm/s atthe piezo.

The force-distance plots of a normal indentation show a very different appear-ance than normal nanoindentation curves in that the unloading part of the curveis above the loading curve. This is most probably due to the fact that the sensoris spring loaded, giving the unusual behavior in the unloading part.

Chapter 5

Discussion

THE results are discussed and contrasted with the literature, treating both thetechnique and the specific materials indented. Conclusions are presented

along with thoughts on future work.

5.1 Nanoindentation in situ TEM

Producing samples for in situ indentation is, at least, in the same order of com-plexity as producing a normal TEM sample, and the time required is about thesame as well. The requirements on equipment and geometry are different, a FIBmicroscope is required to produce good samples and the geometry, while be-ing thin enough for electron transparency, must give support so that plastic de-formation is possible before bending and failure through fracture. Minor et al.[29, 30] used lithography to make silicon wedges onto which the film material tobe studied was deposited, giving an increased stability at the possible expense ofchanged growth conditions which could alter the properties measured.

There is still the question on how the radically different geometry from normalnanoindentation affects the results. Certainly the windows indented are muchless restricted and can bend and flex as a whole. The stress field is also differentas it is constrained between the edges of the sample [30] and this, along with thelarge free surface for dislocations to exit, will influence the results.

Instrumentation needs to improve, as some of the limits mentioned in section4.5 are just a matter of design of the instruments and nothing inherent in theexperiment. With more stable instruments and better cameras it would be easierto more directly, if not conclusively, interpret the videos of the indentations.

In this work the samples have all been hard, with the exception of the Cr/Scmultilayers, and therefore required a high load to induce plasticity and disloca-tion movement. The maximum load of the instrument was reached (500 µN ) andthis should be contrasted to the work on Aluminium by Minor et al. where themaximum load was 80 µN [29].

The most suitable indenter tip was the chisel geometry as it was less prone to

27

28 Discussion

slipping off. It also has the advantage that the contact area is more easily extractedfrom the movie than for a threesided Berkovich tip as the contact length is visible.

5.1.1 Cr/Sc and Sapphire

The deformation of the Cr/Sc multilayers did not yield much information fromthe video due to the small period of the layers, but the post-indent study revealedlayer rotation, bending and slip across layers which is in line with the publishedliterature on multilayers [11]. If the experiment would be done again on a samplewith a larger period these deformation mechanisms should be detectable.

The indentation in sapphire shows that it is possible to indent a thin TEM-foilof a hard brittle material in a meaningful way. Basal slip was observed formingin the unloading phase but the unknown change in bending of the window maypartly be the reason that they were not observed earlier. According to the litera-ture prismatic slip should be active at lower loads than for basal slip [14]. Neitherprismatic nor pyramidal slip was observed, which can be attributed to the geom-etry where slip through the window is likely easier, and to the crystal orientationof the sample.

5.1.2 Ti3SiC2

The motion of dislocations, observed here during the indenation, was along thebasal planes as predicted from the high c/a ratio of the material which is a re-quired feature of the Incipient Kink Band model [20, 31]. Kinks were observedwithout any delamination. The latter is opposed to the incipient to normal kinktransformation outlined in [21]; this may be due to the fact that the kinks formeddownwards into the film and not outwards as observed in [31] and very close tothe indenter.

5.2 Conclusions

The technique of nanoindentation in situ TEM has been proven to work and togive meaningful results, although there are constraints arising from the equip-ment, sample geometry and sample preparation. The most limiting factors inthis work were the resolution of the camera and the stability-derived restrictionon the magnification of the microscope. Indentations have been made in silicon,chromium/scandium multilayers, sapphire (α−Al2O3) and MAX-phase titaniumsilicon carbide (Ti3SiC2), where the Silicon and sapphire were from single crystalsamples and the Cr/Sc and MAX-phase were thin films on Silicon and sapphirerespectively. Basal-plane flow of dislocations in Ti3SiC2 was observed and it wasfound that the film had kinked.

As observation of deformation phenomena is the most direct way to test thetheories of mechanical deformation in a given material the technique needs tobe developed further. The constraint placed by the geometry also needs to be

5.3 Future work 29

analysed theoretically to ensure that comparisons of deformations between thinfilms and bulk materials may be done.

5.3 Future work

There are three main paths for future work; firstly, the instrumentation must beimproved, secondly, the influence of the geometry of the samples must be studiedand thirdly, our knowledge of deformation mechanisms at the nanoscale must beverified and deepened.

It is not just the instrument that can improve, but also the techniques of prepar-ing the samples to lessen the influence of bending.

The deformation and kinking of MAX phases is still not fully understood sincekinking has not yet been observed in situ ; further experiments are thus required.

30 Discussion

References

[1] M.A. Wall and U. Dahmen. An in situ nanoindentation specimen holderfor a high voltage transmission electron microscope. J. Microsc. Res. Techn.,42:248–254, 1998.

[2] M.S. Bobji, C.S. Ramanujian, J.B Pethica, and B.J Inkson. A miniaturizedTEM nanoindenter for studying material deformation in situ. Meas. Sci.Technol., 17:1324–1329, 2006.

[3] A Stach, T Freeman, A.M. Minor, D.K Owen, J. Cumings, M.A. Wall,T Chraska, R. Hull, J.W Jr. Morris, A. Zettl, and U Dahmen. Developmentof a nanoindenter for in situ transmission electron microscopy. Microsc. Mi-croanal., 7:507–517, 2001.

[4] C. Kittel. Introduction to Solid State Physics. Wiley, eighth edition, 2005.

[5] D. A. Porter and K. E. Easterling. Phase Transformations in Metals and Al-loys. Chapman & Hall, second edition, 1992.

[6] J.F. Shackelford. Introduction to Materials Science for Engineers. PearsonPrentice Hall, 2005.

[7] P.E.J. Flewitt and R.K. Wild. Physical Methods for Materials Characteriza-tion. Bristol: Institute of Physics, 1994.

[8] F.C Frank and A.N. Stroh. On the theory of kinking. Proc. Phys. Soc., 65:811–821, 1952.

[9] D. Hull and D.J. Bacon. Introduction to dislocations, volume 37 of Interna-tional Series on Materials Science and Technology. Pergamon Press, 1984.

[10] J. Birch, T. Joelsson, F. Eriksson, N. Ghafoor, and L. Hultman. Single crystalCrN/ScN superlattice soft x-ray mirrors: Epitaxial growth, structure, andproperties. Thin Solid Films, 514:10–19, 2006.

[11] Y. Long, F. Giuliani, S.J. Lloyd, J. Molina-Aldareguia, Z.H. Barber, and W.J.Clegg. Deformation processes and the effects of microstructure in multilay-ered ceramics. Composites Part B: Engineering, 37(6):542–549, 2006.

31

32 References

[12] E. Wallin. Alumina Thin Film Growth: Experiments and Modeling. Lic.thesis, IFM, Linköping University, February 2007.

[13] S.J. Lloyd, J.M. Molina-Aldareguia, and W.J. Clegg. Deformation undernanoindents in sapphire, spinel and magnesia examined using transmissionelectron microscopy. Phil. Mag., 82(10):1963–1969, 2002.

[14] K.P.D. Lagerlöf, A.H. Heuer, J. Castaing, J.P Rivière, and T.E. Mitchell. Slipand twinning in sapphire. J. Am. Ceram. Soc., 77(2):385–397, 1994.

[15] H. Nowotny. Strukturchemie einiger verbindungen der Übergangsmetallemit den elementen C, Si, Ge, Sn. Prog. Solid State Chem., 2:27, 1970.

[16] M.W. Barsoum and T. El-Raghy. Synthesis and characterization of a remark-able ceramic: Ti3SiC2. J. Amer. Cer. Soc., 79:1953, 1996.

[17] M.W. Barsoum. The MN+1AXN phases: A new class of solids. Prog. SolidSt. Chem, 28:201–281, 2000.

[18] J. Emmerlich. MAX Phase Thin Films; Unique Multifunctional Ceramicswith the Elements Ti, Si, Ge, Sn and C. PhD thesis, Linköping University,SE-581 83 Linköping, Sweden, 2006.

[19] P. Eklund, J.-P. Palmquist, J. Höwling, D.H. Trinh, T. El-Raghy, H. Högberg,and L. Hultman. Ta4AlC3: Phase determination, polymorphism, and defor-mation. Unpublished.

[20] M.W. Barsoum, T. Zhen, S.R. Kalidini, M. Radovic, and A. Murugaiah. Fullyreversible, dislocation-based compressive deformation of Ti3SiC2 to 1 GPa.Nat. Mater., 2:107, 2003.

[21] M.W. Barsoum, A. Murugaiah, S.R. Kalindi, and T. Zhen. Kinking nonlinearelastic solids, nanoindentations, and geology. Phys. Rev. Letters, 92(25), June2004.

[22] A.G. Zhou, M.W. Barsoum, S. Basu, S.R. Kalidini, and T. El-Raghy. Incipientand regular kink bands in fully dense and 10 vol.% porous Ti2AlC. ActaMat., 54:1631–1639, 2006.

[23] M.W. Barsoum, A. Murugaiah, S.R. Kalindi, T. Zhen, and Y. Gogotsi.Kink bands, nonlinear elasticity and nanoindentation in graphite. Carbon,42:1435–1445, 2004.

[24] A.C. Fischer-Cripps. Critical review of analysis and interpretation of nano-indentation test data. Surf. Coat. Technol., 200:4153–5165, 2006.

[25] F. Eriksson, G.A. Johansson, H.M. Hertz, and J. Birch. Enhanced soft x-ray re-flectivity of Cr/Sc multilayers by ion-assisted sputter deposition. Opt. Eng.,41(11):2903–2909, 2002.

References 33

[26] J. Emmerlich, J.-P. Palmquist, H. Högberg, J. Molina, Zs. Czigány, Sz. Sasvári,P.O.Å. Persson, U. Jansson, and L. Hultman. Growth of Ti3SiC2 thin filmsby elemental target magnetron sputtering. J. Appl. Phys., 96:4817, 2004.

[27] L. Johnson. Movie of indent in sapphire. http://www.diva-portal.org/diva/getDocument?urn_nbn_se_liu_diva-8333-1__movie_part-2.mpg , 2007.

[28] L. Johnson. Movie of indent in Ti3SiC2. http://www.diva-portal.org/diva/getDocument?urn_nbn_se_liu_diva-8333-1__movie_part-1.mpg , 2007.

[29] A.M. Minor, S.S.A. Asif, Z. Shan, E.A. Stach, E. Cryanowski, T.J. Wyrobek,and O.D. Warren. A new view of the onset of plasticity during the nanoin-dentation of aluminium. Nat. Mat., 5:697–702, 2006.

[30] A.M. Minor, E.T Lilleodden, M. Jin, Stach. E.A., D.C. Chrzan, and J.W. Mor-ris JR. Room temperature dislocation plasticity in silicon. Phil. Mag., 85(2-3):323–330, 2005.

[31] J.M. Molina-Aldareguia, J. Emmerlich, J. Palmquist, U. Jansson, and L. Hult-man. Kink formation around indents in laminated Ti3SiC2 thin films stud-ied in the nanoscale. Scripta Materialia, 49(155-160), 2003.

34 References

Appendix A

Guidelines for using theIndenter-Holder

Using the Indenter-Holder is not always straightforward but if the general guideoutlined here is followed it should not pose any major problems.

After inserting the holder in the microscope connect it to the control box andturn on both the box and the control software “NFControl”. Start measuring,answering “no” to “Use the last tuning values?”. This is to allow for a correctadjustment of the stray capacitances which would otherwise affect the force mea-surement. The next step is to “Tune Without Electron Beam”; this requires thatmore than 100 sample points have been collected (takes approximately ten sec-onds). The data window should now display a force close to zero.

When the sensor is up and running and the vacuum is low enough it is timeto turn on the filament and align the microscope. Here it is important to align theeucentric height to the tip as it is the tip that is fixed and not the sample. So findthe tip, center it on the screen and do all alignments close to it.

Once the microscope is aligned sufficiently well it is time to find the sample.Move the holder lengthwise so that the sample is visible and take note of thearea. Find the thin area for indenting by moving the holder crosswise, and returnto the starting point — in line with the tip. Now use NFControl to move thesample until the thin part is visible. Now bring the thin area into focus using the“up/down” controls in NFControl. If the tip is in focus already this brings thesample into approximately the same height. Now move the sample towards thetip using the “forward” button in NFControl and following with the holder, keepthe thin area in view and in focus as well as possible to avoid large movementsclose to the tip.

During the alignment and approach only the coarse movement controls areused — that is: utilising inertial sliding — and the fine slider controls shouldbe left at zero to allow the greatest possible range of motion when close to thesample. This includes the forward/back control, and it is imperative that it isleft at zero when doing the final approach with the inertial sliding. The problemis that when moving back in with the coarse controls the sample will jump for-

35

36 Guidelines for using the Indenter-Holder

ward before staring to move backwards. The magnitude of this jump dependson the “Amplitude” slider in NFControl, but so does the size of the backwardsmovement, and if left too low the sample will not move backwards at all.

Once the sample is close to the tip it is time to use the sliders for the finalalignment of the height. The easiest way to accomplish this is to use the videocapture screen in NFControl and at the same time watch the force plot and min-imise the overlap between the sample and the tip when the force goes up whenmoving the sample into the tip. This can be done on the TEM screen as well, butthis requires some familiarity with the system.

With everything aligned, only the simple matter of performing an indenta-tion remains. There are two ways to do this, either manually or with NFControl’sbuilt in facility. For either approach there are advantages and disadvantages; au-tomatic indentation is the simplest, giving charging-corrected values, but manualcontrol allows for acting upon how the indentation develops. Another problemwith automatic indentation is that all data is lost if it is stopped before comple-tion. The main problem with manual indentations is that the logged data is notcorrected for charging so data extraction and correction is an additional problem.

As manual indentation was preferred in this work a script was written to ex-tract and perform charge correction on indentation log data. It can be found inappendix B.

Appendix B

Log data extractor

This script extracts data from an window from a NFControl log file, and performsthe second correction step, the charge calibration. This requires the logfile tocontain a hundred data points before the start of the window; for correct valuesthe data must represent a unloaded sensor just prior to the indentation.

# ! / usr / b in / python# −∗− c o d i n g : l a t i n −1 −∗−" " "Performs the charging c a l i b r a t i o n normally done in NFcontrol , on

the raw log data . Uses the formulas displayed in thedocumentation of NFcontrol .

" " "

import sysfrom c o l l e c t i o n s import deque

# C o n s t a n t s used f o r t h ek = 1000 # N/m, Spr ing c o n s t a n t o f s e n s o rSensorArea = 7 . 0 9 e−7 #m^2 , S enso r Areap e r m i t t i v i t y = 8 .8541878 e−12 # F /m, p e r m i t t i v i t y o f f r e e s p a c e

def compare ( d1 , d2 ) :t 1 = [ i n t ( a ) for a in d1 . s p l i t ( ’ _ ’ ) ]t 2 = [ i n t ( a ) for a in d2 . s p l i t ( ’ _ ’ ) ]for i in zip ( t1 , t2 ) :

i f i [ 0 ] < i [ 1 ] :return −1

e l i f i [ 0 ] > i [ 1 ] :return 1

return 0

def e x t r a c t ( l i n e ) :return l i n e . s p l i t ( )

37

38 Log data extractor

def handle ( s t a r t , stop , i n f i l e , o u t f i l e ) :header = i n f i l e . next ( )# p r i n t >> o u t f i l e , h e a d e rheader = header . r s t r i p ( ’\n\r \ t ’ )o u t f i l e . wri te ( header+ ’\ t F c o r r \n ’ )l i n e = ’ ’l i n e l = Nonep r e v i o u s _ l i n e s = deque ( )for l i n e in i n f i l e :

l i n e l = e x t r a c t ( l i n e )i f l i n e . s t a r t s w i t h ( ’ Time ’ ) :

continue #re s = compare ( l i n e l [ 0 ] , s t a r t )res2 = compare ( l i n e l [ 0 ] , stop )i f r es >= 0 and res2 <= 0 : # h i t !

breake l i f re s < 0 :

p r e v i o u s _ l i n e s . append ( l i n e l )

# E x t r a c t c h a r g i n g v a l u e s .C_n = [ ]for i in range ( 1 0 0 ) :

C_n . append ( f l o a t ( p r e v i o u s _ l i n e s . pop ( ) [ 3 ] ) )C_charging = (sum ( C_n ) /100) ∗1e−15p r e v i o u s _ l i n e s . c l e a r ( )print C_charging# Handle f i r s t l i n eC_cont = f l o a t ( l i n e l [ 3 ] ) ∗1e−15F_corr = k∗ p e r m i t t i v i t y ∗SensorArea∗ ( 1/ ( C_charging ) − 1/

C_cont ) ∗1 e3l i n e = l i n e . r s t r i p ( ’\n\r \ t ’ )o u t f i l e . wri te ( l i n e )print >> o u t f i l e , ’\ t %.4 f ’ % ( F_corr )

#The r e s tfor l i n e in i n f i l e :

l i n e l = e x t r a c t ( l i n e )# i f l i n e . s t a r t s w i t h ( ’ Time ’ ) :# c o n t i n u e #r es = compare ( l i n e l [ 0 ] , s t a r t )res2 = compare ( l i n e l [ 0 ] , stop )i f r es >= 0 and res2 <= 0 : # h i t !

#Do s t u f fC_cont = f l o a t ( l i n e l [ 3 ] ) ∗1e−15F_corr = k∗ p e r m i t t i v i t y ∗SensorArea∗ ( 1/ ( C_charging ) −

1/C_cont ) ∗1 e3l i n e = l i n e . r s t r i p ( ’\n \ t ’ )o u t f i l e . wri te ( l i n e )print >> o u t f i l e , ’\ t %.4 f ’ % ( F_corr )

39

e l i f res2 == 1 :# f i n i s h e dreturn

i f __name__ == ’ __main__ ’ :i f len ( sys . argv ) < 3 :

print >>sys . s tderr , " I n v a l i d argument ! "print >>sys . s tderr , " usage : xlogdata . py < s t a r t > <stop > (

i n f i l e ) ( o u t f i l e ) "print >>sys . s tderr , " s t a r t and stop has the form of :

YYYY_MM_DD_HH_SS"print >>sys . s tderr , " i f i n f i l e i s not given s t d i n i s used

. "print >>sys . s tderr , " Outputs to stdout . "sys . e x i t ( 1 )

i n f i l e = sys . s t d i no u t f i l e = sys . s tdouts t a r t = sys . argv [ 1 ]stop = sys . argv [ 2 ]t r y :

i n f i l e = open ( sys . argv [ 3 ] , ’ r ’ )o u t f i l e = open ( sys . argv [ 4 ] , ’w’ )

except IndexError :pass

handle ( s t a r t , stop , i n f i l e , o u t f i l e )i n f i l e . c l o s e ( )

Additionally, a script to add a simple counter to make plotting easier:

# ! / usr / b in / python# −∗− c o d i n g : l a t i n −1 −∗−

import sys

def h a n d l e _ f i l e ( in f , out ) :n = 0for l i n e in i n f :

i f l i n e . s t a r t s w i t h ( ’ Time ’ ) :continue

l i n e = l i n e . r s t r i p ( ’\n ’ )

l i n e += ’\ t%d\n ’ % nout . wri te ( l i n e )n += 1

i f __name__ == ’ __main__ ’ :for arg in sys . argv [ 1 : ] :

i n f = open ( arg , ’ r ’ )

40 Log data extractor

out = open ( arg+ ’ t ’ , ’w’ )h a n d l e _ f i l e ( in f , out )i n f . c l o s e ( )out . c l o s e ( )

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c© Lars Johnson