TEM and nanoindentation studies on sputtered Ti40Ni60 thin films

Materials Chemistry and Physics 97 (2006) 308–314

TEM and nanoindentation studies on sputtered Ti40Ni60 thin films A.K. Nanda Kumar, C.K. Sasidharan Nair, M.D. Kannan, S. Jayakumar ∗ Thin Film Center, Department of Physics, PSG College of Technology, Coimbatore, India Received 10 March 2005; received in revised form 25 July 2005; accepted 9 August 2005

Abstract Ti40 Ni60 thin films were sputter deposited from an equiatomic alloy target onto Si(1 0 0) substrates and the mechanical response of these films have been studied. The compositions of the films were determined by RBS. The transformation of the amorphous films into an ordered B2 Austenite phase on annealing was confirmed by electron microscopy. The effects of second phase segregation and superlattice formation and their implication on the electron diffraction patterns have been discussed. The mechanical response of the films to indentation was studied by a standard Hysitron triboindenter. The deformation behavior of the films was analysed via a series of load–displacement (P–h) curves and the singular occurrence of intermittent discontinuities in the otherwise smooth P–h curves are attributed to the onset of plastic deformation. Mechanical response parameters like hardness and Young’s modulus have been determined from the unloading curve as per the standard Oliver–Pharr analysis. © 2005 Elsevier B.V. All rights reserved. Keywords: TiNi shape memory alloy; Thin film; Austenite phase; Nanoindentation; Oliver–Pharr analysis; Young’s modulus; Hardness

1. Introduction Recently, there has been a growing interest in the tribological and mechanical analysis of surfaces and films owing to relatively new, ingenious materials that have found resourceful applications in various hybrid systems like micro electro mechanical systems (MEMS), that require actuation and sensing mechanisms with high efficiency. Where such high efficacy actuation mechanisms are concerned, one is immediately interested in the physical mechanism of the thin film element and hence knowledge of the mechanical responses becomes quite imperative when devising an actuator that works on a microscale. Thin films of a variety of materials with good mechanical properties are presently being studied for applications such as abrasion-resistant protective coatings as well as micro actuators in micro devices. In our present work, we have analysed thin films of titanium–nickel (shape memory actuator) alloy of specific composition (Ti40 Ni60 ), by nanoindentation studies. Titanium–nickel alloy has been classified as a, “shape memory” material due to the fact that ∗

Corresponding author. E-mail address: s jayakumar [email protected] (S. Jayakumar).

0254-0584/$ – see front matter © 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.matchemphys.2005.08.020

it can recuperate a large amount of plastic deformation on suitable heat treatment, or thermal training. It is hence widely used as an actuator due to its mechanical response to a thermal input. When used as an actuator in microsystems, the shape memory alloy (SMA) provides a maximum work density of 10.6 MJ m−3 , compared to other actuators, like piezoelectric (0.1 MJ m−3 ), bimetallic strip (0.5 MJ m−3 ), etc. [1,2].

2. Preparation of the samples Thin films of TiNi were sputter deposited from an equiatomic target on (1 0 0) oriented Si single crystals maintained at 250 ◦ C by conventional dc magnetron sputtering. The films were analysed for their composition by Rutherford back scattering (RBS) technique. The composition of the films was found to be Ti 40 at.%–Ni. The films are Ni rich due to the larger sputtering yield of Ni. The details of the deposition conditions are given in Table 1. The films thus deposited were annealed at 525 ◦ C to obtain crystallisation under a vacuum of nearly 1 × 10−5 mbar, to prevent any oxidative reaction of Ti. The cooling rate was maintained at 10–12 ◦ C min−1 . The as-deposited (deposited at 250 ◦ C)

A.K. Nanda Kumar et al. / Materials Chemistry and Physics 97 (2006) 308–314 Table 1 Preparation details of the dc magnetron sputtered TiNi films Substrate Substrate temperature (◦ C) Ar gas pressure (mbar) Sputtering power (W) Target–substrate distance (cm) Composition (RBS, at.%) Film thickness (nm)

Si(1 0 0) 250 0.01 100 4 Ti: 40; Ni: 60 500

films and the heat-treated films were then subjected to electron microscopy and nanoindentation analysis. Usually, TiNi films prepared at room temperature are amorphous and no longer behave as SMA. Hence, a proper thermal annealing process is indispensable. Also, if the crystallisation temperature is too high or if the annealing period is too long, there is an interfacial mixing and reaction [3]. The onset temperature of crystallisation also varies with the alloy addition and composition. As a rule, as the x in Tix Ni1−x decreases, the crystallisation temperature also decreases, indicating less stability based on the enthalpy of mixing [4]. The crystallisation temperature is also substantially affected by the composition, since, the boundary on either side of the equiatomic TiNi are Ti2 Ni and Ni3 Ti, which are two different intermetallic compounds with different crystallisation temperatures [5].

3. The structural identification of the Austenite phase by TEM The TiNi alloy system exists in two phases—the high temperature Austenite phase (B2 cubic structure) and the low temperature Martensite (B19’ monoclinic structure). The dominant phase in the annealed films could not be discerned by elementary XRD methods, due to the overlap of certain high intense peaks of both phases. Hence, a TEM analysis was done on the films to identify the crystal structure of the films and to study the microstructure.

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Fig. 1(a) and (b) shows the bright field TEM micrographs of the annealed films after ion beam thinning. The growth of Austenite grains is seen. Fig. 2 shows the selected area diffraction (SAD) patterns from some of the Austenite grains, along different orientations. The grains seem to have 1 1 1 fibre texture. TEM studies also revealed twin defects in certain grains (Fig. 3) that have nucleated while annealing. The results of the TEM analysis suggests that the films are in the Austenite (B2) phase with a substantial amount of segregation of Ni3 Ti. Also, certain grains seem to have inherent twin defects. These results help in understanding certain observations that were obtained while nanoindenting the films.

4. Mechanical studies on the sputtered films To understand the mechanical responses of thin films is difficult, and a survey of the literature during the past one and a half decade illustrates that most of the mechanical tests on such thin films were done using the Bulge test, details of which are provided by the group of Makino et al. [6–9]. Recently, indentation tests have perhaps become one of the most commonly applied means of testing the mechanical properties of thin film materials. Nanoindentation provides accurate measurements of the continuous variation of indentation load, P, as a function of indentation penetration depth, h. In all indentation tests, the indenter is loaded at a specified rate and the load is then held constant for a period of time, while the displacement is monitored. For our present study, all nanoindentation studies were done using Hysitron Triboindenter, with a three-sided pyramidal Berkovich diamond indenter, with an included angle of 143.5◦ between the opposite faces, and the radius of curvature of the tip being 150 nm. The maximum loads were 250, 500, and 750 ␮N and, in each case, the loading was held at the maximum value for a period of 5 s. Also, the indentations were done at four different regions of the films, spaced ∼100 nm apart, so that there is no intervention on

Fig. 1. (a) Bright field image of the annealed films showing Ni3 Ti precipitates in Austenite matrix (88600X); (b) some of the Austenite grains (122000X) of ˚ average grain size 1000 A.

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˚ at 122,000× and SAD pattern viewed along 1 1 1 direction (inset) and along 1 1 0 of B2 Fig. 2. Bright field images from some Austenite grains (∼1000 A) cubic phase. The appearance of the 1 0 0 forbidden reflections (BCC) is due to the ordered superlattice, and superposition of Ni3 Ti.

Fig. 3. TEM micrographs of twinned grains (left) and corresponding SAD pattern (right) from the twinned grains. Note the overlap of the diffraction spots (inset) and the failing intensity from the central spot.

the P–h curves at each spot. All the indentation tests were done at room temperature with the samples firmly fixed to the holder. As the indenter is driven into the material, both elastic and plastic deformation occurs, which results in the formation of a hardness impression that conforms to the shape of the indenter used. Fig. 4 shows the depth profile of the surface of the as deposited and annealed films after the indentation. The projected contour shows higher penetration depth of the as deposited amorphous films, inveterate by the inherent porosity and softness of the films prior to any heat treatment. It is also observed that the roughness profile of the heat-treated films is much corrugated than the as prepared samples without any heat treatment. This is due to the strain relaxation of the films from the substrate at localized centers of the interface where the residual stresses are very high. Studies on the surface topography and optical reflectivity while the system changes from the amorphous to Austenite to the Martensite phase have been reported by Fu et al. [10]. The difference in surface roughness also seems to tell upon the optical reflectivity. Hence, a similar consequence is only logically

expected in the hardness and mechanical properties of the films. An overall representative curve of a nanoindentation study is a parabola with intermittent discontinuities in an otherwise

Fig. 4. Roughness profile of the as deposited and crystallized films. The indentation is much deeper for the amorphous, as deposited films due to porosity and lower strength.

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eventless curve. All indentation analysis are carried out under two basic premises: (i) the deformation upon unloading is purely elastic; (ii) the compliance of the sample and the indenter tip can be combined as a system of springs in series, with the effective (or reduced) modulus Eeff being 1 Eeff

=

1−υi 2 Ei

+

1−υs 2 Es ,

(1)

where E and υ represents the Young’s modulus and Poisson’s ratio of the indenter, i and sample, s, respectively. For a sharp Berkovich indenter with an included angle of 140◦ , the loading path equation takes the general form of [11,12]: P = Ch2

(2)

which, in all respects implies a parabolic fit with a constant of equation, C, which can be determined from the load (P)–displacement (h) curves. The unloading path of an elastic material, for many simple indenter geometries, follow the general power law (Kick’s law) [13]: P = α(h − hf )m ,

(3)

where α and m are constants and h and hf the instantaneous and final displacements of the indenter. All estimations of the elastic property parameters like Young’s modulus and hardness, etc. is done from this curve. The entire load–unload curve represents the overall elasto–plastic response of the film [14]. As a prologue to the analysis, it should be mentioned that the films have been assumed to be of the 1 1 1 fiber texture as seen from the electron diffraction patterns and the results are assumed to be the aggregate mechanical response of a system of polycrystalline grains favorably oriented along this direction. In all totality, this has been taken as a very near approximation to the real system.

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4.1. Analysis of the P–h curve The P–h curve, as is evident in Fig. 5, represents a discrete and continuous deformation process occurring in the material. The response of a nanoindented film shows purely elastic behaviour with intermittent microplasticity [15]. The unloading curve, during indentation withdrawal, is smooth which means that it is purely the recovery on relaxation. This facilitates the invocation of elastic solutions to model the contact process during indenter unloading. Fig. 5 shows the P–h curves obtained for the as deposited amorphous films and the films annealed at 525 ◦ C with loads of 250 and 500 ␮N. The amorphous phase exhibits a lower work output than the crystallised TiNi Austenite phase films. The area enclosed by the curves represents the total work done during the entire load–unload cycle and is an indication of the work lost due to plastic deformation. Considering a load of 500 ␮N, the area enveloped by the as deposited films is 3.95 × 10−12 J, whereas the area enclosed by the heat-treated films is found to be 3.03 × 10−12 J. We have observed a decrease in the mechanical hysterisis for the annealed films, due to better crystallisation by annealing. An important feature to be noticed in the curves is the shift in the total P–h loop for the as deposited and annealed films. It is a common trait exhibited by the films at both the loads, that the loading curve is nearly identical for both the amorphous and annealed films until the first displacement burst occurs at ∼153 ␮N. This first displacement burst seems to be independent of the maximum load. This appears to be the critical load, after which the deformation behavior changes. The width of the plateau region, during loading is indicative of the extent of plastic deformation suffered by the films, on compressive indentation. The width of the flat region is a projection of the amount and the nature of the dislocation configuration that has nucleated at the indenter tip. Attention is also directed to a particular aspect reported in

Fig. 5. Load (P) vs. displacement (h) curves for the as deposited and annealed films at maximum load levels of 250 ␮N (left) and 500 ␮N (right). Insets: Pyramidal indentation profiles of the annealed films.

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[16,17], that the onset of the first displacement burst appears to occur when the computed maximum shear stress at the indenter tip approaches the theoretical shear strength of the material. There is also the phenomenon of stress-induced Martensite phase transformation in the TiNi alloy system. The driving force of Martensite transformation can be either the free energy changes or the stress induced basal shear mechanism. Only shear stresses aid Martensitic transformation. It is still indecisive whether the stress values obtained at the first displacement spike is indicative of a phase change, but it might as well be a critical stress where a combination of both twin dislocation nucleation and phase transformation occur. Such interfering phenomena can only be resolved through suitable electron microscopy analysis of the indented area. A more formidable model has to be developed for a polycrystalline, non-isotropic media. It is also possible that the second and higher discontinuities can be the nucleation of twins at the slip dislocations, which can act as stress concentrated areas. Rapid formation of twins will surely result in sudden increment of the tensile strain. 4.2. Estimation of hardness and Young’s modulus Nanoindentation simplifies to a great extent, the determination of two very critical parameters, namely, the hardness, H, and Young’s modulus, E through a single test. We have used the Oliver–Pharr analysis method [13] to determine the hardness and Young’s modulus of the film. The key parameters that are used for analysis in this method are the maximum load, Pmax and the displacement at peak load, hmax . The Oliver–Pharr data analysis procedure begins by fitting the unloading curve to the power law relation, given earlier in Eq. (3). The effective modulus of elasticity, Eeff , accounts for the fact that elastic deformation of both the specimen and indenter is given by Eq. (1), wherein we have used a diamond indenter, with a Poisson’s ratio, of 0.07 and Ei of 1141 GPa [18].

Hardness can be calculated by the expression [13], H=

Pmax , A

(8)

‘A’ being the effective area of contact. Fig. 6 shows the Young’s modulus variation as a function of the indentation load for the as deposited and annealed films. The nature of the graph is the same for both the films. At loads greater than 500 ␮N, the Young’s modulus saturates to an almost constant value. The elastic modulus and hardness of the annealed films was found to be greater than the as deposited films. The value of the elastic modulus is 150 GPa for the annealed films and 138 GPa for the as deposited amorphous films. Hardness variation of both as deposited and annealed films are also shown in Fig. 6. A similar trend as in the elastic modulus is obtained here also, with a maximum hardness value of H = 7.11 GPa for the as deposited films and H = 8.42 GPa for annealed films. Both the elastic modulus and the hardness are highly non-linear with temperature, and these values are found to be considerably higher for the thin films as compared to the bulk (E = 85 GPa for the Austenite phase [18]). This behaviour is attributed to the compositional change and phase segregation occurring in the films on annealing. Although defect density is reduced by annealing, phase segregation of Ni3 Ti leads to precipitation hardening and hence an increase in the hardness and modulus. It is well known that the strength of metallic films rises dramatically when the thickness falls below about 0.3 ␮m and the Hall Petch relation, which is usually obeyed by thin films, is also dependant on film thickness [5]. It is mostly observed in all indentation analysis, that, at low loads (low penetration depths), the hardness of the material increases [17,19]. This effect has now been confirmed to be due to the nucleation and propagation of dislocations. Insights into the phenomenal occurrence of nucleation of dislocations and defects in thin films can be gained by exploring the local stresses that arise at the tip of the indenter during the loading process.

Fig. 6. Variation of Young’s modulus and hardness with load for the as deposited and amorphous films.

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Table 2 Critical parameters calculated for the as deposited and annealed samples for a maximum load of 250 ␮N Nature of the films

Wt (×10−12 J)

hr (nm)

Wp /Wt = hr /hmax

We /Wt = 1 − (hr /hmax )

As deposited Ti40 Ni60 Annealed Ti40 Ni60

1.71 1.53

7.95 4.89

0.386 0.265

0.614 0.734

4.3. Analysis of the elastic recovery ratio (hr /hmax ) An important parameter used to explain the indentation behaviour of a given material is the ratio of the residual or final depth of penetration, hr , upon complete unloading to the maximum penetration depth, hmax, prior to unloading, and this factor is indicative of the extent of plastic deformation and strain hardening. hr /hmax is determined from the P–h curve. For complete elastic deformation, hr /hmax = 0, and for a purely plastic deformation, the ratio, hr /hmax = 1. Obviously, the limits of this factor is 0 < hr /hmax < 1. Intermediate values indicate different extents of plastic response to deformation. Alternatively, hr /hmax gives the amount of plastic work done during the indentation process. The total work done during the entire load-unload cycle is given by the sum of the elastic work (We ) and plastic work (Wp ): Wt = We + Wp

(9)

An appropriate analysis in this context is divulged in [12], which uses the following relation: Wp hr = hmax Wt

(10)

hr /hmax = 0.895 is the critical stress for which there is no pile up or sink in of the material under the indenter [12]. If the ratio is <0.895, the material is characterized by high yield stresses and/or high work hardening, which seems to be the case in this study. Our calculation of hr /hmax yields a ratio of 0.265 for the annealed films and 0.386 for the as deposited films at a constant load of 250 ␮N. The work hardening of the annealed films is greater due to the combined effect of precipitation hardening and crystallisation. Annealing leads to increase in strength by grain boundaries and dislocation interactions, whereas the as deposited films are almost amorphous. But, from a practical outlook, the direct measurement of the residual depth, hr , is prone to considerable errors, due to a number of factors, like surface roughness (Fig. 4). Hence, another equivalent parameter that can be more reliably estimated for the elastic and plastic components of deformation during indentation, is the area under the P–h curve. The area under the loading portion of the P–h curve, is a measure of the total work, Wt , done by the indenter in deforming the material and is a function of the maximum load (Pmax ) and indentation depth at maximum load (hmax ), approximated by the relation [12]: Wt =

Pmax hmax 3

(11)

The values so calculated for a load of 250 ␮N is given in Table 2. Similar results are obtained at all the maximum loads (500 and 750 ␮N). 5. Conclusions Nanoindentation of Ti40 Ni60 thin films coupled with TEM imaging proves to be a very viable tool for a profound, comprehensive analysis of the deformation behaviour of shape memory alloy thin films on a microscopic scale. Numerous mechanical parameters have been derived through a series of load–displacement curves and the estimation of two very important parameters, namely, the modulus of elasticity and the hardness are determined through a single experiment. Inspection of the P–h curves show marked discontinuities. The discontinuities have been argued to be the indication of the onset of plastic deformation and the nucleation and growth of defects and dislocations. The displacement bursts are found to be fairly independent of the load (Pmax ), and seems to be slightly affected by the heat treatment procedure adopted (there is a marked difference, albeit small, between the critical loads in the as deposited and annealed films). Elastic modulus and hardness values have been determined, which seems to be quite greater than the data available for the bulk. In summary, the present work is a record and analysis of the behaviour of Ti40 Ni60 thin films subjected to nanoindentation. Although the film has not been studied for its shape memory property, more studies are being carried out by the present authors to compare the indentation behaviour of a near equiatomic TiNi alloy thin film with the presently reported composition. Acknowledgements The authors are sincerely grateful to the Defence Metallurgical Research Laboratory (DMRL), Hyderabad, for the TEM studies and to the Mechanical Engineering Department of Indian Institute of Science (IISc), Bangalore, for providing the very valuable NanoIndentation facility. The authors also thank the Department of Science and Technology (DST), India, for supporting the project financially and PSG College of Technology, Coimbatore, for providing the infrastructure. References [1] R.H. Wolf, A.H. Heuer, J. Microelect. Mech. Syst. 4 (1995) 206. [2] C.M. Wayman, Proc. Mat. Res. Soc. Symp. 21 (1984) 657.

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