Combination of Instrumented Nanoindentation and Scanning Probe Microscopy for Adequate Mechanical Surface Testing

Transcription

1 J. Mater. Sci. Technol., Vol.25 No.1, Combination of Instrumented Nanoindentation and Scanning Probe Microscopy for Adequate Mechanical Surface Testing Enrico Tam 1), Mikhail Petrzhik 2), Dmitry Shtansky 2) and Marie-Paule Delplancke-Ogletree 1) 1) Université Libre de Bruxelles, Brussels 1050, Belgium 2) Moscow State Institute of Steel and Alloys, Moscow , Russia [Manuscript received December 12, 2007, in revised form March 3, 2008] The elastic indentation modulus and hardness of standard bulk materials and advanced thin films were determined by using the nanoindentation technique followed by the Oliver- Pharr post-treatment. After measurements with different loading/unloading schemes on chemically polished bulk titanium a substantial decrease of both modulus and hardness vs an increasing loading time was found. Then, hard nanostructured TiBN and TiCrBN thin films deposited by magnetron sputtering (using multiphase targets) on substrates of high roughness (sintered hard metal) and low roughness (silicon) were studied. Experimental modulus and hardness characterized by using two different nanoindenter tools were within the limits of standard deviation. However, a strong effect of roughness on the spread of the experimental values was observed and it was found that hardness and elastic indentation modulus obeyed a Gaussian distribution. The experimental data were discussed together with scanning probe microscopy (SPM) images of typical imprints taken after the nanoindentation tests and the local topography s strong correlation with the results of nanoindentation was described. KEY WORDS: Nanoindentation; Thin films; Coatings; Scanning probe microscopy 1. Introduction Instrumented nanoindentation has become very powerful and is one of the few techniques usable for the quantitative analysis of surface layers of bulk materials and thin films to estimate their mechanical properties [1]. Over the last few years, enormous efforts were made to improve both the theory of the nanoindentation [2 5] and the commercially available instruments [6,7]. The number of publications in which nanoindentation is used to derive the mechanical properties of surfaces is growing year by year, but not all of them contain a thorough analysis of the experimental data [8,9]. There are also some questionable reports [10] about ultra high hardness of advanced films (more than 100 GPa). This seems doubtful, because these films would be harder than diamond, which is the indenter material, and it is perhaps a result of miscalculation of the contact area [11,12]. So, along with a widening use of the nanoindentation technique, an interest in a reliable interpretation of the resulting data is rising. Recently published ISO and ASTM standards [13] state strong rules to conduct indentation tests: (1) the surfaces must be smooth enough compared to the penetration depth, which means that the indentation depth should be at least 20 times greater than the Ra average roughness of the tested material; (2) the maximum indentation depth should not exceed 10% of the thickness of the film to avoid the influence of the substrate. As a result, when indenting films of 1 µm in thickness, the penetration depth must be less than 100 nm, and the surface roughness must be less than 5 nm. Fulfilling these conditions can be difficult in practice. Roughness in the order of one nanometer (which Corresponding author. Prof., Ph.D.; Tel.: ; address: enrico.tam@ulb.ac.be (Enrico Tam). means that the indentation depth in this case should be more than 20 nm), can be easily obtained for metallic alloys like steels and Ti-alloys since they are homogeneous in structure and thus easily polished. In this case, the application of the ISO rules is not a problem. However, this kind of alloy often exhibit creep behavior even during indentation at room temperature, which make the experimental data variable depending on the applied loading-unloading pattern. This means that indentations of ductile materials as well as indentation of thin films deposited on them can be strongly affected by the creep phenomenon. As a result, creep effect should be taken into account especially when analyzing hard film on soft substrate systems. On one hand, these kinds of coatings are very important for many modern medical applications [14] and also for the development of protective films on steel parts [15]. On the other hand, for many wear-resistant applications hard films on hard substrates are generally more suitable. In this case, the substrate (which gives the roughness to the final coating) is usually made of a sintered hard metal and consists of hard micro particles embedded into a metallic matrix. Since the properties of the different components composing the substrate are sometimes strongly different, it is difficult to prepare a smooth surface because the wear during the polishing will be selective. In addition, the surface roughness of the substrate can increase after the ion etching performed before the deposition and which is a mandatory operation before PVD. In this case, when very thin layers are considered (less than 1 µm thick), it becomes difficult to fulfill the two ISO rules simultaneously. For instance, considering again the example of the 1 µm thick film, the penetration depth should be less than 100 nm (first ISO rule) and consequently the surface roughness R a must be less than 5 nm (second ISO rule) which is sometimes very

2 64 J. Mater. Sci. Technol., Vol.25 No.1, 2009 difficult to reach. From a practical point of view it becomes important to know how to derive reasonable results from instrumented nanoindentation when the thickness of the coating does not permit the fulfillment of the new ISO rules. The aim of this work is to conduct a comparative study of the mechanical properties of soft and hard surfaces of bulk and coated materials of various roughnesses using instrumented nanoindentation and scanning probe microscopy and to study the creep and roughness effects on them. 2. Experimental Details The surface preparation technique for polycrystalline, commercially-pure titanium consisted of a mechanical polishing with grinding paper followed by an electro-polishing. The grain size was found to be of the order of µm. TiBN and TiCrBN films of about 1 µm in thickness were formed by magnetron sputtering of TiBN and TiCrBN targets in gaseous mixture of argon and nitrogen [17]. The targets were produced by the combined force SHS-pressing technology, as described in the paper of Levashov et al. [16]. Single crystal silicon (100) and multi-phase hard alloy (TT8K6 brand mark, (W,Ti,Ta)C-6%Co) were used as substrates. The structure of TiCrBN films consisted of a mixture of amorphous and nanocrystalline phases without a strong preferred orientation. The grain size ranged from 3 to 8 nm [17]. Nanoindentaion experiments were performed and compared by using two apparatus, the Hysitron Inc. Triboindenter and the Nanohardness Tester, CSM Instr, both equipped with Berkovich indenters. All tests were done as one cycle loading-unloading indentation with different holding time periods (0, 10 and 50 seconds respectively) while the loading and the unloading segment duration were held constant. Penetration depth was kept about 100 nm in order to remain at a penetration of less than 10% of the film thickness as required by the new ISO rule. Roughness of all the samples was estimated by scanning probe microscopy (SPM) head build-in Hysitron Inc. Triboindenter and by DEKTAK 3030 contact profilometry. The titanium sample was electrochemically polished and characterized by a small Ra roughness (in the order of 1 nm) compared to the penetration depth which was approximately 100 nm. The same Ra roughness was found for the thin film deposited on silicon whereas the roughness of the thin films deposited on the hard metal substrates exceeded 5 nm and therefore would not be acceptable under the new ISO rule. 3. Results 3.1 Nanoindentation of bulk titanium Load-displacement curves of standard materials The fused quartz reference sample was studied at first to verify the calibration of the equipment. As expected, it showed a very consistent and repeatable displacement behavior. Several indentations made at the same maximum load produced a difference in the Fig. 1 Load-displacement curve sets obtained using the Hysitron Inc. Triboindenter after indentations on: (a) the fused quartz specimen at 2 mn and 5 mn load; (b) the titanium sample at 0.5 mn and at 1 mn load maximum depth of penetration below few nanometers (Fig. 1(a)) which is a small fraction (2 3%) of the maximum penetration. Conversely, the titanium sample exhibited a large difference in the maximum depth of penetration (about 10 nm) at the end of the loading branch (Fig. 1(b)). This seems to be due to the creep behavior [9]. In addition, the slope of the final part of the unloading branch varies strongly (see Fig. 1(b)). This effect could be connected with surface modifications caused by the mechanical or chemical polishing due to the chemical activity of titanium. Cold working cannot be an issue of the recovery because of the pronounced pile-up formation (Fig. 2). However, stress can induce surface phase transformations and may be a possible explanation of the behavior due to hydrogen adsorption and to the formation of titanium hydrides near the surfaces during the chemical polishing [18] Elastic indentation modulus and hardness characterization Elastic modulus and hardness data were derived by using the Oliver and Pharr post treatment method [2]. The penetration depth was kept at approximately 100 nm where previous experiments showed the sensitivity was higher and the dispersion lower. Considering a series of ten indentations, elastic modulus mean value and its standard deviation turned out to be 70.8±0.7 and 129.4±6.3 respectively for the fused quartz and titanium. That is found to be in agreement with the literature values [19,20]. Concerning hardness, the mean value was 9.7 GPa for the fused quartz sample and 2.9 GPa for titanium for a standard deviation of 0.1 and 0.3 GPa, respectively. Experimental hardness values are in reasonable agreement with a previous study [21]. The slightly

3 J. Mater. Sci. Technol., Vol.25 No.1, Fig. 2 SPM images of imprints made on titanium in which the pile-up effect is clearly visible Fig. 3 Holding time influence on elastic modulus and hardness for titanium for three holding time (0, 10 and 50 s) higher modulus value for Ti compared to the expected one could be related to the pile-up phenomenon [8] Holding time influence From Fig. 3, it can be seen that increasing the holding time systematically induces a decrease of both elastic modulus and hardness. Elastic modulus significantly decreases as the holding time passes from 0 to 10 s and it gradually decreases in the range of 10 to 50 s duration. In contrast, the hardness of titanium was observed to decrease almost linearly for holding time comprised between 0 and 50 s. It should be noted that if 50 s holding is applied, titanium elastic indentation modulus is almost equal to the macroscopic values known from literature [19, 20] Discussion Consistent and repeatable values were found when indenting on fused quartz and on chemically polished titanium. Experiments performed in order to study the holding time effects showed that both elastic modulus and hardness decrease with an increasing holding time duration. Trying to find a plausible explanation for this behavior, a series of load-displacement curves were examined. If no holding time is applied, titanium experiences the creep phenomenon at the early stages of the unloading segment (see Fig. 4). It means that even if the load is decreased, which is the case during the unloading phase, and the tip continues to penetrate into the specimen for a few nanometers. According to the Oliver and Pharr theory, both reduced modulus and hardness depend on the contact stiffness S which is calculated by differentiating Fig. 4 Examples of the load-displacement curves obtained when investigating the holding time influence on titanium analytically the fit of the unloading curve and which represents the slope at maximum load S= dp dh p=p max, where P is the applied load and h is the penetration depth. As a result, a modification of the slope of the unloading curve induces the resulting contact stiffness to change and implies the alteration of both the elastic modulus and hardness. It is worth noting that after 50 s the effect of creep seems to be attenuated. Finally, it should be said that pile-up (see Fig. 2) certainly plays a role in the accuracy of the resulting data [2,8]. The slightly higher modulus value for Ti compared to the expected one (see part) could be related to the pile-up phenomenon. Pile-up around the indenter causes a smaller penetration and this leads to the underestimation of the real contact area. Consequently, elastic modulus and hardness derived by using the Oliver and Pharr post treatment method will be overestimated [2]. A quantification of its extent should be done in order to take its effects into account. 3.2 Nanoindentation of hard thin films The effects of roughness on elastic modulus and hardness According to surface topography measurements performed under a DEKTAK profilometer, the R a roughness is within 6 nm and 9 nm for thin films deposited on hard metal and nm for those deposited on silicon. Typical thickness of the magnetron sputtered films considered here is in the order of 1 µm. The penetration depth during the nanoindentations was chosen in order to avoid any substrate effects according to the new ISO rules. This meant to perform indentation at a depth of 1/10 of the film thickness (approximately 100 nm). This indentation depth was similar to that used for bulk Ti. Bulk sapphire single crystal was studied at first as a standard hard material imposing a contact depth around 100 nm. Then, hard TiBN and TiCrBN films deposited on hard metal (high roughness) and on single crystal silicon (low roughness) were characterized by using a Nanohardness Tester, CSM followed by Oliver-Pharr post treatment using 9-16 indentations per sample. As expected, Fig. 5 shows that there is a very high dependence of E and H standard deviation on roughness. The higher the R a roughness, the higher the spread of hardness and elastic modulus values. To elucidate the effect of roughness on data, a detailed statistical study was carried out by using

4 66 J. Mater. Sci. Technol., Vol.25 No.1, 2009 Fig. 5 Correlation between roughness and spread of elastic modulus observed during the nanoindentation experiments. The same correlation was observed for the spread of the hardness Fig. 7 Gaussian distribution of the experimental elastic modulus obtained from indentations on the TiCrBN600 thin film. Hardness showed the same kind of distribution. The oval, diamond and triangle marks are associated respectively with the imprint 1, 2 and 3 in next figure operator controlled approach and statistical one. During the operator controlled treatment the average values of hardness and elastic modulus were calculated by the mean square method using only appropriate and consistent curves. This lead to consistent and repeatable elastic modulus and hardness values. Nevertheless, it must be noted that the use of the arithmetic mean value can be appropriate for films deposited on Si substrate, but not for hard metal substrate because of the high standard deviation observed in this case and which is due to the high roughness. Then, a statistical approach based on the analysis of all the experimental data was applied. In addition, the morphology of several imprints was studied just after the test in order to find and explain such a high experimental deviation from the mean value encountered for some indentations. This part will be discussed in detail in the next section. It should also be noted that, independently of the type of nanoindenter and the treatment used, a strong effect of the R a roughness on the spread of the experimental values was observed. Fig. 6 Values of elastic modulus (a) and hardness (b) derived according to the Oliver and Pharr post treatment method using experimental data obtained with the Hysitron Inc. Triboindenter (ULB) and by the Nanohardness Tester, CSM Instruments (MISIS) a triboindenter, Hysitron Inc., and performing up to 100 indentations on each sample. This will be presented in the next paragraph. Experimental data received with both nanoindenter instruments are compared and the results are given in Fig. 6. The difference in terms of elastic modulus and hardness obtained with the two distinct apparatus was found to be within the limits of the standard deviation (cf. Fig. 6). This proves that the equipment was correctly calibrated and confirms the consistency of the Oliver and Pharr post treatment method. Two different approaches were used to determine the average value of hardness and elastic modulus: Statistical distribution of elastic modulus and hardness After a statistical analysis of one hundred indentations it was found that both elastic modulus and hardness follow a statistical distribution (see Fig. 7). Bell-shaped curve was supposed to fit experimental points by a Gaussian model using the following equation: A y = y o + w (x xc) 2 π/2 e 2 w 2 (1) where: w=2σ 2 double dispersion, x c =center of bellcurve, A=area under bell-curve, y 0 =offset. As an example, the resulting fitting curve parameters for the Ti-Cr-B-N thin film deposited on TT8K6 on hard alloy substrate is given in Table 1. It can be appreciated in Fig. 7 the presence of many points well above the Gaussian bell curve with modula up to 670 GPa. The reason for this scatter is that indentations were accidentally performed on surface irregularities which caused the distortion of the load displacement curves leading to a miscalculation of hard-

5 J. Mater. Sci. Technol., Vol.25 No.1, Table 1 Ti-Cr-B-N thin film on TT8K6 hard alloy Averaged experimental values Parameters of Gaussian fitting curve mean Std dev median Center Width Area Offset Height H/GPa E/GPa Fig. 8 Indentation on a peak (imprint 1), on a valley (imprint 2) and on a smooth surface (imprint 3). These three imprints are correlated with the diamond, triangle and circle mark in Fig. 7 ness and modulus (since they are calculated using the Oliver & Pharr post treatment method which is very sensitive to the modifications of the load-displacement curve). In this case (see Fig. 7), the arithmetic mean value derived considering all the experimental data (with the aberrant points taken into account), is different from the maximum of the Gaussian curve (which constitutes the mean value when using this approach). As a result, it can be stated that when a high number of data is collected, the Gaussian fitting curve, which is less sensitive on the aberrant data, is more accurate than the use of the arithmetic mean value and can help to reduce the effects of roughness Discussion about the contact area determination To understand why such a wide dispersion was obtained for the films deposited on hard substrate, the effects of roughness on the penetration process were analyzed combining SPM (Scanning Probe Microscopy) images with experimental results. Indentations performed on hills correspond to lower values of hardness and elastic modulus (oval mark in Fig. 7, imprint 1 in Fig. 8). Due to the limited dimension of the hills, the penetration and deformation are easier and the resulting contact depth and area are higher. This clearly leads to an underestimation of both elastic modulus and hardness knowing that (literature [2] for more details) they are inversely proportional (1/ Ac and 1/Ac respectively) to the contact area. When indentations are performed in valleys (imprint 2 in Fig. 8 and diamond mark in Fig. 7), the penetration will be more difficult because the edges of the indenter interact with the valley sides and the resulting contact depth is smaller, leading to an overestimation of the calculated elastic modulus and hardness. The ideal case is observed when indenting on smooth surfaces (imprint 3 in Fig. 8 and triangle mark in Fig. 7) for which the resulting data in terms of elastic modulus and hardness are in agreement with the maximum of the statistical distribution curve. So, it is shown by analyzing the SPM images of typical imprints that the local topography is strongly correlated with the nanoindentation results. Average values of hardness and modulus corresponded to indentations performed on flat areas of the surface, lower values were found when indenting on hillocks, and higher values were obtained when indenting on valleys. 4. Conclusions Depth-sensing nanoindentation was used to characterize the mechanical properties of standard bulk materials and thin films. Several indentations performed at a contact depth of approximately 100 nanometers confirmed that the Oliver and Pharr post treatment method [2] can be used to reliably evaluate elastic indentation modulus and hardness. Agreement was found between experimental and literature data concerning the elastic modulus of bulk titanium. Holding time influence was taken into account by applying three different holding time period. Elastic modulus and hardness were clearly found to decrease when the holding time was increased. This phenomenon was already observed for other materials [9]. Indentations on TiBN and TiCrBN nanostructured thin films deposited by magnetron sputtering of SHS targets on smooth (silicon) and rough (hard metal) substrates were performed using two different apparatus and the resulting data were in good agreement. However, an increase of the spread of the elastic modulus and hardness experimental values was observed with an increasing roughness. A statistical analysis of the nanoindentation results was proposed and it was shown that the use of the arithmetic mean value can lead to imprecision when aberrant data are present. The correlation between roughness and standard deviation was investigated by special SPM inspection of typical imprints giving different mechanical properties. It was shown that average values of hardness and modulus corresponded to indentations performed on flat areas of the surface, lower values were found when indenting on hillocks, and higher values were obtained indenting on valleys. Acknowledgements This work was supported by the Communauté Française de Belgique ARC 04/ and was done in the context of the EC VI FW international EXCELL Project. Authors are grateful to Mr. A. N. Sheveiko and

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