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SiNx nanomultilayers
Thin Solid Films 534 (2013) 367–372
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SiNx thickness dependent morphology and mechanical properties of CrAlN/SiNx nanomultilayers Wei Li a,⁎, Ping Liu a, Yongsheng Zhao b, Ke Zhang c, Fengcang Ma a, Xinkuan Liu a, Xiaohong Chen a, Daihua He a a b c
School of Materials Science and Engineering, University of Shanghai for Science and Technology, Shanghai, 200093, China School of Mechanical Engineering, University of Shanghai for Science and Technology, Shanghai, 200093, China National Engineering Research Center for Nanotechnology, Shanghai, 200241, China
a r t i c l e
i n f o
Article history: Received 13 September 2012 Received in revised form 7 February 2013 Accepted 11 February 2013 Available online 26 February 2013 Keywords: Multilayers Thin films Superlattices Microstructure X-ray diffraction Thermodynamic modeling Sputtering Nitrides
a b s t r a c t CrAlN/SiNx nanomultilayers with different SiNx layer thicknesses were synthesized by reactive magnetron sputtering. The microstructure and mechanical properties were investigated by X-ray diffraction, high-resolution transmission electron microscopy and nano-indentation techniques. The average crystallite size, microstrain and average dislocation density of CrAlN/SiNx nanomultilayers with different SiNx layer thicknesses were evaluated by X-ray diffraction line profile analysis method. The results indicated that, when SiNx layer thickness was below 0.6 nm, SiNx was forced to crystallize and grew epitaxially with CrAlN layers, resulting in the decrease of average dislocation density and the enhancement in hardness and elastic modulus. As the SiNx layer thickness further increases, the epitaxial growth was firstly interrupted and then crystallized SiNx layers transformed back to amorphous state, leading to the increase of average dislocation density and the decrease of hardness and elastic modulus. An energy balance model was established to explain the microstructure evolution. © 2013 Elsevier B.V. All rights reserved.
1. Introduction As superhard film material, nanomultilayer films have been widely investigated in the past decades [1,2]. The conception of nanomultilayers was firstly theoretically proposed by Koehler [3] in 1970 suggested that solids can be strengthened by forming a laminate structure of thin layers with large shear modulus mismatch. The strengthening effect was firstly experimentally observed in Au–Ni and Cu–Pd superlattice thin films by Yang [4] in 1977, which verified Koehler's theory. The abnormal increase of hardness in nanomultilayer can be called as the superhardness effect. Helmersson et al. [5] later reported that the superhardness effect with a maximum hardness of above 50 GPa was obtained in TiN/VN nanomultilayer, indicating that nanomultilayers had a significant research value and bright application prospect. Subsequently, the superhardness effect has been found in many nanomultilayer systems and its generation prerequisite has been deeply studied [6–12]. Representative nanomultilayers are crystal/amorphous system, such as TiN/SiC [13], TiAlN/SiO2 [14] and CrN/Si3N4 [15] nanomultilayers, etc., in which amorphous SiC, SiO2 and Si3N4 layers can be crystallized below a critical thickness (generally less than 1 nm) under template effects of crystalline TiN, TiAlN and CrN layers, and grown coherently with crystal ⁎ Corresponding author. Tel./fax: +86 21 55271682. E-mail addresses: [email protected] (W. Li), [email protected] (K. Zhang). 0040-6090/$ – see front matter © 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.tsf.2013.02.049
layers. Then superhardness effect appears in nanomultilayers. As the thicknesses of intrinsically amorphous layers increase, these layers cannot keep their crystallization state and change back to amorphous state, leading to the destruction of epitaxial growth structure and decrease of hardness. However, during the disappearance of superhardness effect, the relation and sequence between the damage of the epitaxial growth and the transformation of crystallized layers back to amorphous state, namely, the details of microstructural evolution, have seldom been studied, which is important to figure out the mechanism of superhardness effect. X-ray diffraction line profile analysis is an effective method to evaluate the microstructural parameters for nanocrystalline materials, including the average crystallite size, microstrain, and average dislocation density etc., which has been widely used in microstructural characterization for nanocrystalline materials [16–19]. However, the microstructural characterization for nanomultilayers through X-ray diffraction line profile analysis has rarely been documented, which is helpful to figure out the microstructural evolution for nanomultilayers. To this end, the crystalline CrAlN with fcc (facecentered cubic) structure and amorphous SiNx are selected to form CrAlN/SiNx nanomultilayer in this paper. The microstructural parameters of the CrAlN/SiNx nanomultilayers with different SiNx layer thicknesses are evaluated by X-ray diffraction line profile analysis method. It is worth noting that Tsai and Duh recently reported the microstructure and mechanical properties of CrAlN/SiNx nanomultilayered films [20]. Unfortunately, they did not give the details of
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microstructural evolution during disappearance of superhardness effect. In this investigation, the correlations of microstructure and mechanical property with SiNx layer thickness are systemically studied by means of X-ray diffraction line profile analysis and other experimental techniques. Particular attention is paid to the microstructural evolution during disappearance of superhardness effect, which will be explained from the energy balance aspect. 2. Experimental details CrAlN/SiNx nanomultilayers were fabricated on silicon substrates by reactive magnetron sputtering system. The CrAlN layers were deposited from a Cr50Al50 alloy target (at.%, 99.99%) by DC sputter deposition with balanced magnetic field configuration and the power was set at 120 W; the SiNx layers were deposited from a Si target (99.99%) by RF sputtering with frequency of 13.56 MHz and the power was set at 80 W. Both CrAl and Si targets were 75 mm in diameter. The base pressure was 5.0 × 10 −4 Pa and the substrate temperature was 300 °C. The Ar and N2 flow rate were 15 and 5 sccm respectively, and the working pressure was 0.2 Pa. The sputtering rates for CrAlN and SiNx layers are respectively 0.5 nm/s and 0.2 nm/s. Before deposition, silicon substrates were ultrasonically cleaned in acetone and alcohol, and further sputtered for 15 min for pretreatment. The configuration of CrAlN/SiNx nanomultilayers was designed with CrAlN layers with a fixed thickness of 5.0 nm along with variable SiNx layer thickness ranging from 0.4 to 1.2 nm. The individual layer thickness was controlled by the switch time of alternate shutters which were modulated by a programmable logic control. The thickness of all films was about 2 μm. The microstructures of CrAlN/SiNx nanomultilayers were investigated by field emission high-resolution transmission electron microscopy (HRTEM) using a Philips CM200-FEG. The preparation procedures of cross-section specimen for TEM observation are as follows. The films with substrate were cut into two pieces and adhered face to face, which were subsequently cut at the joint position to make a slice. The slices were thinned by mechanical polishing followed by argon ion milling. The X-ray diffraction (XRD) profiles were measured by a D/max-2550 X-ray diffractometer (18 kW) with Cu Kα radiation. The step size and step time were 0.02° and 5 s respectively. The instrumental correction was made by the powder pattern of a Si standard and Stokes correction procedure [21]. The selected reflection of CrAlN/SiNx nanomultilayer, (200), was evaluated for the average crystallite size (D), root mean square (r.m.s.) microstrain and average dislocation density (ρ) by XRD line profile analysis using the modified Warren–Averbach method [22,23], which is described in detail in Section 3. The hardness (H) and elastic modulus (E) of the films were measured by a MTS G200 nanoindenter by using the Oliver and Pharr method [24]. The measurements were performed by using a Berkovich diamond tip with a load of 5 mN. The indentation depth was about 100 nm, less than 1/10th of the film thickness to minimize the effect of substrate on the measurements. Each hardness or elastic modulus value was an average of at least 10 measurements. 3. Evaluation of the X-ray diffraction line profiles Warren–Averbach Fourier analysis is a most useful method to measure microstructural parameters for nanocrystalline materials, including the average crystallite size, microstrain, and average dislocation density. Based on the Warren–Averbach Fourier method [25], which provides the apparent crystallite size and the mean square strain, Wilkens [26] proposed a model of “restrictedly random” distribution of dislocations to analyze line profiles, the dislocations density ( ρ), configuration parameter ( M), strain field range ( Re ), stored elastic energy ( E V ) could be deduced for monocrystalline-copper.
Considering the idea suggested by Langford [27] that each profile could be regarded as a convolution of several Gaussian functions with Cauchy ones, Wang et al. [28,29] worked out a practical procedure and standard curves for line profile analysis of face-centered cubic (fcc), body-centered cubic (bcc) and hexagonally close packed (hcp) polycrystals. However, two orders of reflections are needed in this method, which makes it unavailable for the case when it can only observe the high-quality first-order reflection due to preferred orientation or broadening of diffraction peak, such as CrAlN/SiNx nanomultilayer prepared by magnetron sputtering in this investigation. As a result, the combination of single-peak Fourier analysis proposed by Mignot and Rondot [30] and Wang et al.'s method [28,29] is used in this work. The fitting procedures are given as follows. By Stokes' deconvolution of XRD profile, the physical broadened profiles and its corresponding Fourier coefficients A(L) are obtained, and then the two components, “particle” coefficients A p(L) and “strain” coefficients A s(L), are separated as: p
A ðLÞ ¼ a−
L : D
D E1 0 2 2 2 −2πm L εL A: A ðLÞ ¼ exp@ d2 s
ð1Þ
ð2Þ
In the analysis of X-ray diffraction line profile, a hook is often been observed in the particle size coefficients (A p(L)) versus Fourier length (L) plots, which is called the “hook effect”. In Eqs. (1) and (2), a is the quantity expressing the “hook” effect, L the reciprocal of the diffraction vector, also called the Fourier length, D the average coherent domain size, m the order of reflection, 〈εL2〉 1/2 the r.m.s. microstrain, d the interplanar spacing of the reflection plane. Combining with the method proposed by Mignot and Rondot [30], D and 〈εL2〉 1/2 can be obtained by non-linear fitting. According to Wang's theory [28], “strain” coefficients A s(L) can be expressed as: s 2 2 A ðLÞ ¼ exp −2βc L−πβg L
ð3Þ
where βc and βg are the Cauchy and Gaussian widths of “strain” broadened profile, which can be gained from Eqs. (1) and (3) by non-linear fitting. Making use of the standard curves and procedures described in [29], the average dislocation density ( ρ) can been obtained. 4. Results and discussion 4.1. Microstructure of CrAlN/SiNx nanomultilayers with different SiNx layer thicknesses Low-angle patterns of CrAlN/SiNx nanomultilayers with different SiNx layer thicknesses (tSiNx) are shown in Fig. 1. The low-angle diffraction peaks are clearly observed when tSiNx was over 0.4 nm, indicating that the nanomultilayers have well modulated structures. The modulation period of nanomultilayers can be calculated using the modified Bragg formula [31]. As the thickness of CrAlN layer (tCrAlN) is kept at 5.0 nm for all the samples, tSiNx can be evaluated, which was verified by the subsequent HRTEM observations and calculated values from depositing time and rate derived from the monolayer, as indexed in the Fig. 1. The high-angle XRD patterns of monolithic film and CrAlN/SiNx nanomultilayers with different tSiNx are shown in Fig. 2. It can be seen that both monolithic CrAlN film and CrAlN/SiNx nanomultilayers exhibit fcc B1-NaCl CrN structure. No other phase except for Si substrate can be detected. The monolithic CrAlN film presents the (111) preferred orientation, while the intensities of (111) and (200)
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Fig. 1. Low-angle patterns of CrAlN/SiNx nanomultilayers with different SiNx layer thicknesses.
Fig. 2. High-angle patterns of CrAlN/SiNx nanomultilayers with different SiNx layer thicknesses as well as monolithic CrAlN film.
diffraction peaks are comparatively close for CrAlN/SiNx nanomultilayers. With the initial increase of tSiNx, the intensity of (111) and (200) peaks gradually increases, and reaches the maximum value when tSiNx is 0.6 nm. With further increase of tSiNx, however, the peak intensities gradually decrease and the peak shapes broaden. When crystal/amorphous nanomultilayer films grow by alternately depositing crystal layer and amorphous layer. Amorphous layers can be crystallized to crystal structure in nanomultilayer and grow epitaxially with crystal layers when amorphous layer thickness is
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below a critical thickness. As the thickness of amorphous layer increases, however, they cannot keep their crystallization state and change back to amorphous state [14–16,32,33]. Therefore, it can be deduced that, when tSiNx is below 0.6 nm, amorphous SiNx layers are inclined to form the crystalline structure and grow epitaxially with CrAlN layers, which improves the crystallization integrity of nanomultilayers and therefore increases the intensities of diffraction peaks. When tSiNx exceeds 0.6 nm, however, the improvement effect on crystallization integrity gradually disappears and the intensities of diffraction peaks accordingly decrease. It can also be seen from Fig. 2 that most nanomultilayer samples show a shift towards high angle in (111) and (200) diffraction peaks compared with the monolithic CrAlN film. The shift amount gradually increases with the initial increase of tSiNx and reaches the maximum value when tSiNx is 0.6 nm. With further increase of tSiNx, however, the shift amount gradually decreases. The shifts correspond to the change of the lattice parameter in CrAlN and SiNx layers. Under the epitaxial growth structure, the originally larger lattice parameter of CrAlN layer is inclined to decrease, leading to generation of interfacial compressive stress in substrate parallel direction, while the lattice parameter of SiNx layer is forced to increase, resulting in formation of interfacial tensile stress in substrate parallel direction. Therefore, the alternating compressive and tensile stress fields create along the growth direction of nanomultilayer, which can hinder the dislocation motion and strengthen the nanomultilayer [34]. Fig. 3 shows the cross-sectional HRTEM images of CrAlN/SiNx nanomultilayers with different tSiNx. It can be clearly seen from Fig. 3(a) that, when tSiNx is 0.6 nm, the lattice fringes continuously go through several modulation layers and interfaces, indicating that SiNx layers have been fully crystallized and grown epitaxially with CrAlN layers. Under the epitaxial growth structure, the originally larger lattice parameter of CrAlN layer is inclined to decrease, leading to generation of interfacial compressive stress, while the lattice parameter of SiNx layer is forced to increase, resulting in formation of interfacial tensile stress. When tSiNx increases to 0.8 nm, the whole epitaxial growth cannot be maintained between the CrAlN layers, while SiNx layer can still be present in a crystallized state, as shown in Fig. 3(b). The different CrAlN layers denoted by A, B and C, exhibit to grow along different orientations. The two different growth orientations are even observed in a CrAlN layer marked C, suggesting that different crystal domains have created in a single CrAlN layer and the epitaxial growth structure between CrAlN layers has been broken. When tSiNx reaches 1.2 nm, as shown in Fig. 3(c), the SiNx layers cannot keep their crystallization and change back to amorphous state. The interfacial morphology between the CrAlN layer and SiNx layer becomes undistinguishable, and more crystal domains with different misorientations are formed in a single CrAlN layer, suggesting that the epitaxial growth structure has been totally broken.
Fig. 3. Cross-sectional HRTEM images of CrAlN/SiNx nanomultilayers with SiNx layer thickness of (a) 0.6 nm, (b) 0.8 nm and (c) 1.2 nm.
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Table 1 The average crystallite size (D), r.m.s. microstrain and average dislocation density (ρ) of CrAlN/SiNx nanomultilayers with different SiNx layer thicknesses, measured by XRD line profile analysis. SiNx layer thickness (nm)
(hkl)
D (nm)
bε2 >
0.4 0.6 0.8 1.0 1.2
200 200 200 200 200
17.2 17.5 13.6 9.4 6.5
6.8 4.7 8.4 9.9 14.4
1/2
× 10−3
ρ (m-2) 5.7 2.7 8.6 12.1 26.6
× × × × ×
1015 1015 1015 1015 1015
4.2. Microstructural characterization of CrAlN/SiNx nanomultilayers by XRD line profile analysis By using our modified Warren–Averbach method, the average crystallite size (D), r.m.s. microstrain and average dislocation density (ρ) of CrAlN/SiNx nanomultilayers with different tSiNx are determined, as shown in Table 1. It can be seen that the average crystallite sizes (D) of various CrAlN/SiNx nanomultilayers are larger than modulation periods (Λ), and with the increase of tSiNx, the average crystallite size gradually decrease to the value of modulation period. The average dislocation density ( ρ) firstly decreases with the initial increase of tSiNx and reaches the minimum value when tSiNx is 0.6 nm, then increases with further increase of tSiNx. The phenomenon for Λ> D can be explained by distinguishing mode of XRD technique. In the measurements of crystallite size, XRD technique distinguishes crystallite with a certain misorientation [35]. In this investigation, however, due to the epitaxial growth between CrAlN and SiNx layer, the misorientation between different modulation periods is so small that it cannot be differentiated by XRD, resulting that the average crystallite size exceeds the modulation period. The initial decrease of the average dislocation density can be attributed to the fact that SiNx layers not only can be fully crystallized under the heteroepitaxy of CrAlN layers, but also the crystallized SiNx layers can improve the crystallization integrity of CrAlN layers [36]. This mutually promoting growth effect is verified by the XRD patterns in Fig. 2, in which the CrAlN/SiNx nanomultilayer with tSiNx of 0.6 nm exhibits a higher peak intensity than that with tSiNx of 0.4 nm. When tSiNx further increases, the epitaxial growth cannot be maintained between the CrAlN layers, and the misorientations between different CrAlN layers become large, resulting in the decrease of the average crystallite size. However, since SiNx layer still keeps the crystallization state, the CrAlN layer can retain the epitaxial growth with neighboring SiNx layer, even with the adjacent CrAlN layer in some regions, as shown in Fig. 3(b), causing that the average crystallite size from XRD still exceeds the modulation period. Due to the destruction of the epitaxial growth between the CrAlN layers and formation of different crystal domains within a CrAlN layer, the
crystallization integrity of CrAlN/SiNx nanomultilayer decreases, leading to the increase of the average dislocation density. When tSiNx reaches 1.2 nm, SiNx layers change back to amorphous state and the epitaxial growth is totally broken. More crystal domains with different misorientations are created within a single CrAlN layer, causing the average crystallite size to further reduce to 6.5 nm, close to the value of modulation period. The interfaces between the crystalline CrAlN layer and amorphous SiNx layer can generate more dislocations, and the misorientations between different crystal domains are also needed to adjust through dislocations [37]. Therefore, the average dislocation density further increases to the high value of 26.6 × 10 15 m −2. 4.3. Energy balance model of microstructure evolution for CrAlN/SiNx nanomultilayers The microstructural evolution of CrAlN/SiNx nanomultilayers with increase of tSiNx can be illustrated by Fig. 4. As tSiNx is less than 0.6 nm, SiNx layer are fully crystallized and grown epitaxially with CrAlN layers, as shown in Fig. 4(a). When tSiNx increases to 0.8 nm, although SiNx layers still present in a crystallized state, the epitaxial growth cannot be maintained between CrAlN layers and crystal domains may create within a CrAlN layer, as illustrated in Fig. 4(b). When tSiNx reaches 1.2 nm, SiNx layers transform into amorphous state, and the epitaxial growth structure is totally broken, and more crystal domains with different misorientations are created within a single CrAlN layer, as presented in Fig. 4(c). Such a microstructural evolution of CrAlN/SiNx nanomultilayers can be explained by an energy balance model including strain-free bulk energy (EB), coherency strain energy (ES) and interfacial energy (Ei), in which the total energy of CrAlN/SiNx nanomultilayer per unit (ET), can be written as: SiN SiN CrAlN CrAlN þ Ei þ ES ET ¼ EB x þ ES x t SiNx þ EB SiN
SiN
ð4Þ
where EB x and ES x are respectively the strain-free bulk energy and coherency strain energy of SiNx layers; EBCrAlN and ESCrAlN are respectively the strain-free bulk energy and coherency strain energy of CrAlN layers, in which ESCrAlN increases with tSiNx; Ei is interfacial energy between CrAlN layer and SiNx layer. During the initial growth of SiNx layer, tSiNx approaches zero, ET is therefore composed of EBCrAlN, ESCrAlN and Ei. Since EBCrAlN is invariable and ESCrAlN is comparatively small when tSiNx approaches zero, ET is dominated by Ei. Formation of coherent interface between CrAlN and SiNx layer can lower Ei. Therefore, SiNx layers are crystallized and grow epitaxially on top of CrAlN layers. When tSiNx increases to 0.8 nm, the total energy of SiNx layer increase. However, the total
Fig. 4. Schematic illustration of microstructural evolution of CrAlN/SiNx nanomultilayers with SiNx layer thickness of (a) less than 0.6 nm, (b) 0.8 nm and (c) 1.2 nm.
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elastic modulus decrease, the average crystallite size gradually decreases according to Table 1. Therefore, the strengthening of CrAlN/ SiNx nanomultilayers is not connected to the grain refinement. 5. Conclusions
Fig. 5. Variation of hardness and elastic modulus of CrAlN/SiNx nanomultilayers with change of SiNx layer thickness.
energy of SiNx layer is not large enough to exceed Ei. Therefore, the SiNx layers still maintain the crystallization state and keep the local epitaxial growth with the adjacent CrAlN layer. However, with increase of tSiNx, the coherency strain energy of the whole CrAlN layers, ESCrAlN, grows. Therefore, the whole epitaxial growth between CrAlN layers cannot be maintained, and even coherent growth cannot be kept within a CrAlN layer, leading to the formation of crystal domains in a CrAlN layer. As tSiNx reaches 1.2 nm, the total energy of SiNx layer further increases, which have more dominant effect on ET than Ei. x x Since ESiN and ESiN of crystalline SiNx are higher than those of amorB S phous SiNx, crystalline SiNx structure has to transform into amorphous state. In addition, ESCrAlN further increases with the increase of tSiNx, resulting in total destruction of epitaxial growth structure and formation of more crystal domains within a CrAlN layer. 4.4. Mechanical properties of CrAlN/SiNx nanomultilayers Fig. 5 shows the dependence of hardness and elastic modulus of CrAlN/SiNx nanomultilayers on SiNx layer thickness. The hardness and elastic modulus of CrAlN monolithic film deposited under the same conditions were measured as 28.7 GPa and 382 GPa, while those of SiNx monolithic film were 18 GPa and 225 GPa. It can be seen that, after SiNx layers are initially inserted, the hardness and elastic modulus of the nanomultilayers increase. When tSiNx is 0.6 nm, the maximum values of hardness and elastic modulus reach 37.6 GPa and 437 GPa, respectively. As tSiNx further increases, the hardness and elastic modulus values gradually decrease, and are below those of monolithic CrAlN film when tSiNx is 1.2 nm. The change of hardness and elastic modulus is closely related to the microstructure of nanomultilayers. When tSiNx is below 0.6 nm, SiNx layers are forced to adopt a the crystalline structure under the template effect of CrAlN layers and grow epitaxially with CrAlN layers. The enhancement in hardness and elastic modulus can be explained by Koehler's modulus-difference strengthening theory [1] and the alternating-stress strengthening theory [38], in which the dislocation motions are hindered at interfaces in nanomultilayers respectively by forces generated from two layers with different shear moduli and the alternating stress fields. When tSiNx exceeds 0.6 nm, however, the hardness and elastic modulus values gradually decrease. Based on the above analysis, the epitaxial growth structure between CrAlN layers is firstly damaged while SiNx layers still maintain the crystallization state, suggesting that the epitaxial growth structure is the key point to achieve the superhardness effect, while the crystallization of SiNx layers is not a requirement. As tSiNx further increases to 1.2 nm, the SiNx layers change to amorphous state, and the epitaxial growth is totally broken, leading to the further decrease of hardness and elastic modulus. It is worth noting that, as the hardness and
In summary, the microstructure and mechanical properties of CrAlN/SiNx nanomultilayers were investigated by XRD, HRTEM, XRD line profile analysis method and nano-indentation techniques. When tSiNx was below 0.6 nm, SiNx layers are forced to crystallize and grow epitaxially on CrAlN layers, resulting in the enhancement in hardness and elastic modulus, which respectively reach maximum values of 37.6 GPa and 437 GPa when tSiNx is 0.6 nm. When tSiNx increases to 0.8 nm, the epitaxial growth structure between CrAlN layers is damaged, while SiNx layers still maintain the crystallized state, resulting in the increase of average dislocation density and decrease of hardness and elastic modulus. As tSiNx reaches 1.2 nm, the SiNx layers change to amorphous state, and the epitaxial growth is totally broken, resulting in further increase of average dislocation density and decrease of mechanical properties. Based on the established energy balance expression, with the joint effects of interfacial energy, strain-free bulk energy and coherency strain energy of CrAlN and SiNx layers, the damage of the epitaxial growth precedes the transformation of the crystallized SiNx layers to amorphous state during the disappearance of superhardness effect. Acknowledgment The present work was financially supported by the National Natural Science Foundation of China under grant no. 51101101, “Shanghai Municipal Natural Science Foundation” under grant no. 11ZR1424600, “Innovation Program of Shanghai Municipal Education Commission” under grant no. 12YZ104, “Shanghai Leading Academic Discipline Project” under grant no. J50503 sponsored by Shanghai municipal education commission in China. We thank Dr. Weizong Xu for calculational assistance in XRD line profile analysis. References [1] S.A. Barnett, M. Shinn, Annu. Rev. Mater. Sci. 24 (1994) 481. [2] G. Abadias, A. Michel, C. Tromas, C. Jaouen, S.N. Dub, Surf. Coat. Technol. 202 (2007) 844. [3] J.S. Koehler, Phys. Rev. B 2 (1970) 547. [4] W.M. Yang, T. Tsakalakos, J.E. Hilliard, J. Appl. Phys. 48 (1977) 876. [5] U. Helmersson, S. Todorova, S.A. Barnett, J. Appl. Phys. 62 (1987) 481. [6] A. Madan, I.W. Kim, S.C. Cheng, P. Yashar, V.P. Dravid, S.A. Barnett, Phys. Rev. Lett. 78 (1997) 1743. [7] C. Ziebert, S. Ulrich, J. Vac. Sci. Technol. A 24 (2006) 554. [8] M. Stueber, H. Holleck, H. Leiste, K. Seemann, S. Ulrich, C. Ziebert, J. Alloy. Compd. 483 (2009) 321. [9] H.C. Barshilia, K.S. Rajam, A. Jain, K. Gopinadhan, S. Chaudhary, Thin Solid Films 503 (2006) 158. [10] J. Musil, Surf. Coat. Technol. 125 (2000) 322. [11] L. Wei, F.H. Mei, N. Shao, M. Kong, G.Y. Li, Appl. Phys. Lett. 86 (2005) 021919. [12] R.A. Andrievski, Surf. Coat. Technol. 201 (2007) 6112. [13] J.J. Lao, N. Shao, F.H. Mei, G.Y. Li, M.Y. Gu, Appl. Phys. Lett. 86 (2005) 011902. [14] W. Li, P. Liu, J.T. Wang, F.C. Ma, X.K. Liu, X.H. Chen, Mater. Lett. 65 (2011) 636. [15] J.H. Xu, K. Hattori, Y. Seino, I. Kojima, Thin Solid Films 414 (2002) 239. [16] L. Li, T. Ungar, Y.D. Wang, J.R. Morris, G. Tichy, J. Lendvai, Y.L. Yang, Y. Ren, H. Choo, P.K. Liaw, Acta Mater. 57 (2009) 4988. [17] J. Gubicza, N.Q. Chinh, T. Csanadi, T.G. Langdon, T. Ungar, Mater. Sci. Eng. A 462 (2007) 86. [18] T. Ungar, O. Castelnau, G. Ribarik, M. Drakopoulos, J.L. Bechade, T. Chauveau, A. Snigirev, I. Snigireva, C. Schroer, B. Bacroix, Acta Mater. 55 (2007) 1117. [19] J. Gubicza, L. Balogh, R.J. Hellmig, Y. Estrin, T. Ungar, Mater. Sci. Eng. A 400/401 (2005) 334. [20] S.H. Tsai, J.G. Duh, Thin Solid Films 518 (2009) 1480. [21] A.R. Stokes, Proc. Phys. Soc. 61 (1948) 382. [22] W. Li, W.Z. Xu, X.D. Wang, Y.H. Rong, J. Alloy. Compd. 474 (2009) 546. [23] W. Li, P. Liu, F.C. Ma, Y.H. Rong, J. Mater. Sci. 44 (2009) 2925. [24] W.C. Oliver, G.M. Pharr, J. Mater. Res. 7 (1992) 1564. [25] B.E. Warren, B.L. Averbach, J. Appl. Phys. 21 (1950) 595. [26] M. Wilkens, Phys. Status Solidi A 2 (1970) 359. [27] J.I. Langford, J. Appl. Crystallogr. 11 (1978) 10.
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SiNx thickness dependent morphology and mechanical properties of CrAlN/SiNx nanomultilayers Wei Li a,⁎, Ping Liu a, Yongsheng Zhao b, Ke Zhang c, Fengcang Ma a, Xinkuan Liu a, Xiaohong Chen a, Daihua He a a b c
School of Materials Science and Engineering, University of Shanghai for Science and Technology, Shanghai, 200093, China School of Mechanical Engineering, University of Shanghai for Science and Technology, Shanghai, 200093, China National Engineering Research Center for Nanotechnology, Shanghai, 200241, China
a r t i c l e
i n f o
Article history: Received 13 September 2012 Received in revised form 7 February 2013 Accepted 11 February 2013 Available online 26 February 2013 Keywords: Multilayers Thin films Superlattices Microstructure X-ray diffraction Thermodynamic modeling Sputtering Nitrides
a b s t r a c t CrAlN/SiNx nanomultilayers with different SiNx layer thicknesses were synthesized by reactive magnetron sputtering. The microstructure and mechanical properties were investigated by X-ray diffraction, high-resolution transmission electron microscopy and nano-indentation techniques. The average crystallite size, microstrain and average dislocation density of CrAlN/SiNx nanomultilayers with different SiNx layer thicknesses were evaluated by X-ray diffraction line profile analysis method. The results indicated that, when SiNx layer thickness was below 0.6 nm, SiNx was forced to crystallize and grew epitaxially with CrAlN layers, resulting in the decrease of average dislocation density and the enhancement in hardness and elastic modulus. As the SiNx layer thickness further increases, the epitaxial growth was firstly interrupted and then crystallized SiNx layers transformed back to amorphous state, leading to the increase of average dislocation density and the decrease of hardness and elastic modulus. An energy balance model was established to explain the microstructure evolution. © 2013 Elsevier B.V. All rights reserved.
1. Introduction As superhard film material, nanomultilayer films have been widely investigated in the past decades [1,2]. The conception of nanomultilayers was firstly theoretically proposed by Koehler [3] in 1970 suggested that solids can be strengthened by forming a laminate structure of thin layers with large shear modulus mismatch. The strengthening effect was firstly experimentally observed in Au–Ni and Cu–Pd superlattice thin films by Yang [4] in 1977, which verified Koehler's theory. The abnormal increase of hardness in nanomultilayer can be called as the superhardness effect. Helmersson et al. [5] later reported that the superhardness effect with a maximum hardness of above 50 GPa was obtained in TiN/VN nanomultilayer, indicating that nanomultilayers had a significant research value and bright application prospect. Subsequently, the superhardness effect has been found in many nanomultilayer systems and its generation prerequisite has been deeply studied [6–12]. Representative nanomultilayers are crystal/amorphous system, such as TiN/SiC [13], TiAlN/SiO2 [14] and CrN/Si3N4 [15] nanomultilayers, etc., in which amorphous SiC, SiO2 and Si3N4 layers can be crystallized below a critical thickness (generally less than 1 nm) under template effects of crystalline TiN, TiAlN and CrN layers, and grown coherently with crystal ⁎ Corresponding author. Tel./fax: +86 21 55271682. E-mail addresses: [email protected] (W. Li), [email protected] (K. Zhang). 0040-6090/$ – see front matter © 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.tsf.2013.02.049
layers. Then superhardness effect appears in nanomultilayers. As the thicknesses of intrinsically amorphous layers increase, these layers cannot keep their crystallization state and change back to amorphous state, leading to the destruction of epitaxial growth structure and decrease of hardness. However, during the disappearance of superhardness effect, the relation and sequence between the damage of the epitaxial growth and the transformation of crystallized layers back to amorphous state, namely, the details of microstructural evolution, have seldom been studied, which is important to figure out the mechanism of superhardness effect. X-ray diffraction line profile analysis is an effective method to evaluate the microstructural parameters for nanocrystalline materials, including the average crystallite size, microstrain, and average dislocation density etc., which has been widely used in microstructural characterization for nanocrystalline materials [16–19]. However, the microstructural characterization for nanomultilayers through X-ray diffraction line profile analysis has rarely been documented, which is helpful to figure out the microstructural evolution for nanomultilayers. To this end, the crystalline CrAlN with fcc (facecentered cubic) structure and amorphous SiNx are selected to form CrAlN/SiNx nanomultilayer in this paper. The microstructural parameters of the CrAlN/SiNx nanomultilayers with different SiNx layer thicknesses are evaluated by X-ray diffraction line profile analysis method. It is worth noting that Tsai and Duh recently reported the microstructure and mechanical properties of CrAlN/SiNx nanomultilayered films [20]. Unfortunately, they did not give the details of
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microstructural evolution during disappearance of superhardness effect. In this investigation, the correlations of microstructure and mechanical property with SiNx layer thickness are systemically studied by means of X-ray diffraction line profile analysis and other experimental techniques. Particular attention is paid to the microstructural evolution during disappearance of superhardness effect, which will be explained from the energy balance aspect. 2. Experimental details CrAlN/SiNx nanomultilayers were fabricated on silicon substrates by reactive magnetron sputtering system. The CrAlN layers were deposited from a Cr50Al50 alloy target (at.%, 99.99%) by DC sputter deposition with balanced magnetic field configuration and the power was set at 120 W; the SiNx layers were deposited from a Si target (99.99%) by RF sputtering with frequency of 13.56 MHz and the power was set at 80 W. Both CrAl and Si targets were 75 mm in diameter. The base pressure was 5.0 × 10 −4 Pa and the substrate temperature was 300 °C. The Ar and N2 flow rate were 15 and 5 sccm respectively, and the working pressure was 0.2 Pa. The sputtering rates for CrAlN and SiNx layers are respectively 0.5 nm/s and 0.2 nm/s. Before deposition, silicon substrates were ultrasonically cleaned in acetone and alcohol, and further sputtered for 15 min for pretreatment. The configuration of CrAlN/SiNx nanomultilayers was designed with CrAlN layers with a fixed thickness of 5.0 nm along with variable SiNx layer thickness ranging from 0.4 to 1.2 nm. The individual layer thickness was controlled by the switch time of alternate shutters which were modulated by a programmable logic control. The thickness of all films was about 2 μm. The microstructures of CrAlN/SiNx nanomultilayers were investigated by field emission high-resolution transmission electron microscopy (HRTEM) using a Philips CM200-FEG. The preparation procedures of cross-section specimen for TEM observation are as follows. The films with substrate were cut into two pieces and adhered face to face, which were subsequently cut at the joint position to make a slice. The slices were thinned by mechanical polishing followed by argon ion milling. The X-ray diffraction (XRD) profiles were measured by a D/max-2550 X-ray diffractometer (18 kW) with Cu Kα radiation. The step size and step time were 0.02° and 5 s respectively. The instrumental correction was made by the powder pattern of a Si standard and Stokes correction procedure [21]. The selected reflection of CrAlN/SiNx nanomultilayer, (200), was evaluated for the average crystallite size (D), root mean square (r.m.s.) microstrain and average dislocation density (ρ) by XRD line profile analysis using the modified Warren–Averbach method [22,23], which is described in detail in Section 3. The hardness (H) and elastic modulus (E) of the films were measured by a MTS G200 nanoindenter by using the Oliver and Pharr method [24]. The measurements were performed by using a Berkovich diamond tip with a load of 5 mN. The indentation depth was about 100 nm, less than 1/10th of the film thickness to minimize the effect of substrate on the measurements. Each hardness or elastic modulus value was an average of at least 10 measurements. 3. Evaluation of the X-ray diffraction line profiles Warren–Averbach Fourier analysis is a most useful method to measure microstructural parameters for nanocrystalline materials, including the average crystallite size, microstrain, and average dislocation density. Based on the Warren–Averbach Fourier method [25], which provides the apparent crystallite size and the mean square strain, Wilkens [26] proposed a model of “restrictedly random” distribution of dislocations to analyze line profiles, the dislocations density ( ρ), configuration parameter ( M), strain field range ( Re ), stored elastic energy ( E V ) could be deduced for monocrystalline-copper.
Considering the idea suggested by Langford [27] that each profile could be regarded as a convolution of several Gaussian functions with Cauchy ones, Wang et al. [28,29] worked out a practical procedure and standard curves for line profile analysis of face-centered cubic (fcc), body-centered cubic (bcc) and hexagonally close packed (hcp) polycrystals. However, two orders of reflections are needed in this method, which makes it unavailable for the case when it can only observe the high-quality first-order reflection due to preferred orientation or broadening of diffraction peak, such as CrAlN/SiNx nanomultilayer prepared by magnetron sputtering in this investigation. As a result, the combination of single-peak Fourier analysis proposed by Mignot and Rondot [30] and Wang et al.'s method [28,29] is used in this work. The fitting procedures are given as follows. By Stokes' deconvolution of XRD profile, the physical broadened profiles and its corresponding Fourier coefficients A(L) are obtained, and then the two components, “particle” coefficients A p(L) and “strain” coefficients A s(L), are separated as: p
A ðLÞ ¼ a−
L : D
D E1 0 2 2 2 −2πm L εL A: A ðLÞ ¼ exp@ d2 s
ð1Þ
ð2Þ
In the analysis of X-ray diffraction line profile, a hook is often been observed in the particle size coefficients (A p(L)) versus Fourier length (L) plots, which is called the “hook effect”. In Eqs. (1) and (2), a is the quantity expressing the “hook” effect, L the reciprocal of the diffraction vector, also called the Fourier length, D the average coherent domain size, m the order of reflection, 〈εL2〉 1/2 the r.m.s. microstrain, d the interplanar spacing of the reflection plane. Combining with the method proposed by Mignot and Rondot [30], D and 〈εL2〉 1/2 can be obtained by non-linear fitting. According to Wang's theory [28], “strain” coefficients A s(L) can be expressed as: s 2 2 A ðLÞ ¼ exp −2βc L−πβg L
ð3Þ
where βc and βg are the Cauchy and Gaussian widths of “strain” broadened profile, which can be gained from Eqs. (1) and (3) by non-linear fitting. Making use of the standard curves and procedures described in [29], the average dislocation density ( ρ) can been obtained. 4. Results and discussion 4.1. Microstructure of CrAlN/SiNx nanomultilayers with different SiNx layer thicknesses Low-angle patterns of CrAlN/SiNx nanomultilayers with different SiNx layer thicknesses (tSiNx) are shown in Fig. 1. The low-angle diffraction peaks are clearly observed when tSiNx was over 0.4 nm, indicating that the nanomultilayers have well modulated structures. The modulation period of nanomultilayers can be calculated using the modified Bragg formula [31]. As the thickness of CrAlN layer (tCrAlN) is kept at 5.0 nm for all the samples, tSiNx can be evaluated, which was verified by the subsequent HRTEM observations and calculated values from depositing time and rate derived from the monolayer, as indexed in the Fig. 1. The high-angle XRD patterns of monolithic film and CrAlN/SiNx nanomultilayers with different tSiNx are shown in Fig. 2. It can be seen that both monolithic CrAlN film and CrAlN/SiNx nanomultilayers exhibit fcc B1-NaCl CrN structure. No other phase except for Si substrate can be detected. The monolithic CrAlN film presents the (111) preferred orientation, while the intensities of (111) and (200)
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Fig. 1. Low-angle patterns of CrAlN/SiNx nanomultilayers with different SiNx layer thicknesses.
Fig. 2. High-angle patterns of CrAlN/SiNx nanomultilayers with different SiNx layer thicknesses as well as monolithic CrAlN film.
diffraction peaks are comparatively close for CrAlN/SiNx nanomultilayers. With the initial increase of tSiNx, the intensity of (111) and (200) peaks gradually increases, and reaches the maximum value when tSiNx is 0.6 nm. With further increase of tSiNx, however, the peak intensities gradually decrease and the peak shapes broaden. When crystal/amorphous nanomultilayer films grow by alternately depositing crystal layer and amorphous layer. Amorphous layers can be crystallized to crystal structure in nanomultilayer and grow epitaxially with crystal layers when amorphous layer thickness is
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below a critical thickness. As the thickness of amorphous layer increases, however, they cannot keep their crystallization state and change back to amorphous state [14–16,32,33]. Therefore, it can be deduced that, when tSiNx is below 0.6 nm, amorphous SiNx layers are inclined to form the crystalline structure and grow epitaxially with CrAlN layers, which improves the crystallization integrity of nanomultilayers and therefore increases the intensities of diffraction peaks. When tSiNx exceeds 0.6 nm, however, the improvement effect on crystallization integrity gradually disappears and the intensities of diffraction peaks accordingly decrease. It can also be seen from Fig. 2 that most nanomultilayer samples show a shift towards high angle in (111) and (200) diffraction peaks compared with the monolithic CrAlN film. The shift amount gradually increases with the initial increase of tSiNx and reaches the maximum value when tSiNx is 0.6 nm. With further increase of tSiNx, however, the shift amount gradually decreases. The shifts correspond to the change of the lattice parameter in CrAlN and SiNx layers. Under the epitaxial growth structure, the originally larger lattice parameter of CrAlN layer is inclined to decrease, leading to generation of interfacial compressive stress in substrate parallel direction, while the lattice parameter of SiNx layer is forced to increase, resulting in formation of interfacial tensile stress in substrate parallel direction. Therefore, the alternating compressive and tensile stress fields create along the growth direction of nanomultilayer, which can hinder the dislocation motion and strengthen the nanomultilayer [34]. Fig. 3 shows the cross-sectional HRTEM images of CrAlN/SiNx nanomultilayers with different tSiNx. It can be clearly seen from Fig. 3(a) that, when tSiNx is 0.6 nm, the lattice fringes continuously go through several modulation layers and interfaces, indicating that SiNx layers have been fully crystallized and grown epitaxially with CrAlN layers. Under the epitaxial growth structure, the originally larger lattice parameter of CrAlN layer is inclined to decrease, leading to generation of interfacial compressive stress, while the lattice parameter of SiNx layer is forced to increase, resulting in formation of interfacial tensile stress. When tSiNx increases to 0.8 nm, the whole epitaxial growth cannot be maintained between the CrAlN layers, while SiNx layer can still be present in a crystallized state, as shown in Fig. 3(b). The different CrAlN layers denoted by A, B and C, exhibit to grow along different orientations. The two different growth orientations are even observed in a CrAlN layer marked C, suggesting that different crystal domains have created in a single CrAlN layer and the epitaxial growth structure between CrAlN layers has been broken. When tSiNx reaches 1.2 nm, as shown in Fig. 3(c), the SiNx layers cannot keep their crystallization and change back to amorphous state. The interfacial morphology between the CrAlN layer and SiNx layer becomes undistinguishable, and more crystal domains with different misorientations are formed in a single CrAlN layer, suggesting that the epitaxial growth structure has been totally broken.
Fig. 3. Cross-sectional HRTEM images of CrAlN/SiNx nanomultilayers with SiNx layer thickness of (a) 0.6 nm, (b) 0.8 nm and (c) 1.2 nm.
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Table 1 The average crystallite size (D), r.m.s. microstrain and average dislocation density (ρ) of CrAlN/SiNx nanomultilayers with different SiNx layer thicknesses, measured by XRD line profile analysis. SiNx layer thickness (nm)
(hkl)
D (nm)
bε2 >
0.4 0.6 0.8 1.0 1.2
200 200 200 200 200
17.2 17.5 13.6 9.4 6.5
6.8 4.7 8.4 9.9 14.4
1/2
× 10−3
ρ (m-2) 5.7 2.7 8.6 12.1 26.6
× × × × ×
1015 1015 1015 1015 1015
4.2. Microstructural characterization of CrAlN/SiNx nanomultilayers by XRD line profile analysis By using our modified Warren–Averbach method, the average crystallite size (D), r.m.s. microstrain and average dislocation density (ρ) of CrAlN/SiNx nanomultilayers with different tSiNx are determined, as shown in Table 1. It can be seen that the average crystallite sizes (D) of various CrAlN/SiNx nanomultilayers are larger than modulation periods (Λ), and with the increase of tSiNx, the average crystallite size gradually decrease to the value of modulation period. The average dislocation density ( ρ) firstly decreases with the initial increase of tSiNx and reaches the minimum value when tSiNx is 0.6 nm, then increases with further increase of tSiNx. The phenomenon for Λ> D can be explained by distinguishing mode of XRD technique. In the measurements of crystallite size, XRD technique distinguishes crystallite with a certain misorientation [35]. In this investigation, however, due to the epitaxial growth between CrAlN and SiNx layer, the misorientation between different modulation periods is so small that it cannot be differentiated by XRD, resulting that the average crystallite size exceeds the modulation period. The initial decrease of the average dislocation density can be attributed to the fact that SiNx layers not only can be fully crystallized under the heteroepitaxy of CrAlN layers, but also the crystallized SiNx layers can improve the crystallization integrity of CrAlN layers [36]. This mutually promoting growth effect is verified by the XRD patterns in Fig. 2, in which the CrAlN/SiNx nanomultilayer with tSiNx of 0.6 nm exhibits a higher peak intensity than that with tSiNx of 0.4 nm. When tSiNx further increases, the epitaxial growth cannot be maintained between the CrAlN layers, and the misorientations between different CrAlN layers become large, resulting in the decrease of the average crystallite size. However, since SiNx layer still keeps the crystallization state, the CrAlN layer can retain the epitaxial growth with neighboring SiNx layer, even with the adjacent CrAlN layer in some regions, as shown in Fig. 3(b), causing that the average crystallite size from XRD still exceeds the modulation period. Due to the destruction of the epitaxial growth between the CrAlN layers and formation of different crystal domains within a CrAlN layer, the
crystallization integrity of CrAlN/SiNx nanomultilayer decreases, leading to the increase of the average dislocation density. When tSiNx reaches 1.2 nm, SiNx layers change back to amorphous state and the epitaxial growth is totally broken. More crystal domains with different misorientations are created within a single CrAlN layer, causing the average crystallite size to further reduce to 6.5 nm, close to the value of modulation period. The interfaces between the crystalline CrAlN layer and amorphous SiNx layer can generate more dislocations, and the misorientations between different crystal domains are also needed to adjust through dislocations [37]. Therefore, the average dislocation density further increases to the high value of 26.6 × 10 15 m −2. 4.3. Energy balance model of microstructure evolution for CrAlN/SiNx nanomultilayers The microstructural evolution of CrAlN/SiNx nanomultilayers with increase of tSiNx can be illustrated by Fig. 4. As tSiNx is less than 0.6 nm, SiNx layer are fully crystallized and grown epitaxially with CrAlN layers, as shown in Fig. 4(a). When tSiNx increases to 0.8 nm, although SiNx layers still present in a crystallized state, the epitaxial growth cannot be maintained between CrAlN layers and crystal domains may create within a CrAlN layer, as illustrated in Fig. 4(b). When tSiNx reaches 1.2 nm, SiNx layers transform into amorphous state, and the epitaxial growth structure is totally broken, and more crystal domains with different misorientations are created within a single CrAlN layer, as presented in Fig. 4(c). Such a microstructural evolution of CrAlN/SiNx nanomultilayers can be explained by an energy balance model including strain-free bulk energy (EB), coherency strain energy (ES) and interfacial energy (Ei), in which the total energy of CrAlN/SiNx nanomultilayer per unit (ET), can be written as: SiN SiN CrAlN CrAlN þ Ei þ ES ET ¼ EB x þ ES x t SiNx þ EB SiN
SiN
ð4Þ
where EB x and ES x are respectively the strain-free bulk energy and coherency strain energy of SiNx layers; EBCrAlN and ESCrAlN are respectively the strain-free bulk energy and coherency strain energy of CrAlN layers, in which ESCrAlN increases with tSiNx; Ei is interfacial energy between CrAlN layer and SiNx layer. During the initial growth of SiNx layer, tSiNx approaches zero, ET is therefore composed of EBCrAlN, ESCrAlN and Ei. Since EBCrAlN is invariable and ESCrAlN is comparatively small when tSiNx approaches zero, ET is dominated by Ei. Formation of coherent interface between CrAlN and SiNx layer can lower Ei. Therefore, SiNx layers are crystallized and grow epitaxially on top of CrAlN layers. When tSiNx increases to 0.8 nm, the total energy of SiNx layer increase. However, the total
Fig. 4. Schematic illustration of microstructural evolution of CrAlN/SiNx nanomultilayers with SiNx layer thickness of (a) less than 0.6 nm, (b) 0.8 nm and (c) 1.2 nm.
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elastic modulus decrease, the average crystallite size gradually decreases according to Table 1. Therefore, the strengthening of CrAlN/ SiNx nanomultilayers is not connected to the grain refinement. 5. Conclusions
Fig. 5. Variation of hardness and elastic modulus of CrAlN/SiNx nanomultilayers with change of SiNx layer thickness.
energy of SiNx layer is not large enough to exceed Ei. Therefore, the SiNx layers still maintain the crystallization state and keep the local epitaxial growth with the adjacent CrAlN layer. However, with increase of tSiNx, the coherency strain energy of the whole CrAlN layers, ESCrAlN, grows. Therefore, the whole epitaxial growth between CrAlN layers cannot be maintained, and even coherent growth cannot be kept within a CrAlN layer, leading to the formation of crystal domains in a CrAlN layer. As tSiNx reaches 1.2 nm, the total energy of SiNx layer further increases, which have more dominant effect on ET than Ei. x x Since ESiN and ESiN of crystalline SiNx are higher than those of amorB S phous SiNx, crystalline SiNx structure has to transform into amorphous state. In addition, ESCrAlN further increases with the increase of tSiNx, resulting in total destruction of epitaxial growth structure and formation of more crystal domains within a CrAlN layer. 4.4. Mechanical properties of CrAlN/SiNx nanomultilayers Fig. 5 shows the dependence of hardness and elastic modulus of CrAlN/SiNx nanomultilayers on SiNx layer thickness. The hardness and elastic modulus of CrAlN monolithic film deposited under the same conditions were measured as 28.7 GPa and 382 GPa, while those of SiNx monolithic film were 18 GPa and 225 GPa. It can be seen that, after SiNx layers are initially inserted, the hardness and elastic modulus of the nanomultilayers increase. When tSiNx is 0.6 nm, the maximum values of hardness and elastic modulus reach 37.6 GPa and 437 GPa, respectively. As tSiNx further increases, the hardness and elastic modulus values gradually decrease, and are below those of monolithic CrAlN film when tSiNx is 1.2 nm. The change of hardness and elastic modulus is closely related to the microstructure of nanomultilayers. When tSiNx is below 0.6 nm, SiNx layers are forced to adopt a the crystalline structure under the template effect of CrAlN layers and grow epitaxially with CrAlN layers. The enhancement in hardness and elastic modulus can be explained by Koehler's modulus-difference strengthening theory [1] and the alternating-stress strengthening theory [38], in which the dislocation motions are hindered at interfaces in nanomultilayers respectively by forces generated from two layers with different shear moduli and the alternating stress fields. When tSiNx exceeds 0.6 nm, however, the hardness and elastic modulus values gradually decrease. Based on the above analysis, the epitaxial growth structure between CrAlN layers is firstly damaged while SiNx layers still maintain the crystallization state, suggesting that the epitaxial growth structure is the key point to achieve the superhardness effect, while the crystallization of SiNx layers is not a requirement. As tSiNx further increases to 1.2 nm, the SiNx layers change to amorphous state, and the epitaxial growth is totally broken, leading to the further decrease of hardness and elastic modulus. It is worth noting that, as the hardness and
In summary, the microstructure and mechanical properties of CrAlN/SiNx nanomultilayers were investigated by XRD, HRTEM, XRD line profile analysis method and nano-indentation techniques. When tSiNx was below 0.6 nm, SiNx layers are forced to crystallize and grow epitaxially on CrAlN layers, resulting in the enhancement in hardness and elastic modulus, which respectively reach maximum values of 37.6 GPa and 437 GPa when tSiNx is 0.6 nm. When tSiNx increases to 0.8 nm, the epitaxial growth structure between CrAlN layers is damaged, while SiNx layers still maintain the crystallized state, resulting in the increase of average dislocation density and decrease of hardness and elastic modulus. As tSiNx reaches 1.2 nm, the SiNx layers change to amorphous state, and the epitaxial growth is totally broken, resulting in further increase of average dislocation density and decrease of mechanical properties. Based on the established energy balance expression, with the joint effects of interfacial energy, strain-free bulk energy and coherency strain energy of CrAlN and SiNx layers, the damage of the epitaxial growth precedes the transformation of the crystallized SiNx layers to amorphous state during the disappearance of superhardness effect. Acknowledgment The present work was financially supported by the National Natural Science Foundation of China under grant no. 51101101, “Shanghai Municipal Natural Science Foundation” under grant no. 11ZR1424600, “Innovation Program of Shanghai Municipal Education Commission” under grant no. 12YZ104, “Shanghai Leading Academic Discipline Project” under grant no. J50503 sponsored by Shanghai municipal education commission in China. We thank Dr. Weizong Xu for calculational assistance in XRD line profile analysis. References [1] S.A. Barnett, M. Shinn, Annu. Rev. Mater. Sci. 24 (1994) 481. [2] G. Abadias, A. Michel, C. Tromas, C. Jaouen, S.N. Dub, Surf. Coat. Technol. 202 (2007) 844. [3] J.S. Koehler, Phys. Rev. B 2 (1970) 547. [4] W.M. Yang, T. Tsakalakos, J.E. Hilliard, J. Appl. Phys. 48 (1977) 876. [5] U. Helmersson, S. Todorova, S.A. Barnett, J. Appl. Phys. 62 (1987) 481. [6] A. Madan, I.W. Kim, S.C. Cheng, P. Yashar, V.P. Dravid, S.A. Barnett, Phys. Rev. Lett. 78 (1997) 1743. [7] C. Ziebert, S. Ulrich, J. Vac. Sci. Technol. A 24 (2006) 554. [8] M. Stueber, H. Holleck, H. Leiste, K. Seemann, S. Ulrich, C. Ziebert, J. Alloy. Compd. 483 (2009) 321. [9] H.C. Barshilia, K.S. Rajam, A. Jain, K. Gopinadhan, S. Chaudhary, Thin Solid Films 503 (2006) 158. [10] J. Musil, Surf. Coat. Technol. 125 (2000) 322. [11] L. Wei, F.H. Mei, N. Shao, M. Kong, G.Y. Li, Appl. Phys. Lett. 86 (2005) 021919. [12] R.A. Andrievski, Surf. Coat. Technol. 201 (2007) 6112. [13] J.J. Lao, N. Shao, F.H. Mei, G.Y. Li, M.Y. Gu, Appl. Phys. Lett. 86 (2005) 011902. [14] W. Li, P. Liu, J.T. Wang, F.C. Ma, X.K. Liu, X.H. Chen, Mater. Lett. 65 (2011) 636. [15] J.H. Xu, K. Hattori, Y. Seino, I. Kojima, Thin Solid Films 414 (2002) 239. [16] L. Li, T. Ungar, Y.D. Wang, J.R. Morris, G. Tichy, J. Lendvai, Y.L. Yang, Y. Ren, H. Choo, P.K. Liaw, Acta Mater. 57 (2009) 4988. [17] J. Gubicza, N.Q. Chinh, T. Csanadi, T.G. Langdon, T. Ungar, Mater. Sci. Eng. A 462 (2007) 86. [18] T. Ungar, O. Castelnau, G. Ribarik, M. Drakopoulos, J.L. Bechade, T. Chauveau, A. Snigirev, I. Snigireva, C. Schroer, B. Bacroix, Acta Mater. 55 (2007) 1117. [19] J. Gubicza, L. Balogh, R.J. Hellmig, Y. Estrin, T. Ungar, Mater. Sci. Eng. A 400/401 (2005) 334. [20] S.H. Tsai, J.G. Duh, Thin Solid Films 518 (2009) 1480. [21] A.R. Stokes, Proc. Phys. Soc. 61 (1948) 382. [22] W. Li, W.Z. Xu, X.D. Wang, Y.H. Rong, J. Alloy. Compd. 474 (2009) 546. [23] W. Li, P. Liu, F.C. Ma, Y.H. Rong, J. Mater. Sci. 44 (2009) 2925. [24] W.C. Oliver, G.M. Pharr, J. Mater. Res. 7 (1992) 1564. [25] B.E. Warren, B.L. Averbach, J. Appl. Phys. 21 (1950) 595. [26] M. Wilkens, Phys. Status Solidi A 2 (1970) 359. [27] J.I. Langford, J. Appl. Crystallogr. 11 (1978) 10.
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