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Effects of periods on the evolution of microstructure and mechanical properties of multilayered Cu-W films during thermal annealing
Surface & Coatings Technology 381 (2020) 125179
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Effects of periods on the evolution of microstructure and mechanical properties of multilayered Cu-W films during thermal annealing
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Jiawei Xuea, Yanhuai Lia, Leiwen Gaoa, Dan Qiana, Zhongxiao Songa, Xiaohua Wangb, ⁎ ⁎⁎ Xiaodong Zhua, , Jian Chenc, a
State Key Laboratory for Mechanical Behavior of Materials, Xi'an Jiaotong University, Xi'an, Shannxi 710049, China State Key Laboratory of Electrical Insulation and Power Equipment, Xi'an Jiaotong University, 710049, China c School of Materials Science and Engineering, Southeast University, Nanjing, Jiangsu 211189, China b
A R T I C LE I N FO
A B S T R A C T
Keywords: Cu-W films Multilayered structure Thermal annealing Modulation period Thermal stability Mechanical properties
Benefiting from the unique structure, the copper‑tungsten (Cu-W) multilayered film is attractive due to its high strength and fracture toughness, diffusion barrier effects and radiation tolerance. However, the influence of period on the interfacial stability and strengthening mechanism during annealing has seldom been touched. Hence, the Cu-W multilayers with the periods from 1.6 to 27 nm were synthesized by altering substrate rotation speed using co-sputtering. In contrast to the findings in literature by sequential sputtering, kinetically constrained deposited atoms contribute to the more amorphous phase near interface in the films with smaller periods and result in the reduced hardness. During annealing, Cu and W atoms are thermal-activated and up-hill diffuse towards Cu- and W-rich layers, respectively, which can not only eliminate the amorphous phases but also enlarge the concentration ratio and result in the spinodal decomposition. In turn, the films with larger periods present higher thermal resistibility against phase separation due to the stronger diffusion barrier effect of thicker W-rich layers. The mechanical properties of the annealed films generally increase with the temperature up to 400 °C due to the reduced amorphous phases and the enhanced interfacial strengthening. At 500 °C, the separation of soft Cu phase and the loss of coherency mainly lead to the reduced mechanical properties. It is striking that the Cu-W film with the moderate thickness exhibits the highest hardness at higher temperature of 500 °C, indicating the significant importance of period to control their high temperature mechanical properties.
1. Introduction The nano-multilayered film has received wide interests as their unique structure leads to the advantages of not only high mechanical properties but also many other attributes [1]. For example, copper‑tungsten (Cu-W) multilayered films are considered as potential candidates for microelectronic [2,3], plasma and nuclear fields [4–6] due to high strength and fracture toughness, diffusion barrier effects and radiation tolerance. In the past years, it has been gained consensus that the modulation period (λ), modulus mismatch and interfacial structure are of great importance on the performance of the multilayered films [7–10]. Co-sputtering is a widely used method to deposit the multilayered films. By altering the substrate rotation speed and/or deposition rate, the bilayer periodicity, individual layer thickness and interfacial structure can be regulated. Müller et al [11] synthesized the Cu-56 at.%
⁎
Ta multilayered films with different periods and found that the period equals to the multiple of the rotation speed and deposition rate. Dong et al [6] reported that the period strongly influenced the irradiation tolerance and mechanical properties of Cu-W multilayers, and the films with smaller periods possessed better morphological stability and resistivity against He+ ions irradiation due to the higher content of interfaces acting as sinks to annihilate the radiation-induced defects. With the decrease of the periods from micrometers to nanometers, the flow strength of Cu-W multilayers with incoherent interfaces is reported to be increased by several studies [12,13]. This is explained by that the interfaces hamper the dislocation movement, leading to complicate deformation mechanisms including dislocation piles-up [14] and confined layer slip [15]. Notably, for specific systems, e.g. Cu-Ru multilayers, the strength may reach a maximum when the layer thickness is down to 1–2 nm due to the formation of coherently interfacial structure between layers [12].
Correspondence to: X.D. Zhu, State Key Laboratory for Mechanical Behavior of Materials, Xi'an Jiaotong University, Xi'an 710049, China. Correspondence to: J. Chen, School of Materials Science and Engineering, Southeast University, Nanjing 211189, China. E-mail addresses: [email protected] (X. Zhu), [email protected] (J. Chen).
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https://doi.org/10.1016/j.surfcoat.2019.125179 Received 23 May 2019; Received in revised form 5 October 2019; Accepted 19 November 2019 Available online 28 November 2019 0257-8972/ © 2019 Published by Elsevier B.V.
Surface & Coatings Technology 381 (2020) 125179
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furnace to room temperature.
The immiscible chemical combination of Cu and W is of great importance to the formation of the multilayers. However, it is usually found that the Cu-W films by co-deposition show smooth interfaces with sinusoidal content transition between Cu- and W-rich layers [11,16]. The existence of internal interfaces will contribute to higher Gibbs energy and hampering dislocation movement, which not only strengthen the films, but also make the multilayers thermodynamically unstable at high temperatures. Zhou et al [17] reported that the cosputtered Cu-62 at.% W multilayered film with λ of 20 nm can resist high temperature up to 600 °C, beyond which the separated Cu phase appeared. As for sequentially deposited Cu-W multilayers with λ of 5 nm, Moszner et al [18] found the degradation of multilayered structure happens at annealing temperature higher than 700 °C, and the degradation process is driven by minimization of the total Gibbs energy via thermal grooving and coarsening of layer fragments. Besides, thermal treatment of Cu-W multilayers above 200 °C results in a drastic reduction of hot hardness [19]. Although a number of literature deal with the formation mechanism and interfacial strengthening effects of the Cu-W multilayered films, limited studies focus on their high temperature stability and its resultant mechanical evolution. Furthermore, during co-sputtering process, the substrate rotation-induced periodicity and corresponding interfacial features dominate the thermal stability and mechanical properties, which has seldom been touched as so far. Therefore, the Cu-W films with different periods have been deposited by altering the substrate rotation speed. Their microstructural evolution by the thermal annealing from 300 to 500 °C has been systematically analyzed by TEM (Transmission Electron Microscope) to investigate the effects of the period on their thermal stability. The corresponding mechanical properties have been also studied and the hardening mechanisms are discussed.
2.2. Characterization The crystallographic information of the films were detected by a Panalytical X'-Pert PRO XRD (X-Ray Diffractometer) operating at 40 kV and 40 mA using Cu Kα irradiation (1.54 Å) and a nickel filter. Diffraction patterns were collected by stepwise scanning in the range of 20° and 80° with a scan rate of 4° min−1. For every sample, the average composition was measured several times at different regions under 200 magnification by Hitachi SU6600 FESEM (Field Emission Scanning Electron Microscope) with an EDS (Energy Dispersion Spectrometer). Detailed microstructures and crystallographic analysis were also carried out by JEM-2100F TEM at a 200 kV accelerating voltage. The crosssection specimens of films for TEM inspection were prepared by means of ion milling on a Leica EM RES102. 2.3. Mechanical testing Hysitron TI950 TriboIndenter with a calibrated Berkovich indenter was used for analyzing the hardness and elastic modulus by nanoindentation. The loading-control mode was applied at a constant rate of 600 μN s−1. The load was held at 3000 μN for 2 s before unloading for all the samples and the maximum depth of indentation was around 80 nm. Each test was repeated 10 times at different points. 3. Results 3.1. Phase structure Fig. 1a shows the XRD patterns of the as-deposited near-to-ratio CuW films (as measured by EDX) in Table 1 with the substrate rotation speed from 10 rpm (D1) to 0.625 rpm (D5) as well as pure Cu and W films sputtered at RT with thickness ~280 ± 5 nm for comparison. The dominated diffraction peaks lie within the range of bcc W110 (40.27°) and fcc Cu111 (43.30°), with much broader FWHM (Full Width at Half Maximum) compared with pure Cu and W film, suggesting the severe alloying effect and nanocrystals in the films [20]. It is found that D1 and D2 possess a main narrow diffraction peak and a cell left to the main peak, indicating the existence of amorphous [20]. While D3-D5 films possess much broader FWHM compared with D1 and D2, which is likely related to the fewer amorphous and different distribution of ordered interplanar spacing in bcc-W based lattice when the incursion thickness of Cu-rich layer increases [21]. The lower angle shift of main reflection for the films with larger period may be related to the increased W content from 45 to 50 at.%. In general, the concentration of Cu solute dissolved in the W phase, XCu, can be simply estimated from the measured lattice parameter using Vegard's law [22]:
2. Materials and methods 2.1. Film deposition and thermal treatment The base pressure of system chamber was evacuated to 3 × 10−4 Pa and working pressure was maintained at 3.0 × 10−1 Pa. Cu-W films were co-deposited using a direct current magnetron co-sputtering system (JPG-450A) with target diameter of 75 mm. The power densities of W (99.99%) and Cu target (99.99%) were 2.26 and 2.71 W·cm−2, respectively. A bias voltage of −50 V and deposition time of 20 min were adopted throughout the deposition. Commercial single-crystalline p-Si (100) substrate with 1 um SiO2 layer was used for as-deposited films and single-crystalline p-Si (100) with 100 nm SiNx diffusion barrier was used for annealing samples to prevent the formation of CuSix. All the substrates before deposition were ultrasonically cleaned by acetone, alcohol, and deionized water in sequence. To obtain the series of Cu-W films with varying periods, only substrate rotation speed was adjusted from 20 to 0.625 rpm with other parameters unchanged and these films are named as D1, D2, D3, D4, and D5 in Table 1, respectively. Besides, after deposition, isothermal annealing temperatures from 300 to 500 °C were applied for 1 h in vacuum environment (< 5 × 10−4 Pa) with heating rate about 25 °C/min, then cooling with
XCu =
αss − αW αCu − αW
(1)
where αss, αCu and αW are the lattice parameters of solid solution Cu-W alloyed films, pure Cu and W [22], respectively. Here we neglect the influence of other internal stresses on the peak position. Taken D1, D3
Table 1 Deposition parameters, thickness, chemical composition and period λ of the Cu-W multilayered films. Sample label
Rotation speed (rpm)
Thickness (nm)
Average composition (at.%)
Predicted λ by the Müller's model (nm)
Experimental λ by TEM (nm)
D1 D2 D3 D4 D5
10 5 2.5 1.25 0.625
260 280 240 260 255
CuW45 CuW45 CuW49 CuW50 CuW50
1.3 2.8 4.8 10.4 20.5
1.6 3.2 6.5 13.2 27.0
± ± ± ± ±
5 5 5 5 5
Other deposition parameters: base pressure 3 × 10−4 Pa, bias −50 V, deposition time and pressure: 20 min and 0.3 Pa. 2
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Fig. 1. XRD patterns of the (a) as-deposited Cu-W films under different substrate rotation speeds (D1, D2, D3, D4 and D5), and (b–d) D1, D3 and D5 films annealed at different temperatures, respectively. Both pure Cu and W films are deposited at RT and shown for comparison.
structure of the as-deposited Cu-W films with different substrate rotation speeds. The period counted statistically from EDS line scanning shows the clear increasing tendency with the decrease of rotation speed, from 1.6 nm (Fig. 2a) to 6.5 nm (Fig. 2b) and to 27 nm (Fig. 2c), which roughly accord with the predicted values in Table 2 calculated by the Müller model (λ = vt), which presents that the co-sputtered modulation period of immiscible binary alloy films (e.g. Cu-Ta or Cu-W) are mainly predictable by multiple of deposition rate v (nm/s) and time t (s) for each rotation based on the ignorance of surface diffusion during deposition [11]. The spot analysis results in Fig. 2b and c certify the W-rich (51–56 at.% W) and Cu-rich (62–64 at.% Cu) region. To be noted, these values only provide estimation of chemical composition in the individual layer due to the extremely small thickness. For the thinnest layer with 1.6 nm thickness in Fig. 2a, it is hard to differentiate the compositions. In addition, when the rotation speed is further increased to 20 rpm, the layer structure is barely seen in the BF and EDS line scanning (Fig. S1). Similar observation has been found in Cu-Ta system [11], indicating there is a threshold thickness, below which no layer forms. All the results indicate that the formation of modulation composition period is caused by that the alternatively arrived and condensed W-rich and Cu-rich flux [17]. It is worthy noted that the EDS measurement in this study was conducted at 200 kV with the probe size of 1 nm. So the electron beam broadening size was estimated as 1.5–3.9 nm (seen details in supplementary material) based on the electron diffraction theory [27,28] (Fig. S2). Thus, if the period is too small, e.g. D1 (1.6 nm) or D2 (3.2 nm), the accuracy of period value from EDS cannot be guaranteed. Fig. 3 shows the schema of deposition geometry in our experimental with two individual confocal targets: Cu and W. In region A, the glow discharge area of W and Cu are severely overlapped and the deposition field is nearly homogeneous, so multilayers cannot be obtained in this
and D5 as examples, the estimated Cu composition within the W phase versus global Cu composition in TEM is shown in Table S1. We can see that the Cu content within W-based phase is lower than its corresponding global content for different period films except for D1. In details, the Cu content within the W based phase of D3 and D5 film is 24.8 at.% and 17.7 at.%, respectively, which is lower than the global content of 51.2 at.% and 50.0 at.%. This phenomenon is also found by Zhao [23]. While, the estimated Cu content within W based phase in D1 film (64.1 at.%) is higher than the global content (55.1 at.%), which is not expected and its result is affected significantly by large amount of amorphous region. Besides, it is also found that the Cu content in W based phase is decreased with larger period, indicating that the wider Cu-rich layer may promote the Cu segregation. After the thermal annealing at different temperatures, it is found that Cu reflections occur at 400 °C for D1, 500 °C for D3 and D5 as indicated by the arrows (seen in Fig. 1b–d), demonstrating the precipitation takes place in the films, especially for those with fast rotation. This implies that there are strong period effects on the thermal stability of co-sputtered Cu-W multilayers. Besides, upon annealing at 300 °C for D1 and 400 °C for D3 and D5, the main diffraction peak shows slightly narrowed peak width compared with as-deposited state, which is likely related to the elimination of film defect [24] and grain growth effect [25]. Until the phase separation occurs, the main diffraction peak width increases obviously, which is likely related to the decreased grain size [18] and elastic strain between newly formed phases [26]. The narrower peak width in D5 at different annealing temperatures also demonstrate that the slower rotation will promote the crystallinity and phase stability of as-deposited multilayers. 3.2. Interfacial structure of the as-deposited Cu-W films Cross-sectional TEM results in Fig. 2 clearly show the multilayered 3
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Fig. 2. Cross-sectional TEM bright field (BF) images of (a) D1, (b) D3 and (c) D5 film. Corresponding EDS line scans (marked with yellow line) and spot analysis (marked as white plus) results are also listed in the figure. The average period value are estimated by EDS line scanning as 1.6 nm (D1), 6.5 nm (D3) and 27 nm (D5). (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
region. In region C, it is too close to W or Cu target during deposition, making the sputtering process more like alternatively deposition. In this experiment, the size of substrate disk includes region A and B and the samples (text wafer) are obtained in region B, where the sputtering distance difference between Cu or W target cannot be ignored and a heterogeneous deposition filed with W-rich and Cu-rich flux areas forms. Thus, when the substrate disk rotates at a certain speed, the text wafer goes through the W- or Cu-rich area alternatively and consequently W- or Cu-rich nanolayer constitute the film.
Table 2 The predicted and experimentally measured (by TEM) value of λ. Sample label
Rotation speed (rpm)
Predicted λ by the Müller's model (nm)
Experimental λ by TEM (nm)
D1 D2 D3 D4 D5
10 5 2.5 1.25 0.625
1.3 2.8 4.8 10.4 20.5
1.6 3.2 6.5 13.2 27.0
Fig. 3. The co-deposition schematic diagram of Cu-W multilayered film. 4
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Fig. 4. Representative cross-sectional TEM images of (a) as-deposited, (b) 400 °C annealed, and (c) 500 °C annealed D3 films. SAED patterns are inserted in BF images (a1, b1&c1). DF images (a2, b2&c2) of W110 are obtained according to diffraction spots (marked as red circles) in SAED patterns. Corresponding HRTEM images and their fast FFT images are shown in (a3, b3&c3). (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
Fig. 1c. The average columnar grain size is also increased with the temperature, from 20 to 40 nm at 400 °C to 30–50 nm at 500 °C in Fig. 4c2, but the period is almost unchanged. EDS spot analysis in Fig. 5a presents that the annealed D3 at 400 °C possess similar contents as the untreated one in Fig. 2b. However, the D3 annealed at 500 °C (Fig. 5b) shows higher Cu (~73 at.% Cu) and W (~54 at.% W) content in Cu- and W-rich layer respectively, meanwhile the increased average peak content from 67.5 at.% Cu and 54.1 at.% W to 70.3 at.% Cu and 56.4 at.% W is clearly observed in the line scan results. Such concentration evolution in individual layers can be attributed to the thermal-activated up-hill diffusion along the film growth direction during annealing [19] (also called spinodal decomposition). Besides, there are also some observable white contrast in between layers in Fig. 5a and b, which can be related to the grain boundary segregation and Cu diffusion during annealing [23]. The HRTEM results in Fig. 5c shows that the coherent interface is preserved after 500 °C annealing, and the interplanar spacing of W (110) is 0.221 nm, slightly less than pure W (110) of 0.224 nm due to the solid solution of Cu. Similarly, the W inclusion in Cu lattice can also
3.3. Effects of annealing temperature on the layered structure The annealing effects on the layered structure have been also investigated by TEM analysis. For the D3 film in Fig. 4a, the as-deposited interfacial structure shows that nanocrystals (Fig. 4a2) coexist with amorphous phase (marked as blue in Fig. 4a3). The formation of amorphous phase is caused by the limited kinetic energy of deposited atoms as reported in literature [21]. A comparison between D3 and D5 in Fig. S2 shows that the mixed amorphous and nanocrystalline phases are barely seen in the D5 film with the slower rotation speed, as the slower rotation speed can promote the nucleation and growth of crystals during deposition [29,30]. Once the D3 film was annealed at 400 °C, the thermal-activated rearrangement of atoms took place and promoted the columnar morphology in Fig. 4b2 with clearer interfaces between W- and Cu-rich regions in Fig. 4b3, and the disappearance of amorphous phase. At higher annealing temperature of 500 °C, the phase segregation is identified by the appearance of Cu reflection in SAED (Selected Area Electron Diffraction) (Fig. 4c1) in accordance with the XRD result in 5
Surface & Coatings Technology 381 (2020) 125179
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Fig. 5. BF images of D3 film annealed at (a) 400 °C and (b) 500 °C, respectively. Their corresponding EDS line scan results (as marked with yellow line), spot analysis (marked as white plus) and the average peak content are also listed in images (red for Cu and blue for W). HRTEM image of (c) 500 °C annealed D3 film are shown. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
as-deposited and annealed Cu-W films. It can be seen that the as-deposited films show increased hardness from 6.6 ± 0.4 to 8.9 ± 0.1 GPa when the period increases from 1.6 to 27 nm. Moreover, all of the H of Cu-W films with various periods increases with the rise of annealing temperature, and reaches the summit at 400 °C except for D3 (500 °C), which is generally accordant with the phase separation temperature as presented in XRD (Fig. 1). This finding means that annealing hardening dominates until the occurrence of separated Cu phase.
result in the continuous expansion of interplanar spacing of Cu (111) from 0.208 to 0.212 nm, which is observed in the enlarged image of local region (marked as yellow square) in Fig. 5c. The cross-sectional TEM and EDS analysis have been also carried out for the annealed D5 film with the largest period of 27 nm at 500 °C. As shown in Fig. S3a and b, similar Cu reflection can be seen in SAED image, and the DF image of Cu (200) presents the laminated Cu phase due to spinodal decomposition during annealing. Periodic composition oscillation is clearly observed in EDS line scanning result in Fig. S3c, and the peak concentrations reach ~91 at.% Cu and ~64 at.% W, which are much higher than ~70 at.% Cu and ~56 at.% W in 500 °C annealed D3 film (Fig. 5b), despite the relatively closer concentration in their asdeposited state (Fig. 2). This can be ascribed by the less interfaces in D5 which can benefit the inter-diffusion of metallic elements. Fig. 6a shows the dominant incoherent interfaces (between region I and II) and few residual coherent interfaces (region III) for annealed D5 film at 500 °C, illustrating the failure of coherency, much severer than that in the annealed D3 at 500 °C. The transformation from coherent to incoherent interface is ascribed to the unbearable elastic strain near interfaces between separated Cu- and W-rich domains during decomposition. The higher elastic strain in the annealed D5 may be related to its high variation of chemical composition as mentioned above. Besides, the fcc-Cu and bcc-W phase with interplanar spacing of 0.209 nm and 0.221 nm are also proved in FFT (Fast Fourier Transform) and inverse FFT images (Fig. 6b–e), showing the similar spinodal decomposition process with annealed D3 film (Fig. 4).
4. Discussion 4.1. Effects of periods on the thermal stability of Cu-W films The above microstructural and mechanical evolution of the Cu-W multilayer films with different periods confirm the significant effects of the periods. To clarify its effects, the phase stabilities of Cu-W films has been summarized in accordance with the annealing temperature and the periods in Fig. 8. It is clearly seen that the Cu-W films with larger periods (D3-D5) show the higher thermal resistibility against phase separation until 500 °C than the films (D1-D2) with smaller periods at 400 °C. It is known that the thermal stability of sequentially sputtered multilayers is generally related to the laminated architecture between layered phases and the degradation of incoherent interfaces plays a decisive role in the overall performance [31]. While for co-sputtered Cu-W multilayers, there is no sharp interfaces between W- and Cu-rich layers and only periodical composition fluctuations can be distinguished in W-based single phase. Thus the thermal stability is strongly controlled by the phase stability and initial laminated
3.4. Mechanical properties of multilayered Cu-W films Fig. 7 shows the measured hardness (H) and elastic modulus (E) of 6
Surface & Coatings Technology 381 (2020) 125179
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Fig. 6. Cross-sectional HRTEM image of D5 film annealed at 500 °C. (a) HRTEM image showing interfacial structure with upright inserted SAED image; (b–c) The fast FFT images of region I and II; (d–e) The inverse FFT image of region I and II.
architecture. Generally, phase separation in immiscible alloys can take place either by nucleation and growth or by spinodal decomposition mechanism [32]. In the case of nucleation and growth the formation of precipitates with fixed composition and sharp interface are expected. For spinodal decomposition, composition fluctuation that is amplifying over time by up-hill diffusion is expected and the interface between formed phases is not very sharp until the finish of phase separation. Then precipitates with fixed composition and sharp interface are formed. Therefore, the final state of nucleation and growth and of spinodal decomposition is indistinguishable and the phase separation mechanism should be identified by studying the early stage during decomposition [32]. In Cu-transition-metal systems, such as Cu-Nb, CuTa or Cu-W etc., some of the phase separations or transformations may be attributed to spinodal decomposition through uphill diffusion of Cu atoms [33]. In this paper, on one hand, we found that the composition ratio in Cu- or W-rich layer increases along with the annealing process (Figs. 4 and 5), e.g. the peak composition in Cu or W-rich layers for as-deposited D3 film are roughly ~62 at.% Cu and ~54 at.% W (Fig. 2b), and after 400 °C annealing (D3@400 °C) the compositions are 67.5 at.% Cu and 54.1 at.% W (Fig. 5a), then 70.3 at.% Cu and 56.4 at.% W (Fig. 5b) for 500 °C annealing (D3@500 °C). This increased composition ratio shows the key feature characteristic for spinodal decomposition. On the other, the sharpness of the interface, which is found visually strengthened by comparing D3 (Fig. 4a3), D3@400 °C (Fig. 4b) and D3@500 °C (Fig. 4c3) when phase separation neared its completion, is another point in favor of the spinodal decomposition hypothesis. Thus, we infer the mechanism of early stage of phase separation in this paper is vertically spinodal decomposition. Such behavior resembles the cases in TiAlN/CrN multilayers [34] and AlCrTiSiN high-entropy alloy nitride coating [35] upon annealing, that the elastic strain near interfaces is the driven force for vertically spinodal decomposition. Notably, the spinodal decomposition only happens at the early stage of phase separation, contributing to the amplifying fluctuation of Cu or W content in individual layers. When the Cu content in Cu-rich layer exceeds an threshold ~70 at.% Cu [36], bcc will transform into fcc structure. Considering the higher diffusivity of Cu than W [37], the separation process is mainly limited by the diffusion of Cu and thus the W-rich layers will act as barriers hampering the diffusion of Cu. Consequently, the thicker W-rich layer (or larger periods) presents better phase stability. Meanwhile, the pre-existing amorphous phases may transform to
Fig. 7. Indentation hardness and reduced modulus of the as-deposited and annealed Cu-W multilayered films under different substrate rotation speeds and annealing temperatures. The dotted square representing the as-deposited values.
Fig. 8. A schematic map showing the summary of the experimental results (shown as squares) on the annealing temperature- and period-dependent phase stabilities.
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bulk approximation [48], where the gas flux during deposition reaches the substrate and then transforms into surface atoms with residual kinetic energy. Before these atoms diffuse to thermodynamically favorable places, next layer of deposited atoms will arrive and frozen the existed atoms, i.e. adatom mobility is constrained. Thus, the faster rotation speed will aggregate the kinetic constrains of deposited atoms and smaller periods will introduce more local mixed (amorphous) regions (Fig. 3a) near the interfaces. In contrast, by slowing the rotation speed, the adatom mobility (especially Cu) is enhanced and benefits the nucleation and growth of crystals, leading to the fewer amorphous phases and better crystallinity. (b) Annealed Cu-W films During annealing, Cu and W atoms are thermal-activated and uphill diffuse towards Cu- and W-rich layers, respectively, enlarging the concentration ratio (Fig. 5), sharpening the interfaces (Fig. 3) and eliminating the amorphous phases (Figs. 3a and S2). At the same time, the columnar grain size slightly increases (Fig. 3). Therefore, the factors including grain size, composition, phase and interface need be considered for their variation of hardness. At stage I from room temperature to 400 °C (Fig. 7), the reduction of amorphous phase is related to the crystallization process, and the formed nanocrystals not only decrease the low-density amorphous phases but also strengthen the films by hampering the dislocation movement [38]. In addition, the interfacial strengthening also contributes to higher hardness by large concentration ratio caused by spinodal decomposition. However, when the temperature increases beyond 400 °C, the grain growth and reduced solid solubility in individual layer become severer, and the large elastic strain at interface can break the coherency and weaken the effects of coherent hardening [49]. The D3 film with moderate period shows the highest H at higher temperature of 500 °C. This can be attributed to the competition between the phase stability and the retained coherency. As mentioned above, the films with smaller periods (D1, D2) show poorer thermal stability due to the restricted barrier effects. Thus the formation of soft Cu phases deteriorates the mechanical properties of the films at 500 °C. On the other hand, although the D5 film with the largest period exhibits the highest thermal stability, the introduced lattice distortion by the highest concentration ratio through spinodal decomposition can break the coherency, leading to the drop of H. The above results indicated that the optimal period can significant improve the microstructural stability and hardness at high temperature. Actually, by rotating the substrate and annealing in this work, the microstructures of co-sputtered Cu-W multilayers varies dramatically and complicatedly, and this evolution may involve some different deformation mechanisms, further affecting the mechanical properties. Therefore, we consider it necessary to perform more in-depth mechanical and microstructural characterizations in the future to reveal its deformation mechanism.
nanocrystals and minimize the total free energy during annealing, the process of which is controlled by atomic energy and interfacial energy between new phase and amorphous matrix [38]. With the further increase of annealing temperature, the decomposition is promoted and the chemical composition gradient across the interface is enlarged, resulting in higher elastic strain due to the amplified lattice parameters of Cu- or W-rich layers, finally leads to the degradation of coherent interfaces. This can explain why the annealed D5 film at 500 °C possess more incoherent interfaces than the annealed D3 film at the same condition. Such trend is similar with the spinodal decomposition process in other systems [39,40]. However, once the annealing temperature rises further, the layered microstructure of Cu-W films can be completed destroyed as demonstrated by annealing D3 at 800 °C in Fig. S4 by severe coarsening effect, with a relatively smooth film surface due to the W-rich protective layer on the surface. In addition, during this non-equilibrium co-sputtering process, the residual internal stress in Cu-W multilayers should also be considered due to its effect on the thermal stability and ultimate performance [41]. The total residual stress in PVD films deposited at room temperature arises from a superposition of coherency and deposition (growth) stresses [42]. The former involves the lattice mismatch at substrate-film interface and the stress generated by the interface between adjacent layers, i.e. interface stress. The latter originates from film growth process (e.g. island nucleation, growth and coalescence) [43]. In this paper, considering the same substrate of p-Si (100) and the unchanged deposition parameters except the substrate rotation speed, the contribution of the substrate-film stress and deposition (growth) stress are thus ignored. Besides, for nanometer thick film, interfaces start to play a dominant role in the definition of the final stress state [44], so only interface stress is involved in the following discussion. Literature have found that the residual stress in Cu or W layer is usually different and the W layers always possess larger compressive stress due to its intrinsic stiffer nature, independent of thickness or period [41], while Cu layer presents different stress states from compressive to tensile, depending on the thickness and the deposition conditions [43]. This effect indicates the available modulation of internal stress in Cu-W multilayers. Druzhinin [44] found that as-deposited Cu/W multilayered films with smaller periods possess lager compressive stress no matter in Cu or W layers, indicating the increasing contribution of interface stress. Besides, during annealing treatment, the high internal stress can promote the outflow of Cu from inner to surface, leading to the relaxion of stress. Thus, for co-sputtered Cu-W multilayers with smaller periods (D1-D2) in this paper, the more contribution of interface stress may promote the up-hill diffusion of Cu atoms to decompose and finally form separated phases, showing the worse thermal stability against annealing. 4.2. Effects of periods on the hardness of Cu-W films (a) As-deposited Cu-W films For sequentially sputtered metallic multilayers, it is often reported that the strength increases with the decrease of period, due to the higher interface density and interfacial barrier for dislocation movement [12]. This seems contradictory to our finding as the Cu-W film with larger period possesses higher hardness. However, the interfacial structure in co-sputtered Cu-W multilayers is smooth and with mixed regions (amorphous) near interfaces, which is different from the sharper heterophase interfaces in sequentially sputtered films. As reported by Rizzo et al [45] and Zong et al [46], the H of Cu-W amorphous phase is lower than their solid solution counterparts, due to the low density of amorphous phases caused by a large amount of vacancylike defects or free volume and large static atomic displacement [47]. According to the TEM observations, higher interface density and more amorphous regions form in the films with smaller periods, leading to their lower mechanical properties. The formation of amorphous phases can be explained by the frozen
5. Conclusion Co-sputtered Cu-W multilayered films with various periods from 1.6 to 27 nm are prepared by altering the substrate rotation speed. The asdeposited films are all composed of single W-based solid solution with alternative Cu- and W-rich layers. With the decrease of the period, kinetically constrained deposited atoms contribute to the more amorphous phases near interface and result in the reduction of hardness. During annealing, Cu and W atoms are thermal-activated and uphill diffuse towards Cu- and W-rich layers, respectively, which can not only eliminate the amorphous phases but also enlarge the concentration ratio and result in the spinodal decomposition. In turn, the films with larger periods present higher thermal resistibility against phase separation due to the stronger diffusion barrier effect of thicker W-rich layers. The mechanical properties of the annealed films show an increasing trend (stage I) followed by a descend (stage II). The strengthening 8
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mechanisms in stage I involve the reduced amorphous phases and the enhanced interfacial strengthening. Afterwards, in stage II, the separation of soft Cu phase and the loss of coherency mainly lead to the reduced mechanical properties. It is striking that the Cu-W film with the moderate thickness exhibit the highest hardness at higher temperature, indicating the significant importance of period to control their high temperature mechanical properties.
[18]
[19]
[20]
Declaration of competing interest [21]
The authors declared that they have no conflicts of interest to this work. We declare that we do not have any commercial or associative interest that represents a conflict of interest in connection with the work submitted.
[22]
[23]
Acknowledgments
[24]
This work was supported by the National Natural Science Foundation of China [grant number 11472080].
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Appendix A. Supplementary data
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Supplementary data to this article can be found online at https:// doi.org/10.1016/j.surfcoat.2019.125179.
[28]
References
[29]
[1] V.V. Bukhanovsky, N.I. Grechanyuk, R.V. Minakova, I. Mamuzich, V.V. Kharchenko, N.P. Rudnitsky, Production technology, structure and properties of Cu–W layered composite condensed materials for electrical contacts, Int. J. Refract. Met. Hard Mater. 29 (2011) 573–581. [2] O. Auciello, S. Chevacharoenkul, M.S. Ameen, J. Duarte, Controlled ion beam sputter deposition of W/Cu/W layered films for microelectronic applications, J. Vac. Sci. Technol., A 9 (1991) 625–631. [3] M. Takeyama, A. Noya, T. Fukuda, Thermal stability of Cu/W/Si contact systems using layers of Cu(111) and W(110) preferred orientations, J. Vac. Sci. Technol., A 15 (1997) 415–420. [4] Y. Gao, T. Yang, J. Xue, S. Yan, S. Zhou, Y. Wang, D.T.K. Kwok, P.K. Chu, Y. Zhang, Radiation tolerance of Cu/W multilayered nanocomposites, J. Nucl. Mater. 413 (2011) 11–15. [5] E.G. Fu, N. Li, A. Misra, R.G. Hoagland, H. Wang, X. Zhang, Mechanical properties of sputtered Cu/V and Al/Nb multilayer films, Mater. Sci. Eng., A 493 (2008) 283–287. [6] L. Dong, H. Zhang, H. Amekura, F. Pan, A. Chettah, M. Hong, W. Qin, J. Tang, L. Hu, H. Wang, C. Jiang, Period-thickness dependent response of Cu/W multilayered nanofilms to ions irradiation under different ion energies, J. Nucl. Mater. 497 (2017) 117–127. [7] H. Zheng, L. Xiong, Q. Luo, S. Lu, Development of multilayer oxidation resistant coatings on Cr–50Nb alloy, Appl. Surf. Sci. 359 (2015) 515–520. [8] I.A. Kopylets, V.V. Kondratenko, E.N. Zubarev, D.L. Voronov, E.M. Gullikson, E.A. Vishnyakov, E.N. Ragozin, Fabrication and characterization of Sb/B4C multilayer mirrors for soft X-rays, Appl. Surf. Sci. 307 (2014) 360–364. [9] H. Khodja, M. Sall, M. Loyer-Prost, C. Cabet, A. Billard, I. Monnet, Radiation tolerance of multilayer coating revealed by IBA, Nucl. Instrum. Methods Phys. Res. 0168583X (2018). [10] L.R. Brandt, E. Salvati, C. Papadaki, H.J. Zhang, S.Q. Ying, E.L. Bourhis, L. Dolbnya, T. Sui, A.M. Korsunsky, Probing the deformation and fracture properties of Cu/W nano-multilayers by in situ SEM and synchrotron XRD strain microscopy, Surf. Coat. Technol. 320 (2017) 158–167. [11] C.M. Müller, A.S. Sologubenko, S.S.A. Gerstl, M.J. Süess, D. Courty, R. Spolenak, Nanoscale Cu/Ta multilayer deposition by co-sputtering on a rotating substrate. Empirical model and experiment, Surf. Coat. Technol. 302 (2016) 284–292. [12] Q. Zhou, Y. Li, F. Wang, P. Huang, T.J. Lu, K.W. Xu, Length-scale-dependent deformation mechanism of Cu/X (X=Ru, W) multilayer thin films, Mater. Sci. Eng., A 664 (2016) 206–214. [13] S. Wen, R. Zong, F. Zeng, Y. Gao, F. Pan, Evaluating modulus and hardness enhancement in evaporated Cu/W multilayers, Acta Mater. 55 (2007) 345–351. [14] Y. Lu, C. Wang, Y. Gao, R. Shi, X. Liu, Y. Wang, Microstructure map for self-organized phase separation during film deposition, Phys. Rev. Lett. 109 (2012) 086101. [15] M.A. Phillips, B.M. Clemens, W.D. Nix, A model for dislocation behavior during deformation of Al/Al3Sc (fcc/L12) metallic multilayers, Acta Mater. 51 (2003) 3157–3170. [16] J.W. Xue, Y.H. Li, L.C. Hao, L.W. Gao, D. Qian, Z.X. Song, J. Chen, Investigation on the interfacial stability of multilayered Cu–W films at elevated deposition temperatures during co-sputtering, Vacuum 166 (2019) 162–169. [17] T. Xie, L. Fu, W. Qin, J. Zhu, W. Yang, D. Li, L. Zhou, Self-assembled metal nano-
[30] [31]
[32]
[33]
[34]
[35]
[36]
[37]
[38]
[39]
[40] [41]
[42]
[43]
[44]
[45]
[46]
9
multilayered film prepared by co-sputtering method, Appl. Surf. Sci. 435 (2018) 16–22. F. Moszner, C. Cancellieri, M. Chiodi, S. Yoon, D. Ariosa, J. Janczak-Rusch, L.P.H. Jeurgens, Thermal stability of Cu/W nano-multilayers, Acta Mater. 107 (2016) 345–353. M.A. Monclús, M. Karlik, M. Callisti, E. Frutos, J. Llorca, T. Polcar, J.M. MolinaAldareguía, Microstructure and mechanical properties of physical vapor deposited Cu/W nanoscale multilayers: Influence of layer thickness and temperature, Thin Solid Films 571 (2014) 275–282. T. Xie, J. Zhu, L. Fu, R. Zhang, N. Li, M. Yang, J. Wang, W. Qin, W. Yang, D. Li, The evolution of hardness in Cu-W alloy thin films, Mater. Sci. Eng., A 729 (2018) 170–177. L.P. Zhou, M.P. Wang, K. Peng, J.J. Zhu, Z. Fu, Z. Li, Structure characteristic and its evolution of Cu-W films prepared by dual-target magnetron sputtering deposition, Trans. Nonferrous Met. Soc. China 22 (2012) 2700–2706. G. Csizar, S.J.B. Kurz, E.J. Mittemeijer, Stability of nanosized alloy thin films: faulting and phase separation in metastable Ni/Cu/Ag-W films, Acta Mater. 110 (2016) 324–340. J. Zhao, J. Zhang, Z. Hou, K. Wu, X. Feng, G. Liu, J. Sun, The W alloying effect on thermal stability and hardening of nanostructured Cu-W alloyed thin films, Nanotechnology 29 (2018) 195705. C.S. Fuh, P.T. Liu, W.H. Huang, S.M. Sze, Effect of annealing on defect elimination for high mobility amorphous indium-zinc-tin-oxide thin-film transistor, IEEE Electron Device Lett. 35 (11) (2014) 1103–1105. Y. Du, L. Li, J.M. Pureza, Y.W. Chung, K.G. Pradeep, S. Sen, J. Schneider, Thermal stability of nanocrystalline grains in Cu-W films, Surf. Coat. Technol. 357 (2019) 662–668. S.K. Sahay, Spinodal decomposition in Co-3 Ti (wt.%) alloy, J. Metall. Mater. Sci. 50 (1) (2008) 33–43. S.J.B. Reed, The single-scattering model and spatial resolution in X-ray analysis of thin foils, Ultramicroscopy 7 (4) (1982) 405–409. D.B. Williams, C.B. Carter, Thickness and bending effects, Transmission Electron Microscopy: A Text Book for Material Science, 2th edn., Springer, Amsterdam, 2004, pp. 369–371. C. Wang, P. Brault, C. Zaepffel, J. Thiault, A. Pineau, T. Sauvage, Deposition and structure of W–Cu multilayer coatings by magnetron sputtering, J. Phys. D 36 (21) (2003) 2709–2713. A. Misra, R.G. Hoagland, H. Kung, Thermal stability of self-supported nanolayered Cu/Nb films, Philos. Mag. 84 (2004) 1021–1028. A. Misra, R.G. Hoagland, Effects of elevated temperature annealing on the structure and hardness of copper/niobium nanolayered films, J. Mater. Res. 20 (2005) 2046–2054. M. Muller, A.S. Claudia, S.S.A. Sologubenko, R.S. Gerstl, On spinodal decomposition in Cu-34 at.% Ta thin films-an atom probe tomography and transmission electron microscopy study, Acta Mater. 89 (2015) 181–192. T. Xie, L. Fu, B. Gao, J. Zhu, W. Yang, D. Li, L. Zhou, The crystallization character of W-Cu thin films at the early stage of deposition, Thin solid films 690 (2019) 137555. I. Povstugar, P.P. Choi, D. Tytko, J.P. Ahn, D. Raabe, Interface-directed spinodal decomposition in TiAlN/CrN multilayer hard coatings studied by atom probe tomography, Acta Mater. 61 (2013) 7534–7542. W. Chen, A. Yan, X.N. Meng, D.Q. Wu, D.C. Yao, D.D. Zhang, Microstructural change and phase transformation in each individual layer of a nano-multilayered AlCrTiSiN high-entropy alloy nitride coating upon annealing, Appl. Surf. Sci. 462 (2018) 1017–1028. K. Chang, D. Music, M. Baben, D. Lange, H. Bolvardi, J.M. Schneider, Modeling of metastable phase formation diagrams for sputtered thin films, Sci. Technol. Adv. Mat. 17 (1) (2016) 210–219. L.T. Kong, X.Y. Li, W.S. Lai, J.B. Liu, B.X. Liu, Interfacial reaction of W/Cu examined by an n-body potential through molecular dynamics simulations, Jpn. J. Appl. Phys. 41 (2002) 4503–4508. G. Wang, B. Jiang, X. Zhang, B. Zhou, L. Meng, Crystallization behavior and mechanical properties of Cu-based bulk metallic glass composites, Mater. Res. 22 (2019) e20190078. B. Derby, Y. Cui, J. Baldwin, R. Arróyave, M.J. Demkowicz, A. Misra, Processing of novel pseudomorphic Cu–Mo hierarchies in thin films, Mater. Res. Lett. 7 (2018) 1–11. F. Findik, Improvements in spinodal alloys from past to present, Mater. Des. 42 (2012) 131–146. B. Girault, P. Villain, E.L. Bourhis, Ph. Goudeau, P.O. Renault, X-ray diffraction analysis of the structure and residual stress of W/Cu multilayers, Surf. Coat. Tech. 201 (2006) 4372–4376. Ph. Goudeau, K.F. Badawi, A. Naudon, G. Gladyszewski, Determination of the residual stress tensor in Cu/W multilayers by X-ray diffraction, Appl. Phys. Lett. 62 (3) (1993) 246–248. C. Cancellieri, F. Moszner, M. Chiodi, S. Yoon, J. Janczak-Rusch, L.P.H. Jeurgens, The effect of thermal treatment on the stress state and evolving microstructure of Cu/W nano-multilayers, J. Appl. Phys. 120 (2016) 195107. A.V. Druzhinin, D. Ariosa, S. Siol, N. Ott, B.B. Straumal, J. Janczak-Rusch, L.P.H. Jeurgens, C. Cancellieri, Effect of the individual layer thickness on the transformation of Cu/W nano-multilayers into nanocomposites, Materialia 7 (2019) 100400. H.F. Rizzo, T.B. Massalski, M. Nastasi, Metastable crystalline and amorphous structures formed in the Cu-W system by vapor deposition, Metall. Trans. A 24 (1993) 1027–1037. R.L. Zong, S.P. Wen, F. Zeng, Y. Gao, F. Pan, Nanoindentation studies of Cu–W alloy
Surface & Coatings Technology 381 (2020) 125179
J. Xue, et al.
[49] J. Yan, L. Ning, F. Xu, Z. Ying, The strengthening effect of spinodal decomposition and twinning structure in MnCu-based alloy, Mater. Sci. Eng., A 618 (2014) 205–209.
films prepared by magnetron sputtering, J. Alloys Compd. 464 (2008) 544–549. [47] R.W. Cahn, Aluminium-based glassy alloys, Nature 341 (1989) 183–184. [48] D. Tabor, The Hardness of Metals (Oxford), (1951).
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Effects of periods on the evolution of microstructure and mechanical properties of multilayered Cu-W films during thermal annealing
T
Jiawei Xuea, Yanhuai Lia, Leiwen Gaoa, Dan Qiana, Zhongxiao Songa, Xiaohua Wangb, ⁎ ⁎⁎ Xiaodong Zhua, , Jian Chenc, a
State Key Laboratory for Mechanical Behavior of Materials, Xi'an Jiaotong University, Xi'an, Shannxi 710049, China State Key Laboratory of Electrical Insulation and Power Equipment, Xi'an Jiaotong University, 710049, China c School of Materials Science and Engineering, Southeast University, Nanjing, Jiangsu 211189, China b
A R T I C LE I N FO
A B S T R A C T
Keywords: Cu-W films Multilayered structure Thermal annealing Modulation period Thermal stability Mechanical properties
Benefiting from the unique structure, the copper‑tungsten (Cu-W) multilayered film is attractive due to its high strength and fracture toughness, diffusion barrier effects and radiation tolerance. However, the influence of period on the interfacial stability and strengthening mechanism during annealing has seldom been touched. Hence, the Cu-W multilayers with the periods from 1.6 to 27 nm were synthesized by altering substrate rotation speed using co-sputtering. In contrast to the findings in literature by sequential sputtering, kinetically constrained deposited atoms contribute to the more amorphous phase near interface in the films with smaller periods and result in the reduced hardness. During annealing, Cu and W atoms are thermal-activated and up-hill diffuse towards Cu- and W-rich layers, respectively, which can not only eliminate the amorphous phases but also enlarge the concentration ratio and result in the spinodal decomposition. In turn, the films with larger periods present higher thermal resistibility against phase separation due to the stronger diffusion barrier effect of thicker W-rich layers. The mechanical properties of the annealed films generally increase with the temperature up to 400 °C due to the reduced amorphous phases and the enhanced interfacial strengthening. At 500 °C, the separation of soft Cu phase and the loss of coherency mainly lead to the reduced mechanical properties. It is striking that the Cu-W film with the moderate thickness exhibits the highest hardness at higher temperature of 500 °C, indicating the significant importance of period to control their high temperature mechanical properties.
1. Introduction The nano-multilayered film has received wide interests as their unique structure leads to the advantages of not only high mechanical properties but also many other attributes [1]. For example, copper‑tungsten (Cu-W) multilayered films are considered as potential candidates for microelectronic [2,3], plasma and nuclear fields [4–6] due to high strength and fracture toughness, diffusion barrier effects and radiation tolerance. In the past years, it has been gained consensus that the modulation period (λ), modulus mismatch and interfacial structure are of great importance on the performance of the multilayered films [7–10]. Co-sputtering is a widely used method to deposit the multilayered films. By altering the substrate rotation speed and/or deposition rate, the bilayer periodicity, individual layer thickness and interfacial structure can be regulated. Müller et al [11] synthesized the Cu-56 at.%
⁎
Ta multilayered films with different periods and found that the period equals to the multiple of the rotation speed and deposition rate. Dong et al [6] reported that the period strongly influenced the irradiation tolerance and mechanical properties of Cu-W multilayers, and the films with smaller periods possessed better morphological stability and resistivity against He+ ions irradiation due to the higher content of interfaces acting as sinks to annihilate the radiation-induced defects. With the decrease of the periods from micrometers to nanometers, the flow strength of Cu-W multilayers with incoherent interfaces is reported to be increased by several studies [12,13]. This is explained by that the interfaces hamper the dislocation movement, leading to complicate deformation mechanisms including dislocation piles-up [14] and confined layer slip [15]. Notably, for specific systems, e.g. Cu-Ru multilayers, the strength may reach a maximum when the layer thickness is down to 1–2 nm due to the formation of coherently interfacial structure between layers [12].
Correspondence to: X.D. Zhu, State Key Laboratory for Mechanical Behavior of Materials, Xi'an Jiaotong University, Xi'an 710049, China. Correspondence to: J. Chen, School of Materials Science and Engineering, Southeast University, Nanjing 211189, China. E-mail addresses: [email protected] (X. Zhu), [email protected] (J. Chen).
⁎⁎
https://doi.org/10.1016/j.surfcoat.2019.125179 Received 23 May 2019; Received in revised form 5 October 2019; Accepted 19 November 2019 Available online 28 November 2019 0257-8972/ © 2019 Published by Elsevier B.V.
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furnace to room temperature.
The immiscible chemical combination of Cu and W is of great importance to the formation of the multilayers. However, it is usually found that the Cu-W films by co-deposition show smooth interfaces with sinusoidal content transition between Cu- and W-rich layers [11,16]. The existence of internal interfaces will contribute to higher Gibbs energy and hampering dislocation movement, which not only strengthen the films, but also make the multilayers thermodynamically unstable at high temperatures. Zhou et al [17] reported that the cosputtered Cu-62 at.% W multilayered film with λ of 20 nm can resist high temperature up to 600 °C, beyond which the separated Cu phase appeared. As for sequentially deposited Cu-W multilayers with λ of 5 nm, Moszner et al [18] found the degradation of multilayered structure happens at annealing temperature higher than 700 °C, and the degradation process is driven by minimization of the total Gibbs energy via thermal grooving and coarsening of layer fragments. Besides, thermal treatment of Cu-W multilayers above 200 °C results in a drastic reduction of hot hardness [19]. Although a number of literature deal with the formation mechanism and interfacial strengthening effects of the Cu-W multilayered films, limited studies focus on their high temperature stability and its resultant mechanical evolution. Furthermore, during co-sputtering process, the substrate rotation-induced periodicity and corresponding interfacial features dominate the thermal stability and mechanical properties, which has seldom been touched as so far. Therefore, the Cu-W films with different periods have been deposited by altering the substrate rotation speed. Their microstructural evolution by the thermal annealing from 300 to 500 °C has been systematically analyzed by TEM (Transmission Electron Microscope) to investigate the effects of the period on their thermal stability. The corresponding mechanical properties have been also studied and the hardening mechanisms are discussed.
2.2. Characterization The crystallographic information of the films were detected by a Panalytical X'-Pert PRO XRD (X-Ray Diffractometer) operating at 40 kV and 40 mA using Cu Kα irradiation (1.54 Å) and a nickel filter. Diffraction patterns were collected by stepwise scanning in the range of 20° and 80° with a scan rate of 4° min−1. For every sample, the average composition was measured several times at different regions under 200 magnification by Hitachi SU6600 FESEM (Field Emission Scanning Electron Microscope) with an EDS (Energy Dispersion Spectrometer). Detailed microstructures and crystallographic analysis were also carried out by JEM-2100F TEM at a 200 kV accelerating voltage. The crosssection specimens of films for TEM inspection were prepared by means of ion milling on a Leica EM RES102. 2.3. Mechanical testing Hysitron TI950 TriboIndenter with a calibrated Berkovich indenter was used for analyzing the hardness and elastic modulus by nanoindentation. The loading-control mode was applied at a constant rate of 600 μN s−1. The load was held at 3000 μN for 2 s before unloading for all the samples and the maximum depth of indentation was around 80 nm. Each test was repeated 10 times at different points. 3. Results 3.1. Phase structure Fig. 1a shows the XRD patterns of the as-deposited near-to-ratio CuW films (as measured by EDX) in Table 1 with the substrate rotation speed from 10 rpm (D1) to 0.625 rpm (D5) as well as pure Cu and W films sputtered at RT with thickness ~280 ± 5 nm for comparison. The dominated diffraction peaks lie within the range of bcc W110 (40.27°) and fcc Cu111 (43.30°), with much broader FWHM (Full Width at Half Maximum) compared with pure Cu and W film, suggesting the severe alloying effect and nanocrystals in the films [20]. It is found that D1 and D2 possess a main narrow diffraction peak and a cell left to the main peak, indicating the existence of amorphous [20]. While D3-D5 films possess much broader FWHM compared with D1 and D2, which is likely related to the fewer amorphous and different distribution of ordered interplanar spacing in bcc-W based lattice when the incursion thickness of Cu-rich layer increases [21]. The lower angle shift of main reflection for the films with larger period may be related to the increased W content from 45 to 50 at.%. In general, the concentration of Cu solute dissolved in the W phase, XCu, can be simply estimated from the measured lattice parameter using Vegard's law [22]:
2. Materials and methods 2.1. Film deposition and thermal treatment The base pressure of system chamber was evacuated to 3 × 10−4 Pa and working pressure was maintained at 3.0 × 10−1 Pa. Cu-W films were co-deposited using a direct current magnetron co-sputtering system (JPG-450A) with target diameter of 75 mm. The power densities of W (99.99%) and Cu target (99.99%) were 2.26 and 2.71 W·cm−2, respectively. A bias voltage of −50 V and deposition time of 20 min were adopted throughout the deposition. Commercial single-crystalline p-Si (100) substrate with 1 um SiO2 layer was used for as-deposited films and single-crystalline p-Si (100) with 100 nm SiNx diffusion barrier was used for annealing samples to prevent the formation of CuSix. All the substrates before deposition were ultrasonically cleaned by acetone, alcohol, and deionized water in sequence. To obtain the series of Cu-W films with varying periods, only substrate rotation speed was adjusted from 20 to 0.625 rpm with other parameters unchanged and these films are named as D1, D2, D3, D4, and D5 in Table 1, respectively. Besides, after deposition, isothermal annealing temperatures from 300 to 500 °C were applied for 1 h in vacuum environment (< 5 × 10−4 Pa) with heating rate about 25 °C/min, then cooling with
XCu =
αss − αW αCu − αW
(1)
where αss, αCu and αW are the lattice parameters of solid solution Cu-W alloyed films, pure Cu and W [22], respectively. Here we neglect the influence of other internal stresses on the peak position. Taken D1, D3
Table 1 Deposition parameters, thickness, chemical composition and period λ of the Cu-W multilayered films. Sample label
Rotation speed (rpm)
Thickness (nm)
Average composition (at.%)
Predicted λ by the Müller's model (nm)
Experimental λ by TEM (nm)
D1 D2 D3 D4 D5
10 5 2.5 1.25 0.625
260 280 240 260 255
CuW45 CuW45 CuW49 CuW50 CuW50
1.3 2.8 4.8 10.4 20.5
1.6 3.2 6.5 13.2 27.0
± ± ± ± ±
5 5 5 5 5
Other deposition parameters: base pressure 3 × 10−4 Pa, bias −50 V, deposition time and pressure: 20 min and 0.3 Pa. 2
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Fig. 1. XRD patterns of the (a) as-deposited Cu-W films under different substrate rotation speeds (D1, D2, D3, D4 and D5), and (b–d) D1, D3 and D5 films annealed at different temperatures, respectively. Both pure Cu and W films are deposited at RT and shown for comparison.
structure of the as-deposited Cu-W films with different substrate rotation speeds. The period counted statistically from EDS line scanning shows the clear increasing tendency with the decrease of rotation speed, from 1.6 nm (Fig. 2a) to 6.5 nm (Fig. 2b) and to 27 nm (Fig. 2c), which roughly accord with the predicted values in Table 2 calculated by the Müller model (λ = vt), which presents that the co-sputtered modulation period of immiscible binary alloy films (e.g. Cu-Ta or Cu-W) are mainly predictable by multiple of deposition rate v (nm/s) and time t (s) for each rotation based on the ignorance of surface diffusion during deposition [11]. The spot analysis results in Fig. 2b and c certify the W-rich (51–56 at.% W) and Cu-rich (62–64 at.% Cu) region. To be noted, these values only provide estimation of chemical composition in the individual layer due to the extremely small thickness. For the thinnest layer with 1.6 nm thickness in Fig. 2a, it is hard to differentiate the compositions. In addition, when the rotation speed is further increased to 20 rpm, the layer structure is barely seen in the BF and EDS line scanning (Fig. S1). Similar observation has been found in Cu-Ta system [11], indicating there is a threshold thickness, below which no layer forms. All the results indicate that the formation of modulation composition period is caused by that the alternatively arrived and condensed W-rich and Cu-rich flux [17]. It is worthy noted that the EDS measurement in this study was conducted at 200 kV with the probe size of 1 nm. So the electron beam broadening size was estimated as 1.5–3.9 nm (seen details in supplementary material) based on the electron diffraction theory [27,28] (Fig. S2). Thus, if the period is too small, e.g. D1 (1.6 nm) or D2 (3.2 nm), the accuracy of period value from EDS cannot be guaranteed. Fig. 3 shows the schema of deposition geometry in our experimental with two individual confocal targets: Cu and W. In region A, the glow discharge area of W and Cu are severely overlapped and the deposition field is nearly homogeneous, so multilayers cannot be obtained in this
and D5 as examples, the estimated Cu composition within the W phase versus global Cu composition in TEM is shown in Table S1. We can see that the Cu content within W-based phase is lower than its corresponding global content for different period films except for D1. In details, the Cu content within the W based phase of D3 and D5 film is 24.8 at.% and 17.7 at.%, respectively, which is lower than the global content of 51.2 at.% and 50.0 at.%. This phenomenon is also found by Zhao [23]. While, the estimated Cu content within W based phase in D1 film (64.1 at.%) is higher than the global content (55.1 at.%), which is not expected and its result is affected significantly by large amount of amorphous region. Besides, it is also found that the Cu content in W based phase is decreased with larger period, indicating that the wider Cu-rich layer may promote the Cu segregation. After the thermal annealing at different temperatures, it is found that Cu reflections occur at 400 °C for D1, 500 °C for D3 and D5 as indicated by the arrows (seen in Fig. 1b–d), demonstrating the precipitation takes place in the films, especially for those with fast rotation. This implies that there are strong period effects on the thermal stability of co-sputtered Cu-W multilayers. Besides, upon annealing at 300 °C for D1 and 400 °C for D3 and D5, the main diffraction peak shows slightly narrowed peak width compared with as-deposited state, which is likely related to the elimination of film defect [24] and grain growth effect [25]. Until the phase separation occurs, the main diffraction peak width increases obviously, which is likely related to the decreased grain size [18] and elastic strain between newly formed phases [26]. The narrower peak width in D5 at different annealing temperatures also demonstrate that the slower rotation will promote the crystallinity and phase stability of as-deposited multilayers. 3.2. Interfacial structure of the as-deposited Cu-W films Cross-sectional TEM results in Fig. 2 clearly show the multilayered 3
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Fig. 2. Cross-sectional TEM bright field (BF) images of (a) D1, (b) D3 and (c) D5 film. Corresponding EDS line scans (marked with yellow line) and spot analysis (marked as white plus) results are also listed in the figure. The average period value are estimated by EDS line scanning as 1.6 nm (D1), 6.5 nm (D3) and 27 nm (D5). (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
region. In region C, it is too close to W or Cu target during deposition, making the sputtering process more like alternatively deposition. In this experiment, the size of substrate disk includes region A and B and the samples (text wafer) are obtained in region B, where the sputtering distance difference between Cu or W target cannot be ignored and a heterogeneous deposition filed with W-rich and Cu-rich flux areas forms. Thus, when the substrate disk rotates at a certain speed, the text wafer goes through the W- or Cu-rich area alternatively and consequently W- or Cu-rich nanolayer constitute the film.
Table 2 The predicted and experimentally measured (by TEM) value of λ. Sample label
Rotation speed (rpm)
Predicted λ by the Müller's model (nm)
Experimental λ by TEM (nm)
D1 D2 D3 D4 D5
10 5 2.5 1.25 0.625
1.3 2.8 4.8 10.4 20.5
1.6 3.2 6.5 13.2 27.0
Fig. 3. The co-deposition schematic diagram of Cu-W multilayered film. 4
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Fig. 4. Representative cross-sectional TEM images of (a) as-deposited, (b) 400 °C annealed, and (c) 500 °C annealed D3 films. SAED patterns are inserted in BF images (a1, b1&c1). DF images (a2, b2&c2) of W110 are obtained according to diffraction spots (marked as red circles) in SAED patterns. Corresponding HRTEM images and their fast FFT images are shown in (a3, b3&c3). (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
Fig. 1c. The average columnar grain size is also increased with the temperature, from 20 to 40 nm at 400 °C to 30–50 nm at 500 °C in Fig. 4c2, but the period is almost unchanged. EDS spot analysis in Fig. 5a presents that the annealed D3 at 400 °C possess similar contents as the untreated one in Fig. 2b. However, the D3 annealed at 500 °C (Fig. 5b) shows higher Cu (~73 at.% Cu) and W (~54 at.% W) content in Cu- and W-rich layer respectively, meanwhile the increased average peak content from 67.5 at.% Cu and 54.1 at.% W to 70.3 at.% Cu and 56.4 at.% W is clearly observed in the line scan results. Such concentration evolution in individual layers can be attributed to the thermal-activated up-hill diffusion along the film growth direction during annealing [19] (also called spinodal decomposition). Besides, there are also some observable white contrast in between layers in Fig. 5a and b, which can be related to the grain boundary segregation and Cu diffusion during annealing [23]. The HRTEM results in Fig. 5c shows that the coherent interface is preserved after 500 °C annealing, and the interplanar spacing of W (110) is 0.221 nm, slightly less than pure W (110) of 0.224 nm due to the solid solution of Cu. Similarly, the W inclusion in Cu lattice can also
3.3. Effects of annealing temperature on the layered structure The annealing effects on the layered structure have been also investigated by TEM analysis. For the D3 film in Fig. 4a, the as-deposited interfacial structure shows that nanocrystals (Fig. 4a2) coexist with amorphous phase (marked as blue in Fig. 4a3). The formation of amorphous phase is caused by the limited kinetic energy of deposited atoms as reported in literature [21]. A comparison between D3 and D5 in Fig. S2 shows that the mixed amorphous and nanocrystalline phases are barely seen in the D5 film with the slower rotation speed, as the slower rotation speed can promote the nucleation and growth of crystals during deposition [29,30]. Once the D3 film was annealed at 400 °C, the thermal-activated rearrangement of atoms took place and promoted the columnar morphology in Fig. 4b2 with clearer interfaces between W- and Cu-rich regions in Fig. 4b3, and the disappearance of amorphous phase. At higher annealing temperature of 500 °C, the phase segregation is identified by the appearance of Cu reflection in SAED (Selected Area Electron Diffraction) (Fig. 4c1) in accordance with the XRD result in 5
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Fig. 5. BF images of D3 film annealed at (a) 400 °C and (b) 500 °C, respectively. Their corresponding EDS line scan results (as marked with yellow line), spot analysis (marked as white plus) and the average peak content are also listed in images (red for Cu and blue for W). HRTEM image of (c) 500 °C annealed D3 film are shown. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
as-deposited and annealed Cu-W films. It can be seen that the as-deposited films show increased hardness from 6.6 ± 0.4 to 8.9 ± 0.1 GPa when the period increases from 1.6 to 27 nm. Moreover, all of the H of Cu-W films with various periods increases with the rise of annealing temperature, and reaches the summit at 400 °C except for D3 (500 °C), which is generally accordant with the phase separation temperature as presented in XRD (Fig. 1). This finding means that annealing hardening dominates until the occurrence of separated Cu phase.
result in the continuous expansion of interplanar spacing of Cu (111) from 0.208 to 0.212 nm, which is observed in the enlarged image of local region (marked as yellow square) in Fig. 5c. The cross-sectional TEM and EDS analysis have been also carried out for the annealed D5 film with the largest period of 27 nm at 500 °C. As shown in Fig. S3a and b, similar Cu reflection can be seen in SAED image, and the DF image of Cu (200) presents the laminated Cu phase due to spinodal decomposition during annealing. Periodic composition oscillation is clearly observed in EDS line scanning result in Fig. S3c, and the peak concentrations reach ~91 at.% Cu and ~64 at.% W, which are much higher than ~70 at.% Cu and ~56 at.% W in 500 °C annealed D3 film (Fig. 5b), despite the relatively closer concentration in their asdeposited state (Fig. 2). This can be ascribed by the less interfaces in D5 which can benefit the inter-diffusion of metallic elements. Fig. 6a shows the dominant incoherent interfaces (between region I and II) and few residual coherent interfaces (region III) for annealed D5 film at 500 °C, illustrating the failure of coherency, much severer than that in the annealed D3 at 500 °C. The transformation from coherent to incoherent interface is ascribed to the unbearable elastic strain near interfaces between separated Cu- and W-rich domains during decomposition. The higher elastic strain in the annealed D5 may be related to its high variation of chemical composition as mentioned above. Besides, the fcc-Cu and bcc-W phase with interplanar spacing of 0.209 nm and 0.221 nm are also proved in FFT (Fast Fourier Transform) and inverse FFT images (Fig. 6b–e), showing the similar spinodal decomposition process with annealed D3 film (Fig. 4).
4. Discussion 4.1. Effects of periods on the thermal stability of Cu-W films The above microstructural and mechanical evolution of the Cu-W multilayer films with different periods confirm the significant effects of the periods. To clarify its effects, the phase stabilities of Cu-W films has been summarized in accordance with the annealing temperature and the periods in Fig. 8. It is clearly seen that the Cu-W films with larger periods (D3-D5) show the higher thermal resistibility against phase separation until 500 °C than the films (D1-D2) with smaller periods at 400 °C. It is known that the thermal stability of sequentially sputtered multilayers is generally related to the laminated architecture between layered phases and the degradation of incoherent interfaces plays a decisive role in the overall performance [31]. While for co-sputtered Cu-W multilayers, there is no sharp interfaces between W- and Cu-rich layers and only periodical composition fluctuations can be distinguished in W-based single phase. Thus the thermal stability is strongly controlled by the phase stability and initial laminated
3.4. Mechanical properties of multilayered Cu-W films Fig. 7 shows the measured hardness (H) and elastic modulus (E) of 6
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Fig. 6. Cross-sectional HRTEM image of D5 film annealed at 500 °C. (a) HRTEM image showing interfacial structure with upright inserted SAED image; (b–c) The fast FFT images of region I and II; (d–e) The inverse FFT image of region I and II.
architecture. Generally, phase separation in immiscible alloys can take place either by nucleation and growth or by spinodal decomposition mechanism [32]. In the case of nucleation and growth the formation of precipitates with fixed composition and sharp interface are expected. For spinodal decomposition, composition fluctuation that is amplifying over time by up-hill diffusion is expected and the interface between formed phases is not very sharp until the finish of phase separation. Then precipitates with fixed composition and sharp interface are formed. Therefore, the final state of nucleation and growth and of spinodal decomposition is indistinguishable and the phase separation mechanism should be identified by studying the early stage during decomposition [32]. In Cu-transition-metal systems, such as Cu-Nb, CuTa or Cu-W etc., some of the phase separations or transformations may be attributed to spinodal decomposition through uphill diffusion of Cu atoms [33]. In this paper, on one hand, we found that the composition ratio in Cu- or W-rich layer increases along with the annealing process (Figs. 4 and 5), e.g. the peak composition in Cu or W-rich layers for as-deposited D3 film are roughly ~62 at.% Cu and ~54 at.% W (Fig. 2b), and after 400 °C annealing (D3@400 °C) the compositions are 67.5 at.% Cu and 54.1 at.% W (Fig. 5a), then 70.3 at.% Cu and 56.4 at.% W (Fig. 5b) for 500 °C annealing (D3@500 °C). This increased composition ratio shows the key feature characteristic for spinodal decomposition. On the other, the sharpness of the interface, which is found visually strengthened by comparing D3 (Fig. 4a3), D3@400 °C (Fig. 4b) and D3@500 °C (Fig. 4c3) when phase separation neared its completion, is another point in favor of the spinodal decomposition hypothesis. Thus, we infer the mechanism of early stage of phase separation in this paper is vertically spinodal decomposition. Such behavior resembles the cases in TiAlN/CrN multilayers [34] and AlCrTiSiN high-entropy alloy nitride coating [35] upon annealing, that the elastic strain near interfaces is the driven force for vertically spinodal decomposition. Notably, the spinodal decomposition only happens at the early stage of phase separation, contributing to the amplifying fluctuation of Cu or W content in individual layers. When the Cu content in Cu-rich layer exceeds an threshold ~70 at.% Cu [36], bcc will transform into fcc structure. Considering the higher diffusivity of Cu than W [37], the separation process is mainly limited by the diffusion of Cu and thus the W-rich layers will act as barriers hampering the diffusion of Cu. Consequently, the thicker W-rich layer (or larger periods) presents better phase stability. Meanwhile, the pre-existing amorphous phases may transform to
Fig. 7. Indentation hardness and reduced modulus of the as-deposited and annealed Cu-W multilayered films under different substrate rotation speeds and annealing temperatures. The dotted square representing the as-deposited values.
Fig. 8. A schematic map showing the summary of the experimental results (shown as squares) on the annealing temperature- and period-dependent phase stabilities.
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bulk approximation [48], where the gas flux during deposition reaches the substrate and then transforms into surface atoms with residual kinetic energy. Before these atoms diffuse to thermodynamically favorable places, next layer of deposited atoms will arrive and frozen the existed atoms, i.e. adatom mobility is constrained. Thus, the faster rotation speed will aggregate the kinetic constrains of deposited atoms and smaller periods will introduce more local mixed (amorphous) regions (Fig. 3a) near the interfaces. In contrast, by slowing the rotation speed, the adatom mobility (especially Cu) is enhanced and benefits the nucleation and growth of crystals, leading to the fewer amorphous phases and better crystallinity. (b) Annealed Cu-W films During annealing, Cu and W atoms are thermal-activated and uphill diffuse towards Cu- and W-rich layers, respectively, enlarging the concentration ratio (Fig. 5), sharpening the interfaces (Fig. 3) and eliminating the amorphous phases (Figs. 3a and S2). At the same time, the columnar grain size slightly increases (Fig. 3). Therefore, the factors including grain size, composition, phase and interface need be considered for their variation of hardness. At stage I from room temperature to 400 °C (Fig. 7), the reduction of amorphous phase is related to the crystallization process, and the formed nanocrystals not only decrease the low-density amorphous phases but also strengthen the films by hampering the dislocation movement [38]. In addition, the interfacial strengthening also contributes to higher hardness by large concentration ratio caused by spinodal decomposition. However, when the temperature increases beyond 400 °C, the grain growth and reduced solid solubility in individual layer become severer, and the large elastic strain at interface can break the coherency and weaken the effects of coherent hardening [49]. The D3 film with moderate period shows the highest H at higher temperature of 500 °C. This can be attributed to the competition between the phase stability and the retained coherency. As mentioned above, the films with smaller periods (D1, D2) show poorer thermal stability due to the restricted barrier effects. Thus the formation of soft Cu phases deteriorates the mechanical properties of the films at 500 °C. On the other hand, although the D5 film with the largest period exhibits the highest thermal stability, the introduced lattice distortion by the highest concentration ratio through spinodal decomposition can break the coherency, leading to the drop of H. The above results indicated that the optimal period can significant improve the microstructural stability and hardness at high temperature. Actually, by rotating the substrate and annealing in this work, the microstructures of co-sputtered Cu-W multilayers varies dramatically and complicatedly, and this evolution may involve some different deformation mechanisms, further affecting the mechanical properties. Therefore, we consider it necessary to perform more in-depth mechanical and microstructural characterizations in the future to reveal its deformation mechanism.
nanocrystals and minimize the total free energy during annealing, the process of which is controlled by atomic energy and interfacial energy between new phase and amorphous matrix [38]. With the further increase of annealing temperature, the decomposition is promoted and the chemical composition gradient across the interface is enlarged, resulting in higher elastic strain due to the amplified lattice parameters of Cu- or W-rich layers, finally leads to the degradation of coherent interfaces. This can explain why the annealed D5 film at 500 °C possess more incoherent interfaces than the annealed D3 film at the same condition. Such trend is similar with the spinodal decomposition process in other systems [39,40]. However, once the annealing temperature rises further, the layered microstructure of Cu-W films can be completed destroyed as demonstrated by annealing D3 at 800 °C in Fig. S4 by severe coarsening effect, with a relatively smooth film surface due to the W-rich protective layer on the surface. In addition, during this non-equilibrium co-sputtering process, the residual internal stress in Cu-W multilayers should also be considered due to its effect on the thermal stability and ultimate performance [41]. The total residual stress in PVD films deposited at room temperature arises from a superposition of coherency and deposition (growth) stresses [42]. The former involves the lattice mismatch at substrate-film interface and the stress generated by the interface between adjacent layers, i.e. interface stress. The latter originates from film growth process (e.g. island nucleation, growth and coalescence) [43]. In this paper, considering the same substrate of p-Si (100) and the unchanged deposition parameters except the substrate rotation speed, the contribution of the substrate-film stress and deposition (growth) stress are thus ignored. Besides, for nanometer thick film, interfaces start to play a dominant role in the definition of the final stress state [44], so only interface stress is involved in the following discussion. Literature have found that the residual stress in Cu or W layer is usually different and the W layers always possess larger compressive stress due to its intrinsic stiffer nature, independent of thickness or period [41], while Cu layer presents different stress states from compressive to tensile, depending on the thickness and the deposition conditions [43]. This effect indicates the available modulation of internal stress in Cu-W multilayers. Druzhinin [44] found that as-deposited Cu/W multilayered films with smaller periods possess lager compressive stress no matter in Cu or W layers, indicating the increasing contribution of interface stress. Besides, during annealing treatment, the high internal stress can promote the outflow of Cu from inner to surface, leading to the relaxion of stress. Thus, for co-sputtered Cu-W multilayers with smaller periods (D1-D2) in this paper, the more contribution of interface stress may promote the up-hill diffusion of Cu atoms to decompose and finally form separated phases, showing the worse thermal stability against annealing. 4.2. Effects of periods on the hardness of Cu-W films (a) As-deposited Cu-W films For sequentially sputtered metallic multilayers, it is often reported that the strength increases with the decrease of period, due to the higher interface density and interfacial barrier for dislocation movement [12]. This seems contradictory to our finding as the Cu-W film with larger period possesses higher hardness. However, the interfacial structure in co-sputtered Cu-W multilayers is smooth and with mixed regions (amorphous) near interfaces, which is different from the sharper heterophase interfaces in sequentially sputtered films. As reported by Rizzo et al [45] and Zong et al [46], the H of Cu-W amorphous phase is lower than their solid solution counterparts, due to the low density of amorphous phases caused by a large amount of vacancylike defects or free volume and large static atomic displacement [47]. According to the TEM observations, higher interface density and more amorphous regions form in the films with smaller periods, leading to their lower mechanical properties. The formation of amorphous phases can be explained by the frozen
5. Conclusion Co-sputtered Cu-W multilayered films with various periods from 1.6 to 27 nm are prepared by altering the substrate rotation speed. The asdeposited films are all composed of single W-based solid solution with alternative Cu- and W-rich layers. With the decrease of the period, kinetically constrained deposited atoms contribute to the more amorphous phases near interface and result in the reduction of hardness. During annealing, Cu and W atoms are thermal-activated and uphill diffuse towards Cu- and W-rich layers, respectively, which can not only eliminate the amorphous phases but also enlarge the concentration ratio and result in the spinodal decomposition. In turn, the films with larger periods present higher thermal resistibility against phase separation due to the stronger diffusion barrier effect of thicker W-rich layers. The mechanical properties of the annealed films show an increasing trend (stage I) followed by a descend (stage II). The strengthening 8
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mechanisms in stage I involve the reduced amorphous phases and the enhanced interfacial strengthening. Afterwards, in stage II, the separation of soft Cu phase and the loss of coherency mainly lead to the reduced mechanical properties. It is striking that the Cu-W film with the moderate thickness exhibit the highest hardness at higher temperature, indicating the significant importance of period to control their high temperature mechanical properties.
[18]
[19]
[20]
Declaration of competing interest [21]
The authors declared that they have no conflicts of interest to this work. We declare that we do not have any commercial or associative interest that represents a conflict of interest in connection with the work submitted.
[22]
[23]
Acknowledgments
[24]
This work was supported by the National Natural Science Foundation of China [grant number 11472080].
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Appendix A. Supplementary data
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Supplementary data to this article can be found online at https:// doi.org/10.1016/j.surfcoat.2019.125179.
[28]
References
[29]
[1] V.V. Bukhanovsky, N.I. Grechanyuk, R.V. Minakova, I. Mamuzich, V.V. Kharchenko, N.P. Rudnitsky, Production technology, structure and properties of Cu–W layered composite condensed materials for electrical contacts, Int. J. Refract. Met. Hard Mater. 29 (2011) 573–581. [2] O. Auciello, S. Chevacharoenkul, M.S. Ameen, J. Duarte, Controlled ion beam sputter deposition of W/Cu/W layered films for microelectronic applications, J. Vac. Sci. Technol., A 9 (1991) 625–631. [3] M. Takeyama, A. Noya, T. Fukuda, Thermal stability of Cu/W/Si contact systems using layers of Cu(111) and W(110) preferred orientations, J. Vac. Sci. Technol., A 15 (1997) 415–420. [4] Y. Gao, T. Yang, J. Xue, S. Yan, S. Zhou, Y. Wang, D.T.K. Kwok, P.K. Chu, Y. Zhang, Radiation tolerance of Cu/W multilayered nanocomposites, J. Nucl. Mater. 413 (2011) 11–15. [5] E.G. Fu, N. Li, A. Misra, R.G. Hoagland, H. Wang, X. Zhang, Mechanical properties of sputtered Cu/V and Al/Nb multilayer films, Mater. Sci. Eng., A 493 (2008) 283–287. [6] L. Dong, H. Zhang, H. Amekura, F. Pan, A. Chettah, M. Hong, W. Qin, J. Tang, L. Hu, H. Wang, C. Jiang, Period-thickness dependent response of Cu/W multilayered nanofilms to ions irradiation under different ion energies, J. Nucl. Mater. 497 (2017) 117–127. [7] H. Zheng, L. Xiong, Q. Luo, S. Lu, Development of multilayer oxidation resistant coatings on Cr–50Nb alloy, Appl. Surf. Sci. 359 (2015) 515–520. [8] I.A. Kopylets, V.V. Kondratenko, E.N. Zubarev, D.L. Voronov, E.M. Gullikson, E.A. Vishnyakov, E.N. Ragozin, Fabrication and characterization of Sb/B4C multilayer mirrors for soft X-rays, Appl. Surf. Sci. 307 (2014) 360–364. [9] H. Khodja, M. Sall, M. Loyer-Prost, C. Cabet, A. Billard, I. Monnet, Radiation tolerance of multilayer coating revealed by IBA, Nucl. Instrum. Methods Phys. Res. 0168583X (2018). [10] L.R. Brandt, E. Salvati, C. Papadaki, H.J. Zhang, S.Q. Ying, E.L. Bourhis, L. Dolbnya, T. Sui, A.M. Korsunsky, Probing the deformation and fracture properties of Cu/W nano-multilayers by in situ SEM and synchrotron XRD strain microscopy, Surf. Coat. Technol. 320 (2017) 158–167. [11] C.M. Müller, A.S. Sologubenko, S.S.A. Gerstl, M.J. Süess, D. Courty, R. Spolenak, Nanoscale Cu/Ta multilayer deposition by co-sputtering on a rotating substrate. Empirical model and experiment, Surf. Coat. Technol. 302 (2016) 284–292. [12] Q. Zhou, Y. Li, F. Wang, P. Huang, T.J. Lu, K.W. Xu, Length-scale-dependent deformation mechanism of Cu/X (X=Ru, W) multilayer thin films, Mater. Sci. Eng., A 664 (2016) 206–214. [13] S. Wen, R. Zong, F. Zeng, Y. Gao, F. Pan, Evaluating modulus and hardness enhancement in evaporated Cu/W multilayers, Acta Mater. 55 (2007) 345–351. [14] Y. Lu, C. Wang, Y. Gao, R. Shi, X. Liu, Y. Wang, Microstructure map for self-organized phase separation during film deposition, Phys. Rev. Lett. 109 (2012) 086101. [15] M.A. Phillips, B.M. Clemens, W.D. Nix, A model for dislocation behavior during deformation of Al/Al3Sc (fcc/L12) metallic multilayers, Acta Mater. 51 (2003) 3157–3170. [16] J.W. Xue, Y.H. Li, L.C. Hao, L.W. Gao, D. Qian, Z.X. Song, J. Chen, Investigation on the interfacial stability of multilayered Cu–W films at elevated deposition temperatures during co-sputtering, Vacuum 166 (2019) 162–169. [17] T. Xie, L. Fu, W. Qin, J. Zhu, W. Yang, D. Li, L. Zhou, Self-assembled metal nano-
[30] [31]
[32]
[33]
[34]
[35]
[36]
[37]
[38]
[39]
[40] [41]
[42]
[43]
[44]
[45]
[46]
9
multilayered film prepared by co-sputtering method, Appl. Surf. Sci. 435 (2018) 16–22. F. Moszner, C. Cancellieri, M. Chiodi, S. Yoon, D. Ariosa, J. Janczak-Rusch, L.P.H. Jeurgens, Thermal stability of Cu/W nano-multilayers, Acta Mater. 107 (2016) 345–353. M.A. Monclús, M. Karlik, M. Callisti, E. Frutos, J. Llorca, T. Polcar, J.M. MolinaAldareguía, Microstructure and mechanical properties of physical vapor deposited Cu/W nanoscale multilayers: Influence of layer thickness and temperature, Thin Solid Films 571 (2014) 275–282. T. Xie, J. Zhu, L. Fu, R. Zhang, N. Li, M. Yang, J. Wang, W. Qin, W. Yang, D. Li, The evolution of hardness in Cu-W alloy thin films, Mater. Sci. Eng., A 729 (2018) 170–177. L.P. Zhou, M.P. Wang, K. Peng, J.J. Zhu, Z. Fu, Z. Li, Structure characteristic and its evolution of Cu-W films prepared by dual-target magnetron sputtering deposition, Trans. Nonferrous Met. Soc. China 22 (2012) 2700–2706. G. Csizar, S.J.B. Kurz, E.J. Mittemeijer, Stability of nanosized alloy thin films: faulting and phase separation in metastable Ni/Cu/Ag-W films, Acta Mater. 110 (2016) 324–340. J. Zhao, J. Zhang, Z. Hou, K. Wu, X. Feng, G. Liu, J. Sun, The W alloying effect on thermal stability and hardening of nanostructured Cu-W alloyed thin films, Nanotechnology 29 (2018) 195705. C.S. Fuh, P.T. Liu, W.H. Huang, S.M. Sze, Effect of annealing on defect elimination for high mobility amorphous indium-zinc-tin-oxide thin-film transistor, IEEE Electron Device Lett. 35 (11) (2014) 1103–1105. Y. Du, L. Li, J.M. Pureza, Y.W. Chung, K.G. Pradeep, S. Sen, J. Schneider, Thermal stability of nanocrystalline grains in Cu-W films, Surf. Coat. Technol. 357 (2019) 662–668. S.K. Sahay, Spinodal decomposition in Co-3 Ti (wt.%) alloy, J. Metall. Mater. Sci. 50 (1) (2008) 33–43. S.J.B. Reed, The single-scattering model and spatial resolution in X-ray analysis of thin foils, Ultramicroscopy 7 (4) (1982) 405–409. D.B. Williams, C.B. Carter, Thickness and bending effects, Transmission Electron Microscopy: A Text Book for Material Science, 2th edn., Springer, Amsterdam, 2004, pp. 369–371. C. Wang, P. Brault, C. Zaepffel, J. Thiault, A. Pineau, T. Sauvage, Deposition and structure of W–Cu multilayer coatings by magnetron sputtering, J. Phys. D 36 (21) (2003) 2709–2713. A. Misra, R.G. Hoagland, H. Kung, Thermal stability of self-supported nanolayered Cu/Nb films, Philos. Mag. 84 (2004) 1021–1028. A. Misra, R.G. Hoagland, Effects of elevated temperature annealing on the structure and hardness of copper/niobium nanolayered films, J. Mater. Res. 20 (2005) 2046–2054. M. Muller, A.S. Claudia, S.S.A. Sologubenko, R.S. Gerstl, On spinodal decomposition in Cu-34 at.% Ta thin films-an atom probe tomography and transmission electron microscopy study, Acta Mater. 89 (2015) 181–192. T. Xie, L. Fu, B. Gao, J. Zhu, W. Yang, D. Li, L. Zhou, The crystallization character of W-Cu thin films at the early stage of deposition, Thin solid films 690 (2019) 137555. I. Povstugar, P.P. Choi, D. Tytko, J.P. Ahn, D. Raabe, Interface-directed spinodal decomposition in TiAlN/CrN multilayer hard coatings studied by atom probe tomography, Acta Mater. 61 (2013) 7534–7542. W. Chen, A. Yan, X.N. Meng, D.Q. Wu, D.C. Yao, D.D. Zhang, Microstructural change and phase transformation in each individual layer of a nano-multilayered AlCrTiSiN high-entropy alloy nitride coating upon annealing, Appl. Surf. Sci. 462 (2018) 1017–1028. K. Chang, D. Music, M. Baben, D. Lange, H. Bolvardi, J.M. Schneider, Modeling of metastable phase formation diagrams for sputtered thin films, Sci. Technol. Adv. Mat. 17 (1) (2016) 210–219. L.T. Kong, X.Y. Li, W.S. Lai, J.B. Liu, B.X. Liu, Interfacial reaction of W/Cu examined by an n-body potential through molecular dynamics simulations, Jpn. J. Appl. Phys. 41 (2002) 4503–4508. G. Wang, B. Jiang, X. Zhang, B. Zhou, L. Meng, Crystallization behavior and mechanical properties of Cu-based bulk metallic glass composites, Mater. Res. 22 (2019) e20190078. B. Derby, Y. Cui, J. Baldwin, R. Arróyave, M.J. Demkowicz, A. Misra, Processing of novel pseudomorphic Cu–Mo hierarchies in thin films, Mater. Res. Lett. 7 (2018) 1–11. F. Findik, Improvements in spinodal alloys from past to present, Mater. Des. 42 (2012) 131–146. B. Girault, P. Villain, E.L. Bourhis, Ph. Goudeau, P.O. Renault, X-ray diffraction analysis of the structure and residual stress of W/Cu multilayers, Surf. Coat. Tech. 201 (2006) 4372–4376. Ph. Goudeau, K.F. Badawi, A. Naudon, G. Gladyszewski, Determination of the residual stress tensor in Cu/W multilayers by X-ray diffraction, Appl. Phys. Lett. 62 (3) (1993) 246–248. C. Cancellieri, F. Moszner, M. Chiodi, S. Yoon, J. Janczak-Rusch, L.P.H. Jeurgens, The effect of thermal treatment on the stress state and evolving microstructure of Cu/W nano-multilayers, J. Appl. Phys. 120 (2016) 195107. A.V. Druzhinin, D. Ariosa, S. Siol, N. Ott, B.B. Straumal, J. Janczak-Rusch, L.P.H. Jeurgens, C. Cancellieri, Effect of the individual layer thickness on the transformation of Cu/W nano-multilayers into nanocomposites, Materialia 7 (2019) 100400. H.F. Rizzo, T.B. Massalski, M. Nastasi, Metastable crystalline and amorphous structures formed in the Cu-W system by vapor deposition, Metall. Trans. A 24 (1993) 1027–1037. R.L. Zong, S.P. Wen, F. Zeng, Y. Gao, F. Pan, Nanoindentation studies of Cu–W alloy
Surface & Coatings Technology 381 (2020) 125179
J. Xue, et al.
[49] J. Yan, L. Ning, F. Xu, Z. Ying, The strengthening effect of spinodal decomposition and twinning structure in MnCu-based alloy, Mater. Sci. Eng., A 618 (2014) 205–209.
films prepared by magnetron sputtering, J. Alloys Compd. 464 (2008) 544–549. [47] R.W. Cahn, Aluminium-based glassy alloys, Nature 341 (1989) 183–184. [48] D. Tabor, The Hardness of Metals (Oxford), (1951).
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