Deposition, Characterization and Performance

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Nano-layered superlattice nitride coatings with two or multi-layer constituents were deposited by unbalanced closed-field reactive magnetron sputtering ...
Surfaces and Interfaces in Nanostructured Materials II Edited by S.M. Mukhopadhyay, N.B. Dahotre, S. Seal, and A.Agarwal TMS (The Minerals, Metals & Materials Society), 2006

DEPOSITION, CHARACTERIZATION AND PERFORMANCE EVALUATION OF NANOLAYERED SUPERLATTICE NITRIDE COATINGS Q. Yang and L.R. Zhao Structures and Materials Performance Laboratory Institute for Aerospace Research National Research Council Canada Ottawa, Ontario, Canada Keywords: Superlattice Coatings, Characterization, Hardness Enhancement, Dry Sliding Wear, Thermal Stability Abstract Nano-layered superlattice nitride coatings with two or multi-layer constituents were deposited by unbalanced closed-field reactive magnetron sputtering technique using pure metallic targets. The modulation periods of the superlattices were measured from the reflection peak positions in lowangle X-ray reflectivity spectra using the modified Bragg’s law. The superlattice structures can lead to remarkably higher hardness than monolayered nitride coatings; and the hardness enhancement is closely related to the superlattice layer constituents, modulation period and coating crystallographic orientation. The results of post-annealing tests indicated that nitride superlattices have a high thermal stability. The results of pin-on-disc dry sliding tests showed that the nitride superlattices exhibit lower friction coefficients and markedly higher wear resistances than commercial TiN hard coatings. These tribological properties make these superlattice coatings good candidates for surface protection against wear in engineering applications. Introduction Superlattices consist of repeating layered structures of different constituent materials with individual layer thickness in the nanometer range. Thin film superlattices, especially transition metal nitrides/nitrides, have drawn a great deal of attention in coating research because their hardness enhancement capability is much higher than predicted by the rule of mixtures [1,2]. In addition to high hardness, superlattice coatings can also offer good oxidation resistance, improved toughness and low coefficient of friction depending the chemistry of layer constituents [3,4,5]. Several studies have shown that nitride superlattices, such as TiAlN/VN [3], TiN/CrN [6], TiN/TaN [7], TiN/NbN and TiN/MoN [8], can outperform single layered TiN in terms of wear resistance. However, the detailed tribological behavior of superlattices and its relation with hardness need further investigation. For applications involving high temperatures, such as surface protection of cutting tools used for high speed machining, the ability of superlattices to maintain structural and mechanical stability is crucial; therefore, the thermal stability of superlattices should be investigated. Although superlattice coatings with two-layer constituents have been successfully applied to cutting and stamping tools, it is challenging to apply these coatings to meet multi-functional requirements for certain applications. As forming multicomponent nitride coatings by alloying can adapt hard coatings to specific tribological applications [9], incorporating multiple layer constituents in superlattice coatings could provide 79

more flexibility in coating design to meet engineering requirements. So far, limited efforts have been made to study superlattice systems with more than two-layer constituents [10,11]. In this study, superlattice coatings with two- and four-layer constituents are deposited. High resolution TEM, low-angle X-ray reflectivity and high-angle X-ray diffraction techniques are used to characterize the superlattice structures. The effects of superlattice layer constituents, modulation period and coating crystallographic orientation on hardness enhancement are discussed. The thermal stability and tribological performances of the superlattices are also evaluated. Deposition of Superlattice Flat substrate specimens of a Ni-base alloy or Ti6Al4V were mechanically polished down to mirror finish with a final polishing using 1Pm diamond paste, followed by ultrasonic cleaning in Varsol and anhydrous ethyl alcohol. Nitride superlattices were deposited on the substrates in a TEER 650 closed-field unbalanced magnetron sputtering coater by reactive sputtering in a mixture of Ar and N2 gases. Fig.1 shows a schematic of the deposition chamber, in which four metallic targets were installed. An optical emission monitor (OEM) was used to regulate nitrogen supply in the chamber through monitoring the luminous emittance of plasma in front of one target. If targets are arranged, in a way such that Fig. 1. Schematic showing the deposition targets 1 and 3 are of the same material chamber of the magnetron sputtering coater. while targets 2 and 4 are of another material, superlattice coatings with twolayer constituents can be deposited. If the four targets are different, superlattices with four-layer constituents can be achieved. To deposit superlattices with alternate layers of similar thickness, the sample rotating stage was either 1) stopped periodically to expose the specimens to each target for predetermined time periods, or 2) continuously rotated to expose the substrate to each target with predetermined currents. In this study, TiN/CrN superlattices and novel multiconstituent superlattices (ZrN/MoN/NbN/AlN and TiN/AlN/ZrN/Mo2N) were deposited on the substrates. The total thickness of the superlattice films was in the range of 4-6 Pm. Superlattice Characterization The nanolayered structures of TiN/CrN and TiN/AlN/ZrN/Mo2N superlattice coatings are revealed by TEM bright-field imaging and energy dispersive X-ray spectra (EDS) mapping (Fig. 2). From the TEM micrographs, the average modulation periods are measured to be 8.3 nm and 21 nm, respectively for the TiN/CrN and TiN/AlN/ZrN/Mo2N superlattices. While TEM is a useful technique to analyze superlattice structures, the method is destructive and timeconsuming. In contrast, the low-angle X-ray reflectivity technique is non-destructive and time efficient, and thus widely used to characterize superlattices.

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c)

Fig. 2. The TEM and EDS elemental mapping images of a TiN/CrN superlattice are shown in a) and b), respectively. The TEM image of a TiN/AlN/ZrN/Mo2N superlattice is shown in c). The low-angle reflectivity spectrum of a superlattice results from the reflection of X-rays from layer interfaces. When X-rays travel across layer interfaces, refraction occurs due to the different refraction indexes of the constituent layers. The reflection peaks of different orders in the lowangle X-ray reflectivity spectra occur at 2T positions, usually given by the modified Bragg’s law by considering the refractions at layer interfaces [12]: mO 2/ sinT m (1 

G sin 2 T m

)

[1]

or in a simplified form:

sin 2 T m

(mO / 2/ ) 2  2G

[2]

where m is the order of the reflection, O is the X-ray wavelength. G is related to the average refractive index of n, and / is the modulation period of a superlattice. Plotting sin2Tm vs. m2 produces a straight line, and / is determined from the slope of the line by /

O

. The modulation period values 2 slope Fig. 3. Low-angle X-ray reflectivity presented in this study were calculated using spectra of superlattcie coatings including this method. Fig. 3 presents the low-angle Xthe two from the superlattices in Fig. 2. ray reflectivity spectra including the two from the superlattices shown in Fig. 2. The strong reflection peaks indicate that these multilayered coatings have abrupt interfaces, which is a beneficial characteristic for hardness enhancement. The modulation periods measured from 81

low-angle X-ray reflectivity spectra show good agreement with the values measured from TEM technique, indicating that the low-angle X-ray reflectivity technique is an accurate method to measure modulation periods. Fig. 4 shows high-angle XRD spectra of superlattices with (111) preferred orientation. Positive and negative satellite peaks on both sides of the main Bragg (111) reflection can be observed with the positions of the m-th order peaks given by [13]: sin T r

sin T B r

mO 2/

[3]

where TB is the position of the main Bragg reflection, / is the modulation period, Ois the X-ray wave length, and Tr are the positions of the m-th order positive (+) and negative (-) satellite peaks. For a given m, the satellite peaks Fig. 4. High-angle XRD spectra of (111) become closer to the main Bragg reflection as preferred oriented TiN/CrN superlattice the superlattice modulation period increases. coatings showing satellite peaks The appearance of the satellite peaks surrounding the main Bragg peak. demonstrates the coatings do have superlattice structures. The superlattice modulation periods can be calculated from the satellite peak positions using Eq. [3]. However, the modulation periods calculated are normally not as accurate as the values obtained from the low-angle X-ray reflectivity technique. Also, since the intensity of the satellite peaks decreases as the intensity of the main Bragg reflection decreases, satellite peaks may not be observed if the main superlattice Bragg peak is weak. Microstructure and Hardness Enhancement

TiN/CrN superlattices with (200) and (111) preferred orientations were deposited by controlling the processing conditions. The influences of modulation period and preferred orientation on hardness are shown in Fig. 5. For both orientations, the hardness first increases with /, followed by a rapid decrease, giving rise to a maximum at /|10 nm. Regarding the effect of preferred orientation, the (200)-oriented superlattices exhibit a much higher hardness enhancement at /|10 nm than the (111)-oriented superlattices. Hardness enhancement in the superlattices can be explained on the basis of dislocation movement within and between layers. When the superlattice modulation period is less than the optimum value for the peak hardness, dislocations can move between layers. The force required to move dislocations across the layer interfaces is related to the difference in shear moduli of the two layer constituents and to the interface width. The maximum hardness enhancement caused by dislocation glide across layer interfaces (i.e. the image effects) can be expressed as [14]:

H max  H A

3(G B  G A ) sin T mS 2

[4]

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where Hmax is the maximum expected hardness of the superlattice, HA is the hardness of the layer with a lower shear modulus, GA and GB are the shear moduli of the A and B layers (GB > GA), m is the Taylor factor, and T is the smallest angle between the slip planes and layer interfaces. When the modulation period is larger than the optimum value, it becomes easier to move dislocations within an individual layer, and the force to move a dislocation decreases as the superlattice period increases. Considering the primary {110}< 110 > slip system and the secondary {111}< 1 1 0 > slip system [15], (200)oriented superlattices have higher T values than (111)-oriented superlattices, which explains the difference in hardness enhancement of the differently oriented superlattices. As indicated by Eq. [4], in order to achieve significant enhancement in hardness, the layer constituents should have a large difference in shear moduli. If the layer constituents have similar shear moduli, little hardness enhancement will be achieved. To verify this assumption, (200) oriented ZrN/NbN superlattices with different modulation periods were deposited, and hardness as a function of modulation period is shown in Fig. 5. As the shear moduli of ZrN and NbN are almost identical (175 and 177 GPa, respectively), ZrN/NbN superlattices show no hardness enhancement at all. Fig. 5. Changes in hardness as a function of modulation period of TiN/CrN and ZrN/NbN superlattices.

Fig. 6. a) Changes in the superlattice structures and b) mechanical properties as a function of modulation period for the ZrN/MoN/NbN/AlN superlattices. “s” in a) indicates the peaks of substrate.

In TiN/CrN and ZrN/NbN superlattices, the layer constituents have the same equilibrium cubic B1 structure. As modulation period changes, the coatings do not show crystallographic changes and retain the same cubic B1 structure. These superlattices are iso-structural. In non-isostructural superlattice coating systems such as TiN/AlN, the two layer constituents have different 83

equilibrium crystallographic structures. With a small modulation period, one layer constituent will grow with a metastable structure that is the same as the other layer constituent, and coherent layer interfaces are formed to reduce the interfacial energy. This phenomenon is called the template effect. However, with a large modulation period, the layers will grow to form their own equilibrium structures. This makes the analysis of the hardness data more complicated. For example, among the four constituents in the ZrN/MoN/NbN/AlN superlattices, ZrN and NbN have an equilibrium rock-salt B1 (cubic) structure, while MoN and AlN an equilibrium hexagonal structure. The superlattices yield a repeated layer sequence of ZrN/MoN/NbN/AlN from substrate to coating surface, in which no adjacent layers have the same equilibrium crystallographic structures. This layer arrangement, with small modulation periods, leads to the growth of cubic MoN and AlN layers with the same B1 structure as the ZrN and NbN layers (Fig. 6a)). The superlattice coatings also have a (200) preferred orientation at smaller modulation periods. At modulation periods greater than 5 nm, (200) preferred orientation diminishes, and several new peaks around the cubic (111) peak are observed and identified as hexagonal MoN and AlN. These results indicate that when the modulation period is larger than a critical value, more stable hexagonal MoN and AlN layers would grow during deposition, as the bulk energy, instead of interfacial energy, becomes dominant. In this case, the loss of collaborative growth between different layers is responsible for the diminishing coating texture. The influence of the modulation period on coating hardness and Young’s modulus is shown in Fig. 6b). Both properties exhibit a dramatic decrease if / increases from 2.5 nm to 5.4 nm. The changes in coating mechanical properties reflect the microstructure evolution with the superlattice modulation period. The strong (200) textured superlattices at small modulation periods (/ 1.2 and 2.5 nm) result in high hardness, while the dramatic decrease in hardness corresponds to the loss of crystallographic preferred orientation and the appearance of stable hexagonal MoN and AlN layers. Thermal Stability

Fig. 7. a) Changes in hardness of the TiN/CrN superlattices as a function of annealing temperature and b) the evolution of the low-angle X-ray reflectivity spectra of the specimen #1 in a) with annealing temperature.

Superlattice coatings are emerging as promising candidates for tribological applications. However, at elevated temperatures, the superlattice structure can become unstable with time, thus degrading coating performance. In miscible superlattice systems, thermally activated interdiffusion of the metal species between adjacent layers at elevated temperatures produces an alloying effect between the layered nitrides as well as altering the layer thickness. This process 84

partially contributes to the gradual vanishing of the layered structure. For immiscible superlattices, the superlattice structures are stable up to temperatures as high as 1000 qC. Above this temperature, the layered structure transforms to an equiaxed two-phase structure within a relatively short period of time [16]. The thermal stability of TiN/CrN superlattices was assessed by annealing the coated specimens in vacuum at temperatures ranging from 250qC to 1050qC for 4 hours. The hardness changes with annealing temperature are shown in Fig. 7a), where 650qC appears to be the critical temperature under which the superlattices can retain the as-deposited hardness. Above 650ºC, the hardness decreases rapidly and approaches the rule-of-mixture value at temperatures above 900qC. No appreciable influences of the modulation period and preferred orientation on the critical temperature have been observed. Fig. 7b) presents the corresponding low-angle X-ray reflectivity curves of the specimen 1 in Fig. 7a) ((200) oriented and /=9.5 nm) with the orders of reflection peaks marked. It is interesting to note that a close correlation exists between changes in the hardness value and reflectivity curve. In terms of the intensities of the reflection peaks, little changes are observed from the reflectivity curves below 600qC. Annealing at the critical temperature (650qC) leads to a slight weakening of the reflection peaks. However, annealing at 700qC causes an intensity decrease of the low order peaks and the disappearance of the high order peaks (4th and 5th peaks). All the reflection peaks eventually disappear as the annealing temperature increases to 900qC and 1050qC. Correspondingly, the hardness drops sharply at 700qC and approaches the rule-of-mixture value at 900qC and 1050qC. When the annealing temperature is sufficiently high, interdiffusion occurs and leads to the formation of diffused interfaces between the layers. This structural change is reflected by the weakening or disappearing of the reflection peaks. It is suggested that the interface broadening above 650qC might be responsible for the sharp decrease in hardness. Dry Sliding Wear Resistance

TiN/CrN superlattice coatings have lower coefficients of friction (0.75~0.95) than the TiN coatings (1.0~1.20). The coefficients of friction also decrease with increasing normal load, a phenomenon reported previously [17,18,19]. Depth profiles of the wear tracks formed at 10N normal load show that TiN coatings exhibit a much deeper wear track (1.85 Pm) than superlattice coatings. The maximum depth of the wear tracks on multilayered superlattices decreases with increasing hardness. For the superlattice of H=35.4 GPa, the maximum wear track depth on is only 0.11 Pm, ~6% of that for TiN. Fig. 8 shows the wear rates of the coatings as a function of Fig. 8. Wear rates of the TiN and TiN/CrN hardness and normal loads. It is evident that superlattice coatings as a function of increase in hardness of the superlattices reduces hardness and the normal loads applied. the wear rate. Furthermore, the wear rates are about 2.5 to 12 times lower than the TiN coatings, depending on the superlattice hardness. At H=35.4 GPa, the wear rates are less than 9% of that for TiN.

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The coefficients of friction of a TiN coating and two multi-component coatings with /=31.1nm and 324nm respectively, as a function of sliding distance, are shown in Fig. 9. The multicomponent coating with /=31.1nm has a relatively stable coefficient of friction with an average value of 0.5, significantly lower than that of the TiN coating (1.06). This tribological behavior is common for various Mo-containing systems. For example, Mo alloying leads to a dramatic reduction in the friction of TiMoN coatings due to the tribo-chemically formed solid lubricant MoO3 on the wear track surface [20]. The specific wear rate of this TiN/AlN/ZrN/Mo2N coating (9.88E-8 mm3/(N*m)) is more than one order of magnitude lower than that of TiN (1.80E-6 mm3/(N*m)). The higher H3/E2 value, together with the lower coefficient of friction, is responsible for this improved wear resistance. For the coating with /=324nm, Mo2N is the top layer and the thickness of each layer is in the Fig. 9. Coefficients of friction as a function sub-micron range. Because of the beneficial role of sliding distance for TiN and two of Mo in reducing the coefficient of friction, the TiN/AlN/ZrN/Mo2N superlattice coatings. coating first shows a very low coefficient of The specific wear rates are also provided in friction in the range of 0.25-0.3. With the exposure of the underlying layers, such as ZrN the plot. and AlN, the coefficient of friction gradually increases. It is clear that both the coefficient of friction and the wear rate can vary substantially during the sliding process if the coating layers are sub-micron thick. Summary

Reactive magnetron sputtering has been used to deposit nanolayered superlattice coatings with two and four different layered constituents. Low-angle X-ray reflectivity and high-angle XRD techniques are effective in characterizing superlattice structures. In this study it has been revealed that there are several important factors affecting the hardness enhancement of a superlattice: layer constituents, modulation period, interface width, layer crystallographic structures and preferred orientation. The results of wear tests show that both TiN/CrN and TiN/AlN/ZrN/Mo2N superlattices demonstrate much higher wear resistance than the monolayered TiN coatings due to the combination of high hardness, high H3/E2 and low coefficients of friction. The authors wish to thank National Research Council Canada for the financial support and Robert McKeller for his technical support in the project. References

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