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Ni3Al multilayers
Thin Solid Films 424 (2003) 93–98
Sputter deposited nanocrystalline Ni-25Al alloy thin films and NiyNi3Al multilayers R. Banerjee*, G.B. Thompson, P.M. Anderson, H.L. Fraser Department of Materials Science and Engineering, The Ohio State University, 177 Watts Hall 2041 College Road, Columbus OH, USA
Abstract Thin films of nominal composition Ni-25at%Al have been sputter deposited from a target of the intermetallic compound Ni3Al at different substrate deposition temperatures. The film deposited on an unheated substrate exhibited a strongly textured columnar growth morphology and consisted of a mixture of metastable phases. Nanoindentation studies carried out on this film exhibited a strong strain hardening tendency. In contrast, the film deposited at 200 8C exhibited a recrystallized non-textured microstructure consisting of grains of a partially ordered Ni3 Al phase. At higher deposition temperatures (;400 8C), larger grains of the bulk equilibrium, long-range ordered, Ll2 Ni3Al phase were observed in the film. Unlike the film deposited on an unheated substrate, the films deposited at elevated temperatures did not exhibit any dependence of the hardness on the indentation depth and, consequently no strain hardening. The average hardness of the film deposited at 200 8C was higher than the one deposited at 400 8C. In addition to monolithic Ni-25Al thin films, multilayered NiyNi3Al thin films were also deposited. Multilayers deposited non-epitaxially on unheated substrates exhibited a strong {111} fiber texture while those deposited epitaxially on (001) NaCl exhibited a {001} texture. Free-standing multilayers of both types of preferred orientations as well as of different layer thicknesses were deformed in tension untill fracture. Interestingly, the {111} oriented multilayers failed primarily by a brittle fracture while the {001} multilayers exhibited features of ductile fracture. 䊚 2002 Elsevier Science B.V. All rights reserved. Keywords: Ni-25Al; Alloy thin films; NiyNi3Al
1. Introduction Thin films of alloys and intermetallic compounds have a wide range of applications. These include conducting barrier layers and passivation layers in microelectronic devices such as integrated circuits, active layers in giant magnetoresistive recording heads in magnetic storage devices and structural coatings for high temperature aerospace applications. Formation of metastable phases in thin films occurs rather commonly because of the rapid cooling rates involved during the transformation from a vapor to a solid film. Nickel-based superalloys are used extensively in the aerospace industry, due to excellent mechanical properties such as yield strength and creep resistance at elevated temperatures w1x. The metal–intermetallic composite alloy consists of a g-Ni(Al) matrix with an *Corresponding author. Tel.: q1-614-292-2553; fax: q1-614-2927523.
intermetallic g9-Ni3Al phase. The latter serves as the primary strengthening phase and often takes the form of cuboidal precipitates with a {001} Ni(Al)yy {001} Ni3Al, N100M Ni(Al)yyN100M Ni3Al interface orientation. In addition to Ni-based superalloys, a large amount of research has been carried out on the intermetallic compound Ni3Al. However, most of the studies have been directed towards the thermo-mechanical processing and study of deformation mechanisms in bulk Ni3Al based alloys. There have been rather limited studies on the thin film form of this compound, especially from the viewpoint of mechanical properties. Some of the earlier reports include investigations of the metastable phases forming in these thin films. Harris and co-workers w2x deposited chemically disordered Ni25Al (all compositions in this paper are in atomic%) thin films on unheated as well as liquid nitrogen cooled polycrystalline copper substrates temperatures using high vacuum evaporation. In a different study, nanocrystalline Ni–Al alloy films were prepared by vacuum
0040-6090/03/$ - see front matter 䊚 2002 Elsevier Science B.V. All rights reserved. PII: S 0 0 4 0 - 6 0 9 0 Ž 0 2 . 0 0 9 2 4 - 0
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evaporation on NaCl (001) substrates at temperatures in the range of 20–400 8C w3x. A number of studies have focussed on the effect of irradiation on Ni–Al alloy thin films. These include co-evaporation of Ni–Al films in the composition range of 18–23 at% Al by Alexander et al. w4x, and deposition of NiyAl multilayers and subsequent ion-beam mixing of the multilayers by Eridon et al. w5x. The ion-beam mixing resulted in the formation of a number of metastable phases including an amorphous phase, a disordered crystalline phase and a hexagonal close-packed (h.c.p.) phase w5x. The formation of a metastable h.c.p. phase has also been reported in irradiated co-evaporated Ni–Al alloy thin films w4x. The present study includes a detailed characterization of the microstructural evolution in sputter deposited Ni25Al thin films and its correlation to initial investigations of the mechanical properties of these films using nanoindentation. In addition, to monolithic thin films, a potentially useful morphology may be a multilayered thin film geometry, since large measured values of hardness have been reported for several material systems w6x. The large hardness values may stem from the confinement of plastic deformation to individual layers, e.g. w7x. Therefore, the study has been extended to sputter deposited NiyNi3Al multilayered thin films with varying layer thicknesses. This paper reports on the fracture behavior of a series of nanolayered NiyNi3Al thin films, with particular attention to the effect of layer thickness and interfacial orientation on the confinement of plastic deformation and on fracture morphology. 2. Experimental procedure All the thin films have been deposited in a custom built UHV magnetron sputtering unit. The base pressure prior to sputtering was ;3=10y8 torr and the argon pressure during sputtering was ;2=10y3 torr. The nominal rate of deposition was 0.3 nmys with the gun being power regulated at 200 W. The monolithic thin films of nominal composition Ni-25Al were sputter deposited from an intermetallic Ni3 Al target. The monolithic films were deposited on oxidized Si (001) wafers that had a 200-nm thick amorphous SiO2 layer. The films were deposited at different substrate temperatures, ;45 8C (i.e. unheated), 200 8C and 400 8C. The nominal thickness of the monolithic films was ;1 mm. EDS analysis of the films, carried out in a scanning electron microscope, revealed the average film composition to be Ni-25at% Al. For depositing the multilayers, in addition to the Ni3Al target, an elemental Ni target was also used. The multilayers were deposited on NaCl (001) substrates at room temperature and at ;450 8C. X-Ray diffraction experiments were conducted on the as-deposited monolithic films and the multilayers using a Scintag PAD V diffractometer using CuKa as the
incident radiation. The films were also characterized in a transmission electron microscope (TEM) using a Philips CM200 TEM operating at an accelerating voltage of 200 kV. The details of the TEM sample preparation procedure are given elsewhere w8x. Nanoindentation studies were carried out on the monolithic Ni-25Al films using a Nano II indentor. For the NiyNi3Al multilayers, the fracture surfaces were imaged using scanning electron microscopy in a Philips XL30 FEG SEM. 3. Results and discussion The microstructure of the monolithic Ni-25Al thin film deposited on the unheated Si substrate (;45 8C) is shown in the bright field micrograph from a planview TEM specimen in Fig. 1a. The grain size ranges from 30 to 50 nm, approximately. Based on electron diffraction (inset in Fig. 1a), the phases present in this film are an f.c.c. phase with a lattice parameter of af.c.c.s0.356 nm designated as Ni(Al) because of the similarity of the lattice parameter with that of Ni, a h.c.p. phase with a lattice parameter of ah.c.p.s0.252 nm, designated as Ni(Al)9 because of the similarity in the nearest neighbor spacing with that of Ni, and a third Al-rich tetragonal phase with lattice parameters atets 0.356 nm and ctets0.521 nm. The details regarding the identification of this phase are discussed elsewhere w9x. Furthermore, the electron and X-ray diffraction results indicate the presence of a {111}f.c.c.yy {0001} h.c.p. preferred orientation in this film The most prominent diffracted intensity arising from the Al-rich tetragonal phase which does not overlap with diffraction rings from either the f.c.c. or h.c.p. Ni-rich phases is observed at an interplanar spacing of ;0.147 nm. When taken together with weaker diffraction rings from this phase, this reflection can be tentatively indexed as the {022} ring of a f.c.c.-like structure with a lattice parameter ;0.416 nm. A lack of any superlattice reflections in the diffraction patterns indicates that these phases are chemically disordered. A dark field image (Fig. 1b) has been recorded using the diffracted intensity at an interplanar spacing of ;0.147 nm from the Al-rich tetragonal phase. These images clearly reveal that this phase is distributed in the form of particles of average size ;5 nm. The particles are in fact distributed within the grains of both the f.c.c. Ni(Al) and h.c.p. Ni(Al)9 phases. For the Ni-25Al film deposited at 200 8C, the microstructure consists of grains of average size ;15 nm (Fig. 1c), smaller as compared to the film deposited on the unheated substrate. Furthermore, both X-ray and electron diffraction suggest the absence of any preferred orientation in this film. Despite the fact that the intensity maxima in the diffraction patterns can be consistently indexed based on a single f.c.c. phase of composition Ni-25Al, partial or short-range chemical ordering cannot be ruled out in this film. It is likely that recrystallization
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Fig. 1. (a) Bright field TEM micrograph from the Ni-25Al film deposited on an unheated Si substrate (;45 8C). A selected area ring diffraction pattern from the same film exhibiting rings from the f.c.c. Ni(Al), h.c.p. Ni(Al)9 and a third Al-rich tetragonal phase. (b) Dark field TEM micrograph imaged using the (022) reflection from the Al-rich tetragonal phase. (c) A bright field TEM micrograph from the Ni-25Al film deposited at 200 8C together with the corresponding diffraction pattern in the inset. (d) A bright field TEM micrograph from the Ni-25Al film deposited at 400 8C. A selected area diffraction pattern from the film is shown in the inset at the top right corner. A w111x zone axis microdiffraction diffraction from a single grain of the L12 Ni3Al phase is shown in the inset at the bottom left corner.
induced by partial chemical ordering has occurred in this film deposited at 200 8C w10x. In addition to the partially ordered f.c.c. phase, a small volume fraction of the metastable Al-rich tetragonal phase also forms in this film. A bright field micrograph from a plan-view specimen of the film deposited at 400 8C is shown in Fig. 1d. The average grain size in this film is ;125 nm. Selected area diffraction patterns from this film (inset at top right corner of Fig. 1d) could be consistently indexed based on a single L12 Ni3Al phase and did not exhibit any diffracted intensities from the Alrich tetragonal phase. A w111x zone axis microdiffration pattern from a grain of this film is shown in Fig. 1d (inset at bottom left corner). The presence of superlattice
reflections indicates long range chemical ordering of the L12 type in the Ni3Al grains. The mechanical properties of the monolithic films have been measured using nanoindentation. The films have been indented upto a maximum depth of 200 nm which is ;20% of the film thickness (;1 mm) in order to minimize substrate effects. Plots of the modulus and hardness of the films as a function of the depth of indentation are shown in Fig. 2a and b, respectively. All three films, deposited at different substrate temperatures, do not exhibit any significant change in the modulus as a function of indentation depth. The moduli of the films are 200, 200 and 180 GPa, for the films deposited at 400, 200 and 45 8C, respectively (Fig. 2a). The hardness
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Fig. 2. (a) Plot of the modulus as a function of depth of indentation for the three Ni-25Al films deposited at different temperatures. (b) Corresponding plot of the hardness as a function of depth of indentation for the three Ni-25Al films deposited at different temperatures.
of the films for the corresponding indentation depths (Fig. 2b) clearly indicates a significant difference between the plots for the 200 and 400 8C deposited films as compared with the film deposited at 45 8C. Thus, while the films deposited at elevated temperatures do not exhibit any variation in hardness with the depth of indentation, the film deposited at 45 8C exhibits a continuous increase in the hardness with indentation depth. For the 45 8C film, the hardness increases from a value of ;2 GPa at a depth of ;30 nm to ;6 GPa at a depth of ;160 nm. Coupling this observation with the fact that there is hardly any change in the modulus over the same range of indentation depths, it can be concluded that there is a strong strain hardening tendency in the Ni-25Al film deposited at ;45 8C. The highest hardness is exhibited by the film deposited at 200 8C (;7.3 GPa) while the film deposited at 400 8C exhibits a hardness ;6.5 GPa. The higher hardness for the film
deposited at 200 8C can be attributed to the smaller grain size in this film as compared with the film deposited at 400 8C. The NiyNi3Al multilayers have been deposited on NaCl (001) substrates at ambient temperatures and also at ;450 8C. The films deposited on unheated substrates exhibited a N111M preferred orientation with the interfaces along the {111} Niyy {111} Ni3Al. On the other hand, the multilayers deposited at 450 8C grew epitaxially and, consequently, exhibited a N001M preferred orientation with the interfaces along the {001} Niyy {001} Ni3Al. Subsequent to deposition, the multilayered thin films have been lifted off the NaCl (001) substrates by floating in water. These free-standing multilayers were fractured in uniaxial, in-plane tension. The macroscopic response of all the samples is brittle, with any plastic deformation occurring only at or near the fracture surface w11x. However, the scanning electron micrographs of these samples reveal striking differences between N001M and N111M oriented multilayers. Fig. 3a and b show the fracture surfaces of N001M samples with 20- and 120-nm layer thicknesses, respectively, viewed approximately normal to the fracture surface. The surfaces display ductile void formation, similar to failure surfaces of ductile monolithic metallic samples. Despite the different layer thickness, both samples display a range of void sizes extending up to 200 to 300 nm in diameter. Clearly, the larger voids have been produced by extensive co-deformation of Ni and Ni3Al layers. Another outcome of the extensive co-deformation is that individual layers cannot be discerned on the fracture surfaces. These results are remarkable since Ni3 Al is an intermetallic compound that displays brittle behavior in bulk form. Fig. 3c and d show SEM fracture surfaces of N111M samples with 20- and 120-nm layer thickness, viewed approximately normal to the fracture surface. The N111M samples appear to have more brittle fracture surface features than the counterpart N001M cases. In the 120nm geometry, for example, co-deformation is so extensive in the N001M orientation that individual layers cannot be discerned (Fig. 3b), yet it is so limited in the N111M orientation that the horizontal bi-layered structure and vertical columnar grain structure are clearly delineated on the fracture surface (Fig. 3d). The 20-nm geometry also supports the observation that the N111M orientation has more brittle fracture surface features. In particular, Fig. 3c shows evidence of ductile void coalescence (lower right) in the N111M orientation, but there are other regions (upper right) where co-deformation is limited and individual layers can be discerned. In contrast, the N001M geometry has profuse co-deformation (Fig. 3a). It is quite interesting to note the marked influence of interfacial orientation on the fracture behavior of these nanolayered metal–intermetallic composites. Samples
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ples may be a result of interfaces that offer little resistance to dislocation transmission and little subsequent work hardening. Furthermore, the marked increase in co-deformation for N001M samples may be a result from interfaces that transmit slip more easily than those in N111M samples. Although experimental results for Niy Ni3Al multilayers have been discussed here, initial observations for g–Ni(Al)yg9–Ni3 Al multilayers are similar. 4. Summary A detailed characterization of monolithic Ni-25at%Al thin films, sputter deposited at different substrate temperatures revealed the formation of novel metastable phases in the film deposited on an unheated substrate (;45 8C), a fine-grained recrystallized microstructure in the film deposited at 200 8C and a coarser grained microstructure consisting of the single L12 Ni3Al phase in the film deposited at 400 8C. Nanoindentation studies carried out on these monolithic films indicate a strong strain hardening tendency in the film deposited at ;45 8C and the highest average hardness (;7.3 GPa) for the film deposited at 200 8C. In addition to monolithic Ni-25Al films, the fracture behavior of sputter deposited NiyNi3Al multilayers has also been investigated. N111M Oriented multilayers primarily exhibited a brittle mode of failure while N001M oriented multilayers primarily exhibited a ductile mode of failure. Furthermore, multilayers with thinner individual layers were found to be more ductile as compared to those with thicker individual layers. Fig. 3. SEM micrographs of the different NiyNi3 Al multilayers, viewed normal to the fracture surface. (a) N001M Oriented 20 nm Niy20 nm Ni3Al multilayer. (b) N001M Oriented 120 nm Niy120 nm Ni3Al multilayer. (c) N111M oriented 20 nm Niy20 nm Ni3Al multilayer. (d) N111M oriented 120 nm Niy120 nm Ni3Al multilayer.
with N001M interface orientation exhibit more ductile fracture features than N111M samples. Corresponding modeling of confined layer slip in multilayers shows that multilayers display remarkable strength when interfaces are able to confine dislocations to individual layers w7x. In such cases, profuse co-deformation of layers is suppressed and more fracture surface features are expected. Corresponding transmission electron microscope studies of in situ deformation in CuyNi multilayers indicate that misfit dislocations serve as pinning sites for glissile dislocations that attempt to transmit across interfaces. An implication is that semi-coherent interfaces are expected to be stronger barriers to dislocation transmission than coherent interfaces. Thus, samples with 20-nm layer thickness and corresponding coherent interfaces are expected to be less resistant to slip transmission than the thicker, 120 nm, counterparts. The marked shear localization displayed by the 20-nm sam-
Acknowledgments The authors would like to acknowledge the support from the Air Force Office of Scientific Research under Grant F49620-96-1-0238 and the Center for the Accelerated Maturation of Materials at the Ohio State University for this research. The authors would also like to thank Dr Douglas Smith and Dr Tim Foecke from the National Institute of Standards and Technology for their assistance in carrying out the nanoindentation studies. References w1x C.T. Simms, W. Hagel (Eds.), The Superalloys, Wiley, New York, 1972. w2x S.R. Harris, D.H. Pearson, C.M. Garland, B. Fultz, J. Mater. Res. 6 (1991) 2019. w3x T. Kizuka, N. Mitarai, N. Tanaka, J. Mater. Sci. 29 (21) (1994) 5599. w4x D.E. Alexander, G.S. Was, L.E. Rehn, J. Appl. Phys. 69 (4) (1991) 2021. w5x J.E. Eridon, G.S. Was, L.E. Rehn, J. Mater. Res. 3 (4) (1988) 626.
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R. Banerjee et al. / Thin Solid Films 424 (2003) 93–98 w6x B.M. Clemmons, H. Kung, S.A. Barnett, MRS Bull. 24 (2) (1999) 20–26. w7x P.M. Anderson, T. Foecke, P.M. Hazzledine, MRS Bull. 24 (2) (1999) 27. w8x G.B. Thompson, M.S. Thesis, The Ohio State University (1997).
w9x R. Banerjee, G.B. Thompson, H.L. Fraser, submitted to Philos. Mag. Lett. (2001). w10x G.B. Thompson, R. Banerjee, X.D. Zhang, P.M. Anderson, H.L. Fraser, submitted to Acta Mater. (2001). w11x R. Banerjee, J.P. Fain, P.M. Anderson, H.L. Fraser, to appear in Scripta Mater. (2001).
Sputter deposited nanocrystalline Ni-25Al alloy thin films and NiyNi3Al multilayers R. Banerjee*, G.B. Thompson, P.M. Anderson, H.L. Fraser Department of Materials Science and Engineering, The Ohio State University, 177 Watts Hall 2041 College Road, Columbus OH, USA
Abstract Thin films of nominal composition Ni-25at%Al have been sputter deposited from a target of the intermetallic compound Ni3Al at different substrate deposition temperatures. The film deposited on an unheated substrate exhibited a strongly textured columnar growth morphology and consisted of a mixture of metastable phases. Nanoindentation studies carried out on this film exhibited a strong strain hardening tendency. In contrast, the film deposited at 200 8C exhibited a recrystallized non-textured microstructure consisting of grains of a partially ordered Ni3 Al phase. At higher deposition temperatures (;400 8C), larger grains of the bulk equilibrium, long-range ordered, Ll2 Ni3Al phase were observed in the film. Unlike the film deposited on an unheated substrate, the films deposited at elevated temperatures did not exhibit any dependence of the hardness on the indentation depth and, consequently no strain hardening. The average hardness of the film deposited at 200 8C was higher than the one deposited at 400 8C. In addition to monolithic Ni-25Al thin films, multilayered NiyNi3Al thin films were also deposited. Multilayers deposited non-epitaxially on unheated substrates exhibited a strong {111} fiber texture while those deposited epitaxially on (001) NaCl exhibited a {001} texture. Free-standing multilayers of both types of preferred orientations as well as of different layer thicknesses were deformed in tension untill fracture. Interestingly, the {111} oriented multilayers failed primarily by a brittle fracture while the {001} multilayers exhibited features of ductile fracture. 䊚 2002 Elsevier Science B.V. All rights reserved. Keywords: Ni-25Al; Alloy thin films; NiyNi3Al
1. Introduction Thin films of alloys and intermetallic compounds have a wide range of applications. These include conducting barrier layers and passivation layers in microelectronic devices such as integrated circuits, active layers in giant magnetoresistive recording heads in magnetic storage devices and structural coatings for high temperature aerospace applications. Formation of metastable phases in thin films occurs rather commonly because of the rapid cooling rates involved during the transformation from a vapor to a solid film. Nickel-based superalloys are used extensively in the aerospace industry, due to excellent mechanical properties such as yield strength and creep resistance at elevated temperatures w1x. The metal–intermetallic composite alloy consists of a g-Ni(Al) matrix with an *Corresponding author. Tel.: q1-614-292-2553; fax: q1-614-2927523.
intermetallic g9-Ni3Al phase. The latter serves as the primary strengthening phase and often takes the form of cuboidal precipitates with a {001} Ni(Al)yy {001} Ni3Al, N100M Ni(Al)yyN100M Ni3Al interface orientation. In addition to Ni-based superalloys, a large amount of research has been carried out on the intermetallic compound Ni3Al. However, most of the studies have been directed towards the thermo-mechanical processing and study of deformation mechanisms in bulk Ni3Al based alloys. There have been rather limited studies on the thin film form of this compound, especially from the viewpoint of mechanical properties. Some of the earlier reports include investigations of the metastable phases forming in these thin films. Harris and co-workers w2x deposited chemically disordered Ni25Al (all compositions in this paper are in atomic%) thin films on unheated as well as liquid nitrogen cooled polycrystalline copper substrates temperatures using high vacuum evaporation. In a different study, nanocrystalline Ni–Al alloy films were prepared by vacuum
0040-6090/03/$ - see front matter 䊚 2002 Elsevier Science B.V. All rights reserved. PII: S 0 0 4 0 - 6 0 9 0 Ž 0 2 . 0 0 9 2 4 - 0
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evaporation on NaCl (001) substrates at temperatures in the range of 20–400 8C w3x. A number of studies have focussed on the effect of irradiation on Ni–Al alloy thin films. These include co-evaporation of Ni–Al films in the composition range of 18–23 at% Al by Alexander et al. w4x, and deposition of NiyAl multilayers and subsequent ion-beam mixing of the multilayers by Eridon et al. w5x. The ion-beam mixing resulted in the formation of a number of metastable phases including an amorphous phase, a disordered crystalline phase and a hexagonal close-packed (h.c.p.) phase w5x. The formation of a metastable h.c.p. phase has also been reported in irradiated co-evaporated Ni–Al alloy thin films w4x. The present study includes a detailed characterization of the microstructural evolution in sputter deposited Ni25Al thin films and its correlation to initial investigations of the mechanical properties of these films using nanoindentation. In addition, to monolithic thin films, a potentially useful morphology may be a multilayered thin film geometry, since large measured values of hardness have been reported for several material systems w6x. The large hardness values may stem from the confinement of plastic deformation to individual layers, e.g. w7x. Therefore, the study has been extended to sputter deposited NiyNi3Al multilayered thin films with varying layer thicknesses. This paper reports on the fracture behavior of a series of nanolayered NiyNi3Al thin films, with particular attention to the effect of layer thickness and interfacial orientation on the confinement of plastic deformation and on fracture morphology. 2. Experimental procedure All the thin films have been deposited in a custom built UHV magnetron sputtering unit. The base pressure prior to sputtering was ;3=10y8 torr and the argon pressure during sputtering was ;2=10y3 torr. The nominal rate of deposition was 0.3 nmys with the gun being power regulated at 200 W. The monolithic thin films of nominal composition Ni-25Al were sputter deposited from an intermetallic Ni3 Al target. The monolithic films were deposited on oxidized Si (001) wafers that had a 200-nm thick amorphous SiO2 layer. The films were deposited at different substrate temperatures, ;45 8C (i.e. unheated), 200 8C and 400 8C. The nominal thickness of the monolithic films was ;1 mm. EDS analysis of the films, carried out in a scanning electron microscope, revealed the average film composition to be Ni-25at% Al. For depositing the multilayers, in addition to the Ni3Al target, an elemental Ni target was also used. The multilayers were deposited on NaCl (001) substrates at room temperature and at ;450 8C. X-Ray diffraction experiments were conducted on the as-deposited monolithic films and the multilayers using a Scintag PAD V diffractometer using CuKa as the
incident radiation. The films were also characterized in a transmission electron microscope (TEM) using a Philips CM200 TEM operating at an accelerating voltage of 200 kV. The details of the TEM sample preparation procedure are given elsewhere w8x. Nanoindentation studies were carried out on the monolithic Ni-25Al films using a Nano II indentor. For the NiyNi3Al multilayers, the fracture surfaces were imaged using scanning electron microscopy in a Philips XL30 FEG SEM. 3. Results and discussion The microstructure of the monolithic Ni-25Al thin film deposited on the unheated Si substrate (;45 8C) is shown in the bright field micrograph from a planview TEM specimen in Fig. 1a. The grain size ranges from 30 to 50 nm, approximately. Based on electron diffraction (inset in Fig. 1a), the phases present in this film are an f.c.c. phase with a lattice parameter of af.c.c.s0.356 nm designated as Ni(Al) because of the similarity of the lattice parameter with that of Ni, a h.c.p. phase with a lattice parameter of ah.c.p.s0.252 nm, designated as Ni(Al)9 because of the similarity in the nearest neighbor spacing with that of Ni, and a third Al-rich tetragonal phase with lattice parameters atets 0.356 nm and ctets0.521 nm. The details regarding the identification of this phase are discussed elsewhere w9x. Furthermore, the electron and X-ray diffraction results indicate the presence of a {111}f.c.c.yy {0001} h.c.p. preferred orientation in this film The most prominent diffracted intensity arising from the Al-rich tetragonal phase which does not overlap with diffraction rings from either the f.c.c. or h.c.p. Ni-rich phases is observed at an interplanar spacing of ;0.147 nm. When taken together with weaker diffraction rings from this phase, this reflection can be tentatively indexed as the {022} ring of a f.c.c.-like structure with a lattice parameter ;0.416 nm. A lack of any superlattice reflections in the diffraction patterns indicates that these phases are chemically disordered. A dark field image (Fig. 1b) has been recorded using the diffracted intensity at an interplanar spacing of ;0.147 nm from the Al-rich tetragonal phase. These images clearly reveal that this phase is distributed in the form of particles of average size ;5 nm. The particles are in fact distributed within the grains of both the f.c.c. Ni(Al) and h.c.p. Ni(Al)9 phases. For the Ni-25Al film deposited at 200 8C, the microstructure consists of grains of average size ;15 nm (Fig. 1c), smaller as compared to the film deposited on the unheated substrate. Furthermore, both X-ray and electron diffraction suggest the absence of any preferred orientation in this film. Despite the fact that the intensity maxima in the diffraction patterns can be consistently indexed based on a single f.c.c. phase of composition Ni-25Al, partial or short-range chemical ordering cannot be ruled out in this film. It is likely that recrystallization
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Fig. 1. (a) Bright field TEM micrograph from the Ni-25Al film deposited on an unheated Si substrate (;45 8C). A selected area ring diffraction pattern from the same film exhibiting rings from the f.c.c. Ni(Al), h.c.p. Ni(Al)9 and a third Al-rich tetragonal phase. (b) Dark field TEM micrograph imaged using the (022) reflection from the Al-rich tetragonal phase. (c) A bright field TEM micrograph from the Ni-25Al film deposited at 200 8C together with the corresponding diffraction pattern in the inset. (d) A bright field TEM micrograph from the Ni-25Al film deposited at 400 8C. A selected area diffraction pattern from the film is shown in the inset at the top right corner. A w111x zone axis microdiffraction diffraction from a single grain of the L12 Ni3Al phase is shown in the inset at the bottom left corner.
induced by partial chemical ordering has occurred in this film deposited at 200 8C w10x. In addition to the partially ordered f.c.c. phase, a small volume fraction of the metastable Al-rich tetragonal phase also forms in this film. A bright field micrograph from a plan-view specimen of the film deposited at 400 8C is shown in Fig. 1d. The average grain size in this film is ;125 nm. Selected area diffraction patterns from this film (inset at top right corner of Fig. 1d) could be consistently indexed based on a single L12 Ni3Al phase and did not exhibit any diffracted intensities from the Alrich tetragonal phase. A w111x zone axis microdiffration pattern from a grain of this film is shown in Fig. 1d (inset at bottom left corner). The presence of superlattice
reflections indicates long range chemical ordering of the L12 type in the Ni3Al grains. The mechanical properties of the monolithic films have been measured using nanoindentation. The films have been indented upto a maximum depth of 200 nm which is ;20% of the film thickness (;1 mm) in order to minimize substrate effects. Plots of the modulus and hardness of the films as a function of the depth of indentation are shown in Fig. 2a and b, respectively. All three films, deposited at different substrate temperatures, do not exhibit any significant change in the modulus as a function of indentation depth. The moduli of the films are 200, 200 and 180 GPa, for the films deposited at 400, 200 and 45 8C, respectively (Fig. 2a). The hardness
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Fig. 2. (a) Plot of the modulus as a function of depth of indentation for the three Ni-25Al films deposited at different temperatures. (b) Corresponding plot of the hardness as a function of depth of indentation for the three Ni-25Al films deposited at different temperatures.
of the films for the corresponding indentation depths (Fig. 2b) clearly indicates a significant difference between the plots for the 200 and 400 8C deposited films as compared with the film deposited at 45 8C. Thus, while the films deposited at elevated temperatures do not exhibit any variation in hardness with the depth of indentation, the film deposited at 45 8C exhibits a continuous increase in the hardness with indentation depth. For the 45 8C film, the hardness increases from a value of ;2 GPa at a depth of ;30 nm to ;6 GPa at a depth of ;160 nm. Coupling this observation with the fact that there is hardly any change in the modulus over the same range of indentation depths, it can be concluded that there is a strong strain hardening tendency in the Ni-25Al film deposited at ;45 8C. The highest hardness is exhibited by the film deposited at 200 8C (;7.3 GPa) while the film deposited at 400 8C exhibits a hardness ;6.5 GPa. The higher hardness for the film
deposited at 200 8C can be attributed to the smaller grain size in this film as compared with the film deposited at 400 8C. The NiyNi3Al multilayers have been deposited on NaCl (001) substrates at ambient temperatures and also at ;450 8C. The films deposited on unheated substrates exhibited a N111M preferred orientation with the interfaces along the {111} Niyy {111} Ni3Al. On the other hand, the multilayers deposited at 450 8C grew epitaxially and, consequently, exhibited a N001M preferred orientation with the interfaces along the {001} Niyy {001} Ni3Al. Subsequent to deposition, the multilayered thin films have been lifted off the NaCl (001) substrates by floating in water. These free-standing multilayers were fractured in uniaxial, in-plane tension. The macroscopic response of all the samples is brittle, with any plastic deformation occurring only at or near the fracture surface w11x. However, the scanning electron micrographs of these samples reveal striking differences between N001M and N111M oriented multilayers. Fig. 3a and b show the fracture surfaces of N001M samples with 20- and 120-nm layer thicknesses, respectively, viewed approximately normal to the fracture surface. The surfaces display ductile void formation, similar to failure surfaces of ductile monolithic metallic samples. Despite the different layer thickness, both samples display a range of void sizes extending up to 200 to 300 nm in diameter. Clearly, the larger voids have been produced by extensive co-deformation of Ni and Ni3Al layers. Another outcome of the extensive co-deformation is that individual layers cannot be discerned on the fracture surfaces. These results are remarkable since Ni3 Al is an intermetallic compound that displays brittle behavior in bulk form. Fig. 3c and d show SEM fracture surfaces of N111M samples with 20- and 120-nm layer thickness, viewed approximately normal to the fracture surface. The N111M samples appear to have more brittle fracture surface features than the counterpart N001M cases. In the 120nm geometry, for example, co-deformation is so extensive in the N001M orientation that individual layers cannot be discerned (Fig. 3b), yet it is so limited in the N111M orientation that the horizontal bi-layered structure and vertical columnar grain structure are clearly delineated on the fracture surface (Fig. 3d). The 20-nm geometry also supports the observation that the N111M orientation has more brittle fracture surface features. In particular, Fig. 3c shows evidence of ductile void coalescence (lower right) in the N111M orientation, but there are other regions (upper right) where co-deformation is limited and individual layers can be discerned. In contrast, the N001M geometry has profuse co-deformation (Fig. 3a). It is quite interesting to note the marked influence of interfacial orientation on the fracture behavior of these nanolayered metal–intermetallic composites. Samples
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ples may be a result of interfaces that offer little resistance to dislocation transmission and little subsequent work hardening. Furthermore, the marked increase in co-deformation for N001M samples may be a result from interfaces that transmit slip more easily than those in N111M samples. Although experimental results for Niy Ni3Al multilayers have been discussed here, initial observations for g–Ni(Al)yg9–Ni3 Al multilayers are similar. 4. Summary A detailed characterization of monolithic Ni-25at%Al thin films, sputter deposited at different substrate temperatures revealed the formation of novel metastable phases in the film deposited on an unheated substrate (;45 8C), a fine-grained recrystallized microstructure in the film deposited at 200 8C and a coarser grained microstructure consisting of the single L12 Ni3Al phase in the film deposited at 400 8C. Nanoindentation studies carried out on these monolithic films indicate a strong strain hardening tendency in the film deposited at ;45 8C and the highest average hardness (;7.3 GPa) for the film deposited at 200 8C. In addition to monolithic Ni-25Al films, the fracture behavior of sputter deposited NiyNi3Al multilayers has also been investigated. N111M Oriented multilayers primarily exhibited a brittle mode of failure while N001M oriented multilayers primarily exhibited a ductile mode of failure. Furthermore, multilayers with thinner individual layers were found to be more ductile as compared to those with thicker individual layers. Fig. 3. SEM micrographs of the different NiyNi3 Al multilayers, viewed normal to the fracture surface. (a) N001M Oriented 20 nm Niy20 nm Ni3Al multilayer. (b) N001M Oriented 120 nm Niy120 nm Ni3Al multilayer. (c) N111M oriented 20 nm Niy20 nm Ni3Al multilayer. (d) N111M oriented 120 nm Niy120 nm Ni3Al multilayer.
with N001M interface orientation exhibit more ductile fracture features than N111M samples. Corresponding modeling of confined layer slip in multilayers shows that multilayers display remarkable strength when interfaces are able to confine dislocations to individual layers w7x. In such cases, profuse co-deformation of layers is suppressed and more fracture surface features are expected. Corresponding transmission electron microscope studies of in situ deformation in CuyNi multilayers indicate that misfit dislocations serve as pinning sites for glissile dislocations that attempt to transmit across interfaces. An implication is that semi-coherent interfaces are expected to be stronger barriers to dislocation transmission than coherent interfaces. Thus, samples with 20-nm layer thickness and corresponding coherent interfaces are expected to be less resistant to slip transmission than the thicker, 120 nm, counterparts. The marked shear localization displayed by the 20-nm sam-
Acknowledgments The authors would like to acknowledge the support from the Air Force Office of Scientific Research under Grant F49620-96-1-0238 and the Center for the Accelerated Maturation of Materials at the Ohio State University for this research. The authors would also like to thank Dr Douglas Smith and Dr Tim Foecke from the National Institute of Standards and Technology for their assistance in carrying out the nanoindentation studies. References w1x C.T. Simms, W. Hagel (Eds.), The Superalloys, Wiley, New York, 1972. w2x S.R. Harris, D.H. Pearson, C.M. Garland, B. Fultz, J. Mater. Res. 6 (1991) 2019. w3x T. Kizuka, N. Mitarai, N. Tanaka, J. Mater. Sci. 29 (21) (1994) 5599. w4x D.E. Alexander, G.S. Was, L.E. Rehn, J. Appl. Phys. 69 (4) (1991) 2021. w5x J.E. Eridon, G.S. Was, L.E. Rehn, J. Mater. Res. 3 (4) (1988) 626.
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R. Banerjee et al. / Thin Solid Films 424 (2003) 93–98 w6x B.M. Clemmons, H. Kung, S.A. Barnett, MRS Bull. 24 (2) (1999) 20–26. w7x P.M. Anderson, T. Foecke, P.M. Hazzledine, MRS Bull. 24 (2) (1999) 27. w8x G.B. Thompson, M.S. Thesis, The Ohio State University (1997).
w9x R. Banerjee, G.B. Thompson, H.L. Fraser, submitted to Philos. Mag. Lett. (2001). w10x G.B. Thompson, R. Banerjee, X.D. Zhang, P.M. Anderson, H.L. Fraser, submitted to Acta Mater. (2001). w11x R. Banerjee, J.P. Fain, P.M. Anderson, H.L. Fraser, to appear in Scripta Mater. (2001).