Strain Effects on the Crystal Growth and Superconducting Properties

Article pubs.acs.org/crystal

Strain Effects on the Crystal Growth and Superconducting Properties of Epitaxial Niobium Ultrathin Films C. Clavero,*,† D. B. Beringer,‡ W. M. Roach,† J. R. Skuza,‡,# K. C. Wong,§,∥ A. D. Batchelor,§,∥ C. E. Reece,⊥ and R. A. Lukaszew†,‡ †

Department of Applied Science and ‡Department of Physics, The College of William & Mary, Williamsburg, Virginia 23187, United States § Department of Materials Science and Engineering and ∥Analytical Instrumentation Facility, NC State University, Raleigh, North Carolina 27695, United States ⊥ Thomas Jefferson National Accelerator Facility, Newport News, Virginia 23606, United States ABSTRACT: Superconducting ultrathin films grown epitaxially onto crystalline substrates exhibit strained epitaxial growth due to lattice mismatch, which can have a significant effect on their superconducting properties. We present a complete correlation of the surface morphology, crystal growth, strain, microstructure, and superconducting properties in single-crystal Nb(110) thin films sputter deposited on a-plane sapphire substrates. Notably, we observe that the lattice mismatch between Nb and sapphire induces the formation of a hexagonal surface structure during the first three atomic layers. This is followed by a strained bcc Nb(110) phase whose in-plane lattice parameter progressively relaxes to bulk value. Similar lattice relaxation was also observed in the direction perpendicular to the interface using X-ray diffraction (XRD) and transmission electron microscopy (TEM). Significant perpendicular strain in films up to 30 nm thick was found to ultimately affect the superconducting properties of the Nb thin films as demonstrated with AC susceptibility measurements, where dissipative effects in the lattice associated with the presence of strain and associated defects were identified.

1. INTRODUCTION Superconducting thin films and multilayers have attracted the attention of the scientific community due to their geometryspecific properties as compared to bulk materials.1−4 Multiple devices have benefited from the unprecedented properties of superconducting thin films such as the Josephson effect5 used in superconducting quantum interference devices (SQUID),6 superconducting single electron transistors,7 and even fundamental elements for quantum computation8 among others. In particular, devices such as stripline detectors for the detection of ions9 and superconducting single-photon detectors10 strongly benefit from the use of ultrathin superconducting thin films. Multiple interesting physical phenomena have been found in superconducting ultrathin films and multilayers such as reentrant superconductivity11 or superconducting proximity effect.12 In addition, current superconducting radiofrequency (SRF) cavities are fabricated using bulk Nb, which limits the maximum acceleration fields achievable due to magnetic field penetration in the superconducting material. Superconducting thin films and multilayers are considered as possible coatings for the interior surface of suitable metallic SRF cavities to shield magnetic fields and thus to push further the accelerator limits.13 Nevertheless, special attention needs to be devoted to the microstructure of such superconducting thin films to achieve © 2012 American Chemical Society

the desired properties. Specifically, multiple aspects such as crystalline quality, film strain, grain size, or the presence of defects need to be considered, all of which are predominantly determined at the nucleation stage. All of these factors can drastically influence the superconducting properties of the films such as transition temperature, critical field, and AC susceptibility. For the case of Nb, epitaxial thin films can be achieved on a broad range of ceramic as well as metallic materials14 such as MgO,15 GaAs, InAs, Ta, Mo, Co, and Cu among others. A good match between the lattices of Nb and sapphire has been established for a number of different substrate orientations, with a unique three-dimensional relationship between the substrate and the overlayer.14 This means that Nb has the same orientation relative to the sapphire substrate for all of the known deposition planes, a property that was described as three-dimensional registry by Claassen et al.16 Growth on sapphire has been largely investigated14,17,18 since it is a prototypical system for metal on ceramic nucleation studies and in addition produces Nb films with high crystalline quality.14 Indeed, high quality films have been obtained on aplane sapphire,19 where Nb grows (110) with its Nb[10̅ 0] Received: February 7, 2012 Revised: March 6, 2012 Published: March 22, 2012 2588

dx.doi.org/10.1021/cg3001834 | Cryst. Growth Des. 2012, 12, 2588−2593

Crystal Growth & Design

Article

3. RESULTS AND DISCUSSION The surface morphology of superconducting thin films is particularly important in SRF applications since it can affect their surface impedance23 and hence their SRF performance. In addition, the surface morphology can also offer further clues regarding possible multiple in-plane lattice orientations in heteroepitaxial growth. AFM topography images (Figure 1)

direction parallel to Al2O3[0001] leading to a lattice mismatch of 10.7% along the Nb[001] direction and of 8.3% along the perpendicular Nb[110] direction. Such a mismatch between substrate and thin film can be accommodated by strain in the epitaxial layer, resulting in a coherent interface, or by the formation of misfit dislocations leading to semicoherent interfaces, resulting in coherent regions separated by misfit dislocations.14,20 Following the Matthews and Blakeslee formalism21 for misfit dislocations, Grier et al.20 predicted a critical thickness of 7.2 nm for Nb, above which misfit dislocations are favorable. They reported the presence of misfit dislocations in Nb(110)∥ a-sapphire films for thicknesses larger than 8 nm using transmission electrom microscopy (TEM). Nevertheless, the strain evolution during the early stages of growth, where coherent interface strain relaxation is expected, remains largely unexplored. Oderno et al.22 showed a transition from an initial hexagonal surface structure for thicknesses below 1.5 nm to bcc Nb(110) for thicker films; nevertheless, they did not report whether or not there is further evolution of the lattice parameters with thickness corresponding to coherent strain relaxation, which we addressed here. The study of Nb thin films nucleation and growth is of paramount importance since the presence of additional phases, coherent relaxation of the strain, misfit dislocations, or other epitaxial growth defects can strongly affect the superconducting properties of the entire system. Here, we present a complete study correlating superconducting properties of Nb thin films epitaxially deposited on a-plane sapphire with their morphology and microstructure. The surface morphology was investigated using atomic force microscopy (AFM), while the structure of the films was investigated using in situ reflection high-energy electron diffraction (RHEED) and ex situ X-ray diffraction (XRD) and TEM. Finally, the superconducting properties of the films were investigated using DC and AC susceptibility measurements.

2. EXPERIMENTAL DETAILS

Figure 1. AFM topography images for the (a) 30, (b) 100, and (c) 600 nm Nb films grown on a-plane sapphire. The insets on the upper right corner show 2D FFT maps showing the evolution of the anisotropy. (d) Profiles for the above-mentioned samples clearly showing the roughening of the surface as the film thickness increases.

Nb thin films were prepared by DC magnetron sputtering deposition in an ultrahigh vacuum (UHV) system with a base pressure in the low 10−10 Torr range. For this study, a-plane sapphire substrates with miscut angles lower than 0.05° were used to minimize the effect of atomic terraces on the growth. The substrates were ultrasonically cleaned and subsequently annealed at 600 °C for 1 h in UHV conditions prior growth, to degas and improve the crystallinity of the surface. Sputtering deposition was carried out from a high-purity (99.95%) Nb target at 5 × 10−3 Torr Ar pressure, resulting in a growth rate of 0.35 Å/s. To improve the deposited film homogeneity, the substrates were azimuthally rotated around the normal axis at a constant speed of 12 rpm during growth. Nb films of thicknesses up to 600 nm were deposited at a substrate temperature of 600 °C, which was determined to favor crystalline ordering.14 In situ RHEED analysis was performed to verify and monitor the crystalline structure of substrate and films during growth. XRD studies were performed ex situ using a standard four-circle diffractometer with Cu Kα radiation to further investigate the crystallinity of the films and the different phases present. High-resolution TEM data were collected with a CCD camera using a cold field emission gun TEM operated at 200 kV. The surface morphology of the samples was characterized ex situ using AFM in dynamic mode using Si N type cantilevers with a resonance frequency of 322.9 kHz and a tip radius around 10 nm. DC and AC susceptibility measurements were carried out with a SQUID that allows measurements at temperatures ranging between 3 and 300 K.

clearly show the evolution of the morphology of Nb films deposited on a-plane sapphire with thicknesses ranging from 30 to 600 nm. Elongated surface features along the two principal directions of the Nb(110) surface, that is, Nb[001] and Nb[110], are observed for thinner films (Figure 1a,b). Twodimensional fast Fourier transform (FFT) maps24 are shown on the upper right corner of the images, depicting the characteristic structure of this 2-fold anisotropy. For thicker films, only one of the anisotropy directions prevails, and grooves along such directions are observed (Figure 1c). A somewhat similar morphology was observed by Zhou et al.19,25 for MBE-grown Nb(110) films on a-sapphire. In their case, fingered structures were found along the Nb[110] direction, and they attributed the observed surface structure to an energy balance where substrate miscut, deposition temperature, faceting, pinholes, and lattice mismatch play fundamental roles. This indicates that different intrinsic growth factors can affect the microstructure and the surface morphology of the films. In addition to the observed anisotropy, it is also worth noticing that in our case the roughness increases linearly with Nb thickness as shown in Figure 1d. The profiles shown along random directions of the samples evolve from 0.5 nm root-mean-square (rms) roughness for a 30 nm thick film to 4.5 nm rms roughness for a 600 nm 2589



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thick film, thus clearly showing that no “saturation” of the roughness occurs for the thicknesses investigated here. The microstructure of the films was investigated in situ with RHEED at various stages during growth and ex situ with XRD and TEM after growth. RHEED is especially sensitive to the surface atomic lattice due to the limited penetration length of electrons in metallic films.26 Thus, it is suitable for probing the atomic spacing of the surface along different in-plane crystallographic directions; therefore, it is very sensitive to the presence of in-plane strain. RHEED patterns were taken for Nb films grown on a-sapphire at 600 °C for thicknesses ranging from 1 to 63 atomic layers (0.23−14.5 nm). The RHEED measurements were performed by interrupting the growth and pumping out the Ar sputtering gas after each particular thickness investigated here. Initially, two different streaks patterns that repeat every 60° are observed when the electron beam is directed along the in-plane [0001] and [11̅ 00] sapphire directions, that is, separated by 30° on the surface. In addition, the interline distances along those two directions have a √3 ratio, indicative of a hexagonal surface structure as previously observed by Oderno et al.22 The lattice parameter for such hexagonal structure was extracted from the pattern found along the Nb[112̅0] direction parallel to the Al2O3[0001] direction, as shown in Figure 2. Because of the hexagonal nature of the

Nb[11̅ 1] direction also parallel to the Al2O3[0001] direction, as shown in Figure 2. Because of its bcc nature, the atomic plane spacing obtained from RHEED was multiplied by 2/√3 to obtain the tetragonal lattice parameter “a”. It is worth noticing that a smooth evolution of the streaks spacing is observed during the transition, which corresponds to a smooth and progressive change in the atomic spacing. This precludes the onset of a strong concentration of defects in the interface between the two phases. Nevertheless, the change in structure from hexagonal to bcc and thus the change of the observed crystallographic direction give rise to an abrupt change in lattice parameter. After the transition, an initial 3.93% expansion of the lattice parameter as compared to bulk is observed for films 5 AL thick due to strain. Such strain is progressively relaxed getting close to the Nb bulk lattice parameter (0.33 nm) after approximately 14 atomic layers (3.22 nm). Our RHEED analysis clearly reveals that the Nb films grown on a-sapphire evolve via two main mechanisms to overcome the initial lattice mismatch: an initial hexagonal phase energetically more favorable for thicknesses up to 3 AL, followed by a strained bcc Nb(110) phase, which progressively relaxes reaching values close to equilibrium after 14 Nb atomic layers. The structure and strain of the Nb films were also investigated ex situ using XRD. It is worth noting that the Xray beam probes the entire Nb film thickness in the range investigated, and thus, the observed reflections are an average of the contributions arising from all of the atomic layers forming the film. Figure 3 shows symmetric XRD scans for 30, 100, and

Figure 2. Evolution of the Nb structure and lattice parameter for films with thicknesses ranging from 1 to 63 atomic layers (0.23 nm/AL). An initial hexagonal structure is observed for the first three atomic layers followed by a strained bcc phase, whose lattice parameter relaxes after 14 atomic layers (3.22 nm).

Figure 3. Symmetric XRD scans for 30, 100, and 600 nm thick Nb films grown on a-plane sapphire. Nb(110) phase is observed in all of them with a clear evolution of the lattice parameter toward bulk as the films get thicker.

lattice, the spacing between atomic planes extracted from the RHEED pattern corresponds to the “a” lattice parameter. We observe that the initial value of 0.294 nm for 1 Nb AL decreases and stabilizes at 0.289 nm for 2 and 3 AL. The hexagonal diffraction patterns are progressively replaced by a new set of streaks corresponding to bcc Nb(001) orientated Nb[1̅11]∥Al2O3[0001], similar in streaks spacing but lacking the hexagonal symmetry described above. The coexistence of both phases is found for Nb thickness ranging from 3 to 5 AL (0.69−1.15 nm). The lattice parameter for the new bcc structure was measured using the pattern found along the

600 nm thick Nb films. All of the films exhibited a single (110) phase confirming the structure observed with RHEED and their crystalline nature. Nevertheless, the Nb(110) reflection for the 30 nm thick film is clearly shifted from the expected bulk position due to the lattice mismatch present between substrate and thin film, and thus, the position of such Nb(110) reflection is closer to the Al2O3(112̅0) peak rather than to the corresponding bulk Nb(110) reflection, revealing an average “a” lattice parameter 1.25% larger than bulk (0.33 nm). On the other hand, the thicker films show a lattice parameter closer to bulk as can be observed for the 100 and 600 nm thick films, with only a 0.36 and 0.2% larger than bulk, respectively. For 2590



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frequency-dependent complex susceptibility can be expressed as χ(ω) = χ′(ω) + iχ″(ω), where the real part χ′(ω) describes the behavior of the system in phase with the incoming AC field and the imaginary part χ″(ω) accounts for the energy losses in the system. Figure 5 shows χ′ (left column) and χ″ (right

such thicker films, the contribution arising from the strained interface is less significant as compared to the subsequently relaxed layers, and thus, the average lattice parameter is very close to bulk. The width of the XRD peaks can be correlated to the X-ray coherence length and thus to the crystalline grain size and in this case exhibits a slight evolution from 30 nm grain size for the 30 nm thick film to 46 nm for the 600 nm thick film, evidencing the columnar growth that developed from Volmer− Weber growth mode during the nucleation stage. The structure close to the interface was further investigated using TEM. Figure 4 shows a focused ion beam (FIB) prepared

Figure 5. Left: Real χ′ and (right) imaginary χ″ parts of the susceptibility for 30, 100, and 600 nm Nb films grown on a-plane sapphire measured with an AC field of 3.5 Oe and 1 Hz superimposed to a continuous 100 Oe DC field. The two steps feature observed in the 30 nm Nb film is associated with the strain in the first atomic layers. The applied 100 Oe DC field causes a slight depression of the transition temperature.

Figure 4. TEM image of the Al2O3(112̅ 0)/Nb(110) interface. The out of the plane lattice parameter of the Nb film mimics that of the substrate across the entire thickness of the film (∼15 nm). FFT maps are shown for different areas of the film.

column) for the 30, 100, and 600 nm thick Nb films measured with an AC field of 3.5 Oe and 1 Hz superimposed to a continuous 100 Oe DC field, both along the plane of the Nb film. For the case of the 100 and 600 nm thick Nb film, a type II superconductor transition is observed at 8.75 and 8.7 K, respectively. It is worth noting that the presence of a 100 Oe DC applied to increase the AC susceptibility signal also decreases the measured transition temperature. A separate fieldfree measurement of Tc on the 600 nm film yielded 9.29 K in agreement with the bulk value and a residual resistance ratio (RRR) of 97, which is one of the largest values obtained for Nb ultrathin films grown on sapphire.17,29 The field used to probe the samples was chosen much smaller than the lower critical field Hc1 for bulk Nb, that is, 1700 Oe,23 to minimize its effects on the superconducting state of the films. This sharp drop of the transition temperature under a small applied field as compared to our field-free measurements hints at the presence of pinning sites correlated with the large surface roughness observed in these films. In addition, for the case of the thinner 30 nm thick Nb film, a χ′ susceptibility transition with two steps can be clearly observed, accompanied by two peaks in the χ″ susceptibility at 7.64 and 8.08 K. Even though the transition temperature is expected to decrease in thinner films, the observed reduction exceeds previous observations.30 The presence of two or more steps in the χ″ vs temperature dependence has been associated with the presence of grains and transport through grain boundaries in superconductors in previous studies.28,31 Such behavior has been explained as the consequence of the onset of intragrain and intergrain currents during the superconducting transition. Nevertheless, we note

cross-sectional TEM image of the Al2O3(112̅0)/Nb(110) interface. A very sharp interface is clearly observed with high crystalline ordering at both sides. A more detailed analysis of the TEM images using FFT24 reveals that the interplane distance in the direction perpendicular to the surface for Nb mimics that of the underlying Al2O3(112̅0) substrate (0.237 nm) across the entire thickness of the film (∼15 nm), similar to the observations by Grier et al.,20 not evidencing any relaxation across this range probably due to the lack of optimal resolution in our case. These results agree with the slow out-of-plane lattice parameter evolution observed with XRD, where the 30 nm thick Nb film was found strained by 1.25%. Once the morphology and structure of the Nb thin films have been completely described, we turned our attention to their superconducting properties and how they are affected by the observed strain near the interface. The susceptibility measurement performed using SQUID magnetometry and DC or AC magnetic fields is a well-established method to investigate the superconducting properties of metals and hightemperature superconductors.27,28 The exclusion of the magnetic field by superconductors below their critical temperature due to the Meissner effect gives rise to a remarkable drop in the susceptibility down to χ = −1 (in SI units) as the films become perfect diamagnets. Thus, the superconducting critical temperature, critical field, field shielding, and susceptibility are key magnitudes that can be used to characterize the films. When AC magnetic fields are applied to a sample, the 2591



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Table 1. In-Plane Lower Critical Magnetic Field Hc1∥ and Upper Critical Magnetic Field Hc2∥ Measured at 4.5 K for the 30, 100, and 600 nm Nb Films Deposited on a-Plane Sapphire, Along with the Bulk Values for Nb at Zero Temperature23

that for the Nb on a-plane sapphire system, we observed similar grain sizes for all of the films, ranging from 30 nm for the 30 nm thick Nb film to only 46 nm for the 600 nm film; thus, we do not expect variations in intergrain contact versus Nb thickness. We have shown with our RHEED and XRD analysis that the Nb films are strained during the early stages of growth. Thus, the observed response for the 30 nm thick Nb film can be attributed to the presence of two predominant phases in the sample: (i) a first one with poor superconducting properties due to strain in the lattice and positioned closer to the interface with the substrate and (ii) a second one corresponding to the relaxed Nb layers that exhibit close to bulk Nb behavior. A characteristic feature of superconducting thin films and multilayers is that they exhibit unusually high critical field values when a magnetic field is applied in the plane of the films.32 The critical field in thin films is known to be dependent on the thickness and also on nonlocal parameters such as the coherence length ξ and the penetration length λL1 of the material. The Ginzburg and Landau theory predicts a strong increase of the critical field in thin films, with extremely high critical fields at low thickness decaying exponentially to bulk value around 600 nm1 films. Thus, we investigated such dependence by measuring the magnetization M in the superconducting state as a function of the externally applied in-plane magnetic field in the present Nb films, as shown Figure 6. The films show the typical behavior for type II super-

Nb thickness

Hc1∥ (Oe)

Hc2∥ (Oe)

30 nm 100 nm 600 nm bulk

157 2000 1000 1700

5000 8250 5182 2400

than expected (Figure 5), and confirms the presence of a phase with poorer superconducting properties at the interface with the substrate due to the presence of strain.

4. CONCLUSIONS In summary, we have presented a complete correlation between morphology and structure with superconducting properties such as critical field, critical temperature, and complex susceptibility for epitaxial Nb(110) thin films sputter deposited on a-plane sapphire substrates. A 2-fold surface anisotropy is found for films up to approximately 100 nm thick turning into uniaxial anisotropy consisting of very elongated surface features in thicker films. RHEED and XRD characterization of the films demonstrated the presence of an initial hexagonal phase energetically more favorable for thicknesses up to 3 AL, followed by a strained bcc Nb(110) phase that relaxes progressively along the in-plane direction reaching values close to bulk after 14 Nb atomic layers. The strain relaxation was also investigated along the perpendicular direction using XRD and TEM. The superconducting properties of such an initial strained layer are found to be poorer than those in the remainder of the significantly relaxed film, exhibiting lower critical temperatures and fields. Thus, our studies provide new insight on the identification of dissipative effects associated with the presence of strain and related defects during the early stages of growth of epitaxial Nb films on a-plane sapphire.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Figure 6. In-plane magnetization vs in-pane applied magnetic field measured at 4.5 K for 30, 100, and 600 nm thick Nb films deposited on a-plane sapphire. The samples were zero-field cooled.

Present Address #

National Institute of Aerospace, Hampton, Virginia 23666, United States. Notes

conductors, where the magnetization initially increases linearly with the externally applied field up to the lower critical magnetic field Hc1∥ after which the film is in the vortex state causing deviation from the linear response as the applied field increases until the films lose completely their superconducting character at the upper critical field Hc2∥. Figure 6 shows such evolution applying the magnetic field parallel to the plane for the 30, 100, and 600 nm thick Nb films zero-field cooled and measured at 4.5 K. The corresponding in plane Hc1∥ and Hc2∥ field values are shown in Table 1. The 100 and 600 nm thick Nb samples follow the predicted behavior, since both critical fields Hc1∥ and Hc2∥ increase considerably when the thickness decreases to 100 nm and are higher than bulk values, that is, Hc1 = 1.7 kOe and Hc2 = 2.4 kOe.23 Nevertheless, the 30 nm Nb film deviates from such trend and exhibits much lower critical fields than expected, especially Hc1∥, which is only 157 Oe. This result also agrees with the observed critical temperature, lower

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS We acknowledge Stuart Wolf and Jiwei Lu for assistance in the XRD measurements, Diefeng Gu for assistance in the TEM sample preparation, and Anne Marie Valente-Feliciano for her valuable suggestions regarding the TEM measurements. This work was funded by the Defense Threat Reduction Agency (HDTRA1-10-1-0072) and the Department of Energy (DEAC05-06OR23177).

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REFERENCES

(1) Toxen, A. M. Phys. Rev. 1962, 127 (2), 382. (2) Chiang, T.-C. Science 2004, 306 (5703), 1900−1901. (3) Yazdani, A. Nat. Phys. 2006, 2 (3), 151−152. (4) Cohn, J. L.; Lin, J. J.; Lamelas, F. J.; He, H.; Clarke, R.; Uher, C. Phys. Rev. B 1988, 38 (4), 2326−2332.

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(5) Josephson, B. D. Phys. Lett. 1962, 1 (7), 251−253. (6) Jaklevic, R. C.; Lambe, J.; Silver, A. H.; Mercereau, J. E. Phys. Rev. Lett. 1964, 12 (7), 159−160. (7) Fulton, T. A.; Gammel, P. L.; Bishop, D. J.; Dunkleberger, L. N.; Dolan, G. J. Phys. Rev. Lett. 1989, 63 (12), 1307−1310. (8) Bouchiat, V.; Vion, D.; Joyez, P.; Esteve, D.; Devoret, M. H. Phys. Scr. 1998, 1998 (T76), 165. (9) Casaburi, A.; Zen, N.; Suzuki, K.; Ejrnaes, M.; Pagano, S.; Cristiano, R.; Ohkubo, M. Appl. Phys. Lett. 2009, 94 (21), 212502. (10) Hadfield, R. H.; Habif, J. L.; Schlafer, J.; Schwall, R. E.; Nam, S. W. Appl. Phys. Lett. 2006, 89 (24), 241129. (11) Zdravkov, V.; Sidorenko, A.; Obermeier, G.; Gsell, S.; Schreck, M.; Müller, C.; Horn, S.; Tidecks, R.; Tagirov, L. R. Phys. Rev. Lett. 2006, 97 (5), 057004. (12) Truscott, A. D.; Dynes, R. C.; Schneemeyer, L. F. Phys. Rev. Lett. 1999, 83 (5), 1014. (13) Gurevich, A. Appl. Phys. Lett. 2006, 88 (1), 012511. (14) Wildes, A. R.; Mayer, J.; Theis-Bröhl, K. Thin Solid Films 2001, 401 (1−2), 7−34. (15) Krishnan, M.; Valderrama, E.; Bures, B.; Wilson-Elliott, K.; Zhao, X.; Phillips, L.; Valente-Feliciano, A. M.; Spradlin, J.; Reece, C.; Seo, K. Supercond. Sci. Technol. 2011, 24 (11), 115002. (16) Claassen, J. H.; Wolf, S. A.; Qadri, S. B.; Jones, L. D. J. Cryst. Growth 1987, 81 (1−4), 557−561. (17) Wu, G.; Valente, A. M.; Phillips, H. L.; Wang, H.; Wu, A. T.; Renk, T. J.; Provencio, P. Thin Solid Films 2005, 489 (1−2), 56−62. (18) McMorrow, D. F.; Cowley, R. A.; Gibaud, A.; Ward, R. C. C.; Wells, M. R. Appl. Phys. Lett. 1993, 63 (16), 2195−2197. (19) Zhou, G. L.; Flynn, C. P. Phys. Rev. B 1999, 59 (12), 7860. (20) Grier, E. J.; Jenkins, M. L.; Petford-Long, A. K.; Ward, R. C. C.; Wells, M. R. Thin Solid Films 2000, 358 (1−2), 94−98. (21) Matthews, J. W.; Blakeslee, A. E. J. Cryst. Growth 1974, 27, 118− 125. (22) Oderno, V.; Dufour, C.; Dumesnil, K.; Mangin, A. M. P.; Marchal, G. Philos. Mag. Lett. 1998, 78 (5), 419−426. (23) Padamsee, H.; Knobloch, J.; Hays, T. RF Superconductivity for Accelerators; Wiley: New York, 1998. (24) Horcas, I.; Fernandez, R.; Gomez-Rodriguez, J. M.; Colchero, J.; Gomez-Herrero, J.; Baro, A. M. Rev. Sci. Instrum. 2007, 78 (1), 013705−013708. (25) Zhou, G. L.; et al. J. Phys.: Condens. Matter 1997, 9 (50), L671. (26) Mahan, J. E.; Geib, K. M.; Robinson, G. Y.; Long, R. G. J. Vac. Sci. Technol., A 1990, 8 (5), 3692−3700. (27) Mitsuhashi, R.; Suzuki, Y.; Yamanari, Y.; Mitamura, H.; Kambe, T.; Ikeda, N.; Okamoto, H.; Fujiwara, A.; Yamaji, M.; Kawasaki, N.; Maniwa, Y.; Kubozono, Y. Nature 2010, 464 (7285), 76−79. (28) Gömöry, F. Supercond. Sci. Technol. 1997, 10 (8), 523. (29) Zhao, X.; Phillips, L.; Reece, C. E.; Seo, K.; Krishnan, M.; Valderrama, E. J. Appl. Phys. 2011, 110 (3), 033523. (30) Gubin, A. I.; Ilin, K. S.; Vitusevich, S. A.; Siegel, M.; Klein, N. Phys. Rev. B 2005, 72 (6), 064503. (31) Singh, R.; Lal, R.; Upreti, U. C.; Suri, D. K.; Narlikar, A. V.; Awana, V. P. S.; Albino Aguiar, J.; Shahabuddin, M. Phys. Rev. B 1997, 55 (2), 1216. (32) Banerjee, I.; Yang, Q. S.; Falco, C. M.; Schuller, I. K. Phys. Rev. B 1983, 28 (9), 5037.

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Strain Effects on the Crystal Growth and Superconducting Properties

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