Investigation on the structure and magnetic properties of co2mnsi heusler alloy for spintronic application

Acknowledgements This author would like to express her heartfelt gratitude to her supervisors, Dr. Chen Jingsheng and Dr. Han Guchang for their guidance, advice, concern and encouragement throughout the course of the project. The author would like to thank Dr Qiu Jinjun for the introduction and maintenance of ULVAC sputter system, and Miss Luo Ping for the help in XPS analysis. In addition, the author is grateful for the friendship and support of Huang Lisen, Yang Yang, Ho Pin, Xu Dongbin, Si Huayan, Chen Chin, Li Huihui, Ko Viloane, Lim Boon Chow, and other research staff and students in the Department of Materials Science and Engineering and Data Storage Institute. Last but not least, the author would like to thank her family for their love, support and understanding during the period of research project. i Table of Contents Acknowledgements i Table of Contents ii Summary v List of Figures vii List of Symbols xi Chapter 1:Introduction 1 1.1 Gaint magnetoresistance (GMR) read head technology .................................... 3 1.1.1 Physics of GMR effect .............................................................................. 3 1.1.2 Current-in-plane (CIP) spin valve............................................................. 6 1.1.3 Current-perpendicular-to-plane (CPP) spin valve ................................... 8 1.2 Tunnel magnetoresistance (TMR) read head technology ................................ 10 1.3 Challenges for next generation read head over 1 Tbits/in 2 .............................. 11 1.4 Literature review of Co2MnSi half-metallic Heusler alloys ............................ 12 1.4.1 Basic properties of Heusler alloys .......................................................... 13 1.4.1.1 Origin of the bandgap.................................................................. 14 1.4.1.2 Slater-Pauling behavior ............................................................... 15 1.4.2 Parameters affecting MR ratio of CPP-GMR head using Co2 MnSi Heusler alloy ...................................................................................................... 17 1.4.2.1 Effect of lattice parameter........................................................... 17 1.4.2.2 Spin-orbit coupling ...................................................................... 19 1.4.2.3 Effect of temperature on spin polarization ................................ 19 1.4.2.4 Effect of disorder, defects, and doping ...................................... 20 1.4.2.5 Other factors from device perspective ....................................... 21 1.5 Objective of thesis .............................................................................................. 22 1.6 Outline of thesis.................................................................................................. 24 ii Chapter 2: Experimental methodologies 25 2.1 Sample fabrication by magnetron sputtering .................................................... 25 2.2 Alternating gradient force magnetometer ......................................................... 26 2.3 X-ray diffraction ................................................................................................. 27 2.3.1 θ - 2θ x-ray scans ..................................................................................... 29 2.3.2 Rocking curve .......................................................................................... 29 2.3.3 Phi-scans ................................................................................................... 30 2.4 Transmission electron microscopy .................................................................... 31 2.5 X-ray photoelectron spectroscopy..................................................................... 32 2.6 Atomic force microscopy .................................................................................. 33 Chapter 3: MgO buffer layer effect on the structure and magnetic properties of Co2MnSi (CMS) films on MgO substrates 34 3.1 Study of CMS thin films on MgO substrates (no buffer layer) ....................... 34 3.1.1 Experimental methods ............................................................................. 35 3.1.2 Results and discussion ............................................................................. 36 3.1.2.1 Structure properties ..................................................................... 36 3.1.2.2 Magnetic properties ..................................................................... 39 3.2 Study of CMS thin film on MgO-buffered MgO substrates ............................ 44 3.2.1 Experimental details ................................................................................ 44 3.2.2 Results and discussion ............................................................................. 45 3.2.2.1 Crystallographic structure ........................................................... 45 3.2.2.2 In-plane epitaxial growth relationship ....................................... 46 3.2.2.3 Heusler L21 texture...................................................................... 48 3.2.2.4 Microstructure ............................................................................. 49 3.2.2.5 Magnetic properties ..................................................................... 50 3.3 Summary ............................................................................................................. 54 iii Chapter 4: Cr buffer layer effect on the texture and magnetic properties of Co2MnSi (CMS) films 56 4.1 Experimental methods 56 4.2 Results and discussion ....................................................................................... 58 4.2.1 Effect of CMS in-situ annealing temperature ........................................ 58 4.2.1.1 Crystallographic structure ........................................................... 58 4.2.1.2 Cr/CMS interface and multi-layer roughness ............................ 63 4.2.1.3 Magnetic properties ..................................................................... 64 4.2.1.4 Microstructure ............................................................................. 68 4.2.2 Effects of Cr in-situ annealing temperature ........................................... 69 4.3 Summary ............................................................................................................. 72 Chapter 5: Conclusions and future work 74 5.1 Conclusions ........................................................................................................ 74 5.2 Future work......................................................................................................... 75 Publication 76 References 76 iv Summary The effects of MgO and Cr buffer layers on the structures and magnetic properties of Co2 MnSi (CMS) thin films have been studied. The CMS thin films were deposited on MgO (001) single crystal substrates with and without MgO or Cr buffer layer at room temperature by magnetron sputtering deposition and annealed at various temperatures. From 2-theta and Phi-XRD analysis, it was found that both MgO and Cr buffer layers could help induce the (001) epitaxial growth of the CMS thin films with B2 or L21 structure, while A2 structure was formed in the films without buffer layer. A four-fold 45ºshift between MgO (022) peaks and Cr (022) or CMS (022) peaks was obtained in the Phi-scan analysis of the post-annealed MgO- and Cr-buffered CMS thin films, which confirmed the epitaxial relationship of CMS [110](001) || MgO [100](001) and CMS [110](001) || Cr [110](001) || MgO [100](001). A smoother MgO buffer layer surface was obtained in an Ar+O2 atmosphere compared to that obtained in Ar atmosphere, which resulted in a stronger CMS (001) texture. Moreover, it was found that the in-situ annealing of the Cr buffer layer at 500°C destroyed the homogenous surface, while the RTdeposited Cr layer showed a very smooth surface and subsequently gave a better CMS crystallinity. The saturated magnetization of the CMS thin films increased v with annealing temperature below 600ºC and then decreased when annealed further at 600ºC. A chemical analysis of inter-diffusion was examined and the results indicated significant diffusion of Pt and Co in the MgO-buffered CMS thin films annealed at 600°C for 1h and Cr in the Cr-buffered CMS thin films annealed at 600°C for 15min. Large initial susceptibilities were found in the Cr buffered CMS thin films. For the MgO/CMS/Pt samples post-annealed at 600ºC, three-stage initial curves were obtained, suggesting a pinning behavior in the magnetic reversal mechanism. vi List of Figures Figure 1.1: Data storage roadmap of HDDs. Adapted from [1]. 1 Figure 1.2: Magnetoresistance magnitude of GMR versus TMR. Adapted from [5]. 3 Figure 1.3: Normalized resistance versus applied magnetic field at 4.2K with CIP current. Adapted from [19]. 4 Figure 1.4: Two current model of GMR effect with F/C multi-layers. 5 Figure 1.5: Schematic representation of an exchange biased spin valve structure, in which the magnetization of one ferromagnetic layer is fixed by exchange coupling to an antiferromagnetic layer. The other ferromagnetic layer has it magnetization direction changed by external magnetic field. 6 Figure 1.6: Different spin valve structures: (a) bottom type; (b) top type; (c) dual type, Cap: capping layer; FL: free layer; NM: non-magnetic layer; RL: reference layer; PL: pinned layer, AFM: anti-ferromagnetic layer; SL: the seedlayer. 7 Figure 1.7: The (a) CIP and (b) CPP GMR read head geometry. Adapted from [11]. 8 Figure 1.8: Schematic density of states for both magnetic electrodes with the parallel and anti-paralle arrangement of magnetizations. The conductivity of each spin channel is proportional to the spin DOS in the emitter and collector electrode. Adapted from [19]. 10 Figure 1.9: Schematic representation of the DOS for a half-metal with respect to normal metals and semiconductors [33]. 13 Figure 1.10: Structure of (a) C1 b half-Heusler alloys; (b) L21 full-Heusler alloys; (c) B2 full-Heusler alloy; (d) A2 full-Heusler alloy. Adapted from [32]. 13 Figure 1.11: Major combination of Heusler alloy formation. Adapted from [32]. 14 Figure 1.12: Origin of the minority band gap in (a) NiMnZ and (b) Co 2MnZ. Adapted from [32]. 15 vii Figure 1.13: Calculated total spin moments, in which dashed line represents the Slater-Pauling behavior. Adapted from [40]. 16 Figure 1.14: DOS for NiMnSb and CoMnSb . Adapted from [40]. 18 Figure 2.1: M-H loop of ferromagnetic materials. Adapted from [69]. 27 Figure 2.2: Illustration of x-ray diffraction. Adapted from [74]. 28 Figure 2.3: X-ray powder scans geometry. Adapted from [75]. 29 Figure 2.4: Phi-scan X-ray diffraction geometry. 31 Figure 3.1: (a) L21 structure of CMS; (b) top view of MgO/Co2MnSi crystal structure. 34 Figure 3.2: Schematic diagram of MgO(sub)/CMS/Pt layered structure. 36 Figure 3.3: AFM image of MgO substrate surface topography. 36 Figure 3.4: 2-theta XRD spectra of CMS films deposited at RT and post annealed at 400°C and 600°C. All peaks labeled in * are from MgO substrate. 37 Figure 3.5: In-plane hysteresis loops of CMS films: (a) deposited at room temperature; (b) post-annealed at 400°C; (c) post-annealed at 600°C; (d) extracted Ms and Hc curves as a function of temperature. 40 Figure 3.6: Line-scan EDX spectra of CMS thin films deposited on MgO substrates at room temperature and annealed at (a) 400°C and (b) 600°C. 41 Figure 3.7: The initial curves of both as-deposited CMS films and CMS films postannealed at 400°C and 600°C. (1), (2), (3) represents for initial stage, second stage, and third step, respectively. 42 Figure 3.8: Schematic of MgO(sub)/MgO/CMS/Pt multi-layer structure. 44 Figure 3.9: The 2-theta XRD spectra of CMS deposited at RT and post-annealed at various temperatures with 7nm MgO buffer layer deposited in (a) Ar+O2 and (b) Ar atmosphere. 45 viii Figure 3.10: The three dimension model of crystal structure of epitaxial grown MgO/CMS layers. 46 Figure 3.11: Phi-scan images of CMS {022} and MgO {022} planes in MgObuffered CMS thin films annealed at (a) 400°C and (b) 600°C. 47 Figure 3.12: Phi-scan spectra for CMS (111) peaks of MgO-buffered CMS thin films annealed at 400°C and 600°C. 48 Figure 3.13: Cross-sectional TEM images and IFFT image of the atomic lattice of MgO-buffered CMS with MgO deposited in (a) Ar; (b) Ar+O 2 gas atmosphere. 50 Figure 3.14: M-H loops of MgO-buffered CMS with MgO deposited in (a) Ar+O 2; (b) Ar gas atmosphere; extracted (c) Ms and (d) Hc curves as a function of temperature. 52 Figure 3.15: The XPS depth profile of MgO(7nm)-buffered CMS films: (a) deposited at room temperature; (b) annealed at 400°C; (c) annealed at 600°C; (d) initial magnetization curves. 53 Figure 4.1: (a) L21 structure of CMS; (b) Top view of MgO/Cr/CMS crystal structure. 56 Figure 4.2: Schematic of MgO(sub)/Cr/CMS/Pt layer structure. 57 Figure 4.3: The spectra of 2-theta XRD of Cr/CMS/Pt multi-layers. 58 Figure 4.4: (a) CMS (002) and (004) Peak-integration and (b) Rocking curve of CMS (002) peak as a function of temperature. 59 Figure 4.5: The 3-D crystal structure model of epitaxial grown MgO(sub)/Cr/CMS layers. 60 Figure 4.6: Phi-scans of MgO, Cr and CMS {022} planes of Cr-buffered CMS thin films annealed at (a) 300°C; (b) 400°C; (c) 500°C; (d) 600°C. 62 Figure 4.7: Phi-scan of CMS (111) peaks of Cr-buffered CMS films annealed at 600°C. 62 ix Figure 4.8: AFM images of as-deposited Cr buffer layer. 63 Figure 4.9: AFM images of Cr/CMS/Pt multi-layers: (a) deposited at room temperature; (b) annealed at 300°C; (c) annealed at 400°C; (d) annealed at 500°C; (e) annealed at 600°C. 64 Figure 4.10: (a) In-plane hysteresis loop of Cr-buffered CMS films deposited at room temperature and post-annealed at various temperatures; (b) extracted Ms and Hc curves as a function of temperature. 65 Figure 4.11: The initial curve of Cr-buffered CMS films. 66 Figure 4.12: XPS depth profile of Cr/CMS/Pt multi-layers in which CMS thin films were: (a) deposited at room temperature; (b) annealed at 300°C; (c) annealed at 500°C; (d) annealed at 600°C. 67 Figure 4.13: Cross-sectional TEM images of Cr/CMS/Pt layers in which CMS films were annealed at (a) 300ºC; (b) 400ºC; (c) 500ºC, (d) high resolution TEM image of Cr/CMS/Pt multi-layers in which CMS films were annealed at 500ºC. 69 Figure 4.14: 2-theta spectra of Cr/CMS/Pt multi-layers with various Cr in-situ annealing temperatures. 70 Figure 4.15: AFM images of Cr(10nm)-buffered MgO substrates with Cr in-situ annealing for 30min at (a) 300°C; (b) 400°C; (c) 500°C; (d) 600°C. 71 x List of Symbols  Wavelength Θ Angle theta ω Angle omega ψ Angle Phi Φ Angle Psi AC Alternating current AD Areal density AFM Atomic force microscopy AMR Anisotropic magnetoresistance AGFM Alternating gradient force magnetometry B Bit length BE Binding energy BF Bright-field CIP Current in plane CPP Current-perpendicular-to-plane d Spacing between pole tip of head and magnetic layer xi dhkl Interplanar spacing DF Dark-field DC Direct current DSI Data Storage Institute DOS Density of States EDX Energy-dispersive X-ray spectroscopy EF Fermi Energy FL Free layer fcc Face-centered cubic FM Ferromagnetic FWHM Full width at half maximum GMR Giant magnetoresistance Hc Coercivity HDD Hard disk drive HR High-resolution Hin Interlayer coupling IFFT Inverse Fast Fourier Transform xii KE Kinetic Energy MTJ Magnetic tunnel junction Mr Remanent magnetization Ms Saturation magnetization Mt Total moments per formula unit N Number of points NM Nonmagnetic PVD Physical Vapor Deposition PL Pinned layer △ RA Resistance Change per Unit Area RAMAC Random Access Method of Accounting and Control RF Radio frequency Ra Surface roughness average Rrms Mathematical Root Mean Square Roughness SV Spin valve SOC Spin-orbit coupling TMR Tunneling Magnetoresistance xiii T Temperature Tc Curie temperature TEM Transmission electron microscopy XRD X-ray diffraction XPS X-ray photoelectron spectroscopy Zt Total valence electrons number Zavg Average value of valence electrons number Zi Current value of valence electrons number xiv Chapter 1: Introduction Hard disk drives (HDD) have a leading position in the digitalized information storage area, accompanied by a dramatic increase in storage density at a rate larger than 40% per year. Fig. 1.1 shows the evolution and roadmap of HDD data storage [1]. The first magnetic HDD called the Random Access Method of Accounting and Control was introduced by IBM with an areal density (AD) of 2000 bits/in2 . To date, HDDs with AD of around 500 Gbits/in2 are widely used. Recently, HDD with AD of around 803 Gbits/in2 has been demonstrated in the lab [2]. Figure 1.1: Data storage roadmap of HDDs. Adapted from [1]. 1 The fast developing speed in HDD industry is closely related to the development in the field of spintronics in read head technology. The first breakthrough came in 1991 when the inductive read head was replaced by the anisotropic magnetoresistive (AMR) read head. Magnetoresistance (MR) refers to the resistance variation, between maximum and minimum resistances, normalized against the minimum resistance value when the magnetic field changed. The readout signal is proportional to the MR ratio. Although the MR ratio of AMR head was only 1%, it was more than double the value demonstrated by inductive read head. The second major development of the read head technology came with the discovery of the giant magnetoresistance (GMR) effect by Fert’s group [3] in France as well as Grunberg’s group [4] in Germany in 1988. In 1997, the first GMR read head was introduced in commercial HDDs with AD of around 2 Gbits/in2. However, this GMR read head with current-in-plane (CIP) geometry possessed some key limitations. As such, both current-perpendicular-to-plane (CPP)GMR and tunneling magnetoresistance (TMR) read heads were investigated as promising candidates to replace this CIP-GMR read head. Fig. 1.2 shows the magnetoresistance of GMR versus TMR [5]. The output signal of MgO and Al2O3 based TMRs indicated a larger magnetoresistance than the CPP-GMR read head. In 2005, Seagate introduced the first TMR read head with both MgO and Al2O3 insulators. 2 Figure 1.2: Magnetoresistance magnitude of GMR versus TMR. Adapted from [5]. 1.1 Gaint magnetoresistance (GMR) read head technology 1.1.1 Physics of GMR effect The giant magnetoresistance (GMR) effect was discovered in iron (magnetic) and chromium (non-magnetic) alternating layers. In these structures, when the magnetization directions in neighbouring iron layers changed from antiparallel to parallel, a significant change in resistance was observed. This phenomenon was known as GMR. For samples with 9 Å Cr layers, the MR (MR = [ρAP –ρP]/ ρP) ratio increased four-fold from 20% at room temperature to 80% at 4.2 K, as shown in Fig. 1.3. 3 Figure 1.3: Normalized resistance versus applied magnetic field at 4.2K with CIP current. Adapted from [19]. In these multilayer structure, the anti-parallel arrangement of magnetization between adjacent ferromagnetic (FM) layers resulted in high resistivity while a parallel arrangement of the magnetization led to an obvious reduction of resistivity magnitude. However, the saturation field Hs required to overcome the antiferromagnetic interlayer coupling effect and align the magnetization of consecutive layers was too large for GMR to be applicable in real devices. On the other hand, further investigation revealed that the anti-ferromagnetic coupling arrangement was not a prerequisite for the GMR effect [6]. The physical origin of GMR effect can be explained by the effect of the electron spin on the electronic transport in ferromagnetic conductors, i.e. spin-dependent scattering. A qualitative understanding of GMR effect was given by Mott’s two-current model [7], as shown in Fig. 1.4. 4 The key point of this model is that two independent conduction carriers – spin up and spin down electrons existed in FM conductors. Figure 1.4: Two current model of GMR effect with F/C multi-layers. As the resistance of the multi-layer structure arose from the scattering processes of spin electrons, strong scattering led to short mean free path, while weak scattering led to long mean free path. When the magnetization directions of two ferromagnetic layers were parallel, the spinup electrons, assumed to be parallel to magnetization, passed through the multi-layers with almost zero scattering. On the other hand, the spin-down electrons were scattered strongly as their spin aligned anti-parallel to the magnetization direction. Thus, parallel configuration resulted in low resistivity. If the magnetizations in the two ferromagnetic layers were anti- 5 parallel, high resistivity would be induced as both spin-down and spin-up electrons would be scattered strongly. 1.1.2 Current-in-plane (CIP) spin valve Figure 1.5: Schematic representation of an exchange biased spin valve structure, in which the magnetization of one ferromagnetic layer is fixed by exchange coupling to an antiferromagnetic layer. The other ferromagnetic layer has it magnetization direction changed by external magnetic field. With the discovery of the GMR effect in the CIP case, S.S. Parkin’s group proposed a CIP spin valve (SV) structure, as shown in Fig. 1.5. This SV structure, composed of four layers, had a much smaller switching magnetic field (Hs) which make the GMR effect practical for spintronic devices [8]. The magnetization of the pinned layer (PL) would be pinned along the anti-ferromagnetic (AFM) layer field cooling direction, while the magnetization of the free layer (FL) would change with the external field. The pinning effect was not only the result of unidirectional anisotropy generated by domains but also that of interface exchange coupling. 6 Figure 1.6: Different spin valve structures: (a) bottom type; (b) top type; (c) dual type, Cap: capping layer; FL: free layer; NM: non-magnetic layer; RL: reference layer; PL: pinned layer, AFM: anti-ferromagnetic layer; SL: the seedlayer. In a read head sensor, it would be crucial to reduce the interaction influence of the pinned layer on the free layer, due to interlayer and magnetostatic coupling H in. On one hand, Hin can be reduced by increasing the NM layer thickness. However, this would reduce the MR signal through scatterings which decreased the flow of polarized conduction electrons [9]. On the other hand, a synthetic anti-ferromagnetically coupled pinned layer [10] can be used, as shown above in Fig. 1.6 (a). This spin valve structure with pinned layer deposited first (bottom type) would reduce or even cancel off the stray field generated from the pinned layer. Fig. 1.6 (b) gives a reversed strucuture (top type) where the free layer was deposited first. The dual spin valve in Fig. 1.6 (c) shows a combination of both top and bottom types. Although this structure gave higher 7 MR ratio, it required more spacing between shields in devices. Therefore, this dual spin valve was limited for ultrahigh density recording. 1.1.3 Current-perpendicular-to-plane (CPP) spin valve Since the discovery of the GMR effect, much works had been done with the CIP spin valve structure. It was relatively easy to measure the resistance of thin films with CIP geometry, while the resistance measurement of CPP spin valves required complicated nano-scale device fabrication process as the sensor length was quite small. However, CIP-GMR had two main limitations, as shown in Fig. 1.7 (a). (a) CIP (b) CPP Figure 1.7: The (a) CIP and (b) CPP GMR read head geometry. Adapted from [11]. 8 (1) Two insulator gaps existed between the sensor and the two shields. A smaller gap would promote higher linear density. As such, the ideal condition was to eliminate the gap. However, this would result in a short circuit between the shield and sensor. (2) The reduction of track width could result in a conflict between higher track density and linearly decreased output signal. Unlike the CIP geometry, the two gaps were not necessary in a CPP spin valve head as the current flowed from the top shield to the bottom shield through sensor stack. In addition, both theoretical [12] and experimental results [13] revealed that the CPP-GMR showed higher intrinsic MR ratio compared to CIP-GMR. The physics involved in the MR measurement of CIP and CPP geometry was significantly different. In a theoretical paper by T. Valet and A. Fert in 1993 [14], they showed that the most important difference between the CPP and CIP was induced by the spin transport process. For CPP, the spin transportation was perpendicular to film interface, which included a spin accumulation effect. This effect allowed the spin transportation in CPP-GMR to be dependent on the long spin diffusion length rather than the short mean free path in CIP geometry. Moreover, a spin dependent interface resistance introduced by specular reflection can also be found in CPP spin valve. The famous Fert-Valet model, which mainly focused on the effect of bulk and interface spin-asymmetry coefficients, was then widely used in the investigation of CPP-GMR devices. 9 Unfortunately, the MR ratio of CPP-GMR still fell below the value required to achieve sufficient bit error rate for areal density exceeding 300 Gbits/in2 [15-18] after several years of investigation. The current used read head is based on CPP-TMR. 1.2 Tunnel magnetoresistance (TMR) read head technology When the metallic NM layer in a tri-layer GMR spin valve is replaced by a thin insulator, the mechanism for MR becomes spin dependent tunneling. This phenomenon is called tunneling magneto-resistance (TMR) and the junction is called magnetic tunnel junction (MTJ). Figure 1.8: Schematic density of states for both magnetic electrodes with the parallel and antiparalle arrangement of magnetizations. The conductivity of each spin channel is proportional to the spin DOS in the emitter and collector electrode. Adapted from [19]. The tunneling process was not only dependent on the available electronic channels in FM electrodes like in GMR, but also dependent on thickness and height of barrier. Julliere reported the first work of TMR measurement in a Co/Oxided-Ge/Fe MTJ in 1975 [20]. In his article, 10 Julliere proposed an explanation of spin polarized tunneling effect. Assuming spin conservation in the tunneling process, the conductance can be understood by the sum of two independent channels, as shown in Fig. 1.8 [19]. The famous Julliere formula relates the relative change of conductance with the density of states of each spin channel: TMR= Di  Di 2 P1 P2 , where Pi = Di  Di 1  P1 P2 D is the DOS of the FM electrodes for spin-up and spin-down direction at Fermi level. Researches on TMR have been very active since 1995 and major breakthroughs were made in 2004 at both Tsukuba (Yuasa et al) and IBM (Parkin et al). It was found that very large TMR ratio (200% at room temperature) could be obtained from MTJ with high quality MgO barrier. TMR ratio of about 600% was reported later [21]. However, the major challenge of TMR head was its high resistance, which limited the working frequency and thus reading speed. Reducing the insulating layer thickness also introduced pin-holes into the structure and deteriorated its performance. 1.3 Challenges for next generation read head over 1 Tbits/in2 As discussed above, TMR-based sensors with higher output signal have replaced CIPGMR sensors in HDD read heads. However, there is increasing interest to replace TMR-heads 11 with CPP-GMR heads for 1 Tbit/in2 read heads. This is mainly due to the low resistance area product in the all-metallic layer structure and the lower capacitance of CPP spin valves, enabling higher data transfer rates. However, the major drawback of the current CPP-GMR device is the low MR signal at room temperature. 1.4 Literature review of Co2MnSi half-metallic Heusler alloys In recent years, half-metallic ferromagnetic materials have attracted much attention due to possible applications in the field of spin-electronics. The existence of these materials was predicted using ab-initio calculations by de Groot et al. 1983 [22]. As shown in Fig. 1.9, spin up and down electrons in the band structure of these materials showed completely different behaviors. Half-metals can essentially be treated as hybrids of metals and semiconductors. As the minority spin band showed semiconductor-like behaviour with a gap at Fermi level, these materials exhibited a 100% polarization at Fermi level. Half-metallic ferromagnets can thus be expected to maximize the efficiency of spin-electronic devices, giving high MR ratios in CPPGMR read heads. Many materials have been predicted to be half-metallic by ab-initio calculations, such as transition metal chalcogenides (e.g. CrSe) and pnictides (e.g. CrAs) [23-26], oxides CrO2 and Fe3 O4 [27], europium chalcogenides (e.g. EuS) [28], double perovskites (e.g. Sr2 FeReO 6) [29], 12 and other kinds of materials [30,31]. However, amongst all these materials, the so-called half and full-Heusler alloys have attracted much more interest due to their high Curie temperatures [32] which is a requirement for practical devices. The basic properties and research progress on Heusler alloys will be discussed below. Figure1.9: Schematic representation of the DOS for a half-metal with respect to normal metals and semiconductors [33]. 1.4.1 Basic properties of Heusler alloys Figure 1.10: Structure of (a) C1b half-Heusler alloys; (b) L2 1 full-Heusler alloys; (c) B2 full-Heusler alloy; (d) A2 full-Heusler alloy. Adapted from [32]. 13 The Heusler alloys can be characterized into half-Heusler alloy (XYZ) in C1 b structure and full-Heusler alloy (X2YZ) in L21 structure, as shown in Fig. 1.10 (a) and (b), respectively. X and Y atoms represent transition metals, while Z is either a non-magnetic metal or a semiconductor, as shown in Fig. 1.11 [22, 32]. The unit cell consists of four interpenetrating face centered cubic (fcc) sublattices, in which the C1 b structure is shaped by removing one of the X sites in L21 structure. In addition, Y-Z atomic disorder in L21 structure of full-Heusler alloy will result in the formation of B2 structure, while A2 structure will form when X-Y and X-Z disorder occur. Figure1.11: Major combination of Heusler alloy formation. Adapted from [32]. 1.4.1.1 Origin of the bandgap According to the calculations by Galanakis et al. [34], the origin of the bandgap in Heusler alloys is caused by the d-d states hybridization of X and Y transition metals, as the DOS 14 in the vicinity of EF is dominated by the d-states. The formation of this gap in half-Heusler alloys (Fig. 1.12(a)) and full-Heusler alloys (Fig. 1.12(b)) is not exactly the same. In the case of halfHeusler alloys, the gap is formed by the hybridization states between elements X and Y directly, while in full-Heusler alloys, the hybridization between the elements X happened before the X-Y elements hybridization, in which only two bonding states among these four X–X orbitals eventually hybridized with the Y element. Figure1.12: Origin of the minority band gap in (a) NiMnZ and (b) Co2MnZ. Adapted from [32]. 1.4.1.2 Slater-Pauling behavior Galanakis et al. [34,35] reported analogous Slater-Pauling behaviour in the Heusler alloys with binary transition metal alloys, which is decribed as Mt = Zt - 18 (half Heusler) and Zt – 24 15 (full Heusler), where Mt represents total moments per formula unit, and Zt represents total valence electrons number. This behavior shown in Fig. 1.13 is a theoretical guide to achieve desired magnetic properties by substituting Y atoms with other transition metals in Heusler alloys. According to Fig. 1.11, there are more than 2000 possibilities to form Heusler alloys. However, there are only tens of alloys which have been reported based on this behavior. For example, a great improvement of T c to about 750 K was successfully made by the substitution of Cr with Fe atoms in Co2CrAl HMFs (T c around RT) with a new composition of Co2Cr0.6Fe0.4 Al [36-39]. Figure 1.13: Calculated total spin moments, in which dashed line represents the Slater-Pauling behavior. Adapted from [40]. 16 1.4.2 Parameters affecting MR ratio of CPP-GMR head using Co2MnSi Heusler alloy Amongst all the Heusler alloys, Co2MnSi, with the full-Heusler L21 structure (space group Fm 3 m) has attracted much attention as it was predicted to be a stable half-metal due to its large band-gap of 0.4 eV [41] in the minority spin band and high Curie temperature of 712°C [42]. Polarization of around 60% of both bulk and thin film full-Heusler alloy Co2 MnSi has been obtained by point contact andreev reflection spectroscopy measurements [43-47]. Unfortunately, experimental attempts on CPP-GMR devices using Heusler alloy by Yakushiji et al. achieved a MR ratio of only 2.4% at RT in a pseudo spin valve structure of Co2 MnSi/Cr/Co2 MnSi [17]. This section serves to give a brief review on the parameters which affect spin polarization in fullHeusler alloys. 1.4.2.1 Effect of lattice parameter Magnetic and electronic properties of both half-Heusler in C1b structure and full-Heusler in L21 structure are dependent on the magnitude of the lattice parameter. Density of States (DOS) in both NiMnSb and CoMnSb with a lattice parameter change of ±2% was calculated [40], as shown in Fig. 1.14. Although the half-metallic property was conserved in both expansion and compression scenarios, shifts of Fermi level occurred. Based on their calculations, the shift of EF 17 was attributed to the larger extension of p states compared to the d states of Sb. Movement of Fermi level towards conduction band took place in compression, while Fermi level moved towards the valence band during expansion. In addition, the dominant strong hybridization between Mn d and Ni or Co d states led to a slight increase in the size of the gap during compression. Similar behaviour is expected of other Co2 MnZ alloy compounds where the halfmetallicity is preserved with lattice constant change of ±2% [34]. In addition, a lattice compression of 4% was reported to lead to a strong increase in band gap energy of 23% [48]. Figure 1.14: DOS for NiMnSb and CoMnSb . Adapted from [40]. 18 1.4.2.2 Spin-orbit coupling Spin-orbit coupling (SOC) was neglected during the calculations of half-metallicity mentioned above. Taking into account the SOC, the electron spin would no longer be a good quantum number. As a result, the electron eigenfunction would not conserve the spin degree, even at 0 K. However, DOS within the gap is expected to be less in materials which have weak SOC effect and its polarization is close to 100% [49, 50]. The Heusler alloys like Co2MnSi, Co2 MnGe and Co2MnSn showed small orbit moments based on the calculations of Galanakis et al. [48, 51]. 1.4.2.3 Effect of temperature on spin polarization Several groups have investigated the temperature effect on the polarization of Heusler alloys qualitatively and quantitatively based on different assumptions and theories, such as tightbinding model, constrained density-function approach, dynamical mean-field way, and doubleexchange theory [52-57]. For example, based on the constrained density-functional theory, studies on NiMnSb showed that the minority spin-bands shifted cross the Fermi level gradually as temperature increased and finally a collapse in polarization occurred at around 0.4 Tc (RT). These calculations were consistent with experimental results of MR ratio loss as temperature changes from 4.2 K to RT. 19 1.4.2.4 Effect of disorder, defects, and doping As we introduced in Fig. 1.10, Heusler alloys can form B2 and A2 structures when X-Y and X-Z/Y-Z disordering occur in L21 structure at temperatures below the melting point. This would happen during the deposition of Heusler alloy thin films in the fabrication process of CPPGMR devices. Investigations showed that some Heusler alloys retained their half-metallicity in B2 structure, while A2 disorder degraded the spin polarization significantly [58, 59]. Picozzi et al. [58] investigated the formation of defects in full-Heusler alloy, in particular Co2 MnGe and Co2 MnSi. They found that the Mn antisites had the lowest formation energy and did not destroy the half-metallicity in contrast to Co antisites. In addition, large formation energies of the Mn-Si and Mn-Co atomic swaps were found. However, these results cannot be generalized to all Heusler alloys. Recently, the investigation of Nd doping effect on the transport and magnetic properties of CMS had been reported by K. Hono et al. [60]. From the resistivity measurements at low temperatures, it was concluded that electron-magnon scattering was suppressed in Nd-rich CMS phase. This was based on the understanding that small density of states near the Fermi level in the spin down mode was related to the mixing of spin up and down DOS caused by inelastic electron-magnon scattering [61]. There were many other factors contributing to the loss of spin polarization of CMS. For example, Suk J. Kim reported that fcc Co precipitated together with Co2MnSi at annealing temperature of 600°C, indicating a meta-stable phase of Co2MnSi. 20 1.4.2.5 Other factors from device perspective According to Valet-Fert model, the use of highly spin-polarized metallic Heusler alloys as FM electrodes of CPP-GMR devices was a feasible way to enhance the MR ratio by increasing the bulk spin asymmetry. However, in a multi-layered CPP-GMR device structure, many other factors may take effect on the output of CPP-GMR devices. Firstly, low resistivity and large spin-diffusion length were required for the space layer to obtain large CPP-GMR values. Low resistivity was also desirable for the buffer layer, which served as the bottom electrode for the measurement of CPP-GMR, to further decrease the total resistance of the CPP structure. A relatively high MR ratio of 14% at 6 K was obtained using Ag buffer layer with Co2 FeAl0.5Si0.5 full-Heusler alloys electrodes [62]. Secondly, the interlayer diffusion between the FM and NM layer caused by high deposition or annealing temperature of full-Heusler alloys had to be reduced in the CPP-GMR devices, as the inter-diffusion can result in the formation of magnetically dead layers. Hence, a compromise must be made between the crystal quality and possible interfacial conditions [63]. Thirdly, the interfacial spin-dependent scattering played an important role in the MR ratio of CPP-GMR devices. Ambrose and Mryasov [33] proposed a selection criterion for maximizing the interface spin asymmetry by changing the ferromagnetic metal and non-magnetic space layers. They pointed out that both band matching for majority spin channel and mismatching for minority spin channel at EF played important roles. The good 21 band matching allowed the spin-up electrons to propagate across the interface. On the other hand, poor matching increased the scattering of spin-down electrons. Seagate introduced an all-Heusler alloy CPP-GMR spin valve using ferromagnetic Co 2MnGe and non-magnetic Rh2CuSn. Based on band structure calculations, the interface spin asymmetry of this structure would be maximized [64]. However, some degree of disorder caused a loss of polarization at E F and hence limited the MR ratio (6.7%) of this system. 1.5 Objective of thesis As discussed in the previous sections, half metallic alloys, Co2 MnSi (CMS), with the fullHeusler L21 structure (space group Fm 3 m), have attracted much attention in the field of spinelectronics due to its large band gap of 0.4eV in the minority spin band and high Curie temperature of 712°C. According to Valet-Fert model, the use of highly spin-polarized metallic Heusler alloys as FM electrodes of GMR devices enhanced the MR ratio mainly by increasing bulk spin asymmetry. However, there are two main issues which could degrade the performance of GMR devices using CMS as FM electrodes. One of the issues involved the loss of spin polarization caused by Co-Mn or Co-Si disorder, which is represented by A2-type structure. The other is surface roughness, which plays an important role in spin-electronics applications. Different buffer layers had been used to induce the (001) texture of CMS for further achieving 22 B2 or L21 structure. For instance, (001) texture of CMS on MgO buffered MgO substrate had been reported, with MgO buffer layer deposited by e-beam evaporation [65]. The main objective of this research work is to investigate the effects of different buffer layers on the structures and magnetic properties of CMS thin films. 1) Co2 MnSi full-Heusler alloy thin films were sputter-deposited onto MgO (001) single crystal substrates without buffer layer to study the structures and magnetic properties of CMS thin films. 2) MgO buffer layers were deposited in different gas atmospheres (Ar and a mixture of Ar and oxygen) on MgO substrate by magnetron sputtering. After which, the structures, magnetic properties and interfacial morphologies of these MgO buffered CMS thin films were studied. 3) Cr was used as the buffer layer. The structures and magnetic properties of the Cr-buffered CMS films were studied. Inter-diffusion of component elements as well as thin films roughness relative to CMS in-situ annealing temperature was examined. Cr buffer layer in-situ annealing effect on the Cr/CMS interface morphology was investigated. The work presented in this thesis is an original work on the structures and magnetic properties of Co2 MnSi (CMS) Heusler alloy thin films for spintronic application. It would 23 provide guidance on the understanding of structural and magnetic properties of CMS Heusler alloys and its further application into the field of spin-electronics. 1.6 Outline of thesis This thesis was organized into 5 chapters. Chapter 1 gave an introduction of HDDs read head technology development and basic principles involved in GMR and TMR read head, and summarized current status of full-Heusler alloy Co2 MnSi as ferromagnetic layer of CPP-GMR read head from both materials and device aspects. Chapter 2 gave the outline of the experimental techniques with regards to sample fabrication, characterization and their corresponding working principles. In chapter 3, the structures and magnetic properties of CMS thin films on MgO single crystal substrates and MgO-buffered MgO single crystal substrates were studied. In chapter 4, the effects of Cr buffer layer on the structural and magnetic properties of CMS films were investigated. Inter-diffusion of Cr element as well as Cr/CMS interface roughness were examined. Conclusion and Future work of the thesis was given in chapter 5. 24 Chapter 2: Experimental methodologies In this Chapter, we introduced the experimental techniques used in this thesis for sample fabrication and characterization. For sample fabrication, thin films (MgO, Co 2MnSi, Cr, Pt) were all deposited on MgO substrates in high vacuum Magnetron sputtering system. X-ray Diffraction (XRD) and Transmission electron microscopy (TEM) were used to analyze the structural properties of thin films. Alternating gradient force magnetometry (AGFM) was used for the measurement of magnetic properties. Atomic force microscopy (AFM) was used to study surface morphologies. Line-scan energy-dispersive X-ray spectroscopy (EDX) and X-Ray photoelectron spectroscopy (XPS) depth profile analysis were used to identify distribution of elements in thin films. 2.1 Sample fabrication by magnetron sputtering Sputtering technique is widely used in the magnetic recording industries for thin films deposition. There are two kinds of collision processes involved in the plasma, elastic and inelastic process: Elastic scattering: e + Ar → e + Ar Inelastic Ionization (conversely, recombination): e + Ar → 2e + Ar + Inelastic Excitation (conversely, relaxation): e + Ar → e + Ar* 25 where Ar* represents the excited state of Ar atom. Elastic collision involves the interchange of kinetic energy only; while inelastic collision involves exchange of internal energies. DC sputter deposition is suitable for the deposition of Co2 MnSi, Pt, Cr, but not MgO with non-conducting property. Initiation of plasma is difficult when applying DC voltage to an insulating MgO target. To avoid this problem, a high frequency alternating voltage is used in place of DC voltage. The RF voltages can be coupled capacitively through the insulating target to initiate the plasma. Detailed discussion can be found in text books on sputtering processes [66, 67]. 2.2 Alternating gradient force magnetometer Alternating gradient force magnetometer (AGFM) is commonly used for characterization of magnetic materials, such as hysteresis loop and initial magnetization curves. It has a high sensitivity of 10-8 emu and small sample size of 3×3 mm2. The working system used in this study was Model 2900 MicroMagTM system. The working principles are summarized below. The testing sample is attached on a fragile glass rod and mounted to a piezoelectric transducer which oscillates as the external magnetic field. The alternating field gradient can apply a force on the sample and this force is proportional to the magnetic moment [68]. AGFM can be used to measure the hysteresis loop including initial magnetization curve, as shown in Fig 26 2.1, in which M represents for magnetization per unit volume and H represents for the external magnetic field [69]. For the measurement of M-H curve, the external magnetic field is applied from Hsat and decreases to (–Hsat) in user defined steps, and then increases back to H sat. Therefore, the whole M-H loop shown in Fig. 2.1 could be measured. For the initial magnetization curve measurement, the testing sample is first AC-demagnetized to a field value of +2 Oe. Then, the testing sample is subjected to an increasing field to H sat. The initial magnetization curves can be used to investigate the magnetization reversal mechanism, as the initial slope reveals the coercivity mechanism [70]. Figure 2.1: M-H loop of ferromagnetic materials. Adapted from [69]. 2.3 X-ray diffraction The X-ray diffraction (XRD) is a technique for phase and texture analysis of crystalline solid thin films. X-ray diffraction measurements are based on the principle of Bragg’s Law. 27 According to Fig. 2.2, X-rays interact with the atoms in a crystal. When there is a phase shift of 2π or a multiple of 2π, constructive interference occurs, as expressed by Bragg's law [71], where, n is an integer, λ is the wavelength of incident wave, d is the spacing between the planes in the atomic lattice, and θ is the angle between the incident ray and the scattering planes. More details about XRD can be found in text books by B.E. Warren and B.D. Cullity [72, 73]. In this thesis, X-ray powder scans and rocking curves of samples were measured from X’Pert Philips diffractometer. Phi-scan analysis was carried out using Bruker D8 Discover High Resolution XRD system. Figure 2.2: Illustration of x-ray diffraction. Adapted from [74]. 28 2.3.1 θ - 2θ x-ray scans In θ - 2θ x-ray powder scans, the sample is fixed in the horizontal position, while the tube and detector moved at an angle of θ and 2θ degree, respectively. The scan geometry of sample, tube and detector is shown in Fig. 2.3. Planes parallel to the film surface would be detected, and intensity pattern of these planes as a function of 2θ would be plotted. The diffraction peaks were then indexed according to patterns from the Joint Committee on Powdered Diffraction Standard (JCPDS) data base as well as literatures. Figure 2.3: X-ray powder scans geometry. Adapted from [75]. 2.3.2 Rocking curve The X-ray rocking curve measurement, also called ω scan, is used to investigate the quality of film in a particular texture. This method is applicable to analyze the orientation 29 distribution of a specific set of planes. During the measurement, the detector is fixed at an angle 2 θ (texture peak position), while ω value is changed ±10°around θ. The full-width at the halfmaximum peak intensity value (FWHM) is a qualitative measurement of the texture of the film. Small FWHM indicates better texture quality. 2.3.3 Phi-scans In θ-2θ XRD measurement, planes parallel to surface would be measured. On the other hand, Phi-scan (ψ) is used to measure planes which are not parallel to the sample surface. The geometry of phi-scan settings is shown in Fig. 2.4. Before the measurement, the sample is rotated an angle of Psi (Φ) such that the plane (hkl) is parallel to the original sample surface position. Then, the sample is moved to the 2-theta position of this plane (hkl). During Phi-scan measurement, the sample is rotated along Phi-axis from -180°to 180°in X-Y plane. Four peaks of (hkl), (-khl), (-h-kl), (k-hl) can be tested if the crystal is of four-fold symmetry in a tetragonal crystal structure, while six peaks would show in the phi-scan pattern for a hexagonal crystal structure with six-fold symmetry. 30 Figure 2.4: Phi-scan X-ray diffraction geometry. 2.4 Transmission electron microscopy Transmission electron microscopy (TEM) is useful of characterization for dislocations and other crystallographic defects, as well as performing chemical analysis of layered microstructures. There are three main imaging modes in TEM analysis: bright-field (BF), darkfield (DF), and high-resolution (HR) TEM. BF-TEM mode shows the transmitted beam information, while DF-TEM mode gives diffracted beam information. HRTEM image includes both transmitted beam information and diffracted beam information and is able to provide atomic resolution [76]. In addition, the inverse fast fourier transform (IFFT) image is useful in calculating the spacing between parallel planes [77]. The Miller index (hkl) can then be extracted based on the relationship between the lattice constant and spacing values of planes. For MgO, 31 which is in a cubic system, the [hkl] are parallel to the normal of (hkl). The inter-planar spacing is given by: 1 (ℎ2 + 𝑘 2 + 𝑙2 ) = 𝑑2 𝑎2 where d is the inter-planar spacing, and a represents the lattice parameter of unit cell. Moreover, line-scan energy dispersive x-ray spectroscopy (EDX) in TEM system is useful in analyzing the element distribution along a specific line. In this study, cross-sectional TEM images, as well as line-scan EDX were studied using JEOL TEM system for sample characterization. Detailed description of TEM could be found elsewhere [76]. 2.5 X-ray photoelectron spectroscopy X-Ray Photoelectron Spectroscopy (XPS), also known as ESCA (Electron Spectroscopy for Chemical Analysis), is a surface analysis technique which is based on the photoelectric effect [78, 79]. The XPS instrument measures the kinetic energy (KE) of both photoelectron and Auger electron, where KE of the electron depends on this equation: KE = hv - BE where BE is Electron Binding Energy, KE is Electron Kinetic Energy. 32 In this thesis, XPS depth profile is used to test the chemical concentration changes as a function of sputter time along depth direction using a spot size of 2х2 mm. Ar sputtering is used to sputter the surface components. A constant sputtering rate is assumed. 2.6 Atomic force microscopy Atomic force microscopy (AFM) is a useful method for surface morphology analysis, with a high resolution of less than nanometer [80, 81]. In this study, tapping mode, also called Dynamic contact mode AFM, was used to measure the surface roughness of thin films. The cantilever moved up and down at the resonance frequency for detection of surface morphology. The value of surface roughness in tapping mode is extracted by software attached in the instrument. The term Ra refers to surface roughness average, while R rms refers to the mathematical root mean square which is an average of peaks and valleys of a materials surface profile, and expressed by the relation n 1 Rrm s   Z i 1  Z avg  2 i N where Zavg is the average value; Zi is the current value and N the number of points. Detailed information of AFM technique can be obtained elsewhere [82]. 33 Chapter 3: MgO buffer layer effect on the structure and magnetic properties of Co2MnSi (CMS) films on MgO substrates The typical crystal structure of L2 1 ordered full-Heusler alloy is shown in Fig. 3.1 (a), while the top view of MgO/CMS structure is presented in Fig. 3.1(b). The lattice mismatch δ between Co2 MnSi (lattice constant a = 0.5670 nm) and 45°in-plane rotation of MgO is ~ 5%. Thus, it is expected that MgO (001) texture could induce the growth of CMS (001) texture through the epitaxial relationship of CMS [110](001) || MgO [100](001). (a) (b) Figure 3.1: (a) L21 structure of CMS; (b) top view of MgO/Co2MnSi crystal structure. 3.1 Study of CMS thin films on MgO substrates (no buffer layer) The L21 or B2 ordered CMS films have been reported to be suitable candidates as the magnetic layer in CPP-GMR devices due to their high spin polarizations, in which the formation 34 of (001) texture and flatten surface of CMS are essential issues. In this section, MgO (001) single crystal substrate was used to induce the epitaxial growth of CMS (001) texture. 3.1.1 Experimental methods The CMS films were deposited on MgO (001) single crystal substrates in an ultrahigh vacuum chamber, with a base pressure better than 5 × 10−6 Pa, by magnetron sputtering. The composition of the CMS target was confirmed to be Co46 Mn32Si22 by XPS-depth profile measurement after pre-sputtering for 40 min to remove surface oxygen. Deposition rates of all materials were calibrated using the average thickness with deposition time of 20 mins. The MgO substrates were first pre-heated at 700oC for 2h to improve the surface condition by decomposing surface contaminations (Mg(OH) 2 and MgCO3) [83] and then cooled down to room temperature (RT). Following which, 50 nm thick CMS film was deposited at RT with Ar pressure of 0.056 Pa. A 4 nm thick Pt capping layer was deposited as a protective layer. Finally, the dual-layer CMS(50 nm)/Pt(4 nm) was post-annealed at high temperatures (400°C and 600°C) for 1 h in vacuum system. The schematic diagram of the MgO(sub)/CMS(50 nm)/Pt(4 nm) layered structure is shown in Fig 3.2. 35 Figure 3.2: Schematic diagram of MgO(sub)/CMS/Pt layered structure. 3.1.2 Results and discussion 3.1.2.1 Structure properties Figure 3.3: AFM image of MgO substrate surface topography. As a smooth interface is one of the key issues for the integration of CMS into spin electronic devices, the surface morphology of pre-heated MgO substrate (700ºC; 2 h) was investigated by AFM. From Fig 3.3, average surface roughness (Ra) of 0.2 nm was observed, 36 indicating a relatively smooth MgO surface topography for CMS thin film deposition. This was much smaller than what Yuya et al. achieved, in which 800°C heated MgO substrate (1 h) showed a roughness of 11.9 Å. MgO (002) Intensity (a.u) CMS (220) Pt (111) o 600 C annealed o 400 C annealed as-deposited MgO sub  30     40  50  60  70 2-theta (deg) Figure 3.4: 2-theta XRD spectra of CMS films deposited at RT and post annealed at 400°C and 600°C. All peaks labeled in * are from MgO substrate. θ - 2θ XRD spectrum of CMS/Pt films deposited at RT and post annealed at various temperatures is shown in Fig. 3.4. No CMS peaks were shown in the as-deposited films, compared to the spectra of the MgO (001) single crystal substrate. After annealing at 400°C for 1h, the CMS (220) peak appeared. The intensity of CMS (220) peak increased with the increase in annealing temperature to 600°C. As seen, the as-deposited CMS was of amorphous phase, while both 400°C and 600°C annealed films were of (110) texture. An additional peak at around 37 40.4°appeared in the 600°C annealed sample. To confirm the source of this peak, a similar sample without the Pt capping layer was prepared and analyzed by XRD. No peak at 40.4°was detected. Therefore, it can be deduced that this peak is due to Pt crystallization at high temperature, and is described as Pt (111). As mentioned in the literature review, the presence of A2-type structure of CMS can be confirmed by (220) peak, B2 (partially ordered) structure by the (002) and (220) diffraction peaks, while L21 (fully ordered) structure by the (111) diffraction peak [84]. Besides, A2 structure can degrade the polarization of CMS half-metallic films due to complete site disorder of Co-Mn and Co-Si. Therefore, B2 or L2 1 structure, rather than A2 structure, is desirable in this study. To study the crystal structure of post-annealed CMS thin films deposited on MgO substrates, Phi-scan XRD measurements were carried out to analyze the CMS (200) and (111) peaks. Based on the geometry of 2-theta XRD scans, it can be inferred that the CMS (220) plane was parallel to the film surface. As the angle between CMS (200) and (220) planes is 45°, according to the principles of Phi-scan introduced in section 2.3.3, the CMS (200) peak can be detected by setting Psi to be 45°and 2-theta to be 31.8°in phi-scan XRD analysis. However, no CMS (200) peak was detected from both 400°C and 600°C annealed samples. 38 The absence of (200) peak and the presence of (220) peak in the Phi- and 2-theta scans can confirmed the presence of A2-type structure in both CMS thin films annealed at 400°C and 600°C. However, A2 structure should be avoided in spin electronics applications, as it would cause an obvious loss in spin polarization value compared to L2 1 structure [58, 59]. Although CMS (001) film was expected due to little lattice mismatch of around 5%, the preferred growth of CMS (220) film could possibly be due to the lower energy required to form the close packed (220) plane. 3.1.2.2 Magnetic properties The hysteresis loops with magnetic field applied in the in-plane direction of the asdeposited and post-annealed CMS films on MgO single crystal substrate are shown in Fig. 3.5 (a) to (c). The saturation magnetization Ms and coercivity Hc, were then extracted from the hysteresis curves, assuming that the CMS thin film volume did not change during high temperature annealing process, as shown in Fig 3.5 (d). All the error bars in Fig. 3.5 (d) were added with the value of 10% of experimental data, in which some of the error bars were too small to be observed. 39 15 1000 (a) (b) 500 M (emu/cm ) 5 3 3 M (emu/cm ) 10 0 As-deposited -5 0 o 400 C annealed -500 -10 -200 0 200 -1000 -400 400 -200 0 H (Oe) 200 400 H (Oe) 1000 (c) 800 800 Ms 3 500 M (emu/cm3) 1000 (d) Ms (emu/cm ) 1000 0 o 600 C annealed 600 600 Hc 400 400 200 200 Hc (Oe) -15 -400 -500 0 0 0 -1000 -400 -200 0 200 400 200 o T ( C) 400 600 H (Oe) Figure 3.5: In-plane hysteresis loops of CMS films: (a) deposited at room temperature; (b) postannealed at 400°C; (c) post-annealed at 600°C; (d) extracted M s and Hc curves as a function of temperature. The as-deposited film showed a ferromagnetic hysteresis loop, with a rather low Ms of 10.69 emu/cm3 and a small Hc of 3.1 Oe at 300 K. This weak ferromagnetic loop in the amorphous CMS films is attributed to Co-rich clusters, as reported by V. A. Oksenenko et al. [85]. They also proposed that the crystallization of Heusler alloy films would result in a 40 significant increase in Ms and Hc, as well as the formation of a more homogeneous magnetic ordering. The Ms of the CMS films annealed at 400°C increased to a much higher value of 827 emu/cm3 (the value of bulk L2 1 is 1040 emu/cm3), as the crystal structure changed from amorphous to (110) texture. However, with annealing at 600°C, the Ms (666 emu/cm3) did not show further increment compared to the films annealed at 400°C, even though the intensity of the CMS (220) peak increased. Moreover, Hc increased from 39.5 Oe to 140 Oe when annealing temperature increased from 400°C to 600°C. In order to further understand the drop in Ms and increase in Hc after annealing at 600°C, both line-scan EDX and initial magnetization analysis were carried out. o 400 C annealed 0 10 20 30 D (nm) 40 50 (a) o 600 C annealed Si Co Pt Mn Mg Counts (a.u) Counts (a.u) Mg Co Mn Si Pt 0 10 20 30 40 50 D (nm) (b) Figure 3.6: Line-scan EDX spectra of CMS thin films deposited on MgO substrates at room temperature and annealed at (a) 400°C and (b) 600°C. 41 Fig 3.6 (a) and (b) show the line-scan EDX spectra of CMS thin films annealed at 400°C and 600°C, from which we carried out the qualitative analysis of elemental distribution. The linescan analysis in this study was carried out along the depth (thickness) direction of all multi layers from top surface. As shown in Fig 3.6 (b), counts of Co reached the maximum value at a depth of around 32 nm, which is much larger than that of 25 nm in Fig 3.6 (a). A similar trend was observed for Si. The difference in these depths may be due to the diffusion of Co and Si to the top surface layer, which contributed to the drop in Ms in CMS thin films annealed at 600°C. Normalized Magnetization 1.0 (3) 0.5 (2) As-deposited o 400 C annealed o 600 C annealed (1) 0.0 0 100 200 300 400 500 H (Oe) Figure 3.7: The initial curves of both as-deposited CMS films and CMS films post-annealed at 400°C and 600°C. (1), (2), (3) represents for initial stage, second stage, and third step, respectively. Coercivity Hc is defined as the magnetic field needed to reduce the magnetization from remanence magnetization (Mr) to zero. The mechanism of this magnetic reversal process is subjected to the mechanism during initial magnetization [70]. Here, qualitative analysis of the magnetic reversal mechanism was studied by initial magnetization curves to account for the 42 trend of Hc in CMS thin films. All samples were AC-demagnetized to initial states with values less than 1% of the saturation magnetizations first. The initial curves are shown in Fig. 3.7. In the as-deposited and 400°C annealed samples, relatively high initial susceptibilities were observed. According to D. Givord et al. [86], this kind of initial magnetization behavior revealed a free domain wall displacement mechanism. The initial curve of 600°C annealed CMS films showed much lower susceptibility, and can be divided into (1), (2), (3) three stages, as shown in Figure 3.7: (1) Initial stage, with weak low-field susceptibility; (2) Second stage, starting at some threshold value, with a sharp increase in magnetization; (3) Third step, magnetization started to saturate. This three-stage behavior in initial curve is the same with a pinning-based switching mechanism in hard magnetic materials [86, 87]. This pinning in 600°C annealed films may be caused by diffusion of Pt protection layer and further crystallization of Pt phase, which could similarly be understood by analysis of XPS depth profile with sputtering time less than 14min in Figure 3.15 (in which all the deposition conditions were kept similar) and 2theta scan analysis in Figure 3.4. The behaviour of the initial curves, explained for the low Hc value (less than 50 Oe) in CMS films deposited at room temperature and annealed at 400°C, while large Hc value (180 Oe) in CMS films annealed at 600°C. 43 3.2 Study of CMS thin film on MgO-buffered MgO substrates 3.2.1 Experimental details In order to induce the epitaxial growth of Co2MnSi (001) texture on MgO (001) single crystal substrates, 7nm of fresh MgO buffer layer was introduced on top of the MgO substrates. The deposition and post-annealing parameters were kept similar as described in section 4.1.1, except that the 7nm MgO buffer layer was deposited in a separate oxidation chamber. Two kinds of gas atmosphere (Ar and Ar+O2) were used in the deposition of MgO buffer layer. The thin films were subsequently transferred to another chamber for CMS deposition without breaking of vacuum atmosphere. The layered structure of MgO(sub)/MgO(7nm)/CMS(50nm)/Pt(4nm) films is as follows: Figure 3.8: Schematic of MgO(sub)/MgO/CMS/Pt multi-layer structure. 44 3.2.2 Results and discussion 3.2.2.1 Crystallographic structure 30000 o 600 C annealed 20000 o 400 C annealed CMS (004) CMS(004) Intensity (counts) 30000 (b) Ar Pt(111) CMS (002) (a) Ar+O2 Pt (111) CMS(002) Intensity (counts) 40000 600C annealed 20000 400C annealed as-deposited as-deposited 10000 10000 MgO sub 30 40 50 2-theta 60 70 20 MgO sub 30 40 50 60 70 2-theta Figure 3.9: The 2-theta XRD spectra of CMS deposited at RT and post-annealed at various temperatures with 7nm MgO buffer layer deposited in (a) Ar+O 2 and (b) Ar atmosphere. To study the effect of MgO buffer layer on the formation of CMS (001) texture, 2-theta XRD spectra of the MgO-buffered (Ar and Ar+O2) CMS films deposited at RT and postannealed at 400°C and 600°C are illustrated in Fig. 3.9 (a) and (b). No crystalline peak from CMS was observed in the as-deposited sample. When annealed at 400°C and 600°C, the (002) peak of CMS appeared. This indicated a successfully epitaxial growth of CMS on the MgO buffered substrate. Therefore, the addition of 7nm MgO buffer layer helped to induce the (001) texture of CMS, compared to the (110) texture obtained without MgO buffer layer. Much stronger CMS (002) and (004) peak intensities were observed in post-annealed CMS films with MgO buffer layer deposited in Ar+O2 gas atmosphere, compared to the case without oxygen gas. 45 3.2.2.2. In-plane epitaxial growth relationship Figure 3.10: The three dimension model of crystal structure of epitaxial grown MgO/CMS layers. The in-plane epitaxial relationship between cubic MgO and CMS was investigated by phi-scans measurements based on the MgO (022) and CMS (022) reflections. Fig. 3.10 shows the expected three dimensional epitaxial model of MgO buffered CMS structure, where the black lattice represents MgO films, while the blue one represents CMS films. Before the Phi-scan analysis, (002) surface alignments of MgO substrates in all the samples were carried out. Thereafter, the samples were positioned at Psi = 45°and 2-theta = 62.2°to detect the diffraction peaks from MgO {022} planes. The sample was then rotated along phi-axis from -180°to 180°, namely a 360°rotation in x-y plane shown in Fig. 3.10. As the cubic MgO substrate is of fourfold symmetry, it is expected that four peaks from MgO (022), (0-22), (202), (-202) will be detected in sequence during Phi-scan. Similarly, to detect the diffraction peaks of CMS {022} 46 planes, the samples were moved to position settings of Psi = 45°and 2-theta = 45.15°with phirotation from -180°to 180°. o MgO {022} CMS {022} 600 C(Ar+O2) MgO {022} CMS {022} o o 400 C(Ar+O2) MgO {022} CMS {022} Intensity (a.u) 45 45o Intensity (a.u) o 600 C (Ar) o 400 C (Ar) MgO {022} CMS {022} o 45 45o -150 -100 -50 0 50 100 Phi (Deg) (a) 150 -150 -100 -50 0 50 100 150 Phi (Deg) (b) Figure 3.11: Phi-scan images of CMS {022} and MgO {022} planes in MgO-buffered CMS thin films annealed at (a) 400°C and (b) 600°C. Fig. 3.11 (a) and (b) illustrates the phi-scan spectrums for the post-annealed CMS films with MgO buffer layer deposited in Ar+O 2 and Ar gas atmospheres, respectively. Four-fold MgO (022) peaks as well as CMS (022) peaks were obtained. In particular, all CMS (022) peaks displayed a shift of 45°from the MgO (022) peaks. As shown in Fig 3.10, the angle between the projections in x-y plane of normal vector of MgO (022) plane (represent by Y direction) and that of CMS (022) plane (represent by Y´direction) was 45°, which was consistent with the 45°shift of MgO (022) and CMS (022) in phi-scan analysis. This results showed that the expected epitaxial relationship of CMS [110](001) || MgO [100](001) was successfully achieved. 47 3.2.2.3. Heusler L21 texture o Intensity of CMS (111) peaks (a.u) 600 C (Ar) o 600 C (Ar+O2) o 400 C (Ar+O2) -150 -100 -50 0 50 100 150 Phi (Deg) Figure 3.12: Phi-scan spectra for CMS (111) peaks of MgO-buffered CMS thin films annealed at 400°C and 600°C. The CMS (111) texture of post-annealed MgO-buffered CMS films was investigated by Phi-scan. Samples were moved to a Psi degree of 54.73°and a 2-theta degree of 27.19°to detect the diffraction peaks from CMS {111} planes. As illustrated in Fig 3.12, four and three peaks from CMS {111} planes were clearly shown in the 600°C and 400°C annealed CMS with MgO buffer layer deposited in Ar+O2 gas, respectively. In contrast, (111) peaks were only observed in 600°C annealed CMS in the case of Ar gas environment. It can be deduced that the MgO buffer layer deposited with Ar would delay the formation of L2 1 Heusler structure, compared to the Ar+O2 gas environment. 48 3.2.2.4 Microstructure In order to study MgO buffer layer effect with different deposition gas (Ar+O 2; Ar), interface condition of MgO/CMS were examined by Cross-sectional TEM, as shown in Fig 3.13. Layered structure and epitaxial lattice of MgO buffer layer on MgO substrate were observed in both cases. Smooth MgO/CMS interface was observed in the case of MgO buffer layer deposited in Ar+O2 gas environment, while waved interface between MgO and CMS layer was found in Ar gas environment, with [001] direction of MgO buffer layer slightly deviated from the film normal. The non-uniform and rough morphology of MgO and CMS interface, could have resulted in the smaller peak intensities of CMS (002) and (004) with MgO buffer layer deposited in Ar gas atmosphere. Furthermore, the lattice constants of both MgO substrate and MgO buffer layer were calculated using the Inverse Fast Fourier Transform (IFFT) image of selected areas. The lattice constants (2.23 Å; 2.20 Å) of MgO buffer layer, for Ar+O2 and Ar case respectively, were less than that of MgO substrate (2.26 Å). This may be due to the difference in stoichiometry of the MgO buffer layers compared to the MgO substrate introduced in the deposition process. The smaller lattice constant of MgO buffer layer could have bought a relaxation of the larger lattice mismatch between CMS layer and MgO substrate, which helped to induce the CMS (001) texture. 49 Figure 3.13: Cross-sectional TEM images and IFFT image of the atomic lattice of MgO-buffered CMS with MgO deposited in (a) Ar; (b) Ar+O2 gas atmosphere. 3.2.2.5 Magnetic properties The in-plane hysteresis loops of as-deposited and post-annealed CMS films with MgO buffer layer deposited in Ar+O2 and Ar gas atmosphere were measured at 300K by AGFM, as shown in Fig. 3.14 (a) and (b). Fig. 3.14 (c) and (d) show the extracted Ms and Hc values as a function of temperature for the Ar+O2 and Ar gas atmosphere respectively. All the error bars in Fig. 3.14 were added with the value of 10% of experimental data, in which some of the error bars were too small to be observed. A comparison of magnetic properties with different MgO deposition atmosphere (Ar+O2; Ar) revealed that: (1) both as-deposited samples showed quite small Ms, indicating a weak ferromagnetic order, regardless of the MgO deposition atmosphere. 50 (2) The trend of Ms and Hc of CMS films remained the same, in spite of changes in the MgO deposition atmosphere. With annealing at 400°C, Ms increased to a high value of about 3.75 µB/f.u., consistent with the transformation of the CMS phase from amorphous to crystalline (001) texture. However, further annealing at 600°C resulted in a drop in Ms and a big increase in Hc. (3) In particular, multi-layers with MgO buffer layer deposited in Ar atmosphere showed lower M s and higher Hc in samples post-annealed at 400°C and 600°C compared to that in the presence of Ar+O2 atmosphere. This was consistent with the structural properties in terms of lower peak intensities of CMS (002) and (004), as well as the delayed appearance of L2 1 structure in Ardeposited-MgO buffered CMS films. (a) Ar+O2 0 M (emu/cm 3) 15 -500 10 5 0 As-deposited 500 As-deposited o 600 C annealed o 400 C annealed (b) Ar 0 15 M (emu/cm 3) 3 M (emu/cm ) 500 1000 As-deposited o 400 C annealed o 600 C annealed M (emu/cm3) 1000 -500 -5 -400 -200 0 -1000 0 H (Oe) 0 As-deposited -5 200 -400 -200 200 400 H (Oe) -200 5 -10 -10 -400 10 400 0 200 400 H (Oe) -1000 -400 -200 0 H (Oe) 200 400 51 250 800 Ar Ar+O2 200 HC (Oe) 3 MS (emu/cm ) Ar Ar+O2 150 400 100 50 (c) 0 (d) 0 0 0 400 400 o o T ( C) T ( C) Figure 3.14: M-H loops of MgO-buffered CMS with MgO deposited in (a) Ar+O2; (b) Ar gas atmosphere; extracted (c) Ms and (d) Hc curves as a function of temperature. To further understand the variation trend of Ms and Hc in (Ar+O2)-MgO buffered CMS films, XPS depth profile and initial curves analysis were carried out. In Fig 3.15 (a) to (c), the atomic concentration variations of component elements in the multi-layered films were shown by XPS depth profile analysis: 100 o (a) RT deposition 90 Atomic Concentration (%) Atomic Concentration (%) 100 80 70 Si2s.ls1 Mn2p3.ls1 Co2p.ls1 Pt4f.ls1 60 50 40 30 20 10 (b) 400 C annealed 90 80 70 O1s.ls2 Si2s.ls1 Mn2p3.ls1 Co2p.ls1 Pt4f.ls1 60 50 40 30 20 10 0 0 0 2 4 6 8 Sputter Time (min) 10 12 0 2 4 6 8 10 12 14 Sputter Time (min) 52 80 o Normalized Magnetization Atomic concentration (%) (c) 600 C annealed 70 60 50 O1s.ls1 Si2s.ls1 Mn2p3.ls1 Co2p.ls1 Pt4f.ls1 40 30 20 10 0 (3) 1.0 0.8 0.6 As-deposited o 400 C annealed o 600 C annealed (2) 0.4 0.2 0.0 (1) (d) Initial Curve Onset field -0.2 0 2 4 6 8 10 Sputter Time (min) 12 14 0 100 200 300 400 500 H (Oe) Figure 3.15: The XPS depth profile of MgO(7nm)-buffered CMS films: (a) deposited at room temperature; (b) annealed at 400°C; (c) annealed at 600°C; (d) initial magnetization curves. (1) Co represented by green line: In the as-deposited sample, slight diffusion of Co was observed. This may due to the gradient of concentration in layered thin film. After annealing at 400°C for 1h, slightly deeper diffusion was observed as the saturation concentration of Co was observed 1.5 min later than that of the as-deposited sample. However, after annealing at 600°C for 1h, significant diffusion was detected as the saturation of Co had not been reached even after 14 minutes of sputtering. (2) Mn represented by blue line: Obvious diffusion was seen only in the films annealed at 400°C, while diffusion in both as-deposited and 600°C annealed sample was negligible. (3) Si represented by red line: Only slight diffusion was monitored regardless of annealing temperature. (4) Pt represented by pink line: Slight diffusion was detected in both as-deposited and 400°C 53 annealed thin films. Significant diffusion of Pt was observed in 600°C annealed films as the maximum position of Pt concentration moved inwards, towards the substrate, resulting in surface oxidation of multi-layers. Compositions of CMS films can be estimated from the saturated concentration of Co, Mn and Si. By calibration of target composition Co46 Mn32Si22, the CMS thin films were found to have compositions of Co2 Mn0.99Si1.11, Co2 Mn0.70Si0.81, Co2 Mn0.96Si1.00 , in the as-deposited, 400°C annealed and 600°C annealed samples, respectively. Initial magnetization curves illustrated in Fig. 3.15 (d) indicated similar trend as the CMS films deposited directly onto MgO substrate, where large initial susceptibilities of both as-deposited and 400°C annealed CMS were obtained, while 600°C annealed films showed a three stage variation with a much lower initial susceptibility. Based on the results of the XPS depth profile and initial curves, the drop of Ms in 600°C annealed CMS films may be due to the diffusion of Co and Pt. The high threshold field in the 600°C annealed sample may be due to pinning effects caused by the diffusion of Pt protection layer and surface oxidation. 3.3 Summary Systematic study of MgO buffer layer effect on structural (the epitaxial growth of Co2 MnSi (001) texture and formation of L2 1 phase) and magnetic properties of CMS thin films was carried out. CMS films were deposited at room temperature (RT) and post-annealed at 54 various temperatures. Epitaxial growth of CMS (001) texture was successfully realized by introducing 7nm MgO buffer layer on MgO substrate. Desirable L21 structure of CMS was also formed when annealed at 600°C. Without a buffer layer on the MgO substrate, only the A2 structure can be obtained. MgO buffer layer deposited in Ar+O2 atmosphere helped to induce much higher peak intensities of CMS (002) and (004), and promoted the formation of L21 Heusler structure of CMS, compared to that deposited in Ar atmosphere. Both 600°C annealed CMS films with MgO buffer layer deposited in Ar+O2 and Ar atmosphere showed 4 (111) peaks, suggesting the L21 structure of CMS. Ms of CMS thin films showed significantly improvement as the CMS phase changed from amorphous to crystalline after post-annealing at 400°C. However, it decreased with further annealing at 600°C. In addition, Hc increased with increasing annealing temperature. The decrease in Ms and increase in Hc may be attributed to the diffusion of Pt protection layer. 55 Chapter 4: Cr buffer layer effect on the texture and magnetic properties of Co2MnSi (CMS) films In this chapter, Cr was introduced as the buffer layer on MgO substrates to promote epitaxial growth of (001) oriented CMS. Compared with the MgO buffer layer, Cr (lattice constant a=0.2884nm) has a much smaller lattice mismatch (1.7%) with Co2 MnSi. The lattice mismatch between a 45°in-plane rotation of Cr buffer layer and MgO substrates is 3.2%. The expected in-plane epitaxial growth of Cr/CMS film on MgO single crystal substrates is shown in Fig. 4.1. (a) (b) Figure 4.1: (a) L21 structure of CMS; (b) Top view of MgO/Cr/CMS crystal structure. 4.1 Experimental methods To investigate the effects of in-situ annealing temperatures of CMS and Cr on the properties of Cr-buffered CMS thin films, 2 series of samples were deposited on MgO (001) 56 single crystal substrates using magnetron sputtering, with the layer structure illustrated in Fig. 4.2. Figure 4.2: Schematic of MgO(sub)/Cr/CMS/Pt layer structure.  During the fabrication process, the MgO substrates were pre-heated at 700oC for 2h before deposition to improve the surface condition by decomposing surface contaminants (Mg(OH)2 and MgCO3) [83] and subsequently cooled to room temperature (RT). The composition of the CMS target used was Co46 Mn32 Si22 , same as previous work done on MgO buffer layers. A magnetron sputtering chamber with a base pressure better than 4 × 10−8 Pa was used. Cr buffer layers were deposited at a deposition rate of 0.69 Å· s−1 with Ar gas pressure of 0.112 Pa, while CMS layers were deposited at a deposition rate of 0.46 Å· s−1 with Ar gas pressure of 0.056 Pa. The Cr buffer layer thickness was fixed at 10 nm, while that of CMS layer was fixed at 30 nm. A 4 nm thick Pt capping layer was deposited at ambient temperature less than 80°C. 57  In series I: 10 nm of Cr buffer layer, followed by 30 nm of CMS, was deposited on MgO substrates at room temperature. Following which, in-situ annealing at different temperatures (300°C; 400°C; 500°C; 600°C) was carried out for 15 mins.  In series II: Cr buffer layer was first deposited at RT and in-situ annealed at different temperatures (300°C; 400°C; 500°C; 600°C) for 30min. After cooling to RT, the CMS layer was deposited and in-situ annealed at 600°C for 15min. 4.2 Results and discussion 4.2.1 Effect of CMS in-situ annealing temperature o 600 C CMS(004) CMS(002) Intensity (a.u) 4.2.1.1 Crystallographic structure o 500 C o 400 C o 300 C as-deposited Cr (002) MgO sub 20 30 40 50 60 70 2-theta (deg) Figure 4.3: The spectra of 2-theta XRD of Cr/CMS/Pt multi-layers. 58 Fig 4.3 shows the 2-theta XRD spectra of CMS films in series I. No peaks from the Pt capping layer were detected. For the as-deposited films, the 2-theta XRD results clearly show the Cr (002) peak which is emphasized by an arrow in Fig 4.3, suggesting the crystallized Cr buffer layer. After post-annealing, except for the peaks from the MgO substrates, Cr (002) peak and CMS (002) and (004) peaks were detected, indicating the formation of CMS (001) texture. Fig 4.4 (a) shows the calculated integration of CMS (002) and (004) peak intensities as a function of temperature. There was a gradual increase of integrated peak intensity with post-annealing temperature. The rocking curves of CMS (002) peaks at different annealing temperatures are shown in Fig 4.4 (b). It can be seen that the FWHM of rocking curves of CMS (002) peaks decreased slightly with increasing annealing temperature, and attained a value of less than 1° 300C FWHM =1.23 400C FWHM =1.05 500C FWHM =0.82 600C FWHM =0.79 (002) peak (004) peak Intensity (a.u) Peak Integration of CMS (a.u) when post annealing at 500°C and 600°C for 15 min. (b) (a) 300 400 o 500 Temperature ( C) 600 10 12 14 16  (deg) 18 20 22 Figure 4.4: (a) CMS (002) and (004) Peak-integration and (b) Rocking curve of CMS (002) peak as a function of temperature. 59 To investigate the in-plane epitaxial growth relationship and L21 Heusler structure of CMS films in series I, Phi-scan XRD measurements were carried out. Fig 4.5 illustrates the expected epitaxial growth structure of Cr/CMS films on MgO (001) single crystal substrates. As discussed in section 3.2.2.2, MgO (022) and CMS (022) peaks should have a four-fold 45°shift in Phi-scan spectra if the epitaxial growth structure drawn in Fig 4.5 is well formed. Figure 4.5: The 3-D crystal structure model of epitaxial grown MgO(sub)/Cr/CMS layers. Before the measurement, the alignment of MgO (002) surface was carried out. During the phi-scan (from -180° to 180°) analysis, samples were fixed at specific setting positions of Psi=45.6°and 2-theta=62.2°for measurement of MgO {022} planes, while Psi angle of 45.6° and 2-theta of 44.3°for Cr {022} planes, and Psi=45.6°and 2-theta =45.15°for CMS {022} 60 planes. Fig 4.6 shows the Phi-scan spectrums of multi-layers in series I. In 300°C in-situ annealed CMS thin films, only two sharp Cr (022) and CMS (022) peaks were detected, suggesting an incomplete crystal structure formation. When annealed at temperatures above 400°C, four Cr and CMS (022) peaks were obtained, revealing complete crystallization with four-fold symmetry. The 45° shift of Cr and CMS (022) with respect to MgO (022) peaks suggested the epitaxial growth of CMS [110](001) || Cr [110] (001) || MgO [100] (001), as shown in Fig 4.5. -60 0 60 -180 -120 180 (a) -60 0 60 CMS (022) 120 180 (b) Cr (022) o Cr (022) o Intensity of 300 C annealed sample(a.u) CMS (022) 120 Intensity of 400 C annealed sample(a.u) -180 -120 MgO (022) o 46 -180 -120 -60 0 60 Phi (deg) 120 180 MgO (022) o 45 -180 -120 -60 0 60 120 180 Phi (deg) 61 0 60 120 CMS (022) 180 -180 -120 (c) 0 60 120 180 (d) Cr (022) o Cr (022) -60 CMS (022) Intensity of 600 C annealed sample (a.u) -60 o Intensity of 500 C annealed sample (a.u) -180 -120 MgO (022) o 45 -180 -120 -60 0 60 120 180 MgO (022) o 45 -180 -120 Phi (deg) -60 0 60 120 180 Phi (deg) Figure 4.6: Phi-scans of MgO, Cr and CMS {022} planes of Cr-buffered CMS thin films annealed at (a) 300°C; (b) 400°C; (c) 500°C; (d)600°C. CMS L21 structure was also examined by phi-scan with setting of Psi = 54.74°and 2theta = 27.19°. CMS (111) peaks were only observed in the 600°C annealed films, with four-fold symmetry, as shown in Fig 4.7. Intensity(a.u) CMS (111) -180 -120 -60 0 60 120 180 Phi(deg) Figure 4.7: Phi-scan of CMS (111) peaks of Cr-buffered CMS films annealed at 600°C. 62 From the 2-theta and Phi-XRD results, the appearance of (002) peaks and the missing (111) peaks showed that these annealed (300°C; 400°C; 500°C) CMS films grown on Crbuffered MgO substrate had the B2 structure; while the existence of both (002) and (111) peaks indicated that 600°C annealed CMS films had the L2 1 structure, which was the Hesuler structure of CMS. 4.2.1.2 Cr/CMS interface and multi-layer roughness The roughness of the Cr/CMS interface plays an important role on the surface scattering, which will affect the output signal in CPP spin valve devices. Fig. 4.8 shows the roughness of Cr(10nm)-buffered MgO substrate with a Ra of 1.2 Å, which was much smaller than the results of as-deposited Cr (40 nm) buffer layer on MgO substrates (8 Å) reported by Yuya Sakuraba et al. [84]. Figure 4.8: AFM images of as-deposited Cr buffer layer. 63 Figure 4.9 shows the mean surface roughness of Cr/CMS/Pt multi-layers in series I. Ra values of all samples were less than 2 Å, suggesting that post-annealing did not affect the surface roughness of CMS, as reported by Yuya Sakuraba et al. [84]. Figure 4.9: AFM images of Cr/CMS/Pt multi-layers: (a) deposited at room temperature; (b) annealed at 300°C; (c) annealed at 400°C; (d) annealed at 500°C; (e) annealed at 600°C. 4.2.1.3 Magnetic properties The in-plane hysteresis loops of Cr/CMS/Pt multi-layers in series I at 300K are illustrated in Fig. 4.10 (a). Fig. 4.10 (b) summarizes the extracted Ms and Hc values as a function of annealing temperature. The amorphous phase of as-deposited CMS films showed an extremely small saturation magnetization (Ms). When annealed at 300°C, saturation magnetization of CMS 64 films increased to about 50 times that of the as-deposited value. The Ms continued to increase with higher annealing temperature, reaching the maximum value at 500°C. In addition, all the Hc values were all less than 4 Oe, and decreased slightly as annealing temperature increased. The increase in Ms and decrease in Hc of CMS thin films when in-situ annealed below 500°C is consistent with the crystal structure improvement obtained from the XRD results. However, when the annealing temperature was increased to 600°C, the Ms value dropped. As-deposited o 300 C annealed o 400 C annealed o 500 C annealed o 600 C annealed 1000 (a) 4.0 (b) 3.5 0 M (emu/cm 3) 20 -500 500 10 0 as-deposited Ms 3.0 Hc 2.5 2.0 Hc (Oe) 3 3 M (emu/cm ) 500 Ms (emu/cm ) 1000 -10 -20 -400 1.5 -200 0 200 H (Oe) -1000 -400 -200 0 H (Oe) 200 400 400 0 1.0 0 300 600 o T ( C) Figure 4.10: (a) In-plane hysteresis loop of Cr-buffered CMS films deposited at room temperature and post-annealed at various temperatures; (b) extracted Ms and Hc curves as a function of temperature. To investigate the magnetic reversal mechanism in the Cr-buffered CMS thin films with different CMS in-situ annealing temperature, initial magnetization curves were measured by AGFM, as shown in Fig 4.11. No pinning effect was obtained here, which was opposite to the results shown in Figure 3.15. This difference may be due to the lower deposition temperature of 65 Pt layer in the former (Figure 4.11), while the Pt layer was annealed together with CMS layer in the latter (Figure 3.15). All initial curves showed quite large initial susceptibilities and small saturation fields. This was consistent with the small Hc values (< 4 Oe) obtained from hysteresis Normalized Magnetization loops. 1 300C annealed 400C annealed 500C annealed 600C annealed 0 0 100 200 300 400 500 H (Oe) Figure 4.11: The initial curve of Cr-buffered CMS films. XPS depth profile was used to investigate the inter-diffusion in the multi-layers and the chemical composition of CMS thin films. Fig 4.12 (a) to (d) show the atomic concentrations of different elements as a function of sputter time from the top surface of Cr/CMS/Pt multi-layers. 100 100 o (b) 300 C annealed Atomic Consentration (%) Atomic Concentration (%) (a) RT deposition Si2s.ls1 Mn2p3.ls1 Co2p.ls1 Pt4f.ls1 50 0 Si2s.ls1 Mn2p3.ls1 Co2p.ls1 Pt4f.ls1 50 0 0 2 4 6 8 Sputter Time (min) 10 12 0 2 4 6 8 10 12 Sputter Time (min) 66 o Atomic Concentration (%) (c ) 500 C annealed Si2s.ls1 Mn2p3.ls1 Co2p.ls1 Pt4f.ls1 50 0 100 o (d) 600 C annealed Atomic Concentration (%) 100 Si2s.ls1 Cr2p.ls1 Mn2p3.ls1 Co2p.ls1 Pt4f.ls1 50 0 0 2 4 6 Sputter Time (min) 8 10 0 2 4 6 Sputter Time (min) 8 10 Figure 4.12: XPS depth profile of Cr/CMS/Pt multi-layers in which CMS thin films were: (a) deposited at room temperature; (b) annealed at 300°C; (c) annealed at 500°C; (d) annealed at 600°C. (1) In as-deposited films, Co, Mn, Si elements slightly diffused out to Pt capping layer due to concentration gradient, while Pt diffused towards CMS. The diffusion of Co, Mn, Si elements increased slightly as the annealing temperature increased. (2) Si element slightly diffused out to the surface when annealed at 300°C and 500°C, as the surface concentration of Si was much higher than that of as-deposited films. In the sample annealed at 600°C, Si diffused out further more heavily than that annealed at 500°C, as the saturation concentration appeared slightly later. This diffusion of Si may have contributed to the drop of Ms when annealing at 600°C. (3) In the CMS films annealed at 600°C, Cr appeared after 1 min of sputtering, indicating significant diffusion of Cr compared to the samples annealed at 500°C and below, where there was no sign of Cr for a sputter time of up to 10min. This significant diffusion could 67 result in the formation of magnetic dead layers which may be the key reason for the loss of ferromagnetism at 600°C. (4) From the atomic concentration curves, the composition of CMS films was estimated. After calibration with the Co46 Mn32Si22 target, the compositions of CMS films have been found to be Co2 Mn0.99 Si1.11, Co2 Mn1.01 Si1.10, Co2Mn1.46Si0.76, Co2Mn0.92Si1.02, in the as-deposited sample, sample post annealed at 300°C, 500°C and 600°C, respectively. 4.2.1.4 Microstructure Fig. 4.13 (a) to (c) show the cross-sectional bright field TEM images of Cr-buffered CMS with in-situ annealing temperatures from 300°C to 500°C. Layered structures of Cr/CMS/Pt in samples annealed at different temperatures (300°C; 400°C; 500°C) were all clearly observed. Fig. 4.13 (d) shows the high resolution (HR) TEM image of the 500°C annealed thin film. Epitaxial lattice growth of Cr/CMS on MgO substrate was obtained. No considerable lattice defects could be seen in the Cr and CMS layers, while misfit dislocations was seen in both Cr/CMS and MgO/Cr interfaces due to the small lattice mismatches in adjacent layers. 68 Figure 4.13: Cross-sectional TEM images of Cr/CMS/Pt layers in which CMS films were annealed at (a) 300ºC; (b) 400ºC; (c) 500ºC, (d) high resolution TEM image of Cr/CMS/Pt multi-layers in which CMS films were annealed at 500ºC. 4.2.2 Effects of Cr in-situ annealing temperature Fig 4.14 illustrates the 2-theta XRD spectra of Cr/CMS/Pt multi-layers with and without Cr in-situ annealing. It was found that Cr in-situ annealing resulted in a decrease in CMS (002) and (004) peak intensities, compared to the one without in-situ annealing. This suggested that the two-time annealing fabrication process gave weaker crystalline quality of CMS films in (001) 69 texture. In particular, the lowest intensities of CMS (002) and (004) peaks were observed in the sample with Cr in-situ annealing at 500°C. Figure 4.14: 2-theta spectra of Cr/CMS/Pt multi-layers with various Cr in-situ annealing temperatures. Surface morphologies of Cr(10nm)-buffered MgO substrates after in-situ annealing at various temperatures were investigated by AFM, as shown in Fig 4.15. When annealed at 300°C or 400°C for 30min, the Cr surface remained smooth, with a small surface roughness (R a) of around 1.2 Å. However, when annealed at 500°C for 30min, the Cr film surface became inhomogeneous, with a sudden increase in surface roughness to 16.6 Å. When annealed further at 600°C, although surface roughness increased to 42.3 Å, surface uniformity improved compared to the 500°C annealed sample. To summarize, Cr in-situ annealing at 300°C and 70 400°C for 30min preserved the smooth Cr buffer layer surface, while annealing at 500°C and 600°C increased the surface roughness dramatically. Combining the results from the 2-theta XRD spectrums and AFM results, it can be concluded that: the large decrease in the intensity of the CMS (004) peaks after in-situ annealing at 500°C was due to the inhomogeneous and rough surface morphology of Cr. Figure 4.15: AFM images of Cr(10nm)-buffered MgO substrates with Cr in-situ annealing for 30min at (a) 300°C; (b) 400°C; (c) 500°C; (d) 600°C. 71 4.3 Summary CMS thin films were deposited on Cr-buffered MgO substrate. The effects of in-situ annealing of CMS and Cr on the structural and magnetic properties of Cr-buffered CMS thin films were studied. Firstly, epitaxial growth of CMS (001) texture with 45°in-plane rotation with respect to MgO substrates was successfully realized after CMS in-situ annealing at high temperatures. Extremely small FWHM of CMS (002) peaks of less than 1°were achieved with CMS in-situ annealing at 500°C and 600°C. According to Phi-scan analysis, thin films with CMS in-situ annealing below 600°C had B2 structure, while L21 Heusler structure was observed in the 600°C annealed films. Epitaxial lattice of the MgO/Cr/CMS layers in the 500°C annealed sample was observed by cross-sectional HRTEM. No appreciable defects from lattices were detected, except for misfit dislocations in MgO/Cr and Cr/CMS interfaces. The saturation magnetization of CMS thin films increased with annealing temperatures below 500°C, consistent with the increase in CMS (002) and (004) peak intensities. However, Ms decreased when annealed temperature was raised to 600°C, which could be mainly due to the significant Cr diffusion, as well as Si diffusion revealed by the XPS depth profile analysis. AFM results showed a smooth surface of Cr-buffered MgO substrate, with a roughness of 1.2 Å, ensuring a smooth interface between the Cr and CMS layer. In addition, CMS in-situ annealing could conserve the smooth surface. 72 Secondly, the in-situ annealing process had important effects on the surface morphology of the Cr layer and subsequently on the structural properties of the CMS thin films. In-situ annealing of 10 nm thick Cr at 300°C and 400°C for 30min preserved the smooth surface of the Cr buffer layer, while annealing at 500°C and 600°C increased the surface roughness dramatically. In particular, the Cr layer in-situ annealed at 500°C showed the most inhomogeneous surface morphology, consistent with the lowest peak intensity of CMS (004) in the 2-theta XRD analysis. 73 Chapter 5: Conclusions and future work 5.1 Conclusions Epitaxial relationship of CMS [110](001) || MgO [100](001) and CMS [110](001) || Cr [110](001) || MgO [100](001) were successfully realized by the introduction of MgO and Cr buffer layers on MgO substrate. CMS thin films deposited directly on MgO substrate displayed the (110) texture. Both MgO-buffered and Cr buffered CMS films showed L21 Heusler structure when post-annealed at 600°C. Micro-structural studies showed that the MgO buffer layer helped to relax the lattice mismatch and a smooth buffer layer surface was obtained in an Ar+O 2 gas atmosphere. Cr-buffered CMS (30 nm) films showed much larger (004) peak intensity than MgObuffered CMS (50 nm) films after high temperature annealing, although the thickness of CMS films was much thinner in the former. This was most likely due to the smaller lattice mismatch between Cr and CMS. As the rationale for the study of Cr buffer layers was to find better deposition or annealing condition which promoted the growth of Heusler or B2 structure of CMS thin films, it was found that Cr in-situ annealing above 500°C destroyed the surface uniformity, while the as-deposited Cr layer had a smooth and uniform surface and gave better crystalline quality of CMS, suitable for spin-electronic applications. 74 Magnetic properties of CMS thin films showed a large increase in M s when annealed below 600ºC, consistent with the transformation of amorphous CMS to (110) or (001) texture. However, Ms decreased when further annealed at 600ºC. Chemical analysis (line-scan or XPSdepth profile) results revealed Pt and Co inter-diffusion in MgO-buffered CMS thin films and significant Cr inter-diffusion in Cr-buffered CMS thin film, which most probably contributed to the reduction in Ms. Large initial susceptibilities were found in Cr-buffered CMS thin films, consistent with the small Hc observed in M-H loops. However, for the MgO/CMS/Pt multi-layers post-annealed at 600ºC, three stage initial curves were obtained, suggesting a pinning behavior in the magnetic reversal mechanism. 5.2 Future work Based on this study, Co2MnSi (CMS) thin films with L2 1 or B2 structure had been successfully achieved by introducing MgO or Cr buffer layer on MgO substrates. Particularly, the FWHM of Cr-buffered CMS thin films annealed at 500 oC was less than 1º.It is expected that Cr-buffered CMS thin films after annealing at 500 oC could be applied into GMR read head devices to be the lower ferromagnetic layer to improve the MR ratio in the future. 75 Publication 1. X. Ji, J.J. Qiu, P. Luo, G.C. Han, J.S. Chen, “ Effect of buffer layer MgO and Cr on structure and magnetic properties of Co2 MnSi thin films” (Submitted to Journal of Alloys and Compounds). 76 References 1 Elizabeth A. Dobisz, Zvonomir Z. Bandic, Tsai-Wei Wu, Thomas Albrecht, Proceedings of the IEEE, 96, 1836 (2008). 2 http://www.dssc.ece.cmu.edu/news/story/2008/10/dssc_affiliate_tdk/ 3 Baibich, M.N., J.M. Broto, A. Fert, F. Nguyen Van Dau, F. Petroff, P. Etienne, G. Creuzet, A. Friederich, and J. Chazelas, Phys. Rev. Lett. 61. 2472 (1988). 4 Binasch G, Grunberg P, Saurenbach F and Zinn W, Phys. Rev. B. 39, 4828 (1989). 5 H.J.M. Swagten, Chapter 1: spin dependent tunneling in magnetic junctions, Hand book of Magnetic Materials, edited by K.H.J. Buschow. 6 C. Dupas, P. Beauvillain, C. Chappert, J.P. Renard, F. Triqui, P. Veillet, E. Velu, and D. Renard, J. Appl. Phys. B. 67, 5680 (1990). 7 Eiichi Hirota, Hirosi Sakakima, Koichiro Inomata, Giant Magnetoresistance Device. 8 B. Dieny, V.S. Speriosu, S.S.P. parkin, B.A. Gurney, D.R. Wilhoit, and D. Mauri, Phys. rev. B. 43. 1297 (1991). 9 B. Dieny, V.S. Speriosu, S. Metin, S.S.P. Parkin, B.A. Gurney, P. Baumgart and D.R. Wilhoit, J. Appl. Phys. 69, 4774 (1991). 10 C. -L. Lee, J. A. Bain, S. Chu and M. E. McHenry, J. Appl. Phys. 91, 7113 (2002). 77 11 http://www.google.com.sg/images?um=1&hl=en&tbs=isch%3A1&sa=1&q=CIPGMR+read+head&aq=f&aqi=&aql=&oq=&gs_rfai= 12 A. Vedyayev, M. Chshiev, N. Ryzhanova, B. Dieny, C. Cowache, and F. Brouers, J. Mag. magn. Mater. 172, 53 (1997). 13 W.P. Pratt, Jr., S.F. Lee, J.M. Slaughter, R. Loloee, P.A. Schroeder and J. Bass, Phys. Rev. Lett. 66, 3060 (1991); W.P. Pratt, Jr., S.-F. Lee, P. Holody, Q. Yang, R. Loloee, J. Bass and P.A. Schroeder, J. Magn. Magn. Mater. 126, 406 (1993); Q. Yang, P. Holody, R. Loloee, L.L. Henry, W.P. Pratt, Jr., P.A. Schroeder and J. Bass, Phys. Rev. B 51, 3226 (1995); M. A. M. Gijs, S. K. J. Lenczowski and J. B. Giesbers, Phys. Rev. Lett. 70, 3343 (1993). 14 A. Fert and T. Valet, Phys. Rev. B, 48, 7099 (1993). 15 Y. Jiang, S. Abe, T. Nozaki, N. Tezuka, and K. Inomata, Appl. Phys. Lett. 83. 2874 (2003). 16 S. Maat, J. Checkelsky, M.J. Carey, J.A. Katine and J.R. Childress, J. Appl. Phys. 98, 113907 (2006) 17 K. Yakushiji, K. Saito, S. Mitani, K. Takanashi, Y.K. Takahashi, and K. Hono, Appl. Phys. Lett. 88. 222504 (2006). 18 Y. Sakuraba, T. Iwase, K. Saito, S. Mitani, and K. Takanashi, Appl. Phys. Lett. 94. 012511 (2009). 78 19 Alain Achuhl, Daniel Lacour, C.R. Physique. 6, 945-955 (2005). 20 M. Julliere, Phys. Lett. 54A, 225 (1975). 21S. Ikeda, J. Hayakawa, Y. Ashizawa, Y.M. Lee, K. Miura, H. Hasegawa, M. Tsunoda, F. Matsukura, and H. Ohno, Appl. Phys. Lett. 93, 082508 (2008). 22 de Groot R A, Mueller F M, van Engen P G and Buschow K H J 1983 Phys. Rev. Lett. 50. 2024. 23 I. Galanakis: Phys. Rev. B 66,012406 (2002); I. Galanakis and Ph. Mavropoulos: Phys. Rev. B 67, 104417 (2003); Ph. Mavropoulos and I. Galanakis: J. Phys.: Condens. Matter 16, 4261 (2004). 24 S.Sanvito and N.A. Hill: Phys. Rev. B 62, 15553 (2000); A. Continenza, S. Picozzi, W.T. Geng, and A.J. Freeman: Phys. Rev. B 64, 085204 (2001); B.G. Liu: Phys. Rev. B 67, 172411 (2003); B. Sanyal, L. Bergqvist, and O. Eriksson: Phys. Rev. B 68, 054417 (2003); W.-H. Xie, B.-G. Liu, and D.G. Pettifor: Phys. Rev. B 68, 134407 (2003); W.-H. Xie, B.-G. Liu, and D.G. Pettifor: Phys. Rev. Lett. 91, 037204 (2003); M. Zhang et al.: J. Phys: Condens. Matter 15, 5017 (2003); C.Y. Fong, M.C. Qian, J.E. Pask, L.H. Yang, and S. Dag: Appl. Phys. Lett. 84, 239(2004); J.E. Pask, L.H. Yang, C.Y. Fong, W.E. Pickett, and S. Dag, Phys. Rev. B 67, 224420(2003); J.-C. Zheng and J.W. Davenport, Phys. Rev. B 69, 144415 (2004). 79 25 H. Akinaga, T. manago, and M. Shirai, Jpn. J. Appl. Phys. 39, L1118 (2000); M. Miizuguchi, H. Akinaga, T. Manago, K. Ono, M. Oshima, and M. Shirai, J. Magn. Magn. Mater. 239, 269 (2002); M. Mizuguchi, H. Akinaga, T. Manago, K.Ono, M. Oshirai, M, Yuri, H.J. Lin, H.H. Heieh, and C.T. Chen: J. Appl. Phys. 91, 7917 (2002); M. M. Mizuguchi, M.K. Ono, M. Oshima, J. Okabayashi, H. Akinaga, T. Manago, and M. Shirai: Surf. Rev. Lett. 9, 331 (2002); M.Nagao, M. Shirai, and Y. Miura: J. Appl. Phys. 95, 6518 (2004); K. Ono, J. Okabayashi, M.Mizuguchi, M. Oshima, A. Fujimori, and H. Akinaga: J. Appl. Phys. 91, 8088 (2002); M. Shirai: Physica E 10, 143 (2001); M. Shirai: J. Appl. Phys. 93, 6844 (2003). 26 J.H. Zhao, F.Matsukura, K. Takamura, E. Abe, D. Chiba, and H. Ohno: Appl. Phys. Lett. 79, 2776 (2001); J.H. Zhao, F. Matsukura, K. Takamura, E. Abe, D. Chiba, Y. Ohno, K. Ohtani, and H. Ohno: Mat. Sci. Semicond. Proc. 6, 507 (2003). 27 R.J. Soulen Jr., J.M. Byers, M.S. Osofsky, B.Nadgorny, T. Ambrose, S.F. Cheng, P.R. Broussard, C.T. Tanaka, J. Nowak, J.S. Moodera, A. Barry, and J.M.D. Coey: Science 282, 85 (1998). 28 M. Horne, P. Strange, W.M. Temmerman, Z. Szotek, A. Svane, and H. Winter: J. Phys.: Condens. Matter 16, 5061 (2004). 29 H. Kato, T. Okuda, Y. Okimoto, Y. Tomioka, K. Oikawa, T. kamiyama, and Y. Tokura: Phys. Rev. B 69, 184412 (2004). 80 30 A. Stroppa, S. Picozzi, A. Continenza, and A.J. Freeman: Phys. Rev. B 68, 155203 (2003). 31 H. Akai: Phys. Rec. Lett. 81, 3002 (1998). 32 P. J. Webster and K.R.A. Zieback, In: Alloysand Compounds of d-Elements with Main Group Elements. Part 2., Landolt-Bornstein, New Series, Group III, vol 19c, edited by H.R.J. Wijn, (Springer, Berlin 1988) pp75-184; A. Hirohata, M. Kikuchi, N. Tezuka, K. Inomata, J. S. Claydon, Y. B. Xu, G. van der Laan, Current Opinion on Solid State and Materials Science 10 (2006) 93-107. 33 T. Ambrose and O. Mryasov, U.S. Patent No.6, 876, 522 (5 April 2005); T. Ambrose and O. Mryasov, Half-Metallic Alloys: Fundamentals and Applications, Lecture Notes in Physics Vol.676, edited by I. Galanakis and P. Dederichs (Springer, Berlin-New York, 2005) 34 Galanakis I, Mavropoulos P, Dederichs PH, J. Phys. D: Appl. Phys. 39, 765 (2006). 35 Galanakis I, Dederichs PH, Papanikolaou N. Phys. Rev. B. 66, 174429 (2002). 36 Ishida S, Fujii D, Kashiwagi S, Asano S. J. Phys. Soc. Jpn. 64, 2151 (1995). 37 Buschow KHJ, van Engen PG. J. Magn. Magn. Mater. 25, 90 (1981). 38 Block T, Felser C, Jakob G, Ensling J, Muhling B, Gutlich P, et al. J. Solid State Chem. 176, 646 (2003). 39 Inomata K, Okamura S, Goto R, Tezuka N. Jpn. J. Appl. Phys. 42, L419 (2003). 40 I Galanakis and Ph Mavropolous, J. Phys.: Condens. Matter. 19, 315213 (2007). 81 41 Ishida S, Fujii D, kashiwagi S, Asano S, Jpn. J. Phys. Soc. 64. 2152 (1995). 42 P.J. Brown, K.V. Neumann, P.J. Websteand, and K.R.A. Zieback, J. Phys.: Condens. Mater 12, 1827 (2000). 43 S.F. Cheng, B. Nadgorny, K. Bussman, E.E> Carpenter, B.N. Das, G. Trotter, M.P. Raphael, and V.G. harris, IEEE Trans. Magn. 37, 2176 (2001). 44 L. Ritchie, G. Xiao, Y. Ji, T.Y. Chen, C.L. Chen, M. Zhang, J. Chen, Z. Liu, G. Wu, and X.X. Zhang, Phys. Rev. B 68, 104430 (2003). 45 L.J. Singh, Z.H. Barber, Y. Miyoshi, W.R. Branford, and L.F. Cohen, J. Appl. Phys. 95, 7231 (2004). 46 L.J. Singh, Z.H. Barder, Y. Miyoshi, Y. Bugoslavsky, W.R. Branford, and L.F. Cohen, Appl. Phys. Lett. 84, 2367 (2004). 47 L.J. Singh, Z.H. Barder, A. Kohn, A.K. Petford-Long, Y. Miyoshi, Y. Bugoslavsky, and L.F. Cohen, J. Appl. Phys. 99, 013904 (2006). 48 Picozzi S, Continenza A, Freeman A, J. Phys. Rev. B. 66, 094421(2002). 49 Mavropoulos Ph, Sato K, Zeller R, Dederichs P H, Propescu V and Ebert H, Phys. Rev. B. 69, 054424 (2004); 50 Mavropoulos Ph, Galanakis I, Popescu V and Dederichs P H, J. Phys.: COndens. Matter. 16, S5759 (2004). 82 51 Galanakis I, Ostanin S, Alouani M, Dreysse H, and Wills J M, Phys. Rev. B. 61, 4093 (2000). 52 Skomski R, Dowben P A, Europhys. Lett. 58, 544 (2002). 53 Dowben P A, Skomski R, J. Appl. Phys. 93, 7948 (2003). 54 Dowben P A and Skomski R, J. Appl. Phys. 95, 7453 (2004). 55 Lezaic M, Mavropoulos Ph, Enkovaara J, Bihlmayer G and Blugel S, Phys. Rev. Lett. 97, 026404 (2006). 56 Chioncel L, Katsnelson MI, de Groot R A and Lichtenstein A I, Phys. Rev. B.68, 144425 (2003). 57 Itoh H, Ohsawa T, Inoue, Phys. Rev. Lett. 84, 2501 (2000). 58 Picozzi S, Continenza A and Freeman A, Phys. Rev. B. 69, 094423 (2004). 59 Y. Miura, K. Nagao, M. Shirai, Phys. Rev. B. 69, 144413(2004). 60 A. Rajanikanth, Y. K. Takahashi, K. Hono, J. Appl. Phys, 105, 063916 (2009). 61 A. Rajanikanth, Y.K. Takahashi, K. Hono, J. Appl. Phys. 105, 063916 (2009). 62 T. Furubayashi, K. Kodama, H. Sukegawa, Y.K. Takahashi, K. Inomata, K. Hono, Appl. Phys. Lett. 93, 122507 (2008). 63 W.H. Wang, M. Przybylski, Phys. Rev. B. 71, 144416 (2005). 64 akonstantin Nikolaev, Paul Kolbo, Tras Pokhil, Xilin Peng, Yonghua Chen, Thomas Ambrose, and Oleg Mryasov, Appl. Phys. Lett. 94, 222501 (2009). 83 65 H. Kijima, T. Ishikawa, T. Marykame, H. Koyama, K. Matsuda, T. Uemura, M. Yamamoto, IEEE. Trans. Magn. 42, 10(2006). 66 D.L. Smith, Thin Film Deposition, McGraw-Hill Professional (2001). 67 J.E. Mahan, Physical Vapor Deposition of Thin Films, John Wiler&Sons (2000). 68 P.J. Flanders, J. Appl. Phys., 63, 3940 (1988). 69 B.D. Cullity, Introduction to magnetic materials, Addison-Wesley Pub. Co. (1972). 70 K. H. J. Buschow, F. R. De. Boer, Physics of Mangnetism and Magnetic Materials (2004). 71 W.L. Bragg, "The Diffraction of Short Electromagnetic Waves by a Crystal", Proceedings of the Cambridge Philosophical Society, 17 (1913), 43–57. 72 B.D.Cullity, S.R.Stock, Elements of X-Ray Diffraction. Prentice-Hall Inc (2001). 73 B.E. Warren, X-Ray Diffraction. Addison-Wesley Publishing Company (1969). 74 http://www.google.com.sg/images?hl=en&q=brag%20law&um=1&ie=UTF-8&source =N&tab=wi 75http://www.google.com.sg/images?um=1&hl=en&tbs=isch%3A1&sa=1&q=XRD&aq=f&aqi =g3g-m1&aql=&oq=&gs_rfai=. 84 76 D.B. Williams, and C.B. Carter, Transmission Electron Microscopy, A Textbook for Materials Science, Plenum Press, New York and London, (1996). 77 M.J. Yacaman, G, Diaz, and A. Gomez, Catal. Today, 23, 161(1995). 78 H. Hertz, Ann. Physik 31, 983 (1887). 79. A. Einstein, Ann. Physik, 1921 Nobel Prize in Physics. 80 Giessibl, Franz J, Rev. Mod. Phys, 75, (2003), 949. 81 Roiter, Y, Minko, S, J. Am. Chem. Soc. 127 (45): 15688–9. 82 G. Binnig, H. Rohrer, ISM J. Res. Dev. 30, 279(2000). 83 D. K. Aswal, K.P. Muthe, S. Tawde, S. Chodhury, N. Bagkar, A. Singh, S.K. Gupta and J. V. Yakhmi: J. Cryst. Growth, 70, 191 (2001). 84 Yuya SAKURABA, Jun NAKATA, Mikihiko OOGANE, Hitoshi KUBOTA1, Yasuo ANDO, Akimasa SAKUMA and Terunobu MIYAZAKI, Jpn. J. Appl. Phys., 44, (2005) 6535–6537. 85 V. A. Oksenenko, V. A. Kulagin, Y. V. Kudryavtsev, J. Dubowik, I. Goscianska, Y. N. Troshchenkov, J. Mag. magn. Mater. 316 (2007). 86 D. Givord, Q. Lu, M. F. Rossignol, P. Tenaud, T. Viadieu, J. Mag. magn. Mater. 83 (1990) 183-188. 85 87 G. Bertotti, Hysteresis in Magnetism: For Physicists, Materials Scientists and Engineers, Academic Press, New York, 1998. 86 [...]... et al [34], the origin of the bandgap in Heusler alloys is caused by the d-d states hybridization of X and Y transition metals, as the DOS 14 in the vicinity of EF is dominated by the d-states The formation of this gap in half -Heusler alloys (Fig 1.12(a)) and full -Heusler alloys (Fig 1.12(b)) is not exactly the same In the case of halfHeusler alloys, the gap is formed by the hybridization states between... in which the C1 b structure is shaped by removing one of the X sites in L21 structure In addition, Y-Z atomic disorder in L21 structure of full -Heusler alloy will result in the formation of B2 structure, while A2 structure will form when X-Y and X-Z disorder occur Figure1.11: Major combination of Heusler alloy formation Adapted from [32] 1.4.1.1 Origin of the bandgap According to the calculations by... coupling effect and align the magnetization of consecutive layers was too large for GMR to be applicable in real devices On the other hand, further investigation revealed that the anti-ferromagnetic coupling arrangement was not a prerequisite for the GMR effect [6] The physical origin of GMR effect can be explained by the effect of the electron spin on the electronic transport in ferromagnetic conductors,... requirement for practical devices The basic properties and research progress on Heusler alloys will be discussed below Figure1.9: Schematic representation of the DOS for a half-metal with respect to normal metals and semiconductors [33] 1.4.1 Basic properties of Heusler alloys Figure 1.10: Structure of (a) C1b half -Heusler alloys; (b) L2 1 full -Heusler alloys; (c) B2 full -Heusler alloy; (d) A2 full -Heusler alloy. .. effect and its polarization is close to 100% [49, 50] The Heusler alloys like Co2MnSi, Co2 MnGe and Co2MnSn showed small orbit moments based on the calculations of Galanakis et al [48, 51] 1.4.2.3 Effect of temperature on spin polarization Several groups have investigated the temperature effect on the polarization of Heusler alloys qualitatively and quantitatively based on different assumptions and theories,... Although the half-metallic property was conserved in both expansion and compression scenarios, shifts of Fermi level occurred Based on their calculations, the shift of EF 17 was attributed to the larger extension of p states compared to the d states of Sb Movement of Fermi level towards conduction band took place in compression, while Fermi level moved towards the valence band during expansion In addition,... across the interface On the other hand, poor matching increased the scattering of spin-down electrons Seagate introduced an all -Heusler alloy CPP-GMR spin valve using ferromagnetic Co 2MnGe and non -magnetic Rh2CuSn Based on band structure calculations, the interface spin asymmetry of this structure would be maximized [64] However, some degree of disorder caused a loss of polarization at E F and hence... antisites In addition, large formation energies of the Mn-Si and Mn-Co atomic swaps were found However, these results cannot be generalized to all Heusler alloys Recently, the investigation of Nd doping effect on the transport and magnetic properties of CMS had been reported by K Hono et al [60] From the resistivity measurements at low temperatures, it was concluded that electron-magnon scattering was... scattering led to long mean free path When the magnetization directions of two ferromagnetic layers were parallel, the spinup electrons, assumed to be parallel to magnetization, passed through the multi-layers with almost zero scattering On the other hand, the spin-down electrons were scattered strongly as their spin aligned anti-parallel to the magnetization direction Thus, parallel configuration resulted... Investigations showed that some Heusler alloys retained their half-metallicity in B2 structure, while A2 disorder degraded the spin polarization significantly [58, 59] Picozzi et al [58] investigated the formation of defects in full -Heusler alloy, in particular Co2 MnGe and Co2 MnSi They found that the Mn antisites had the lowest formation energy and did not destroy the half-metallicity in contrast to ... films for spintronic application It would 23 provide guidance on the understanding of structural and magnetic properties of CMS Heusler alloys and its further application into the field of spin-electronics... [34], the origin of the bandgap in Heusler alloys is caused by the d-d states hybridization of X and Y transition metals, as the DOS 14 in the vicinity of EF is dominated by the d-states The formation... properties of Heusler alloys Figure 1.10: Structure of (a) C1b half -Heusler alloys; (b) L2 full -Heusler alloys; (c) B2 full -Heusler alloy; (d) A2 full -Heusler alloy Adapted from [32] 13 The Heusler