Strain-Controlled Giant Magnetoresistance in Spin Valves Grown on

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Strain Controlled Giant Magnetoresistance in Spin Valves Grown on Shape Memory Alloys Mengxi Wang, Xiaoguang Xu, Shijie Hao, Feng Yang, Libai Zhu, Yong Wu, Kangkang Meng, Jikun Chen, Jun Miao, and Yong Jiang ACS Appl. Electron. Mater., Just Accepted Manuscript • DOI: 10.1021/acsaelm.9b00106 • Publication Date (Web): 13 May 2019 Downloaded from http://pubs.acs.org on May 21, 2019

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ACS Applied Electronic Materials

Strain Controlled Giant Magnetoresistance in Spin Valves Grown on Shape Memory Alloys

Mengxi Wang†, Xiaoguang Xu*, †, Shijie Hao‡, Feng Yang‡, Libai Zhu†, Yong Wu†, §, Kangkang Meng†, Jikun Chen†, Jun Miao†, Yong Jiang*, †

† Beijing

Advanced Innovation Center for Materials Genome Engineering, School of

Materials Science and Engineering, University of Science and Technology Beijing, Beijing 100083, China. ‡

Department of Material Science and Engineering, China University of

Petroleum-Beijing, Beijing 102249, China § Institute

of Physics, Johannes Gutenberg-University Mainz, 55128 Mainz, Germany

ACS Paragon Plus Environment

ACS Applied Electronic Materials 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

ABSTRACT: We report a strain mediated giant magnetoresistance (GMR) in spin valves (SPVs) grown on shape memory alloys (SMAs). The SPVs with a stacking structure of Al2O3/Co90Fe10/Cu/Co90Fe10/IrMn/Pt were deposited on NiTi and NiTiNb SMA substrates with and without predeformation. The GMR of the SPVs on the precompressed NiTi substrates changes after annealing, whereas the one on SMAs without predeformation remains unchanged under the same annealing treatment. Especially, the GMR of the SPVs on the NiTiNb substrates can be reversibly modulated under a thermo-cycling treatment on account of the shape memory effect caused by the phase transitions of NiTiNb. The study demonstrates a new strain mediated approach to reversibly control GMR through phase transitions, which provides a potential way to design strain sensitive spintronic devices. KEYWORDS: Spin valve, Giant magnetoresistance, Spintronic device, Shape memory alloy, Strain engineering.


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1. INTRODUCTION Strain engineering is an effective method to manipulate magnetic electrical

6-11

1-5

and

properties of thin films and spintronic devices. Piezoelectric materials

are often used as strain sources due to their reversible piezoelectric effect. However, the maximum strain induced by piezoelectric materials is limited to be less than 1% in a specific direction under large voltage.

12,13

To study large strain manipulating

effect on spintronic devices, a few reports turn to shape memory alloys (SMAs). 14,15 Since SMAs usually show large strain originating from martensite phase transition under thermal treatment, such as the maximum deformation of 8% for NiTi based SMAs16-20, it is reasonable to expect strain mediated spintronic devices manipulated by SMA substrates. In addition, compared with piezoelectric substrates

21-23,

the

SMA ones provide a new driving method of spintronic devices by strain, which could be possible to take the variation of natural temperature as driving force in some applications. Spin valves (SPVs) had been extensively used in magnetic read heads of hard disk drives due to their giant magnetoresistance (GMR).

24

The GMR of SPVs is

greatly dependent on the properties of magnetic thin films.

25-27

According to

previous studies, a strain transferred to magnetic thin films can induce the change of their magnetic anisotropy. According to Mott’s model,

28

the change of effective

magnetic anisotropy constant could affect the transport of electrons inside ferromagnets, therefore the strains should affect the GMR of SPVs. SMA substrates have been employed as strain sources to control the properties of magnetic thin films,



such as the coercivity (Hc) of L10-FePt film

18,19

and the dynamical manipulation of

oxygen ion (O2−) in Pt/Co/MgO heterostructures.

20

However till now there is no

report on the directly strain controlled GMR on SMA substrates. In this work, we use SMA substrates to tune the GMR values of the SPVs deposited on them. For precompressed NiTi substrates, due to the martensite (M) to austenite (A) phase transition at higher temperature, the precompressed state is relaxed, and it leads to a tensile strain in the magnetic thin films. For precompressed NiTiNb substrates, it leads to tensile strain states in the magnetic thin films at room temperature with the phase transitions from M to A (room temperature to 375 K and cooling to room temperature) or A to M (room temperature to 200 K and warming to room temperature). During the A to M phase transition, the substrates can still induce a tensile strain, because NiTiNb cannot revert to the martensite phase completely, so that the tensile strain cannot be entirely relaxed. It is demonstrated that the strain generated from the SMA substrates drives reversible GMR variations at the transition temperatures between martensite and austenite phases.

2. EXPERIMENTAL SECTION First, The Ni50Ti50 (NiTi) and Ni45Ti45Nb10 (NiTiNb) substrates were precompressed to 0%, 2%, 4% and 6% respectively along the x-axis to induce orientated martensitic phase in SMAs. Then, the surfaces of the NiTi and NiTiNb substrates were polished until their roughness reduces to less than 1 nm, as demonstrated by the atomic force microscope images shown in Figure S1a, and the sizes of the substrates are both 3×3×1 mm. The two kinds of SMA substrates are all


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paramagnetic based on their in-plane hysteresis loops shown in Figure S1b. Resistance strain gauges were applied to measure the detailed strain information with the phase transitions of the SMAs. The strain gauges were pasted closely to the surfaces of the SMA substrates with different amount of predeformation. The resistances of the strain gauges were measured to reflect the variation of strains for the SMA substrates. In order to systematically study the effect of strain on magnetic thin films. We have deposited the multilayers of NiTi/Al2O3 (30)/Co90Fe10 (5)/Pt (3) (denoted as Sample I), Ni45Ti45Nb10 (NiTiNb)/Al2O3 (30)/Co90Fe10 (5)/Pt (3) (Sample II), NiTi/Al2O3 (30)/Co90Fe10 (5)/IrMn (12)/Pt (3) (Sample III) and NiTiNb/Al2O3 (30)/Co90Fe10 (5)/IrMn (12)/Pt (3) (Sample IV) (thickness in nanometer) with different precompression of substrates Two samples of SPVs were also fabricated on the precompressed NiTi (Sample V) and NiTiNb (Sample VI) substrates with a stacking structure of SMA/Al2O3 (30)/Co90Fe10 (5)/Cu (3)/Co90Fe10 (5)/IrMn (12)/Pt (3). To improve the quality of our samples, we have optimized the stacking structure on Si/SiO2 substrates before fabricating the SPVs on SMA substrates. One-way and two-way regulations on the SPVs were performed through the shape memory effect of the SMAs. The films were deposited on the substrates by magnetron sputtering with an external field of 100 Oe along the 30o direction of the strains to unify the easy magnetization orientation of the magnetic layers. The base pressure was better than 1 ×10-5 Pa. During deposition, the Ar pressure was 0.5 Pa. Since the SMA substrates are conductive, an Al2O3 layer was deposited on them as an insulating layer to ensure that the current only flows through the SPVs. It can also block the magnetic coupling



between the substrates and the magnetic layers. The SPVs deposited on the NiTi substrates are for one-way regulation while those on the NiTiNb ones are for two-way regulation. The process of heating or cooling is required to control the strain transferred to the samples before measurements. The samples and strain regulations are schematically illustrated in Scheme 1. The samples with the precompressed NiTi substrates were heated up to 375 K and kept for 20 minutes. After the thermal treatment, a tensile strain was transferred to the thin films deposited on them, because the substrates tend to revert to their original shapes due to the one-way memory effect. As for the NiTiNb substrates with two-way effect, after experiencing a cycle of heating and cooling treatment, they were heated to 375 K then hold for 20 minutes to ensure a complete phase transition into austenite phase and then cooled down to room temperature (300 K). The SMA substrates can “remember” their shapes at high temperatures. Next, the substrates were cooled down to 200 K then hold for 20 minutes and returned to room temperature, in this case the substrates can transit to martensite phase and “remember” their shapes at low temperatures. The resistances of the strain gauges for the low-temperature and high-temperature states were measured to reflect the variation of strains for the SMA substrates. The states of the samples are denoted as A1 and A2 with the SMAs taking on the austenite phase for the first and the second time during the thermal treatment, and M1 and M2 for those taking on the martensite phase, respectively. The annealing and cooling processes were realized in vacuum with the pressure better than 0.1 Pa. Predeformation is a method to induce orientated martensitic phase in SMAs.


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The images of surface roughness of the substrates are obtained by atomic force microscope (AFM, Bruker). The magnetic properties of the multilayers were measured by vibrating sample magnetometer (VSM, Quantum Design). Differential scanning calorimeter (DSC, NETZSCH) was also used to observe the phase transition process of the SMAs. GMR signals were measured at room temperature using Keithley 2400 and 2182A. And all the measurements of magnetic and electrical properties were done at room temperature (300 K). To further understand the effect of the strains on the magnetic properties of the SPV/SMA heterostructures, we have also calculated the M-H loops of the free and pinned layers by using Stoner-Wohlfarth model. 29

3. RESULTS AND DISCUSSION To use the one-way memory effect of SMA, we compressed the NiTi substrates at room temperature. Upon heating above the phase transition temperature, the substrates recover to their original shapes. When cooling down to 300 K, they will remain their high-temperature shapes. For the NiTiNb substrates with the two-way memory effect, they were treated with heating and cooling. NiTiNb shows shape memory effect during both heating and cooling processes. The strains induced by the shape changes transfer to the films and manipulate the magnetic and electrical properties of the SPVs. In this study, the phase transition temperatures of the SMAs were determined by DSC. The austenite transition temperatures of NiTi and NiTiNb are about 324 K and 313 K respectively, while the martensitic transition temperature of NiTiNb is about 270 K, as shown in Figures 1a and b.



We have measured the strain information of the SMA substrates by the strain gauges. The resistance of the strain gauges changes with the deformation of the SMA substrates. The deformation recovery rate ε can be obtained from the formula 𝛥R /R=KS × ε, where the sensitivity coefficient KS is 2.0, and ΔR and R represent the variation and initial value of the strain gauges’ resistance, respectively. The dependences of ε on the predeformation and phase state are presented in Figures 2a and b for the NiTi and NiTiNb substrates, respectively. It can be seen that ε increases with the increased predeformation for the NiTi substrate. For the NiTiNb substrate, the austenite phase shows similar variation of ε with that of the NiTi substrate, while the martensite phase only slightly increases ε. Moreover, ε almost keeps constant for the same phase during the thermal treatment cycles, indicating the stable shape memory effect of the SMA substrates. The magnetic properties of the free and pinned layers of the SPVs were studied under different strains induced by thermal treatment. Using the method shown in Figure 3a, we have measured the M-H loops of Sample I to IV under different deformation with an external magnetic field along the x-axis. All the hysteresis loops were obtained at room temperature. The M-H loops vary with the deformation and phase state of the substrates. For Sample I, the tensile strain along the x-axis caused by the phase change can be transferred to the upper Co90Fe10 layer, and increase the coercivity (Hc) of the magnetic film, as shown in Figure 3b. For Sample II shown in Figure 3d, the precompressed NiTiNb substrate went through the cycle of heating and cooling to realize the reversible regulation of strain due to the reversible phase


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transition of NiTiNb. The NiTiNb substrate reaches the austenite transition temperature at 375 K. The tensile strain generates along the direction of the deformation and passes to the films. And the austenite phase (A1) can maintain after returning to room temperature. Upon cooling down to 200 K, the NiTiNb substrate reaches the martensitic transition temperature, and then tend to be back to the martensitic phase (M1) before annealing. In the next heating and cooling cycle, the sample undergoes the same strain cycle from high-temperature A2 to low-temperature M2. By measuring the M-H loops during the thermal treatment cycles, we found similar tendencies of Hc changes for the pinned layers of Sample III and IV, as shown in Figure 3c and Figure 3e. In addition, although the 6% precompressed samples can provide a lager tensile strain, the surface roughness becomes worse. As a result, the strain cannot be transferred to the top films effectively, leading to the smaller changes of Hc for the 6% precompressed samples than those for the 4%. Therefore, the changes of Hc for the 6% precompressed samples are comparable or smaller than those of the 4% precompressed samples. Moreover, we also measured the M-H loops for SPVs used to compare with the GMR results. The detailed M-H loops of Sample I and Sample II are shown in Figure S2, the loops of Sample III and Sample IV are shown in Figure S3 in the supplementary material and the loops of SPVs are shown in Figure S4. In order to further study the change of magnetic properties of the magnetic layers on the two-way SMA substrates during the phase transition from austenite to martensite, the M-H loops of Sample II with different amount of predeformation were



measured from 300 K to 200 K, after the heating and cooling cycle between 375 K and room temperature to form austenite phase. As shown in Figure 5, the Hc of the samples grown on the precompressed substrates initially decreases and then increases with the decreasing temperature. However, the Hc of the sample deposited on the substrate without deformation treatment keeps increase with the decreasing temperature. The detailed M-H loops are shown in Figure S5 in the supplementary material. The phase transition from austenite to martensite occurs in the precompressed SMA substrates when the temperature decreases from 300 K. Correspondingly, the Hc of the Co90Fe10 layer decreases with the cooling process on the precompressed substrates. The change of Hc increases with the increased precompression strain. Below the temperature of the phase transition, the substrates are in martensite phase, and Hc increases slightly with the decreasing temperature, which is consistent with the trends of normal magnetic films during cooling process below Curie temperature. For all the samples from I to IV, the SMA substrates without deformation were also employed for comparison. The results demonstrate that the Hc values of both the free and pinned layers do not change after the same thermal treatment processes. Since both the free and pinned layers can be regularly modulated by strain, it is reasonable to expect the strain-controlled GMR of the SPVs deposited on the SMA substrates. Before preparing the SPVs on SMA substrates, we have fabricated SPVs on Si/SiO2 substrates to optimize their structure and technological parameters. The GMR values of SPVs on the Si/SiO2 substrates are about 3.4% as shown in Figure S6


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in the supplementary material. Figure 5a is the schematic diagram of the structure and measurement for the SPVs fabricated on the SMA substrates. All the GMR measurements were carried out at room temperature after heating or cooling treatment. The GMR values are less than 2% on the SMA substrates, which may be due to the relatively large surface roughness of the SMA substrates. It is reasonable since the GMR values are sensitive to the substrate roughness. As shown in Figure 5b, the GMR of the SPVs grown on the precompressed NiTi substrates (Sample V) is enhanced by annealing, while that on NiTi without predeformation remains unchanged after the same annealing process. More interesting is that GMR increases linearly with the predeformation of NiTi. However, GMR is stable for the SPVs deposited on the substrates without predeformation. To demonstrate the reversible modulation of strain on the GMR of the SPVs, we have measured the GMR curves after a heating or cooling treatment of the SPVs deposited on the 4% precompressed NiTiNb substrates (Sample VI), which are presented in Figure 5c. The GMR values are in high states on the heated NiTiNb substrates in austenite phase in different thermal treatment cycles. When the NiTiNb substrates were cooled down to take on the martensite phase, the GMR values decrease to low states. We repeated the thermal treatment cycles, the GMR values repeated with the phase transition of SMA substrates. We also observed the similar phenomenon on the 2% precompressed NiTiNb substrates, while the GMR values show no obvious changes on NiTiNb without predeformation (See Figure S7 in the supplementary material). Therefore, the SPVs are sensitive to the strain from the SMA substrates, so as to “remember” the



GMR values corresponding to the phases of the two-way effect SMA substrates. To further understand the effect of strain on the SPV/SMA heterostructures, we have studied the magnetic properties of the free and pinned layers by using Stoner-Wohlfarth model. 29 The total energy E1 of the free layer and E2 of the pinned layer are given by Equation (1) and (2), respectively. 𝑬𝟏 = 𝑲𝒖𝒔𝒊𝒏𝟐 (𝝋 ― 𝟑𝟎𝒅𝒆𝒈) ― 𝑲𝝈𝒔𝒊𝒏𝟐 𝝋 ― 𝝁𝟎𝑴𝒔𝑯𝒄𝒐𝒔 𝝋 (1) 𝑬𝟐 = 𝑲𝒖𝒔𝒊𝒏𝟐 (𝝋 ― 𝟑𝟎𝒅𝒆𝒈) ― 𝑲𝝈𝒔𝒊𝒏𝟐 𝝋 ― 𝑲𝒆𝒃𝐜𝐨𝐬 (𝝋) ― 𝝁𝟎𝑴𝒔𝑯𝐜𝐨𝐬 (𝝋), (𝟐)

where 𝜃 is the angle between the easy axis and the magnetic field, 𝜑 represents the angle between the magnetization direction and the magnetic field. H and 𝑀𝑠 are the magnetic field and the saturation magnetization, respectively. 𝐾𝑢 is the uniaxial anisotropy calculated by 𝐾𝑢 = 𝑀𝑠𝐻𝑘/2, in which 𝐻𝑘 represents the magnetic anisotropy field. The unidirectional anisotropy 𝐾𝑒𝑏 equals 𝑀𝑠𝐻𝑒𝑏, where 𝐻𝑒𝑏 is the exchange bias field. The external strain can induce magnetic anisotropy and be expressed as 𝐾𝜎 = ― 3/2 𝜆𝜎 and 𝜎 = 𝑌𝜀, where 𝜎 and 𝜀 are the stress and strain corresponding to the deformation amount from the substrates, respectively. The saturation magnetostriction coefficient 𝜆 is about 83 ppm for Co90Fe10, and the Young’s modulus Y is about 210 GPa for Co90Fe10. 30,31 The detailed derivation of the formulas is shown in supplementary. In the simulation section, we have simplified our simulation by considering the free and pinned layers to be single domain. The simulated magnetization curves of the free and pinned layers based on the NiTi and NiTiNb substrates are shown in Figures 6 a-h, together with the schematic diagram of the effect of strain on the spin transport


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of the SPVs based on two current model. 28 Phase transition occurs after the heating or cooling treatment of the SMA substrates. The consequent strain passes on to the magnetic thin films deposited on them, which drives the changes of 𝐾𝜎 and Hc of the magnetic layers. For the NiTi substrates, they tend to recover their original shapes after annealing. The strain drives the magnetic layer into a tensile state in the direction of x-axis and leads ferromagnetic moments to orientate along the x-axis. As a result, K σ and Hc increase with the increased predeformation as shown in Figures 6a and b. For the two-way SMA, the NiTiNb substrates can remember two different shapes at high and low temperatures. The magnetic layer grown on them will also obtain two strain states, thus Hc can be reversibly tuned by the phase transitions of NiTiNb as shown in Figures 6c-h. For the free layer on the SMA substrates with larger predeformation, Hc shows larger changes, as shown in Figures 6a and c-e. For the pinned layer, the magnitude of the exchange bias field is related to the spin alignment of the magnetic atoms at the ferromagnetic/antiferromagnetic (FM/AFM) interface. 32-36

As shown in Figures 6b and f-h, when the external magnetic field is parallel to

the x-axis at room temperature, strain has no obvious influence on the exchange bias. Comparing the M-H loops between the simulated and experimental ones, the change of Hc is smaller for the later. Another reason for this is that our simulation assumes the complete transfer of the strain to the magnetic layers. However, in fact the strain should be released when it is transferred to the Al2O3 layer. Therefore, the multi-domain states and interactions in the actual systems will decrease the efficiency of the strain transmission.



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As shown in Figure 6i, M0 represents the magnetization direction of the original state after depositing on the precompressed substrates. MA represents the magnetization direction of the substrates in austenite phase after annealing at 375 K and then cool down to room temperature. MM represents the magnetization direction of the substrates in martensite phase state after cooling down to 200 K then back to room temperature. During the phase transition of the SMA substrates, the strain will change and lead to the rotation of the magnetization. The directions given in the diagram schematically represent relative trends. For one way effect, tensile strain is transferred to the magnetic layers after the annealing of NiTi substrates at 375 K, which rotates the magnetization from M0 to MA. For two way effect, the NiTiNb substrates change between martensite and austenite phases and cannot be completely restored to their original shapes, so the magnetization shall rotate between MA and MM. The above changes of magnetic properties of the thin films can explain the strain mediated GMR of the SPVs grown on the SMA substrates, which is schematically illustrated in Figure 6j based on the two-current model.

28

When the direction of the

external magnetic field is parallel to the x-axis, the tensile strain parallel to the x-axis increases the uniaxial anisotropy of the free and pinned layers. So, the magnetic moments are much easier to orientate in the direction of x-axis under the magnetic field. Under the tensile strain, the scattering of the minority spins at the ferromagnetic/non-magnetic (FM/NM) interface increases, therefore, the antiparallel resistance increases, resulting in a larger GMR.

4. CONCLUSION


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We have studied the effect of strain on the magnetic layers and SPVs deposited on different SMA substrates. The strains from the SMAs induced by thermal treatment can be introduced to the heterostructures on the substrates, and thereafter affect their magnetic properties. Hc increases with the tensile strain for both the free and pinned layers, while the exchange bias is not sensitive to the strain. GMR shows a regular increase under the tensile strain with the increased predeformation of the SMA substrates. Moreover, both Hc and GMR can be tuned reversibly for the SPVs deposited on the NiTiNb substrates due to the two-way memory effect. According to our study, the SMAs could be employed as potential materials to design strain mediated spintronic devices.



■ ASSOCIATED CONTENT Supporting Information AFM images for the surfaces of the NiTi and NiTiNb substrates, in-plane hysteresis loops of the NiTi and NiTiNb substrates; One-way regulation of the M-H loops of Sample I. “M” represents the initial state after deposition, “A” represents the austenite phase after annealing from 375 K, one-way regulation of the M-H loops of Sample II; Two-way regulation of the M-H loops of Sample II. A1 and A2 represent the austenite phases after annealing from 375 K for the first and second times, and M1 and M2 represent the corresponding martensite phases after cooling down to 200 K, respectively, two-way regulation of the M-H loops of Sample IV; M-H loops for the SPVs grown on the 0% and 4% precompressed SMA substrates; M-H loops of Sample II deposited on the NiTiNb substrates with different deformations during the cooling treatment; The GMR curve of a SPV deposited on a SiO2 substrate with the stacking structure of Co90Fe10 (5)/Cu (3)/Co90Fe10 (5)/IrMn (12)/Pt (3); The dependence of GMR on different phases for the SPV on the 0% and 4% deformed NiTiNb substrates; The detailed derivation of the formulas.

■ AUTHOR INFORMATION Corresponding Author * E-mail: [email protected] (X.X.). * E-mail: [email protected] (Y.J.). Notes The authors declare no conflict of interest.


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■ ACKNOWLEDGEMENTS This work was partially supported by the National Basic Research Program of China (Grant No. 2015CB921502), the National Science Foundation of China (Grant Nos. 51671019, 51731003, 11574027, 51602022, 61674013).



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Magnetoelectric

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using significant strain induced by shape memory substrates. Sci. Rep. 2016, 6, 20199. (20)Feng, C.; Wang, S.; Yin, L.; Li, X. J.; Yao, M. K.; Yang, F.; Tang, X. L.; Wang, L.; Mi, W. B.; Yu, G. H. Significant Strain‐Induced Orbital Reconstruction and Strong Interfacial Magnetism in TiNi(Nb)/Ferromagnet/Oxide Heterostructures via Oxygen Manipulation. Adv. Funct. Mater. 2018, 1803335. (21)Zhao, W. B.; Huang, W. C.; Liu, C. C.; Hou, C. M.; Chen, Z. W.; Yin, Y. W.; Li, X, G. Electric-Field-Controlled Nonvolatile Magnetization Rotation and Magnetoresistance Effect in Co/Cu/Ni Spin Valves on Piezoelectric Substrates. ACS Appl. Mater. Interfaces 2018, 10, 21390−21397. (22)Rizwan, S.; Zhang, S.; Yu, T.; Zhao, Y. G.; Han, X. F. Piezoelectric enhancement of giant magnetoresistance in spin-valves with different magnetic anisotropies. Journal of Appl. Phy. 2013, 113, 023911. (23)Lei, N.; Devolder, T.; Agnus, G.; Aubert, P.; Daniel, L.; Kim, J. V.; Zhao, W. S.; Trypiniotis, T.; Cowburn, R. P.; Chappert, C.; Ravelosona, D.; Lecoeur, P.; Strain-controlled

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Multiferroic


Figure Captions Scheme 1. Schematic diagram of the samples and strain regulations. The substrates were precompressed first. Then the films were deposited under an external field of 100 Oe along the 30o direction of the strains. The magnetic properties of the free layers and pinned layers and the GMR of the SPVs were measured at room temperature (300 K) after annealing or cooling process. The magnetization states and GMR values change along with the phase transitions of SMA substrates controlled by thermal treatment.

Figure 1. DSC curves of (a) the NiTi SMA obtained during heating and (b) the NiTiNb SMA during heating (red curve) /cooling (blue curve) at a heating/cooling rate of 5 oC/min.

Figure 2. The deformation recovery rate ε of the SMA substrates after thermal treatment cycling measured by the strain gauges. (a) ε of a NiTi substrate before and after annealing with different predeformation. The inset shows the schematic diagram of the SMA substrate with a strain gauge attached closely and the arrows represent the direction of strain. (b) ε of a NiTiNb substrate after annealing from 375 K and cooling down to 200 K.

Figure 3. Coercivity modulations of the free layer Al2O3/Co90Fe10 (Sample I), pinned layer Al2O3/Co90Fe10/IrMn (Sample III) via strains induced by the NiTi substrates, the


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free layer Al2O3/CoFe (Sample II) and pinned layer Al2O3/CoFe/IrMn (Sample IV) via strains induced by the NiTiNb substrates. (a) Schematic diagram of the measurement for M-H loops. The black solid arrows represent the direction of the strain after thermal treatment, and the black dashed arrow shows the direction of external magnetic field. (b) One-way regulation of the Hc of Sample I before and after annealing. and (c) One-way regulation of the Hc of Sample III before and after annealing. (d) Two-way regulation of the Hc of Sample II. (e) Two-way regulation of the Hc of Sample IV. A1 and A2 represent the austenite phases after annealing at 375 K for the first and the second times, and M1 and M2 represent the corresponding martensite phases after cooling down to 200 K, respectively.

Figure 4. The relationship between the Hc of Sample II and temperature during the cooling treatment.

Figure 5. GMR of the SPVs deposited on SMA substrates with different amount of predeformation. (a) Schematic diagram of the structure and GMR measurement. (b) The left part shows the GMR curves of the SPVs deposited on the NiTi SMA substrates with different amount of predeformation before and after annealing. The right part shows the dependence of GMR on predeformation for the SPVs on the NiTi substrates. The bottom part represents the relationship different strain states under different phases of NiTi substrates. (c) The left part shows GMR curves of the SPVs deposited on the 4% deformed NiTiNb substrates, which were measured after heating



(red curves) and cooling process (blue curves). The right part shows the dependence of GMR on different phases for the SPV on the 4% deformed NiTiNb substrate. The bottom part represents the relationship different strain states under different phases of NiTiNb substrate.

Figure 6. (a) and (b) are the simulated M-H loops of the free and pinned layers on the NiTi substrates, respectively. [(c) to (e)] and [(f) to (h)] correspond to the simulated M-H loops of the free and pinned layers on the NiTiNb substrates with different predeformation and phases of the substrates, respectively. (i) Schematic diagram of the effect of strain on the magnetic layers of the SPVs on the SMA substrates. The red arrow shows the strain state of the substrates after annealing from 375 K. The blue one shows the strain state after cooling down from 200 K. M0, MM and MA illustrate the change of magnetization under different strain states of the substrates. (j) Schematic diagram of spin scattering in an antiparallel SPV under different strain states.


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Scheme 1. (Wang M. X. et al.)



Figure 1. (Wang M. X. et al.) 0.5

DSC (mW/mg)

0.4

NiTi from 300 K to 400 K

(a)

0.3 0.2

Heating

0.1 0.0 280

300

320

340

360

380

400

Temperature (K) 0.5 0.4

DSC (mW/mg)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

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NiTiNb from 400 K to 200 K NiTiNb from 200 K to 400 K

(b)

0.3 0.2

Heating

0.1 0.0

Cooling

-0.1 200

240

280

320

Temperature (K)


360

400

Page 29 of 34


(a)

 (%)

1.5 1.0 0.5 0.0 0%

2%

4%

6%

Pre-strained stress 1.2

(b) 0.8

 (%)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


0% 2% 4% 6%

0.4

0.0 A1

M1

A2

Phase


M2


Figure 3. (Wang M. X. et al.)

(b)

(c) 40

C (Oe)

C (Oe)

60 NiTi-Free Layer 40 20

NiTi-Pinned Layer

20

0

0 0%

2%

4%

6%

0%

Pre-strained stress 0% 2% 4% 6%

20 0

NiTiNb-Free Layer

4%

6%

(e) 20 0

0% NiTiNb-Pinned Layer 2% 4% 6%

-20

-20 -40

2%

Pre-strained stress

C (Oe)

(d)

C (Oe)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

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A-1

M-1 A-2 Phase

M-2

-40

A-1


M-1 A-2 Phase

M-2

Page 31 of 34


0% 2% 4% 6%

160 150

Hc (Oe)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Cooling process

140 130 120 110

200

220

240

260

280

Temperaterature (K)


300




Page 32 of 34

Page 33 of 34


(a) NiTi-free layer

1.0

M/Ms

0.0

0.0

-0.5

0% 2% 4% 6%

-1.0 -40

-20

0

(c) 2% NiTiNb-free layer

0% 2% 4% 6%

-1.0

40

-40

-20

0

1.0

0.5

M/Ms

M/Ms

M/Ms

Strain-free state Martensite phase Austenite phase

-1.0 -30

-20

-10

0

H/Hk

10

20

-0.5

(f) 2% NiTiNb-pinned layer

-30

1.0

-20

-10

0

H/Hk

10

20

(g) 4% NiTiNb-pinned layer

M/Ms Strain-free state Martensite phase Austenite phase

-1.0 -20

-10

0

H/Hk

10

20

30

-10

0

H/Hk

10

20

30

(h) 6% NiTiNb-pinned layer

0.0

-0.5 Strain-free state Martensite phase Austenite phase

-1.0 -30

-20

0.5

-0.5

-0.5

-30

1.0

0.0

0.0


-1.0

30

0.5

0.5

-30


-1.0

30

(e) 6% NiTiNb-free layer

0.0

-0.5

-0.5

40

0.5

0.0

0.0

20

H/Hk

(d) 4% NiTiNb-free layer

1.0

0.5

1.0

20

H/Hk

-0.5

M/Ms

1.0

NiTi-pinned layer

(b)

0.5

M/Ms

0.5

M/Ms

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


-20

-10

0

H/Hk

10

20

30



-1.0 -30

-20

-10

0

H/Hk

10

20

30


Table of Contents Graphic. (Wang M. X. et al.)


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Strain-Controlled Giant Magnetoresistance in Spin Valves Grown on

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