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STUDIES OF CURRENT-PERPENDICULAR-TO-PLANE
MAGNET ORESISTANCE (OPP-MR) AND CURRENT-
INDUCED MAGNETIZATION SWITCHING (CIMS)

presented by

HUSEYIN KURT

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STUDIES OF CURRENT-PERPENDICULAR—TO-PLANE
MAGNETORESISTANCE (CPP-MR) AND CURRENT-INDUCED
MAGNETIZATION SWITCHING (CIMS)

By

Huseyin Kurt

A DISSERTATION
Submitted to
Michigan State University
in partial fulﬁllment of the requirements
for the Degree of

DOCTOR OF PHILOSOPHY
Department of Physics and Astronomy

2005

ABSTRACT
STUDIES OF CURRENT-PERPENDICULAR—TO-PLANE
MAGNETORESISTANCE (OPP-MR) AND CURRENT—INDUCED
MAGNETIZATION SWITCHING (CIMS)
By

Huseyin Kurt

We present two CPP-MR studies of Spin-valves based upon ferromagnetic/ non-
magnetic/ ferromagnetic (F/ N/ F) trilayers.

We measure the Spin-diffusion lengths of N = Pd, Pt, and Au at 4.2K, and both
the speciﬁc resistances (sample area A times resistance R) and spin-memory-loss of
N /Cu interfaces. Pd, Pt and Au are of Special device interest because they give
perpendicular anisotropy when sandwiching very thin Co layers. Comparing our
spin—memory-loss data at Pd/ Cu and Pt/ Cu interfaces with older data for Nb/Cu
and W/ Cu gives insight into the importance of spin-orbit coupling in producing such
loss.

We reproduce and extend prior studies by Rid of ‘magnetic activity’ at the inter-
face of Co and N-metals (or combinations of N-metals), when the other side of the
N-metal contacts a superconductor (S). Our data suggest that magnetic activity may
require strong spin-ﬂipping at the N / S interface.

We present ﬁve studies of a new phenomenon, CIMS, in F1 / N / F2 trilayers, with
F1 a thick ‘polarizing’ layer and F2 a thin ‘Switching’ layer.

In all prior studies of CIMS, positive current caused the magnetization of F2 to
switch from parallel (P) to anti-parallel (AP) to that of F1— ‘normal’ switching. By

judicious addition of impurities to F-metals, we are able to controllably produce both

‘normal’ and ‘inverse’ switching— where positive current switches the magnetization
of F2 from AP to P to that of F1. In the samples studied, whether the switching is
normal or inverse is set by the ‘net polarization’ produced by F1 and is independent
of the preperties of F2. AS scattering in the bulk of F1 and F2 is essential to pro-
ducing our results, these results cannot be described by ballistic models, which allow
scattering only at interfaces.

Most CIMS experiments use Cu as the N-layer due to its low resistivity and long
spin-diffusion length. We Show that Ag and Au have low enough resistivities and long
enough spin-diffusion lengths to be useful alternatives to Cu for some devices.

While most technical applications of CIMS require low switching currents, some,
like read-heads, require high switching currents. We Show that use of a synthetic
antiferromagnet can increase the switching current.

Manschot et al. recently predicted that the positive critical current for switching
from P to AP could be reduced by up to a factor of ﬁve by using asymmetric current
leads. In magnetically uncoupled samples, we ﬁnd that highly asymmetric current
leads do not signiﬁcantly reduce the switching current.

A CIMS equation given by Katine et a]. predicts that lowering the demagnetiza-
tion ﬁeld should reduce the switching current. To test this prediction, we compare
switching currents for Co/ Au/ Co(t) / Au nanopillars with t = 1 to 4 nm (where the
easy axis should be normal to the layer planes at least for t=1 and 2 nm) with those
for 00/ Cu/ Co(t) / Au nanopillars (where the easy axis should be in the layer planes).

We do not ﬁnd signiﬁcant differences in switching currents for the two systems.

To my parents

iv

ACKNOWLEDGEMENTS

I would like to thank my thesis advisor professor Jack Bass for his professional guid-
ance and assistance. I would also like to thank professor William Pratt Jr. for his
assistance in my experiments, technical teachings and instrument designing. I am
honored to work with both Dr. Bass and Dr. Pratt.

My special thanks to Dr. Reza Loloee for his help and technical support. I owe
most of my experimental and technical expertise in thin ﬁlms and vacuum systems
to Reza.

I would like to thank my friends and colleagues Antonio Zambano, Ion Moraru,
Khalid Eid, Michael Crosser, Mustafa Al-Haj Darwish, and Sergei Urazhdin for their
friendships, help and suggestions. Each of them made my life in the basement of BPS
building enjoyable and memorable.

I want to thank Thomas Palazzolo for setting up a machine shop class, in which
I learned the basics of machining and designing. I also want to thank the machine
shop crew Thomas Hudson and James Muns for helping me design and make some
instruments I built.

Finally I would like to acknowledge the ﬁnancial support of MSU CF MR, CSM,
Keck Microfabrication Facility, NSF Grants No. DMR 02-02476, No. 98-09688 and

Seagate Technology for the research presented in this thesis.

Contents

1 Introduction 1
1.1 Homogenous metals ............................ 1
1.2 Giant Magnetoresistance (GMR) in Magnetic Multilayers ....... 2

1.2.1 Producing P and AP states ................... 4
1.2.2 Current directions in multilayers ................. 5
1.2.3 GMR device applications ..................... 8
1.3 Spin Transfer Torque & Current-Induced Magnetization Switching . . 9
1.4 This thesis ................................. 14

2 Theory 16
2.1 Origin of GMR .............................. 16
2.2 Spin-polarized transport ......................... 17

2.2.1 Normal MR ............................ 18

2.2.2 Inverse MR ............................ 20

2.3 Two-Current Series Resistor (2CSR) model ............... 20
2.3.1 Analysis of a F1 / N / F2 trilayer in AP and P states using 2CSR

model ............................... 22

2.3.2 Validity of 2CSR model ..................... 24

2.4 Valet-Fert (VF) model .......................... 25

2.4.1 Overview ............................. 25

2.4.2 VF equations for an F/ N multilayer ............... 26

2.4.3 Finite spin-diﬁ’usion-length effects in a trilayer ......... 27

2.5 Real Fermi Surface models of CPP-MR ................. 27

2.6 Spin ”ﬁansfer Torque & Current-Induced Magnetization Switching . . 28

2.6.1 Overview ............................. 28

2.6.2 Ballistic transport equations ................... 30

2.6.3 Diffusive transport equations ................... 33

2.6.4 Summary ............................. 36

3 Sample Fabrication and Measurement 37
3.1 Overview .................................. 37
3.2 Making and measuring Nb—sandwiched magnetic multilayers for CPP-

MR studies ................................ 38
3.2.1 Preparation ............................ 38
3.2.2 Sputtering ............................. 38
3.3 Measuring the Resistances of CPP samples ............... 41

vi

3.3.1 Pinning .............................. 41

3.3.2 Resistance measurement at 4.2K ................. 42
3.4 Making and measuring nanopillars for CIMS studies .......... 44
3.4.1 Preparation of substrates using photolithography and evaporation 45
3.4.2 EBL for the multilayer and sputtering ............. 48
3.4.3 EBL for nanopillars and evaporation of A1 ........... 52
3.4.4 Ion-milling and Insulation .................... 55
3.4.5 Low angle Ion-milling ....................... 59
3.4.6 Lift-off and wet etch ....................... 60
3.4.7 Top contacts ........................... 61
3.4.8 Measuring the Resistances of Nanopillars ............ 62
4 Spin-memory-loss in sputtered Pd,Pt and Au and at their interfaces
with Cu at 4.2K 66
4.1 Introduction ................................ 66
4.2 Spin-memory-loss detection technique .................. 68
4.3 Experiment ................................ 71
4.3.1 Examples of AR(H) with X inserts ............... 72
4.4 Spin-Memory-Loss at 4.2K in Sputtered Pd, Pt, and at Pd/Cu and
Pt/ Cu interfaces ............................. 74
4.5 Spin-Memory-Loss and CPP-MR at 4.2K in Sputtered Multilayers with
Au ..................................... 80
4.6 Summary ................................. 84
4.7 Comparison of measured ABS with theory ............... 85
5 Changes in magnetic scattering anisotrOpy at a ferromagnetic/ su-
perconducting interface 87
5.0.1 Overview ............................. 87
5.0.2 Introduction ............................ 88
5.0.3 Experiment and Results ..................... 89
5.0.4 Discussion ............................. 93
5.0.5 Summary ............................. 97
6 Controlled normal and inverse CIMS and MR in magnetic nanopil-
lars 98
6.1 Introduction ................................ 98
6.2 Experiment and Results ......................... 100
6.3 Summary ................................. 105
7 CIMS in permalloy-based nanopillars with Cu, Ag, and Au spacers 107
7.1 Introduction ................................ 107
7.2 Experiment and Results ......................... 107

vii

8 Reduction of spin polarization and spin transfer by synthetic anti-

ferromagnets 112
8.1 Introduction ................................ 112
8.2 Experiment and Results ......................... 113
8.3 Summary ................................. 116
9 Effect of asymmetric leads on the switching currents of nanopillarsll7
9.1 Introduction ................................ 117
9.2 Experiment ................................ 118
9.3 Summary ................................. 123

10 Effect of lower demagnetizing ﬁelds on the switching currents of
nanopillars 125
10.1 Introduction ................................ 125
10.1.1 Perpendicular anisotropy in Au / Co/ Au systems ........ 126
10.2 Experiment ................................ 127
10.3 Summary ................................. 132
11 Summary & Conclusions 134
A Determining the resistivities of metals and alloys 138
A.1 Overview .................................. 138
A2 Technique ................................. 139
A21 Samples and Measurement .................... 141
A.3 Results ................................... 141
A.3.1 Discussion ............................. 143
LIST OF REFERENCES 145

viii

List of Figures

1.1

1.2

1.3

1.4

1.5

2.1

2.2

2.3

2.4

2.5

(a) Antiferromagnetically coupled F / N multilayer, where the local mag-
netizations lie antiparallel in the absence of magnetic ﬁeld H (b) At
high H the magnetic moments are aligned parallel along H. .....
Normalized resistance vs. H for Fe/ Cr multilayers from Baibich et al.
The resistance decreases as the external ﬁeld increases. ........
Different current ﬂow directions for multilayers, H is applied in-plane.
CIP current ﬂows through area A = wt and CPP current ﬂows through
area A 2 wt .................................
A representative cartoon of (a) longitudinal and (b) perpendicular
recording. Perpendicular recording allows much higher areal densities.
A schematic phase diagram of a magnetically uncoupled nanopillar.
The regions of hysteretic switching (Hys), parallel (P), antiparallel(AP)
and ‘between’(B) are labeled. ......................

A representative diagram of density of states that are available to elec-
trons in (a) a normal metal and (b) a ferromagnetic metal whose spin-
up and spin-down states are not equally ﬁlled giving a non-zero net
magnetization ................................
Normal MR: A schematic representation of spin-dependent scatter-
ing in a F/N/F trilayer with p} < p)... (a) In the AP state the net
scattering in both spin channels is the same. (b) In the P state the
net scattering in the down-channel is greater than the one in the up—
channel. Therefore the resistance in the AP state is higher. Breaks in
the dashed lines indicate strong scattering. ...............
Inverse MR: A schematic representation of spin-dependent scattering
in a F1 /N/F2 trilayer with p}, < p}, and p112 > pin. (a) In the AP
state the net scattering in the down-channel is greater than the one in
the up-channel. (b) In the P state the net scattering in both channels
is the same. Therefore the resistance in the P state is higher. Breaks
in the dashed lines indicate strong scattering. .............
A schematic drawing of a F1 / N/ F2 trilayer showing two channels of
electron ﬂow in a) AP alignment b) P alignment. Block arrows repre-
sent the direction of local magnetization in each ferromagnetic layer.
p1,“) is the resistivity for electrons with moment up (down) .......
A representative cartoon of an elliptical nanopillar. ..........

ix

12

17

19

19

22

2.6

3.1
3.2

3.3

3.4
3.5

3.6
3.7
3.8
3.9

3.10

3.11
3.12
3.13

3.14

3.15

3.16

Spin-transfer torque acting on the magnetization of a 1 nm thick nano-
magnet as a function of 0 .........................

A picture of a rotary CPP mask and a CPP sample ..........
(a) A CPP sample with lumps of indium on Nb contacts. (b) Center
of a CPP sample .............................
A picture of the photomask. The open areas have the shape of the
leads and the rest of the mask is coated with Or to block UV light.

A cartoon of a good undercut after developing ..............
A picture of a 3" wafer after lift-off, which has 21 0.5” x 0.5” square
substrates ..................................
A single substrate (left) and its center (right) with the gold leads evap-
orated ....................................
The DesignCAD drawing for the shape of the multilayer. .......
An opening for the multilayer created by EBL .............
(a) A cross-sectional cartoon of a multilayer. (b) An SEM image of a
multilayer after lift-off. ..........................
(a) An optical image of 5 elliptical Openings fabricated using EBL, the
rest of the substrate is covered with e-beam resist. Since the dimensions
of ellipses are below optical resolution we only see some light coming
out of these holes. Scale as in Fig. 3.9. (b) A cartoon of the cross-
section of multilayer + A1 mask (c) An SEM image of an elliptical Al
mask (top view). .............................
(a) An SEM image of a Au pillar. (b) An SEM image of a Ag pillar. .
The DesignCAD drawing of a window for ion-milling. .........
A special sample holder for ion-milling. It has a heat sink to remove
heat from the substrate and a magnetic disk to hold it using a magnetic
arm. ....................................
(a) An opening over the multilayer for ion-milling, created by EBL. (b)
A cartoon of the cross-section of a magnetically uncoupled nanopillar.
(c) A cartoon of the cross-section of a magnetically uncoupled nanopil-
lar insulated by SiO evaporation. ....................
A special sample holder for low-angle ion-milling with a circular hole
in the middle and a nearly vertical (3°) slit for the substrate.(a) Side

36
42

47
48

49
50
51
52

53

54
56
57

59

view. (b)Top view. (c) A representative cartoon of low-angle ion-milling. 61

(a) An optical image after SiO evaporation and lift-off. Scale as in Fig.
3.9. (b) A cross-sectional cartoon after SiO evaporation lift-off. (c) A
cross-sectional cartoon after wet etch with KOH .............

3.17 The DesignCAD drawing for top contacts. The outer rectangle is ex-

posed to less e—beam than the inner one to get a good undercut. . . .

3.18 (a) Openings for 5 electrodes for top contacts created by EBL. (b)

An optical image showing the multilayer, SiO insulating layer and 5
electrodes for top contacts. (c) A cartoon of a cross-section of ﬁnished
nanopillar device. .............................

61

63

63

3.19 (a) The sample is mounted on the quick-dipper and leads soldered with
indium. (b) Quick dipper is inserted into the electromagnet through
the bored pole pieces and the sample is positioned at the center of the
electromagnet ................................

3.20 A diagram of resistance measurement using lock-in detection ......

3.21 Differential resistance (dV/dI) vs. H of a Py/Cu/Py nanopillar while
I=0 (on the left). dV/dI vs. the dc current (I) while H = 0 for the
same sample. ...............................

4.1 Examples of AR vs. H for Py-based EBSVs (a) with no insert X, (b)
X=Pd(10), (c) X=Pt(10) and (d) X=Au(10) ..............
4.2 ARs vs. H for some EBSV samples with X=[Pd(3)/Cu(3)]~ inserts. .
4.3 ABS vs. H for some EBSV samples with X=[Pt(3)/Cu(3)]~ inserts. .
4.4 AAR vs. tpd or tpy for Pd (ﬁlled triangles and solid curve), Pt (ﬁlled
diamonds and solid curve), Nb (open triangles and dashed curve), and
W (open diamonds and dashed curve). The dotted curves represent
the behaviors of AAR expected for no spin-memory-loss in either the
bulk Pd or Pt or at the Pd/ Cu or Pt/ Cu interfaces ...........
4.5 AR(AP) vs. N for [Pd/Cu]~ and [Pt/Cu]N inserts compared with
[Nb/Cu]N and [W/Cu]N. The symbols are the same as in Fig. 4.4.
The lines are least square ﬁts. ......................
4.6 AAR vs. N for [Pd/Cu]N, [Pt/Cu]~, [Nb/Cu]N, and [W/Cu]N inserts.
Note the similarity of behaviors of [Pd/Cub; and [Nb/Cub, and of
[Pt/Cu]N and [W/Cu]N. The symbols are the same as in Fig. 4.4.
The lines are least square ﬁts. ......................
4.7 AAR vs. tN for N = Au (ﬁlled circles and thick solid curve), V (dot-
ted curve), Nb (dashed curve), Pd (broken curve), and Pt (thin solid
curve). The datum at t Au = 15 nm is only a lower bound, as its AR(H)
curve did not give a ﬂat AR(AP) region. That point is not included in
the ﬁt to the Au data. ..........................
4.8 AR(AP) vs N for [Au/ Cu] N inserts compared with [V/ Cu] N , [Nb/Cu]N,
[Pd/ Cu] N, and [Pt/ Cu] N. The symbols and lines are the same as in
Fig. 4.7 ...................................

4.9 AAR vs. Nfor [Au/Cu]N, [V/Cu]N, [Nb/Cu]~, [Pd/Cu]N, and [Pt/Cu]N

inserts. The symbols and lines are the same as in Fig. 4.7. AAR for
the N = 15 datum dropped by a factor of ten over three successive
measurements. We show the average and ﬂuctuations, but neglect this
point in our ﬁt. In contrast, repeated measurements of the N = 10
and 20 data remained stable to within the uncertainty bars shown. . .
4.10 AR(H) vs. H for representative Co/Ag, Co/Au, and Co/Cu EBSVs
with tea = 6 nm and tN = 10 nm for N = Ag, Au, or Cu. The data
around H = 0 overlap for minor and major loops; the former extend
up to about + 250 Oe, the latter to + 1500 Oe. ............

xi

64
65

66

74
75
76

‘78

79

80

82

83

84

5.1

5.2

5.3

5.4

6.1

6.2

6.3

7.1

7.2

8.1

8.2

8.3

ABS for Nb/X / Nb sandwiches with X =Co(20), Ru(4)/Co(20)/Ru(4),
Au(6)/Co(20)/Au(6), and FeMn(2)/Co(20)/FeMn(2). Thicknesses are
in nanometers. Inset: The circular sample with bottom (V-, I-) and top
(V+, I+) crossed Nb strips. Current (I) and voltage (V) connections
are indicated. A is the area through which the current flows at 4.2K. 91
AR vs. H for Nbl50/Cu1/Py24/Cu20/Ru2/CoZ/X/Nb150 with no
X, X =Ru(2) and X =Au(2). The AP states are labeled for each case. 93
AAR for samples of the form Nb(150)/Cu(1)/Py(24)/Cu(20)/Ru(2)/
Co(2)/X/Nb(150) for a variety of X, including no X. Samples with
italic labels leave "yea/X5 = 0, whereas samples with bold labels give
l7Co/XSl > O ................................. 95
AAR vs. N for inserts X =[Ru(3) /Nb(3)] N, X =[Au(3) /Nb(3)] N and
X =[Cu(3) / Nb(3)] N in exchange-biased spin valves of the form Nb(150) /
Cu(10)/FeMn(8)/Py(24)/Cu(10)/X/Cu(10)/Py(24)/Cu(10)/Nb(150). 97

Py(24)/Cu(10)/Py(6) data at 295K (top) and 4.2K (bottom) showing
normal MR (dV/dI vs. H at I = 0) in the insets and normal CIMS for
dV/dI vs. I in the main ﬁgures at H = 0 0e for 295K and at H = 20

Oe for 4.2K. In all ﬁgures, I > 0 represents electron ﬂow from F2 to F1.103
Ni(Cr)(20)/Cu(20)/Py(10) data at 295K (top) and 4.2K (bottom) show-

ing inverse MR (dV/dI vs. H at I = 0) in the insets and inverse CIMS

for dV/dI vs. I at H = 0 in the main ﬁgures. ............. 104
Py(24)/Cu(10)/Ni(Cr)(4) data at 4.2K showing inverse MR (dV/dI

vs. H at I = 0) in the inset and normal CIMS for dV/dI vs. I at

H = 0 in the main ﬁgure. ........................ 104

MR (top) and CIMS (bottom) at 295K for Py/ N / Py nanopillars with
N =Cu (left), Ag (middle) and Au (right). Layer thicknesses are in run. 109
MR (top) and CIMS (bottom) plots at 4.2K for the samples in Fig.
7.1. Layer thicknesses are in nm. .................... 110

A cross-sectional cartoon of the structure of a magnetically uncoupled
nanopillar with a synthetic antiferromagnet (SAF). All thicknesses are
in run. ................................... 115
MR curve for a Co(20)/Ru(0.6)/Co(4)/Cu(10)/Co(4) nanopillar at
zero current at room temperature. The highest resistance antipar-
allel state represents antiparallel alignments of two Co(4) layers. The
parallel alignments of Co(4) layers represent a pseudo parallel state,
which gives a higher resistance than the complete parallel alignments
of all Go layers. .............................. 115
CIMS curve for the nanopillar in Fig. 8.2 at zero ﬁeld at room temper-
ature. The positive switching current is not seen because it is higher
than the limit (18 mA) of our current supply. ............. 116

xii

9.1

9.2

9.3

10.1

10.2

10.3

10.4

10.5

A.l
A.2

Cross-sectional cartoons of (a) X=Co(20) / Cu( 10) / Co(4), (b) A=Pt(20) / X
/B=AgSn(30), and (c) A=AgSn(30)/X/B=Pt(20) magnetically un-
coupled nanopillars. All thicknesses are in nm .............. 120
Examples of MR and CIMS curves for A=Pt(20) / X / B=AgSn(30) nanopil-
lars at room temperature. ........................ 122
Examples of MR and CIMS curves for A:Ag8n(30) / X / B=Pt(20) nanopil-
lars at room temperature. ........................ 123

The magnetizations of Au(10)[Co(1)/Au(10)]x5 (ﬁlled symbols) and
Au(10)[Co(2)/Au(10)] x5 (open symbols) samples as a function of an-
gular position 6 in a magnetic ﬁeld of H = 1000 Oe at room tempera-
ture. 0 is the angle between H and the sample plane. ........ 128
The magnetization of a Au[Co(1)/Au(10)] x5 ﬁlm as a function of an-
gular position 0 in external magnetic ﬁelds of H = 500,1000 and 5000
Oe at room temperature. 0 is the angle between H and the sample
plane. ................................... 129
A comparison between AAR/AAR(CO(4)) of mm2 Nb cross-stripped
CPP samples with Co(t) sandwiched between Au and Cu layers at 4.2K.130

MR vs. ta“ of Co(20)/Au(10)/Co(t)/Au(115) nanopillars at room
temperature. The MR increases as t increases up to 3-4 nm thick Co
and then saturates. ............................ 132
A comparison of AJ vs. t5? for Co(20)/Au(10)/Co(t)/Au(115) (ﬁlled
circles) and Co/Cu(10)/Co(t)/Au (open circles) nanopillars. ..... 133
The van der Pauw geometry for a thin ﬁlm metal ........... 140
The graph of correction factor f vs. R1/R2 ............... 140

xiii

List of Tables

3.1
3.2

4.1
4.2

4.3

5.1

6.1

7.1
7.2

8.1

9.1
9.2

A.1

A2
A3
A.4

Typical sputtering rates for Nb, Cu, Py, Au and Co ........... 41
Argon ion etch rates for the metals we etch at an ion current density
of 1 mA/cm2 and an ion energy of 500 eV. ............... 58

Typical sputter deposition rates of Nb, Py, FeMn, Cu, Pd, Pt and Au. 73
A table of resistivities and spin-diffusion lengths of Pd, Pt and Au, and
the spin-memory—losses at Pd/ Cu, Pt/ Cu and Au/ Cu and the speciﬁc
resistances of these interfaces at 4.2K ................... 85
A comparison between the experimental and theoretical values of ABS
at 4.2K. Calculations agree with the experimental values if the lattice
parameters of the metals constituting the interface are close. ..... 87

AARs and the number of samples used to determine the error bars for
inserts X in Fig. 5.3 ........................... 96

Scattering anisotropies (+=positive, -=negative) of F1, F1 / N, and
their net anisotropy F1 (net), those of F2, N / F2, and their net anisotropy

F2(net), and the observed MR and CIMS (+=normal, -=inverse). . . 103
Approximate sputter deposition rates of Py, Cu, Ag and Au. ..... 109
Average values of R, AR, MR(%), AI, and H, at 295K and at 4.2K

of Py/N/Py with N=Cu, Ag and Au ................... 110
The approximate sputter deposition rates for Au, Co, Cu and Ru. . . 114

Approximate sputter deposition rates of Cu, Pt, Au, Co and AgSn(5%).121
Average values of T1" = r‘B/rz, R(AP), MR(%), AI 2 1+ — I", I"
and I + for our samples of the form Cu(80) / A/ X/ B / Au(110) with X =
Co(20 )/Cu(10)/Co(4), and A, B = Pt(20), AgSn(30), or nothing. The
AlHajDarwish samples are similar to our ones without Pt and AgSn,

except that the thickness of the ﬁxed Co layer is 30 nm ......... 125
A comparison between the CIP and CPP resistivities of Co and CoZr

at 4.2 K ................................... 140
The resistivities of sputtered gold (Au) samples ............ 143
The resistivities of sputtered platinum (Pt) samples .......... 143
The resistivities of sputtered palladium (Pd) samples ......... 143

xiv

A5 The resistivities and superconducting transition temperatures of sput-

tered niobium (Nb) samples ....................... 142
A6 The resistivities of sputtered silver-tin (AgogsSnoﬂs) samples ..... 143
A7 A comparison between the Aps of pure Au, Pt, Pd, Nb and Ag and

their sputtered counterparts ........................ 144

XV

Chapter 1

Introduction

1 .1 Homogenous metals

Magnetoresistance (MR) is deﬁned as the percentage change in the resistance of a
conductor in response to an external magnetic ﬁeld H. The MR for a homogenous

conductor is usually deﬁned as;

R(H) - 12(0)
13(0)

 

MR(%) = [ ] x 100, (1.1)

where R(H) is the resistance in H and R(0) is the resistance in zero magnetic ﬁeld.

For normal (N) metals or alloys, the external magnetic ﬁeld exerts a Lorentz force
on the conduction electrons that gives rise to a change in resistance, in which the
resistance increases with increasing H [1, 2]. This effect is known as Lorentz MR.
Lorentz MR increases with H 2 at small ﬁelds (where the electrons are scattered many
times before completing a cyclotron orbit) and can level off or continue to grow at high
ﬁelds (where electrons complete several cyclotron orbits between scattering events).
For N-metals or alloys the MR is typically less than 0.1% at room temperature at
H = IT [3]. However, for high purity compensated metals, where the numbers of
electrons and holes are the same, the Lorentz MR increases indeﬁnitely with H 2 at

very low temperatures[4]. Since it is large only in very pure metals and at very low

temperatures, it is impractical to use in devices.

Ferromagnetic (F) metals or alloys, on the other hand, display a different kind
of MR, which depends on the angle between the current and H [5]. This effect is
called anisotropic MR (AMR)and arises from the spin-orbit coupling of electrons in
F-metals [6]. The AMR for a thin ﬁlm of the alloy permalloy (Py=Ni0_8Fe0_2) is
relatively large (~ 2%) at room temperature in a small ﬁeld of ~ 10 Oe. Due to this
2% MR, Py ﬁlms were used in the read heads of hard drives until the introduction of

“Giant Magnetoresistive” read heads in 1997 [7].

1.2 Giant Magnetoresistance (GMR) in Magnetic
Multilayers

GMR was discovered in 1988 in epitaxially grown, antiferromagnetically coupled,
Fe/Cr/ Fe trilayers [8] and Fe/ Cr multilayers [9], in which the magnetizations of the
Fe layers were aligned anti-parallel (AP)(Fig. 1.1(a)) at zero ﬁeld due to the anti-
ferromagnetic exchange coupling [10]. As shown in Fig. 1.2, the resistances of such
Fe/ Cr samples decrease as the moments of the Fe layers aligned parallel (P) (see Fig.
1.1(b)) upon applying a high enough ﬁeld. A dramatic reduction (up to 50% at 4.2K)
in the resistances of these samples was observed. This big reduction in the resistances
of the Fe/ Cr multilayers led to the name “Giant” Magnetoresistance.

GMR depends on the relative orientation of the magnetizations of the F -metals
separated by an N-metal and results from spin polarization in the bulk F layers and

at the F/ N interfaces [11, 12]. The MR for F/ N multilayers or F/ N/ F trilayers is

(b)

   

“H

Figure 1.1: (a) Antiferromagnetically coupled F/ N multilayer, where the local mag-
netizations lie antiparallel in the absence of magnetic ﬁeld H (b) At high H the
magnetic moments are aligned parallel along H.

R/R(H=0)

 

(Fe 3011/0“),0

 

 

 

0.5» HS
1 l l l * 1 J L L + H «6)
«40 m -20 -10 0 10 20 30 4o

 

Figure 1.2: The resistance vs. H graph for Fe/Cr multilayers by Baibich et a1. [9].
The resistance decreases as the external ﬁeld increases.

usually deﬁned as;

R(AP) — R(P)
R(P)

 

MR(%) = [ ] x 100, (1.2)

where R(AP) is the resistance in the anti-parallel (AP) state, and R(P) is the resis-
tance in the parallel (P) state. GMR is usually positive (normal) i.e. the resistance

is higher in the AP state, but, as we shall see below, there are cases in which GMR

can be negative (inverse). Theoretical explanations for the sign of GMR will be given

in Chapter 2.

1.2.1 Producing P and AP states

P states can be easily achieved by increasing the external magnetic ﬁeld H to above
the saturation ﬁeld(s), H,, of the F-metals or alloys. AP states however are harder
to achieve for wide ranges of thicknesses of both F and N-metals. For AP states to
be achieved, either adjacent F-layers must be antiferromagnetically coupled or their

switching ﬁelds must be different. This can be done in 3 ways:

1. Using antiferromagnetically coupled F-metals. The exchange coupling between
two F-metals separated by an N-metal oscillates with increasing N-metal thick-
nesses between ferromagnetic coupling (which aligns the magnetizations of the
two F-metals in the same direction) and antiferromagnetic coupling (which
aligns the magnetizations in opposite directions). The strength of the cou-
pling decreases with increasing thickness of N and becomes negligible beyond a

certain thickness [13, 14].

2. Using an exchange-biased spin valve (EBSV) in which one of the F-metals is in
contact with an antiferromagnetic (AF) metal. A simple EBSV has the form
AF / F1 / N / F2, where F1 and F2 can be same or different and N is thick enough
so that exchange coupling between F1 and F2 is negligible. Heating the sample
to above Néel temperature of the AF-metal and cooling in the presence of a
high enough external ﬁeld H can pin F1 due to the strong exchange interaction

between AF and F1 . The hysteresis loop of the pinned F1 shifts away from

H = 0, changing its switching ﬁeld dramatically [15, 16]. Therefore F1 and
F2 can have very different switching ﬁelds. This conﬁguration enables the free
layer to switch back and forth over many cycles while the pinned layer stays

ﬁxed. For this reason, EBSVs are used in read heads of hard drives.

3. Using a hybrid spin valve, with two F-metals separated by a thick N-metal layer,

one can achieve an AP state in two ways:

0 Using different F-metals with different switching ﬁelds.

0 Using one type of F -metal with different thicknesses, giving different switch-

ing ﬁelds.

1.2.2 Current directions in multilayers

There are two main geometries for current ﬂow direction used in GMR experiments.

0 Current-in-plane (CIP), in which the current ﬂows parallel to the plane of the

layers as in Fig. 1.3. The MR for this geometry is called “CIP-MR”.

o Current-perpendicular-to—plane (CPP) geometry, in which the same current
ﬂows through every layer as in Fig. 1.3. The MR for this geometry is called

“CPP-MR”.

Both in the CIP—MR and CPP-MR experiments the external magnetic ﬁeld is usually
applied in the plane of the samples as shown in Fig. 1.3.
Consider a metallic thin ﬁlm with dimensions of l = 1 cm, w = 1 mm and t = 100

nm and resistivity of 200 an. If the current ﬂows through the area wt, the CIP

I ICPP

N“
If” \ICIP

Figure 1.3: Different current flow directions for multilayers, H is applied in-plane.
CIP current ﬂows through area A = wt and CPP current ﬂows through area A = wl.

  

' H

resistance is;

l 107nm
= _ = Q . —— = .
RCIP pwt 20011 m 106nm x 100nm 209 (1 3)

and the CPP resistance is;

t 100nm
= _ = Q . — = 1.
RCPP pwl 200n m 107nm x106nm 2nQ ( 4)

The CIP resistance of this ﬁlm is much larger than its CPP resistance. Resistances of
ohms can easily be measured using a current source and either a lock-in ampliﬁer or
a digital voltmeter. On the other hand, measuring resistances of nanoohms requires a
sophisticated measurement setup including a superconducting quantum interference
device (SQUID) [17] or a low temperature ac ampliﬁer [18]. The CPP resistance of a
metallic sample can be on the order of ohms if the sample area is reduced to smaller
than 0.01pm2, which can be done using lithography (see Section 3.2) or making F/ N
multilayer nanowires [19, 20, 21] . The early studies of GMR were done using the
CIP geometry [9, 8], because CIP resistances are easier to measure for thin ﬁlms. For

the same reason, today’s GMR devices use CIP geometry.

Theoretically CIP-MR is more complex than CPP-MR due to the complications in
the distribution of current through the layers and there are more parameters involved,
including especially the mean free path of the 1 -metal [11]. GMR is only possible
if the electrons are able to distinguish between the P and AP states of the F-layers,
which requires the electrons to travel between different F—layers without losing their
spin polarization. Thus, for CIP-MR the thicknesses of the N-layers separating the
F -layers should not be much longer than their respective mean free paths. Current
flow in the CIP geometry is non-uniform and also complicated, i.e. more current
passes through lower resistivity layers. Since our research doesn’t involve CIP-MR
we won’t go into further details, referring the interested reader to Ref. [11].

Since the same current passes through every layer in the CPP geometry, there turn
out to be fewer parameters involved, making it easier to analyze the experimental data
and derive the underlying physics [22, 23, 12, 24]. The ﬁrst CPP-MR experiments
were done by our group by using crossed superconducting Nb strips as leads to get a
uniform current ﬂow through a sample [25]. All measurements in this thesis are done

using CPP currents. We can divide the theories of CPP-MR into two categories.

1. Free-Electron Based Models: Semi-classical models neglecting Fermi surface

effects in F and N layers and at F/ N interfaces.

2. Realistic Fermi Surface Models, which take the different Fermi surfaces of the

constituents of an F/ N system into account.

These theoretical models will be reviewed in Chapter 2.

1.2.3 GMR device applications

The discovery of GMR stimulated great interest in spin-dependent electronics and
physics, nowadays called “spintronics”, which has the potential to change the present
technology based on only charge transport.

Different uses of GMR sensors include determining the speed and position of
electric motor shafts, general ﬁeld detection in implantable medical devices, and wheel
speed sensing for ABS brake applications [26]. But their biggest industrial application
is as the read heads of hard drives, where they opened a new era in the ﬁeld of high
density magnetic storage. In 1997 IBM introduced the ﬁrst GMR read head in hard
drives [27] with 6% MR which replaced the AMR (2% MR) read heads. Higher MR
of these ﬁrst GMR read heads allowed higher areal densities (3 Gbits/in2) of bits on
a hard disk. Further increases in the areal densities are achieved by improving the
MR of the GMR read heads.

Conventionally, the magnetic bits are recorded in the plane of the disk (Fig. 1.4
(a)). A new method of magnetic recording suggests storing the bits in perpendicular
magnetic domains which allows more dense packing of bits (see Fig. 1.4 (b)). In De-
cember 2004 Toshiba announced the ﬁrst commercial hard drive using perpendicular
recording with an areal density of 133 Gbits/in2 [28]. More research is under way to
increase the areal densities beyond 1 Tbits/in2 [29, 30, 28].

As the areal densities on a hard disk continue to increase, more sensitive read
heads are needed to detect weaker magnetic ﬁelds from smaller bits. Today’s modern

hard drives use CIP-MR read heads. It has been shown that CPP-MR could be

 

 

/ / / / j
[s—pu ]u<—s[n<—s[s—su L

 

(a) Longitudinal recording
/

/
N S S N S N N S [S
Iii. It I I H.
(b) Perpendicular recording

 

 

 

 

 

 

 

 

 

 

 

 

 

 

m—p:
:<—m

 

Figure 1.4: A representative cartoon of (a) longitudinal and (b) perpendicular record-
ing. Perpendicular recording allows much higher areal densities.

several times larger than the CIP-MR of the same structures [25, 31, 32], and thus
could be used for more sensitive read heads. However, for the CPP-MR devices with
standard metals or alloys to have big enough resistances for read heads, the area of

the devices should be smaller than 0.01pm2.

1.3 Spin Transfer Torque & Current-Induced Mag-
netization Switching

In 1996 Slonczewski [33] and Berger [34] independently predicted that a polarized
high-density CPP current could produce a spin—transfer—torque (STT) on the magne-
tization of a nanomagnet (a thin (1-10 nm thick) ferromagnet with S, 100 nm lateral
dimensions). Such STT acting on the magnetization of a nanomagnet is predicted
to produce a steady precession of magnetization upon applying a high enough cur-
rent in a high enough external magnetic ﬁeld, and a novel type of magnetization
switching upon applying a current higher than the threshold. Slonczewski focused on
switching of that layer’s magnetization and Berger focused upon excitation of coher-

ent magnons, but subsequent works showed that their models were almost equivalent

[35, 36]. For simplicity, we focus mostly on the Slonczewski analysis and outgrowths
thereof.

Slonczewski [33] proposed a trilayer system of F 1 /N /F 2, where F1 is a thick
ferromagnet with magnetization M1 (therefore its magnetization doesn’t ﬂip with
current), N is a nonmagnetic spacer thinner than its spin-diffusion length, and F2 is
a thin nanomagnet with magnetization M; with S, 100 nm lateral dimensions. When
electrons pass through F1 they get polarized, therefore carry angular momentum.
If M; is non-collinear with M1, the electrons polarized by F1 and transmitted into
F2 exert a torque (spin-transfer-torque or STT) on M; in an azimuthal direction
to align or anti-align M2 with M1. If a current greater than a threshold (~ 107’8
A/cm2) is applied, M2 ﬂips. For the current applied in the opposite direction, the
electrons transmitted into F1 cannot ﬂip M1 because F1 is thick, but the electrons
reﬂected from the F1 / N interface can ﬂip M2 when they are injected back into F2.
This gives an asymmetric switching because only M2 can be ﬂipped. Due to the
relative alignments of M1 and M2 (P or AP), the resistance of the trilayer system
changes as the alignments change from P to AP or vice versa, just like in the case of
GMR, but with a different switching mechanism, in which the relative alignments of
magnetizations are determined by the direction of the current ﬂow. This phenomenon
is now called “Current-Induced Magnetization Switching (CIMS)”.

In 1998, Tsoi et al. for the ﬁrst time, experimentally observed the effect of STT by
measuring the variations in the resistance of a Co/ Cu multilayer, induced by a high
current density z 108 A / cm2 injected into the multilayer through a mechanical point

contact [37]. The precession of magnetization by STT has opened a new ﬁeld of dc-

10

generated GHz Spin waves, which has potential uses in high-frequency communication
devices. We won’t go into details of the microwave studies, as they are irrelevant to
the present thesis, referring the interested reader to some appropriate papers [38, 39,
40,41,42,43]

In 1999, Sun reported the ﬁrst magnetic switching in a trilayer manganite junction
[44] and then Myers et al. reported the ﬁrst measurements showing a magnetization
reversal using lithographically deﬁned point contacts to a Co/Cu/Co trilayer [45].
In 2000, Katine et al. [46] reported the ﬁrst magnetization reversal in Co/Cu/Co
nanopillar devices fabricated by electron beam lithography. Later, several groups
fabricated nanopillars, reproducing and extending the results of Katine et al. (see
e.g. refs. [47] and [36]). At MSU we fabricate our nanopillars in such a way that both
dipolar and exchange couplings are minimal, which are referred to as ‘magnetically
uncoupled nanopillars’ in this thesis. The details of the fabrication of our nanopillars
are given in Chapter 3.

The discussion so far has neglected any applied magnetic ﬁeld, H. In the absence
of H, the magnetization of the switching layer F2 reverses at a positive current I j
from parallel (P) to that of F1 to AP to that of F1, and at a (not necessarily the
same magnitude) negative current I 37‘ from AP to P. A large H in either direction
tends to drive the F 1 / N / F2 nanopillar to the P state. Increasing-magnitude negative
H and negative I reinforce each other in driving the system to a P state. As shown
schematically in Fig. 1.5, for increasing-magnitude negative I, the switching ﬁelds
gradually decrease in magnitude until a certain value of I, after which the system

stays in the P state for all larger-magnitude I and all H. For increasing positive I,

11

in contrast, the current and ﬁeld tend to Oppose each other. For a given H, large
enough I should eventually drive the system to the AP state. But for large ﬁxed I ,
increasing H beyond a certain value is found to drive the system from the AP state
to a ‘between’ (B) state involving excitation of one or more magnons. Urazhdin et
a1. [48] showed that, in the vicinity of the AP to B transition, the system displays
telegraph noise, and a peak in dV/dI that is located approximately where the rates

for switching from the AP to B state and vice versa become equal.

 

 

 

-H Hys 4. +H

 

 

 

-I

Figure 1.5: A schematic phase diagram of a magnetically uncoupled nanopillar. The
regions of hysteretic switching (Hys), parallel (P), antiparallel(AP) and ‘between’(B)
are labeled.

A phase diagram of H vs. I is obtained by measuring MR at different currents
and measuring CIMS at diﬂerent external ﬁelds. This diagram allows one to easily
distinguish between the regions of no switching, switching and precession. In this
thesis we only studied the MR at zero current and CIMS at zero (or very low) ﬁeld.

All of the CIMS experiments in this thesis use the F1/N/F2 trilayer geometry
just described, in which F1 generates the polarized current that exerts a torque on
the magnetization of F2, thereby generating magnetic excitations. In 2001, Heide

[49, 50, 51] predicted that such excitations could be produced without the polariz-

12

ing ferromagnet, albeit only at larger exciting currents. Using a mechanical point
contact, Ji et al. [52] observed spin-wave excitations in single Co layers for electrons
ﬂowing only from the tip into the Co ﬁlm in the presence of very high (2-9 Tesla)
external ﬁelds applied perpendicular to the layers for currents ~ 5 x 109 A/cm2.
However, their critical currents were higher for thicker F-layers, inconsistent with
Heide’s model. Ji et a1. explained their excitations adapting Slonczewski’s model to
a single F-layer [38], deﬁning two regions in the F-layer, a region near the contact
where serves as the free layer and the rest of the layer which remains aligned with
the external ﬁeld and serves as the ﬁxed layer. Polianski and Brouwer [53] and Stiles
et al. [54] subsequently argued that an unpolarized current injected from an N-metal
into an F -metal can produce STT excitations for one polarity of current, provided
that the nonmagnetic contacts to the F-layer are asymmetric. Ozyilmaz et al. [55]
reported current-induced excitations in single Co layer nanopillars with asymmetric
(Cu(thin)/Co/Cu(thick)) leads for only one current polarity, currents ~ 109 A/cm2,
and ﬁelds > 1.5 T perpendicular to the sample plane. Similar currents produced no
excitations in nanopillars with symmetric leads. Excitations in single F-layers are
also outside the sc0pe of the present thesis.

We conclude this introduction by noting that CIMS has several potential techno-
logical applications, mainly in non-volatile magnetic random access memories (MRAMS).
Using CIMS devices in MRAMs has the advantage that they don’t require an external
magnetic ﬁeld for switching. However, the high critical switching currents of present
devices make them inappropriate for MRAMs due to self-heating. For CIMS devices

to be used in MRAMS, the switching currents should be reduced to ~ 105“6 A/cm2

13

[56]. Therefore much of the research on CIMS aims to reduce the switching currents.
Another possible technological application is to use nanopillars as read heads of hard
drives. Since the nanopillars are very small, the current used to measure the resistance
of a nanopillar might switch the magnetization of the free layer or at least generate
noise. This problem can be solved by increasing the currents needed for switching.
Chapter 8 discusses a means of increasing the switching currents by reducing the net

spin polarization of the polarizing layer.

1.4 This thesis

We brieﬂy describe the theoretical models of CPP—MR and CIMS in Chapter 2.
Chapter 3 describes the experimental techniques and procedures we used to fabricate
and measure our samples. The fabrication of nanopillars is given in detail.

In Chapter 4 we describe CPP-MR experiments involving spin-memory-loss at
4.2K in sputtered Pd, Pt, Au and at their interfaces with Cu to see whether these
materials could be used in devices without signiﬁcant spin-memory-loss. We also
compare our results with previously published data to understand the effect of spin-
orbit interaction in spin-memory loss.

In Chapter 5 we describe a new phenomenon initially observed by Bid [57], which
involves changes in magnetic scattering anisotropy at a ferromagnetic / superconducting
(F / S) interface. We extend his measurements to include new metals and alloys, and
combinations of metals and alloys to try to understand the source of such changes in
the scattering anisotropies.

In Chapter 6 we describe experiments involving normal and inverse CIMS or MR

14

to investigate the effect of positive and negative bulk and interface anisotropies.

In Chapter 7 we compare the switching currents and MRs of nanopillars with
Cu, Ag and Au spacers to see if Ag and Au are potential competitors of Cu for
nonmagnetic spacers.

In Chapter 8 we discuss the effect of a synthetic antiferromagnet (SAF) polarizing
layer on the net polarization and switching currents of nanopillars, to get an insight on
how the switching currents and MR are affected by reduction in the spin polarization.

In Chapter 9 we test a recent prediction by Manschot et al. [58] that the positive
switching current of a nanopillar (from P to AP state) could be reduced by up to a
factor of 5 by using asymmetric nonmagnetic leads to a nanopillar.

In Chapter 10 we study the effect of lower demagnetizing ﬁelds on the switching
currents of a nanomagnet by spin transfer torque to see whether a lower demagnetizing

ﬁeld noticeably reduces the switching currents as predicted by Katine et al. [46].

15

Chapter 2

Theory

In this chapter we brieﬂy review those theoretical models of CPP-MR in magnetic
multilayers that are relevant to our data. Earlier theses from our group [59, 60, 57]
and reviews [24, 61, 2] cover CPP-MR theory more completely. We also brieﬂy review
the theory of spin-transfer and current-induced magnetization switching (CIMS) in
section 2.6. The theoretical models of ST T are given in more detail in the review by

Stiles and Miltat [36].

2.1 Origin of GMR

GMR is a result of the spin-dependent scattering of electrons within bulk F-metals and
also at F / N interfaces. Such scattering occurs in F-metals due to the spin imbalance
of electrons in their 3d sub-bands. Spin imbalance occurs in F-metals because their
energy levels are shifted, resulting in unequal ﬁlling of energy bands (Fig. 2.1), which
is the source of non-zero net magnetization. Therefore F-metals can be used as a
source of spin-polarized currents. Whereas in N-metals the energy levels are not
shifted and therefore the ﬁlling of the energy levels is the same for both spins, which

gives zero net magnetic moment. In F / N structures, it is the spin-polarization of the

16

F-metals that gives rise to GMR.

 

 

 

 

4s 3d 43 3d
E "55-
D(E)
<——>
(a) Normal Metal (b) Ferromagnetic Metal

Figure 2.1: A representative diagram of density of states that are available to electrons
in (a) a normal metal and (b) a ferromagnetic metal whose spin-up and spin-down
states are not equally ﬁlled giving a non-zero net magnetization. E is the energy of
the electron, Ep is the Fermi level, and D(E) is the density of states.

2.2 Spin-polarized transport

Since the F-metals have a net magnetization, the conduction electrons entering the
F-metals will be affected differently depending on the relative orientations of their
moments and the local magnetization of the ferromagnet. The resistivity for electrons
with moments along (opposite) the local magnetization is denoted by pT (pl) using
the Valet-Fert notation [12]. Depending on the ferromagnet, pT can be larger than pf
or vice-versa. In different papers and reviews [61] spin-up is sometimes used instead of
moment-up. In reality the spin of an electron is opposite to its moment because of the
negative charge of the electron. More appropriate terms are majority and minority

electrons. Majority electrons (T) have their moments along the local magnetization

I7

and minority electrons (l) have their moments opposite to local magnetization.

For F/N/F trilayers and F/N multilayers, the total resistance changes as the
magnetizations of F layers switch from parallel (P) to anti-parallel (AP) alignments.
The change in resistance is deﬁned as AR = R(AP) -— R(P). This change can be
positive or negative depending on the bulk and interface spin-dependent properties.

If AR > 0, it is called “normal” MR and if AR < 0, it is called “inverse” MR.

2.2.1 Normal MR

Consider a simple picture of a F / N/ F trilayer where the minority electrons scatter
more strongly than the majority electrons (pT < pi) in the F-metal. To simplify the
process of spin dependent scattering we assume only bulk scattering in F layers. If we
assume no spin-ﬂipping throughout the sample we can treat the transport with two
independently propagating channels of moment up and down electrons (two-current
model). Current-mixing due to scattering from magnons and phonons becomes neg-
ligible at low temperatures. Therefore the assumption of no current-mixing is valid
at low temperatures.

In the AP state (Fig. 2.2(a)) the moment-up electrons scatter weakly in the ﬁrst F
layer and strongly in the second, whereas the moment-down electrons scatter strongly
in the ﬁrst F layer but weakly in the second. Therefore the total scattering rates for
each channel are the same in the AP state. On the other hand, in the P state (Fig.
2.2(b))the moment-up electrons scatter weakly in each F layer, whereas moment-down
electrons scatter strongly in each F -layer. Therefore in the up—channel we have two

weak scattering events and in the down—channel we have two strong scattering events.

18

A weak scattering represents a low resistance and a strong scattering represents a
high resistance. Since we have two different channels of transport, we can treat these
channels as two resistors in parallel. The resistance of the trilayer is higher in AP
state in this case. This type of MR is known as “normal” MR. Although we assumed
only bulk scattering, the interface scattering is also important in determining MR. A

more detailed analysis will be given in Sections 2.3 and 2.3.1.

Anti-parallel (AP) Parallel (P)

-»V

    
 

Figure 2.2: Normal MR: A schematic representation of spin-dependent scattering in
a F/ N/ F trilayer with p]; < pin (a) In the AP state the net scattering in both spin
channels is the same. (b) In the P state the net scattering in the down—channel is
greater than the one in the up—channel. Therefore the resistance in the AP state is
higher. Breaks in the dashed lines indicate strong scattering.

Antijparallel (AP) Parallel (P)

 

Figure 2.3: Inverse MR: A schematic representation of spin-dependent scattering in
a F1/N/F2 trilayer with pTF1 < p}, and p}? > pt? (a) In the AP state the net
scattering in the down-channel is greater than the one in the up—channel. (b) In the
P state the net scattering in both channels is the same. Therefore the resistance in
the P state is higher. Breaks in the dashed lines indicate strong scattering.

19

2.2.2 Inverse MR

Consider a more complex trilayer of F1 / N / F2 where F 1 and F2 are different F-metals.
We assume that p}, < p}, and p22 > pi? Again we assume only bulk scattering in
F -layers and no spin ﬂipping.

In the AP state (Fig. 2.3(a)), the moment-up electrons scatter weakly in F1 and
as well as in F2 because p22 > pf”, whereas moment-down electrons scatter strongly
in F1 and also strongly in F2 due to the same reason. Therefore we have two weak
scattering events in the up—channel and two strong scattering events in the down-
channel. On the other hand, in the P state (Fig. 2.3(b)) the moment-up electrons
scatter weakly in F1 and strongly in F2, whereas moment-down electrons scatter
strongly in F1 and weakly in F2. Using the same idea of resistors in parallel we can
see that the resistance in the P state is higher in this case. This type of MR is known
as “inverse” MR. Although we assumed only bulk scattering the interface scattering is
also important in determining MR. A more detailed analysis will be given in Sections

2.3 and 2.3.1.

2.3 Two-Current Series Resistor (2CSR) model

The ﬁrst phenomenological explanation of GMR was given by Baibich et al. [9] using
a two—current model. Zhang and Levy [22] developed the ﬁrst theoretical model for
CPP transport in F/ N multilayers, assuming diffuse scattering and no spin ﬂipping.
In this model the electrons with up and down moments are divided into two parallel
channels and the resistance of each channel is merely the series resistor sum of the

resistances of bulk and interfaces constituting the multilayer.

20

Later, Lee et a1. [23] and Valet and Fert [12] wrote the ﬁrst 2CSR equations using:

0 Bulk scattering anisotropy parameter BF
0 Interface scattering anisotropy parameter 7pm
0 Renormalized resistivity of ferromagnet pf;

o Renormalized interface speciﬁc resistance ARE/N, where A is the area through

which the current ﬂows.

All of these parameters in the 2CSR model are independent of the layer thick-
nesses, which are the only lengths in the 2CSR model. 61: determines the relation
between the resistivities for electrons with moments along ([0].) and opposite (pt) to

the local magnetization and is deﬁned as

l _ T
5.. = p—f—pg. (2.1)
PF +PF

7pm determines the relation between the interface speciﬁc resistances for electrons
with moments along (ARg/N) and opposite (ARé/N) to the local magnetization and

is deﬁned as

ARfr/N — ARP/N

 

 

 

”VF/N = . (2.2)
ARRN + AR},N
The renormalized resistivity for F-metals is deﬁned as
T l
1. PF + PF PF
= = _, 2.
The renormalized interface speciﬁc resistance for a F / N interface is deﬁned as
ART + ARi AR
ARE/N = “N “N — —-“"—” (2.4)

4 _1_7123‘/N.

21

2.3.1 Analysis of a F1/N/F2 trilayer in AP and P states
using ZCSR model

Consider a simple trilayer with two different F-metals separated by a N -metal, F 1 and
F2 with different thicknesses tm and tF2 as shown in Fig. 2.4 and the CPP current is
passing through an area A. The total speciﬁc resistance of each channel is given by

the series resistor sum of bulk and interfaces.

 

a) Antiparallel (AP)

 

b) Parallel (P)

Figure 2.4: A schematic drawing of a F1/N/F2 trilayer showing two channels of

electron ﬂow in a) AP alignment b) P alignment. Block arrows represent the direction

of local magnetization in each ferromagnetic layer. p1,“) is the resistivity for electrons

with moment up (down).

The total speciﬁc resistances of the “up” and “down” channels in the AP state

are given by
ARM"(AP) = Pfutm + ARf-‘i/N + PfltN + ARf‘z/N + Pfi‘thZ' (2-6)

22

Similarly, the total speciﬁc resistance of the “up” and “down” channels in the P state

are given by
AR""(P) = pmF1 + AR;1 ,N + pgtN + .412ng + pigs, (27)

ARdow"(P) = ptltpl + 2412;”N + pttN + Alli.”N + 4.2a, (2.8)

The electron scattering in an N-metal is spin-independent. Therefore

pi = pix = 2m (29)

where the factor 2 comes from the equal splitting of the total number of electrons
into up and down channels. Since the “up” and “down” channels are in parallel the

equivalent resistances of the trilayer in the AP and P states are given by;

_ AR"P(AP, P) x AR"°“’"(AP, P)
AR(AP’ P) ‘ AR"P(AP, P) + ARdown(AP, P) (2'10)

 

Using Eqns. 2.5, 2.6, 2.7, 2.8, 2.9 and 2.10 one obtains the difference in the speciﬁc
resistances between AP and P states as:
AAR = AR(AP) — AR(P) =

[ARPl/N — ARf‘i/N + (Pin - va1)tFlllARfv2/N " ARfrz/N + (Pin " PfF2)tF2l
(Pfu + Pf‘i)tFl + (Pf? + Pliny” + 4PNtN + ARfﬂ/N + ARPI/N + ARf‘z/N + ARfm/N
(2.11)

 

Substituting renormalized resistivities p" (Eqn. 2.3), interface speciﬁc resistances
AR" (Eqn. 2.4), bulk anisotropy parameters 5 (Eqn. 2.1) and interface anisotropy

parameters 7 (Eqn. 2.2) for F1 and F2 in Eqn. 2.11, one obtains:

(Pripfqtrl + ’YPI/NAREI/lerzpfrgtrz + 7F2/NARE2/N)

AAR = 4
pﬁltpl + p§2tp2 + pNtN + ARIN/N 'i' ARP2/N

(2.12)

 

23

where the denominator of Eqn. 2.12 equals AR(AP).

One can make the sign of AAR and therefore MR negative by making one of the
parentheses in the numerator of Eqn. 2.12 negative. If for an F-metal 51: > 0 and for
the interface of this F-metal with an N-metal 7pm < 0 or vice-versa, then one can
adjust tp to make the MR negative.

Assume that this trilayer is sandwiched between two superconducting (S) leads.
Experiments showed that the S / F interfaces are spin independent [61], thefore ARg/F =
ARé/F = 2ARs/p. Then, the S/ F1 and F2/ S interfaces add only a constant term to
each channel. Such constant terms only appear in the denominators of Eqns. 2.11
and 2.12, reducing AAR.

If F1=F2=F and tpl = tpg = tp then Eqn. 2.12 becomes

 

 

 

 

*t+ AR’ 2 *t+ AR" 2
AAR : 4 (PFPF F ’YF/N F/N) : (IBFPF F 7F/N F/N) , (2.13)
Zpiétp + pNtN + 2ARP/N + ZARs/F AR(AP)
where for completeness we now include the 2ARS )1: term.
2.3.2 Validity of 2CSR model
For an F / N multilayer with N bilayers, AAR is
N zlﬂrpfvtr‘ + 2’i’F/IW‘UQE Nl2
AA = / , 2.14
R AR(AP) ( )
which can also be rewritten as

 

This equation predicts that for ﬁxed tp a plot of ﬂAAR][AR(AP)] vs. N should be

a straight line passing through origin. Since the N-metal resistivity doesn’t appear in

24

Eqn. 2.15, data for alloys of a given host metal should fall on the same line, so long as
the spin-diffusion length in the alloy is long (i.e. the impurity has a weak spin—orbit
interaction in the host). Demonstrations that data for Ag and AgSn [23] (and later
for Cu and CuGe [62]) fell on the same line to within experimental uncertainties,
although their residual resistivities differed by factors of 20, led to acceptance of the

2CSR model as at least a good ﬁrst approximation.

2.4 Valet-Fert (VF) model

2.4.1 Overview

Valet and Fert (VF) developed the ﬁrst theoretical model of CPP-MR that includes
spin-ﬂipping in the layers of a F / N multilayer [12]. Using a Boltzmann equation model
VF showed that the CPP-MR transport equations reduce to macroscopic transport
equations if the spin-diffusion length (the distance an electron diffuses before its spin
ﬂips) is much longer than the mean free path ls; >> A (independent of the ratio of
the layer thicknesses to the A).

To simplify the analysis, VF assumed single parabolic conduction bands for F
and N layers with the same effective mass and Fermi velocity throughout the F/N
multilayer. The VF analysis is valid at low temperature, where electron-magnon
and electron-phonon scattering are weak, leaving spin-orbit interaction, defects or
impurities and exchange interaction by paramagnetic moments diluted in N-layers as
the active sources of scattering. In the limit A << 13", which generally holds for metals,
the macroscopic transport equations, independent of the layer thicknesses, reduce to:

8 6J8 _ 213 _ﬂ—s

_ _ 2.16
a, ('92 I?! ’ ( )

 

25

033;?

J3: e 02

(2.17)

where e is the electron charge, a, is the conductivity for spin .3, Ps,(—s) is the chemical
potentials for spins s, (—s), J, is the current density and z is the direction of current
ﬂow. Eqn. 2.16 states that, in the steady state, the spin-accumulation (i.e. the
difference between the spin-up and spin-down Fermi energies) due to the spin-current

divergences is balanced by spin-ﬂipping and Eqn. 2.17 is just the Ohm’s law [12].

2.4.2 VF equationsfor an F / N multilayer

VF derived the total speciﬁc resistances of an F/ N multilayer with N bilayers in P

and AP magnetic conﬁgurations as

AR(P’AP) = N(r0 + 2.411"P ,AP)), (2.18)

where To = (1 — ﬁOpHp + PNtN + 2(1 - 72)AR;‘/N’

 

 

 

 

 

 

 

 

 

 

 

AR?! =
B _ ,7 2 t 2 t 2
Plef coth El?!- +FngCOth 51E; +ng§
——1N—coth 1%— —,lp—coth 1%- + Kilt-— AN- coth 1%- + -,—1p-coth Fiﬁ-
Plef .213}. pFlsf .2131. F/N .Plef .213]. pFlsf .21.!“
(2.19)
and A1250 =

Wtamhb ]+1—;L:p-coth[§tl?— ]+71-1,%:/—N

pst .

1 * 1 t 1 t
m 1...}. [2);] mm" [27%;] + m l‘irtan" [Pill + 711.. 6°” [55?“
(2.20)

In the limit, where lg”, >> tN and 1:} >> tp these VF equations reduce to the 2CSR

equations [12].

26

2.4.3 Finite spin-diffusion-length effects in a trilayer

The 2CSR equations do not include effects of spin-ﬂipping anywhere. The only lengths
appearing in Eqn. 2.12 are the thicknesses of the F- and N-layers. If tp >> If, then
the 2CSR equation should be corrected by replacing the tp with If, because the spin
dependent scattering within F-layers happens up to the spin diffusion length of F-
layer and the remaining thickness of the F-layer (tp — lff) is treated as a constant
spin-independent term, which drops out of AAR. If tp < 1:} then Eqn. 2.12 remains
approximately valid. Assuming both tpl >> If} and tpg >> If} the 2CSR equation 2.12

becomes (see e.g. ref. [61])

(PFIPi'nlf)l + VFI/NARIPI/N)(ﬂp2pf‘2l§f2 + 7F2/NARP2/N)

AAR = 4
pf‘llf; + P1321512 + PNtN + Ang‘l/N + ARPl/N

 

(2.21)

2.5 Real Fermi Surface models of CPP-MR

We noted in Sections 2.3 and 2.4 that the” 2CSR and VF models were derived assuming
identical spherical Fermi surfaces for both the F- and N-metals. Some CPP—MR
models have expanded the 2CSR model to include effects of real Fermi surfaces. Such
effects might, or might not, modify the parameters of the 2CSR and VF models.
For example, interfacial speciﬁc resistances derived using the VF model can agree
well with values independently calculated based upon the real Fermi surfaces (see
Section 4.7). Of more importance are claims that incorporating real Fermi surfaces
can cause the CPP-MR to show ‘mean-free—path’ eﬂects, i.e., behaviors that vary

with the ratios of the mean-free—paths to layer thicknesses of the F- or N-layers [63,

27

64]. As noted in Section 2.3.2, the initial search for such mean-free-path effects, by
comparing square-root plots for Ag and AgSn and for Cu and CuGe, did not show
any to within experimental uncertainties. Subsequent claims of observations of mean-
free-path effects [65, 66, 67] have been disputed on experimental grounds [68, 69].
At the present time, one cannot rule out mean-free-path effects on the borderline
of experimental uncertainties. But we do not know of any data that cannot also
be explained by including plausible spin-ﬂipping without any mean-free-path effects

([69, 70], and references therein).

2.6 Spin Transfer Torque & Current-Induced Mag-
netization Switching

2.6.1 Overview

The spin-transfer phenomenon is treated theoretically assuming either ballistic (in
which only the interface scattering is important) or diffusive transport (in which
scattering in the bulk and spin accumulation at F/ N interfaces are also included).
Initial theories assume a macro-spin (single domain) structure for the nanomagnets.
Micromagnetic (multi-domain) simulations also describe the underlying physics of the
spin transfer excitations.

The original circuit theory of Slonczewski [33] assumed ballistic transport. Slon-
czewski [33] assumed that upon injection of polarized current into a nanomagnet the
transverse component of the electron spin is absorbed by the free layer, providing
the source of the torque. By adding this torque term to the Landau-Lifshitz—Gilbert

(LLG) equation, Slonczewski solved for the motion of magnetization of a single do-

28

main nanomagnet, from which he predicted a steady precession of magnetization and
a new type of magnetic switching.

On the other hand, assuming diffusive transport, Berger [34, 71] predicted that
upon injection of polarized current, coherent magnons are created due to spin ac-
cumulation at an F/ N interface so that a spin can ﬂip from minority to majority
giving rise to the torque and excite a spin wave while conserving energy and angular
momentum. Berger showed that the total spin current and spin accumulation are in
ﬁxed ratio showing that the Slonczewski model [33] based on spin current and Berger
model [34, 71] based on spin accumulation are more or less equivalent [35, 72].

Later, Heide et. al [50, 51, 49] using diffusive transport equations argued that
it is the non-equilibrium exchange interaction between s- and d-electrons driven by
the current that promotes the torque. In Heide’s s-d model the similarities of Fermi
surfaces greatly reduces or even eliminates the spin transfer effect. In this case the
transverse spin current is not absorbed by the nanomagnet but a transverse spin ac-
cumulation develops in the nanomagnet. This transverse spin accumulation precesses
relatively slowly around the magnetization giving rise to a torque on the magnetiza-
tion.

Stiles et al. [73], using a matrix Boltzmann equation formalism, also solved the
LLG equation for a nanomagnet. Fert et a1. [74] and later Barnas et a1. [75] showed
that the LLG equations can be solved using diffusive transport for the entire nanopil-
lar structure using Boltzmann equation formalism, which was also used in the theory
of CPP-MR [12]. All of these semiclassical models describing the spin transfer phe-

nomenon give qualitatively the same results [36].

29

The theories described above ignored temperature. Stimulated by their experi-
mental results on dynamical responses of nanOpillars, including telegraph noise, We-
growe et al. [76, 77], Myers et al. [78] and Urazhdin et al. [48, 79], proposed thermally
activated pictures of magnetic excitations (Wegrowe and Urazhdin with the dc current
I producing incoherent magnons, the effects of which they described via an effective
temperature equal to the sum of the ambient temperature and a current dependent
temperature). Urazhdin et a1. [80] subsequently used a thermally activated model
to explain differences in switching behaviors of magnetically uncoupled, ferromagnet-
ically coupled, and antiferromagnetically coupled trilayers. Li and Zhang [81] and
Apalkov and Visscher [82, 83] showed that the temperature dependence of the STT
could be described as an activated process with an effective temperature that de-
pends upon the exciting current I, but their effective temperatures were proportional
to the ambient temperature. Micromagnetic simulations by Lee et al. [84] provided
qualitative support for an effective temperature as an approximate way to describe
the broad range of excitations generated by the STT current. Krivorotov et a1. [85]
subsequently argued that the switching rates in their samples depended only upon
the true sample temperature, adding support for a simple proportionality of effective

and ambient temperatures.

2.6.2 Ballistic transport equations

The initial theoretical model of STT by Slonczweski assumed ballistic transport
throughout the magnetic trilayer structure, i.e. the mean free paths of each layer are

much longer than their respective thicknesses, therefore bulk scattering is neglected.

30

Let’s consider an F1 / N/ F2 elliptical nanopillar structure with ~ 100 nm lateral
dimensions (see Fig. 2.5). Conventionally the positive CPP current indicates electron
ﬂow from the thin F2 (with magnetization M2) to thick F1 (with magnetization M1).
The equation of motion of M2 is given by the LLG equation, which is modiﬁed to

include the STT term.

 

Figure 2.5: A representative cartoon of an elliptical nanopillar, where F1 is a thick
(ﬁxed) ferromagnet, N is a nonmagnetic spacer thinner than its spin diffusion length,
and F2 is a thin (free) ferromagnet. The angle between the magnetizations M1 and
M2 is 0.

Slonczewski [33] derived the STT as

% = (Jng) [M2 X (M2 >< Mlll (2-22)

where M, is a unit vector along Mi, J is the current density, and 0 is the angle

between M1 and M2. The scalar function g(l9) is given by

 

—1
_ _ 3%
9(0) _ [ 4 + (1+ P) 4133/2 (2.23)
Here P is the polarizing factor of the ferromagnet which is deﬁned by
nT — 71*
= 2.24
nT + 714' ( )

31

where nm) is the Fermi level majority (minority) spin densities in a ferromagnet.
Notice that in Eqn.2.22 and 2.23 there is no thickness dependence. This arises from
averaging the ﬂuxes with respect to the phase of the exponent 6”“ occurring therein,
involving phase differences across the thickness (t) of the nonmagnetic spacer layer.
Consider an effective uniaxial anisotropy ﬁeld H“, which includes the effect of
magnet shape and the phenomenological Gilbert damping coefficient a. The LLG
equation for such a single domain ferromagnet, modiﬁed to include the torque term

(Eqn. 2.22) is [33]

dM2 .

dt = M2 x (711,2 . M26 — aid—ME

dt

 

where 7 is the gyromagnetic ratio and a ﬁxed frame is deﬁned by orthogonal unit
vectors 6,6 and 6 of which 6 is the symmetry axis of anisotropy. In the case of
H, > 0, :11“: are the easy directions for M2.

The motion of nanomagnet F2 is found by substituting the solution
M2 = (sin 6)(& cos wt + 68in wt) + écos 6 (2.26)

into the LLG equation in 2.25 under the assumption of weak damping. Applying the

condition [6] < [w] yields the lowest order relations w = 7H,, cos 6 and

6 = —(oryH,, cos 6 + M) sin 6 (2.27)
M28

where g(6)(> 0) is given by Eqn. 2.23 with M1 - M2 = cos 6.
Under some conditions , time dependent solutions to the LLG equation 2.25 de-
scribe switching with 6(t) varying between orientations near easy directions 6 = 0

and 6: 7r.

32

Based on this ﬁrst theory of Slonczewski, Katine et al. [46] derived the critical
currents for switching (assuming 1+ ﬂows from ﬁxed to free layer) from parallel (P)

to antiparallel (AP) states as

 

I: = %[H.,,(0) + 2an (2.28)
and from AP to P as
_ a 65'
c = 927T) [H,,,(7r) — 27rM] (2.29)

where M is the magnetization of the switching layer, S is the total spin of the switch-
ing layer, and He” = Han cos(6) + Hm — He“ is the sum of anisotropy, external and
exchange (but not demagnetizing) ﬁelds. 9(0) and g(rr) depend on the spin-dependent
transmission probabilities of the F / N interfaces, the relative orientation (0=parallel,
7r=antiparallel) of the nanomagnet and any spin-ﬂip scattering that may occur in the

system.

2.6.3 Diffusive transport equations

Although the circuit theory of Slonczewski [33] describes much of the experimental
data obtained at least qualitatively, the ballistic transport assumed in his theory does
not explain all the data. In 2002, Slonczewski [86] extended his model to include dif-
fusive transport in the ferromagnetic layers, with the spacer still treated as ballistic
because it is usually thin compared to its mean-free-path. In the diffusive trans-
port calculations an additional phenomenon ‘spin-accumulation’ enters the transport
equations [12]. At an F / N interface, far from the interface to the F side the incoming
electrons are polarized, where there are more majority electrons and far from the in-

terface to the N side the outgoing electrons are unpolarized. This implies that within

33

the F and N spin diffusion lengths of the F/ N interface majority electrons must ﬂip
their spins to give equal majority and minority channels well into the N-metal. The
splitting between the chemical potentials (Fermi energies) of majority and minority
channels, is called “spin-accumulation” [87].

Later Xiao et al. [88] derived the torque using matrix Boltzmann equation for-
malism assuming diffusive transport. The torque exerted on the magnetization M2

using micromagnetic simulations is expressed as

 

Is; = ———--e—’P[M2 x (M2 x 1911)], (2.30)

where 113 is the Bohr magneton, J is the charge current density normal to the inter—
faces, M 5 is the saturation magnetization [36] and d is the thickness of the nanomag-
net. One evident difference of this torque term is that it includes the thickness of
the switching layer which is inversely proportional to the torque. M2 is the magnetic
moment of the nanomagnet and M1 is a unit vector in the direction of the magnetic
moment of the ﬁxed layer (see Fig. 2.5). By convention 9,, is negative for free elec-
trons. Since the orbital moment in transition metals is small, the Lande’ g-factor
for transition metals is close to -2. ’P is a polarization function that contains all the
information originating from the structure of the sample and depends on the relative

orientations of the magnetizations and is given by.

(1+ Q-
: 2.31
P Bo+Blcos6+Bo-Blcos6’ ( )

 

 

where cos6 = M1 - M2. The parameters Bo, B1, q+ and q- depend on the geometry
of the sample and the physical properties of the bulk and interfaces including the spin

accumulations in the N Spacer layer at both F1 / N and N / F2 interfaces. In particular

34

they depend on the effective spin-dependent resistances from the spacer layer in either
direction over one spin diffusion length [88, 36]. ’P is usually positive but as we will
show later in Chapter 6, it can be made negative using judicious choice of materials

[89]. Inserting this STT into the LLG equation one obtains the equation of motion

 

Of M2 as
32:“) = -1[M2(r) x Hear) — %%§§P<r)lmz(r> x (szl‘) x 1911)]
+0: (M2(r) x Wag-Z) . (2.32)

Fig. 2.6 shows a comparison of the torques calculated using Slonczewski’s diffusive
model [86] with Boltzmann equation model and Slonczewski’s ballistic model [33]. As
it is clear from Fig. 2.6 the Boltzmann equation model of Xiao et al. [88] agree very
well with Slonczewski’s diffusive circuit theory [86] but not so well with the ballistic
Slonczewski model [33].

Fert et al. [74] and Barnas et al. [75] derived the STT at small angle assuming
diffusive transport for the entire nanopillar structure including spin-accumulation and

spin-current as

v mP(AP) jPMP)
7PM?) = -5 [(L + "’1” ) (1- e“tN/)‘N) (2.33)

 

8 2

P(AP)
v m , _ .. . ..
+ (——F :1 + Jﬂfg’f’) e ‘N/AN] M2 x (M2 x M1)

where M1 and M; are unit vectors along the magnetizations of ﬁxed and free layers

respectively, tN is the thickness of N-spacer, AN is the mean-free-path in the N-metal,

and up is the Fermi velocity. mSMP’ and jig?) are the spin-accumulation density and

spin current density in the N spacer layer at the N / F2 interface, whereas mam” and

35

 

' ' ' ‘ 1

0,5. Boltzmann o
Slonczewski—

 
 
  

0.4 -

0.2 -

0.l -

 

Spin transfer torque (In/29)
4:
w

 

J

o 30 60 90 e 120 150 180

Figure 2.6: Spin-transfer torque acting on the magnetization of a 1 nm thick nano-
magnet as a function of 6. Solid circles are Boltzimann equation results. The solid
curve is the torque derived from Slonczewskis later theory [86] including bulk scat-
tering. The dashed curve is the torque with all bulk scattering removed. The ﬁgure
is taken from ref. [88].

ngP’ are the spin-accumulation density and spin-current density in F1 at the N / F 1

interface. The ﬁrst parenthesis in Eqn. 2.33 dominates in the limit tN >> AN, and
the second parenthesis in Eqn. 2.33 dominates in the limit tN << AN. One advantage
of this Fert equation for torque is that it directly includes the spin-accumulation and
spin-current terms, which would provide a more straightforward analysis using the

parameters obtained from CPP-MR experiments.

2.6.4 Summary

In the simple ballistic picture described in the introduction, F1 polarized the current
and the component of that polarized spin-current perpendicular to M2 of F2 exerted
a torque on M2. In the more complex picture arising from analysis of diffusive
scattering, both spin-current and spin-accumulation play important roles in STT,

and both must be determined self-consistently for the whole sample— i.e., taking into

36

account the properties of all components of the sample-layers F1, N, F2, the leads to

both F1 and F2, and the interfaces between these layers.

37

Chapter 3

Sample Fabrication and
Measurement

3. 1 Overview

In this chapter, we describe the experimental procedures used to make and measure
the samples for both CPP-MR and CIMS studies. For the CPP-MR studies we make
Nb—sandwiched magnetic multilayers. To get a uniform current density through the
multilayer, we sandwich the multilayer between two crossed Nb leads, which super-
conduct at our measuring temperature of 4.2K [59]. For CIMS studies we fabricate
magnetically uncoupled elliptical nanopillars of ~ 130 x 70 nm lateral dimensions.
Since both studies involve sputtering, our sputtering procedures are described under
CPP-MR sample fabrication (in section 3.2.2) and then referenced for nanopillar fab-
rication (section 3.4), in which sputtering is only one step in a complex fabrication

process.

38

3.2 Making and measuring Nb-sandwiched mag-
netic multilayers for CPP-MR studies

3.2.1 Preparation

We begin with 3” diameter (100) Si wafers, which have insulating Si02 layers on top.
For CPP-MR studies we dice the wafer into 0.5" x 0.5” square substrates. We clean
these Si substrates in a dilute alconox solution and rinse them thoroughly with deion-
ized (DI) water, then in acetone and ethanol respectively using ultrasonic agitation
in each step. Then, we blow-dry the substrates with pure nitrogen gas to get rid of
any residues from ethanol.

We clean the substrate holders, masks and the sputtering gun parts by washing
them in a nitric acid solution to remove any previously deposited metals. These parts
are then cleaned with acetone and ethanol using ultrasonic agitation in each step,
and then we dry them using a heat gun.

Clean substrates are loaded into sample holders and placed on a rotating sample
positioning and movement assembly (SPAMA) plate [59]. The SPAMA plate has two
computer-controlled ﬁlm thickness monitors (FTM) placed opposite to each other
and 8 circular openings for sample holders and masks. The SPAMA plate can hold
8 substrates for Nb—sandwiched CPP samples or 16 substrates for single ﬁlm samples

[59, 31].

3.2.2 Sputtering

We make the samples by sputtering metal or alloy targets. Four triode and two

dc-magnetron guns allow us to make complex multilayers. The vacuum chamber is

39

pumped down by ﬁrst a mechanical and then a cryo pump. TO lower the pressure
and partly eliminate the water vapor inside the chamber, the chamber is baked for 8
hours right after closing the system. It takes 2 nights to pump down to 3 — 4 x 10‘8
Torr base pressure for the ﬁrst run. If we want to make more samples using the same
targets, we just load new substrates into the sample holders and close the chamber.
In this case it takes 1 night to pump down to 3 — 4 x 10‘8 Torr.

Before starting-up the sputtering guns, we ﬁll a cold trap in the sputtering chamber
with liquid nitrogen. The ﬂow rate of the liquid nitrogen is adjusted to maintain a
constant cold trap temperature of 82K (we cannot reach 77K because of the pressure
we apply for the liquid nitrogen to ﬂow). This allows us to freeze most of the water
vapor inside the chamber and get 2 — 3 x 10‘8 Torr base pressure. An argon gas
puriﬁer is used to get ultra pure argon gas for sputtering. We sputter the metals or
alloys in 2.5 mTorr argon pressure.

While all masks and the shutter are closed, we pre-sputter all targets inside the
chamber (except Au and Pt because they don’t oxidize) for 10-20 minutes tO remove
oxide surface layers and to stabilize their surfaces. Each target has a chimney to
collimate the sputtered atoms and prevent contamination from nearby targets.

The deposition rates of sputtering targets are individually measured just before
making each sample. Table 3.1 shows typical sputtering rates for some materials.
The actual sputtering rates will be given with the experiments. A computer controls
the shutter and the SPAMA plate. The computer initially moves the sample over the
chosen target by rotating the SPAMA plate, Opens the shutter, and leaves the sample

over the target for a time calculated by the computer using the deposition rate from

40

a FTM reading. After the deposition is completed, the shutter is closed or kept open
depending on the next target’s position in the chamber. Then the sample is moved

to the next target in the multilayer sequence and so on.

 

Target Nb Cu Py Au CO
Deposition rate (X/s) 4.5-5 9-11 4.5-5 4-5 2-4

 

 

 

 

 

 

 

 

 

Table 3.1: Typical sputtering rates for Nb, Cu, Py, Au and CO.

During the sputtering Of multilayers, the substrate temperature is kept between
-30 and +30°C. Making one CPP sample usually takes 25-30 minutes and this causes
the substrate temperature to increase by only a few degrees centigrade. To achieve
this approximately constant substrate temperature during sputtering, we cool the
SPAMA plate constantly by a capillary tube that is in contact with both the liquid
nitrogen cold trap and the SPAMA plate. We run pressurized nitrogen gas at 1000
psi through the capillary tube to enable the necessary heat exchange between the cold
trap and the SPAMA plate [59].

Each substrate has its own rotary CPP mask, attached to the sample holder [59].
A CPP mask has 4 positions as shown in Fig. 3.1(a), which is initially in the closed
position (1). We turn the mask manually (using a vacuum-sealed wobble stick) to:
ﬁrst, position (2) and sputter 150-200 nm Nb as a bottom lead, then we turn the
mask to position(3) and sputter the multilayer and then we turn it to position (4)
and sputter 150—200 nm Nb as the top lead. Finally, the mask is turned to position
(1) to protect it from contamination during the sputtering for other substrates in the

chamber. A ﬁnished CPP sample is shown in Fig. 3.1(b).

41

   

(b)

Figure 3.1: (a) A rotary CPP mask: position (1) is closed to protect the sample
before and after sputtering, position (2) is for the bottom Nb contact, position (3)
is for the multilayer and position (4) is for the top Nb contact. (b) A ﬁnished CPP
sample. The width Of each Nb contact is ~ 1.1 mm.

3.3 Measuring the Resistances of CPP samples
The CPP samples for this thesis have two different structures as described earlier:
1. Exchange-Biased Spin Valves (EBSV).
2. Hybrid spin valves.
3.3.1 Pinning

In an exchange-biased spin valve (EBSV) one Of the ferromagnetic (F) layers is sput-
tered next to an antiferromagnetic (AF) layer (FeMn in our samples). To pin the
magnetization of this F layer, we heat the sample to above the Néel temperature of
the antiferromagnet in a vacuum chamber and cool the sample back to room tempera-
ture in an external magnetic ﬁeld. This process pins the magnetization of the F layer

that is next to the AF layer. Pinning a ferromagnetic metal shifts its hysteresis curve

42

toward the opposite direction Of the applied magnetic ﬁeld, changing its switching

ﬁeld dramatically.

3.3.2 Resistance measurement at 4.2K

TO make good electrical contacts with very low contact resistances to the Nb leads,
we put indium on the Nb leads by using an ultrasonic soldering iron before pinning
the ferromagnetic layer. It is necessary to use an ultrasonic soldering iron to make
a good electrical contact to Nb, because Nb oxidizes very quickly. The sample is
mounted on a measurement stick (quick dipper) and the current and voltage leads
are pressed on indium at room temperature. We avoid using any hot soldering tO
connect the leads because it can unpin the pinned ferromagnetic layer in an EBSV.

All CPP samples are measured at 4.2K in liquid helium, where the Nb leads
become superconducting. Since the top and bottom Nb contacts superconduct at
4.2K, our standard CPP measuring current Of 100 mA passes uniformly through a
big area A~ 1.2 mm2 like a parallel plate capacitor (see Figure 3.2). The area A, is
the product of the widths Of two Nb strips. The widths Of Nb strips are measured
using a Dektak surface proﬁlometer [59, 60]. The uncertainties in the areas Of the
samples are ~ 5%.

The resistances of these CPP samples are in the order of 10‘s!) because the sam-
ples are very thin (t ~ 100 nm) and the areas of the samples are large (A~ 1.2 mm2).
TO measure these very small resistances, we use a Superconducting Quantum Inter-
ference Device (SQUID) based circuit. The details Of this 4-probe measurement are

described elsewhere [59, 31]. The sample is placed inside a 1” diameter superconduct-

43

   

(b)

Figure 3.2: (a) A CPP sample with lumps of indium on Nb contacts. (b) Center Of
a CPP sample; The current passes only through the area Of intersection of two Nb
leads at 4.2K.

ing solenoid magnet, which is capable of producing ﬁelds up to 1.5 T. The sample
is positioned in the center Of the solenoid in such a way that the ﬁeld produced is
parallel to the sample plane. Then, the stick is slowly lowered in a liquid helium
storage dewar, until the sample and the SQUID are under the liquid helium surface.
It takes ~ 20 minutes to cool down to 4.2K and ~ 20 minutes to warm up to room
temperature. Due to this very short time for cooling of the sample compared to other
methods [90], the measurement stick is called “quick-dipper”. At 4.2K we measure
the resistances of CPP samples at different external magnetic ﬁelds produced by the
superconducting solenoid. For Py—based EBSVs we sweep the ﬁeld in steps between
-300 Oe and +400 Oe, and for Py-CO hybrid spin valves we sweep the ﬁeld in steps

between -1 kOe to +1 kOe. Resistance vs. magnetic ﬁeld plot gives us the change in

resistances between parallel and anti-parallel states Of the sample.

44

3.4 Making and measuring nanopillars for CIMS
studies

Making nanopillars is a multi-step process, which requires a lot of care and patience.
Starting from scratch it takes about a week to fabricate nanOpillars, assuming no
problems. We make magnetically uncoupled F1/N/F2 elliptical nanopillars of ~
130 x 70 nm lateral dimensions, where F1 is a thick and wide (ﬁxed under CIMS)
polarizing ferromagnet, N is a normal metal thinner than its spin diffusion length
but thick enough (~ 10 nm) to minimize exchange coupling and F2 is a thin (free
under CIMS) nanomagnet. In this system, since F1 is several microns wide and F2
is a nanomagnet, the dipolar coupling between F1 and F2 is minimal. We make 5
nanopillars (with the same structure) on a single 0.5” x 0.5” Si substrate. In a single
run we process 4-8 substrates and therefore 20-40 nanopillars. Following is an outline

of the fabrication of nanopillars.

1. Preparation Of substrates using photolithography and evaporation.
2. Electron beam lithography (EBL) for the multilayer and sputtering.

3. EBL for nanOpillars and evaporation of A1.

A

. Ion-milling and insulation.

01

. Low-angle ion-milling.

6. Top contacts.

Photolithography and EBL steps are done in a class-100 clean room.

45

3.4.1 Preparation of substrates using photolithography and
evaporation

We start with single crystal 3” diameter (100) Si wafers. Because the Si wafers are
packed in a clean room environment, and we Open the wafer box in a clean room,
they usually don’t need further cleaning. However, if the wafer is visibly dirty, it is
cleaned in acetone, isopropyl alcohol (IPA) using ultrasonic agitation in each step.

Then the wafer is immediately blow-dried.

Photolithography

Before sputtering the multilayer, Au current and voltage leads must be evaporated
onto the wafer. Photolithography is used to Open the areas on the wafer for the leads.

Following is a recipe for photolithography.

1. Spin 81805 (propylene glycol monomethyl ether acetate) photoresist (PR) on a

Si wafer at 3500 rpm for 40 seconds to get a uniform PR thickness Of ~ 500 nm.

2. Bake the wafer at 90°C for 1 hour to evaporate the solvent and harden the PR

layer.

3. Place the wafer on a vacuum chuck in a mask aligner and align with the pho-
tomask, which has the shape of the leads transparent to ultraviolet (UV) light
and the rest coated with Cr to block UV light (see Fig. 3.3). Bring the wafer
and the mask into complete contact leaving almost no air between them and
expose the wafer to the UV light for 2.5 seconds. UV light breaks chemical
bonds in the PR polymer by allowing the exposed PR to be dissolved in photo

developer (KOH)

46

 

Figure 3.3: A picture of the photomask. The open areas have the shape Of the leads
and the rest of the mask is coated with Cr to block UV light.

4. Harden the surface of the PR by dipping the exposed wafer in chlorobenzene for
20 seconds . Immediately after taking out the wafer from chlorobenzene, rinse
the wafer under running DI water for 3-5 minutes until the wafer is completely
free Of chlorobenzene residues. Then blow-dry the wafer with pure nitrogen gas.
Hardening the surface of PR is crucial to get a good undercut after developing.
This process makes the PR layer behave like a bilayer i.e. the surface of the
unexposed PR reacts with KOH slower than the rest of the PR giving a good
undercut (see Figure 3.4). A good undercut ensures a trouble-free lift-off. If the
undercut is not very good, the lift—off is irregular because acetone (the chemical

used for lift-Off) cannot easily penetrate into the PR layer.

01

. Finally, we develop the wafer in photo-developer (452, KOH) for ~ 45 seconds,
until the entire exposed area Of PR is dissolved.
Evaporation Of leads

1. Load the wafer into a thermal evaporator within the clean room and close the

chamber to pump down to 1 — 2 x 10‘7 Torr. Fill the cold trap with liquid

47

undercut

  
 

PR PR

 

- Si substrate ]

._.. _..._...__.....A.

Figure 3.4: A cartoon Of a good undercut after developing. The top part of the PR
reacts very slowly with KOH.

nitrogen to get to 1 — 2 X 10‘7 Torr faster.

2. First, evaporate 3 nm Ti to get a strong adhesion to the Si wafer, then evaporate
80 nm Cu tO decrease the resistance of the leads and ﬁnally, evaporate 100 nm
Au to protect it from oxidation. A 100nm thick Au top layer is crucial to see

the leads better in SEM during electron beam lithography.

3. Take the wafer out of the evaporator and place it in acetone facing the evapo-
rated side down to lift-off. PR dissolves in acetone and since the PR is every-
where on the wafer except the areas for leads, all excess metal layers will be
removed in acetone leaving the metals in the areas for the leads created by

photolithography.

4. Rinse the wafer with IPA and DI water respectively immediately after taking
it out Of acetone and then blow-dry it with pure nitrogen gas. Fig. 3.5 shows a

3” Si-wafer after lift-Off with the leads evaporated.

48

.
.
nu
n.<
>
I‘.
"i
I

-
‘l.nl'n-..l

1 ta

:7] a l

 

Figure 3.5: A picture of a 3” wafer after lift-Off, which has 21 0.5” x 0.5” square
substrates.

3.4.2 EBL for the multilayer and sputtering

E-beam Lithography (EBL)

EBL works much the same way as photolithography except the resists are different
and there is no mask involved. In EBL, a computer controlled electron beam in an
SEM scans the desired shape and the e-beam resist in the scanned area becomes
soluble in e-beam developer, which is 1:3 mixture of MIBK (4—methyl-2—pentanone)
and IPA). Using EBL one can make features smaller than 50 nm, depending on
the type and thickness of the e—beam resist [91]. For EBL we use a modiﬁed SEM
(JEOL 840A with tungsten hairpin electron source)with a beam blanker. The SEM
is connected to a computer through an interface which allows the computer to control
the e-beam using NPGS (Nano Pattern Generation System) software. NPGS software
scans the e—beam over the shape and size of our desired feature, which we draw using

DesignCAD software.

49

Preparation for EBL

To prepare the substrate for EBL we have to coat it with e-beam resists. To achieve
a uniform thickness of the resist we spin the wafer during coating. To get a good
undercut we spin a bilayer of e-beam resists. Following are the steps for preparation

for EBL:

1. Spin 9 % MMA (9 % ethyl lactate) e-beam resist on the wafer at 4000 rpm for

40 seconds and bake the wafer for 5 minutes at 170°C.

2. Spin 2 % PMMA (2 % chlorobenzene) e—beam resist on top Of MMA layer at

4000 rpm for 40 seconds and bake the wafer for 30 minutes at 170°C.

We bake the wafers to evaporate the solvent in the e—beam resists. Then we dice
the wafer along the boundaries of the individual 0.5” x 0.5” substrates. We get 21
substrates out of a 3” Si wafer. Fig. 3.6 shows a single substrate and its center.
The crosses are the alignment marks used to align the sample in a Scanning Electron

Microscope (SEM) for EBL.

Alignment marks
'+

’v
a
‘—
’—

----->

   

V
------ 9 Au leads

Figure 3.6: A single substrate (left) and its center (right) with the gold leads evapo-
rated.

50

EBL for the multilayer

We need to Open a window for the multilayer at the center of our substrate, which
is totally covered with an e—beam resist bilayer. The following is a recipe of EBL for

the multilayer.

1. Load the substrate into the SEM and write the pattern in Figure 3.7. The
beam current is 25 pA. The area dose is 130/1C /cm2 for the inner pattern and
70nC/cm2 for the outer one. Therefore the outer shape is exposed to less e-
beam than the inner shape. This gives us a bigger undercut which is needed
for sputtering, because sputtered atoms go in every direction unlike the more

collimated evaporated ones.

I
-I

-----------------------------------------------------

I
\-----c--- ’

'---—----‘
--------

—————————————————————————————————————————————————————

 

 

.......

Figure 3.7: The DesignCAD drawing for the shape Of the multilayer.

2. Take the sample out Of SEM and dip it in the e-beam developer for 30 seconds

to develop. The developing time depends on the beam current and dose.

3. Take the substrate of the e-beam developer and immediately dip it in IPA for
10 seconds to remove the residues from the developer. Then take it out Of the

IPA and blow-dry immediately.

4. Check the development using an optical microscope.

51

If the feature does not develop perfectly within the ﬁrst 30 seconds develop 5
seconds more and so on until it is fully developed. Fig. 3.8 shows a developed area

for the multilayer.

 

Figure 3.8: An Opening for the multilayer created by EBL. Scale as in Fig. 3.7.

Sputtering

We load these substrates into the sample holders behind stainless steel masks, which
have small circular holes at their centers to prevent excess heating Of the e-beam
resists. The details of our sputtering procedure are the same as described in section
3.2.2 of this chapter, except we use different sample holders and rotary masks that
have only open and closed positions.

For the present current-induced magnetization switching (CIMS) studies, we usu—
ally make a trilayer structure of F1/N/F2. The polarizing F1 layer is thick enough so
that current doesn’t ﬂip its magnetization. The N layer is thick enough to eliminate
exchange coupling but thinner than its spin diffusion length. The F2 layer is thin
enough so that a high enough dc current can ﬂip its magnetization.

To get a nearly uniform current through a nanopillar, the sheet resistance, p/t, of
the contacts should be small compared to the resistance Of the sample. For this reason

we sputter 80nm Cu as a bottom contact which has 0.259 sheet resistance at room

52

temperature whereas the resistance Of a typical nanopillar is ~ 19. Then, we sputter
a F l / N / F2 trilayer, and then sputter 15-20nm Au to protect it from oxidation. Since
we will do several more processes on this multilayer, it is important to keep it from
oxidation. After the sputtering, we put the samples in acetone to lift-off. Fig. 3.9

shows a multilayer after this step.

Inc“ ' "a Eu".'-2‘.‘:' ﬁﬁ"h.‘* v . -: -~.- ~ .w-wrv- amen-(nasal?119*» r-r.m<-x-ﬂ~~.-rt:'
D
N ..

F1 1 = a
t '1 .1

CU ~
I n
'~.
I3? .. .‘.. ....t....- - .. _ '1- 'u._.-i.-..~o..: a:.~.'_u'..'.- ..1. '53»;

Figure 3.9: (a) A cross-sectional cartoon Of a multilayer. (b) An SEM image Of a
multilayer after lift-Off.

3.4.3 EBL for nanopillars and evaporation of Al
EBL for nanopillars

Now we have our multilayer sputtered, and we need to create 5 nanopillars on this
“big” sample (~ 8 X 50pm). The following is a recipe for opening nano—sized holes in

the e-beam resist.
1. Spin 9% MMA at 6000 rpm for 30 seconds and bake it for 5 minutes at 170°C.
2. Spin 2% PMMA at 6000 rpm for 30 seconds and bake it for 30 minutes at 170°C.

3. Load the sample into SEM and write 5 ellipses (~ 130 x 70 nm lateral dimen-
sions) 6pm apart from each other on the midline of the multilayer (see Fig. 3.10

(a)). The beam current is 9 pA and the area dose for the features is 475uC/cm2.

4. Develop the ellipses in e-beam developer as described in section 3.4.2.

53

5. Check the development using an Optical microscope. This step in the fabrication
is very crucial to get a well deﬁned elliptical nanopillar. Since their dimensions
are below Optical resolution, it is impossible to see the details of these ellipses
using an Optical microscope. However, when these features develop it is possible
to see some light coming out of them. There is no way of seeing whether these
ellipses are over or under developed, therefore there is always an uncertainty
in the dimensions of these holes. Fig. 3.10 (a) shows an optical image of the

multilayer with 5 elliptical Openings after this step.

 

Figure 3.10: (a) An optical image Of 5 elliptical Openings fabricated using EBL, the
rest of the substrate is covered with e-beam resist. Since the dimensions of ellipses
are below Optical resolution we only see some light coming out of these holes. Scale
as in Fig. 3.9. (b) A cartoon of the cross-section of multilayer + Al mask (c) An
SEM image Of an elliptical Al mask (top view).

54

Evaporation of Al

Now we have to evaporate Al into these elliptical openings to serve as a mask for

ion-milling.

1. Load the substrates into a thermal evaporator in a class-10000 clean room and
close the chamber to pump down to 1 — 2 x 10‘7 Torr. Fill the cold trap with

liquid nitrogen to get to l — 2 x 10‘7 Torr faster.

2. Evaporate 60 nm of Al. When Al is evaporated, it goes through the elliptical
Openings created by the EBL and sticks to the Au layer, making an elliptical

Al mask on the surface of the multilayer.

3. Put each substrate in a test tube and ﬁll the tube with acetone to lift off. TO
speed up the process Of lift-Off, heat the test tube on a hot plate at 80°C. It
takes about half an hour to lift-Off. Fig. 3.10 (b) shows a cartoon of the cross-
section of the multilayer + Al mask after lift-off and Fig. 3.10 (c) shows a top

view SEM image of an Al elliptical mask on top of the sputtered Au layer.

At this stage we can check the shape Of Al pillar masks using an SEM. As we
can see from Fig. 3.10 (c) the shape Of the Al mask is not perfectly elliptical and
the surface Of the Al pillar doesn’t seem to be ﬂat. The possible reasons for the

inaccuracies in the shape are:

a The resolution Of EBL is comparable to the size Of our nanOpillar.

o It is impossible to check the develOpment using an Optical microscope.

55

All of our Al masks appear like the one in Fig. 3.10(c), i.e. we don’t see a ﬂat
surface. We also made Au and Ag pillars Of the same dimensions and checked them
with an SEM to see whether they give a similar surface roughness. It seems from
Fig. 3.11 that if we evaporate Au or Ag in to the ~ 130 x 70 nm elliptical Openings
they tend to have a ﬂat surface. But Au and Ag are not good candidates to be used
as a mask for ion-milling because their ion—milling rates are much larger than the
ion-milling rate of Al (see Table 3.2). Although the surfaces of our Al masks are very
rough, this shouldn’t affect the ion-milling of the multilayer because the Al mask is

thick enough so even thinner parts don’t mill completely through.

 

(a) ""3606” (b) ""léotlt'”

Figure 3.11: (a) An SEM image of a Au pillar. (b) An SEM image of a Ag pillar.

3.4.4 Ion-milling and Insulation
EBL for Ion-milling

To ion-mill through the multilayer, we need to Open a window over the multilayer
using EBL, because we don’t want to ion-mill the leads. The following is a recipe for

opening a window using EBL.

1. Spin 9% MMA at 4000 rpm for 30 seconds and bake for 5 minutes at 170°C.
2. Spin 2% PMMA at 4000 rpm for 30 seconds and bake for 30 minutes at 170°C.

56

3. Load the sample in the SEM and write the pattern in Fig. 3.12. The beam cur-
rent is 25 pA. The area dose is 130/1C/cm2 for the inner pattern and 70/1C/cm2

for the outer one.

 

 

Figure 3.12: The DesignCAD drawing of a window for ion-milling. Outer rectangle
is exposed to less e-beam than the inner one to get a good undercut.

4. Develop the sample as described in section 3.4.2. Fig.3.14 (a) shows an image

Of a window over a multilayer after develOpment.

Ion-milling

Ion-milling is a process used to dry-etch metals with inert gas ions (we use argon).
The process is the Opposite of sputtering, in which the target atoms are knocked Off by
accelerated argon ions and deposited On a substrate. We ion-mill our samples using
a different vacuum chamber, which has a broad-beam ion source, a dc-magnetron
sputtering gun, an evaporation boat to evaporate silicon monoxide (SiO), and a ﬁlm
thickness monitor (FTM). This vacuum system also has a load-lock, which allows us
to insert and take out samples without venting the chamber.

Since we ion-mill through a multilayer Of different materials, we need to estimate
the ion—milling rates Of different materials from the ion-milling rate Of Au. TO measure

the ion-milling rate for Au, we:

1. Sputter ~ 400A Au on the FTM crystal.

57

2. Turn-on the ion-mill and let it run for 5 minutes to stabilize.

3. Move the FTM over the ion-mill and read the Au removal rate from the F TM

controller.

Once the ion-milling rate for Au is measured, we estimate the rates for other
metals using Table 3.2 and calculate the total time needed to ion-mill. Then the

sample is moved over the ion-mill and kept there for the duration calculated.

 

Metal Al Fe Co Ni Py Ru Ag Cu Au
Rate (A/min.) 730 530 550 660 634 610 2200 1100 1700

 

 

 

 

 

 

 

 

 

 

 

 

 

Table 3.2: Argon ion etch rates for the metals we etch at an ion current density Of 1
mA/cm2 and an ion energy of 500 eV. Values are taken from Ref. [92] except for Py,
which is calculated from the Ar ion etch rates of Ni and Fe.

Ion-milling around the Al masks

Fabrication of a F/ N/ F nanopillar requires ion-milling around the A1 masks which
are on top Of the multilayer. TO ion-mill we put the substrate in a special sample
holder (see Fig. 3.13) which has a stainless steel mask with a tiny hole at the center.
Therefore we expose only a small area at the center of the substrate to the argon
ions, which prevents the accelerated ions from burning the e-beam resist.

When we ion-mill around the Al masks, we etch the parts Of the multilayer which
are not covered with Al. The areas below the Al masks stay intact because the
Al masks get ion-milled instead of the multilayer, leaving a F/ N/ F nanopillar right
under the masks. Therefore these nanopillars have approximately the elliptical shape
and the size (~ 130 x 70 nm) of the Al masks. To eliminate the dipolar coupling

between two ferromagnets, we try to ion-mill approximately half-way through the

58

center of the substrate

 

Figure 3.13: A special sample holder for ion-milling. It has a heat sink to remove
heat from the substrate and a magnetic disk to hold it using a magnetic arm.

nonmagnetic metallic spacer layer (see Fig. 3.14 (b)) such that the top ferromagnet is
an elliptical nanomagnet and the bottom F-layer is left untouched as a big rectangular
(~ 8 x 50pm) ferromagnet. Since the magnetic poles of the ferromagnets are far apart,
there is no dipolar coupling.

Insulation

Right after ion milling, without breaking vacuum, we evaporate SiO tO insulate the
nanopillar as well as the top and bottom contacts to the nanopillar. To get a uniform
thickness of SiO we hold the sample holder using a magnetic arm and spin it at 60
rpm while evaporating SiO. We adjust the total thickness of evaporated SiO such
that ~ 10 nm Au is left outside Of the insulating layer for the top contact. However,
since we evaporate SiO all over the Opening for the ion-milling (see Fig. 3.14 (a)), we

actually insulate the nanopillar completely as seen in the cartoon of Fig. 3.14 (c).

59

 

Figure 3.14: (a) An Opening over the multilayer for ion-milling, created by EBL.
(b) A cartoon of the cross-section Of a magnetically uncoupled nanopillar. (c) A
cartoon of the cross-section Of a magnetically uncoupled nanopillar insulated by SiO
evaporation. Scale as in Fig. 3.9.

3.4.5 Low angle Ion-milling

Since the whole pillar region is insulated with SiO we must ion-mill away the excess
SiO that is insulating the top contact region of the nanopillar. To do this we use a
special sample holder (see Fig. 3.15 (a, b)), which has a cylindrical hole at the center
and a nearly vertical (3 degrees with the vertical) slit for the sample. The ion milling
time is estimated based on the thickness of the SiO layer and the ion-milling rate.
Fig. 3.15 (c) shows a cartoon Of the low angle ion-milling from one side. Later, the

substrate is ﬂipped and ion-milled from the other side as well.

60

Slit for the substrate

 

. .: We; Q3; ’
a) side view b) top view c) Low-angle ion-milling

Figure 3.15: A special sample holder for low—angle ion-milling with a circular hole

in the middle and a nearly vertical (3°) slit for the substrate.(a) Side view. (b)Top
view. (c) A representative cartoon Of low—angle ion—milling.

3.4.6 Lift-off and wet etch

After the low—angle ion-milling, we lift-off the SiO as described in section 3.4.3. Fig.

3.16 (a) shows a picture of a multilayer and SiO layer after lift off.

 

Figure 3.16: a) An Optical image after $10 evaporation and lift—Off. Scale as in Fig.
3.9. (b) A cross-sectional cartoon after SiO evaporation lift-Off. (c) A cross-sectional
cartoon after wet etch with KOH.

61

After lift-Off we have some Al left on top Of the nanopillar (see Fig. 3.16 (b)). We
etch this Al by dipping the substrate in photodeveloper (452, KOH) for ~ 4 minutes
and rinse with DI water and blow-dry immediately. Fig. 3.16 (c) shows a cartoon of

the cross-section of a nanopillar at this stage.

3.4.7 Top contacts

The ﬁnal step is to put top contacts to the nanopillars. The following is a recipe for

top contacts.

1. Spin 9% MMA at 4000 rpm for 30 seconds and bake for 5 minutes at 170°C.
2. Spin 2% PMMA at 4000 rpm for 30 seconds and bake for 30 minutes at 170°C.

3. Load the sample in the SEM and write the pattern in Figure 3.17. The beam
current is 25 pA. The area dose is 130pC/cm2 for the inner patterns and

70)uC/cm2 for the outer ones.

4. DevelOp the sample as described in section 3.4.2. Figure 3.18 (a) shows an

image Of the Openings for the top contacts after development.

5. After developing put the sample in the ion-milling chamber and ion-mill the pro-

truding Au pillars for 5 seconds to clean up the contact region Of the nanopillars.
6. Without breaking the vacuum sputter 100nm Au.

7. Lift-Off as described in section 3.4.3. Fig. 3.18 (b) shows an Optical image Of
the center Of the substrate after lift-Off. Fig. 3.18 (c) shows a cartoon of a

cross-section Of a ﬁnished nanopillar device.

62

||___'l|l—_.'I||—_.1||—__1||__1
(I I“ ll.| ll,l llll ||
| | l
I: III: II]: III: III: ll
I! II“ ll,I III III ll
'l ll, ll| ll,‘ llll ll
l
l | l | |
(I llll ll,l lllI llll ll
'1 IllI ”ll ”II ”II H
,l ll| ll| III ll. II
I
| l | l
'l lllI IIII Illl ll,l H
'l llll Illl ll,| ll,| II
I III Ill ll.I ll, ll
I: III: ”II ll, ll,I ll
| ”l
l I

ll ll
:I ll l
I

Figure 3.17: The DesignCAD drawing for top contacts to 5 nanopillars. The outer
rectangles are exposed to less e-beam than the inner ones to get a good undercut.

 

Figure 3.18: (a) Openings for 5 electrodes for top contacts created by EBL. (b) An
optical image showing the multilayer, SiO insulating layer and 5 electrodes for top
contacts. (c) A cartoon Of a cross-section Of ﬁnished nanopillar device. Scale as in
Fig. 3.9.

3.4.8 Measuring the Resistances of Nanopillars

To measure the resistance of the nanopillars, we have to put indium on the Au pads
of the substrate using a sharp tip soldering iron. Initially all the leads are shorted

to protect the nanopillar from burning out in the case Of a static electric discharge

63

from the operator. We mount the sample on a quick-dipper stick and solder the
current and voltage leads as shown in Fig. 3.19 (a). To measure a speciﬁc nanopillar
on the substrate, we scratch the shorting leads using a scriber. During soldering,
and scribing, we ground ourselves by wearing a grounding wrist strap. For another
level Of protection, the quick-dipper has shorting switches, which are shorted until
the measurement starts. We insert the sample in a room temperature electromagnet
with magnetic poles bored out, which produces 820 Oe/A magnetic ﬁeld. We place

the sample at the center of the electromagnet as seen in Fig. 3.19 (b).

electromagnet

pole piece

quick-dipper

 

: . ’ 145;.“
'sample at the center.’

(b)

 

Figure 3.19: (a) The sample is mounted on the quick—dipper and leads soldered with
indium. (b)Quick dipper is inserted into the electromagnet through the bored pole
pieces and the sample is positioned at the center of the electromagnet.

We connect the current leads to the dc current source and the voltage leads to the
pre—ampliﬁer. We built our own battery-powered dc current source the magnitude of
the output of which is set at 1.98 mA dc current per volt supplied by the lock-in (see
Fig. 3.20). The preampliﬁer sends the ampliﬁed ac voltage signal from the sample

to the lock-in ampliﬁer’s input. The differential resistance dV/ d1 is measured by the

lock-in ampliﬁer at frequency 8 kHz, ac measuring current ~ 2011A, and time constant

64

100 ms.

 

dc signal

 

 

dc current source _
ac Slgnal

 

 

 

 

 

Lock-in

 

 

 

+ input signal

 

 

Pre-amplifier

 

 

 

 

Figure 3.20: A diagram of resistance measurement using lock-in detection.

A computer running Labview controls the power supply for the electromagnet and
the lock-in. The measurement is fully computer controlled. We ﬁrst measure dV/dI
vs. the external magnetic ﬁeld H (applied along the easy axes Of nanopillars) of the
nanOpillar to get an idea how gOOd the sample is, i.e. if it has well deﬁned parallel
(P) and anti-parallel (AP) states. For our speciﬁc samples we usually sweep the ﬁeld
in steps between -1 kOe and +1 kOe. Then we sweep the dc current in steps passing
through the nanopillar to check for the current driven phenomenon. dV/dI vs. I
is measured with an effective sweep rate about 0.1 mA/s, calculated from a current
step of 0.1 mA and step time 10 times the lock-in time constant. Positive current 1+
represents current ﬂow from thick F1 to thin F2 i.e. electron ﬂow from F2 to F1. We

do two kinds Of measurement:

1. We keep the current ﬁxed and sweep the magnetic ﬁeld, which gives CPP-MR

data.

65

2. We keep the magnetic ﬁeld ﬁxed and sweep the dc current, which gives CIMS

data.

For example, Fig. 3.21 shows the CPP-MR and CIMS data of a Py/ Cu / Py nanopillar
at room temperature. Noticeably, the change in the resistance, AR 2 R(AP)-R(P),
of the sample is the same for both types of switching. The upward curvatures in the

CIMS data for large current are due tO joule heating.

 

1.5 2-..-.-WA

1.48} AP

1.
7'3

 

 

 

 

 

 

 

 

 

l A A A l A A A l A A A

0.4 L012“ 0 0.2
H(kOe)

 

 

Figure 3.21: Differential resistance (dV/dI) vs. H of a Py/Cu/Py nanopillar while
I=0 (on the left). dV/dI vs. the dc current (I) while H = 0 for the same sample.
We can also measure at 4.2K by dipping the quick-dipper into liquid helium. In
this case we use a superconducting solenoid magnet with a persistent switch to provide
the external magnetic ﬁeld. The persistent switch Of the superconducting magnet is

also controlled by the computer.

66

Chapter 4

Spin-memory—loss in sputtered
Pd,Pt and Au and at their

interfaces with Cu at 4.2K

4.1 Introduction

Spin-dependent bulk and interface scattering in ferromagnetic/nonmagnetic (F/N)
structures are the sources of Giant Magnetoresistance (GMR). Such scattering also
results in spin-memory-loss through spin-ﬂipping both in the layers and at interfaces.
The main parameter characterizing the spin-memory-loss and therefore determining
the MR is the spin-diffusion (spin-ﬂipping) length Is, that is the distance an electron
diffuses before its spin ﬂips. Currently there is no experiment that can directly give
the amount of spin-ﬂipping at a F/ N interface (see, however, [68]). However, hoping
to get insight into the spin-ﬂipping at F/ N interfaces, Park et al. [93] developed a
method to measure the spin-ﬂipping at N 1 / N2 interfaces. Park et al. [93] studied
spin-memory-losses in sputtered Ag, V, Nb and W and at Ag/Cu, V/Cu, Nb/Cu
and W/ Cu interfaces at 4.2K. They also measured the interface speciﬁc resistances of
N1 / N2 to look for patterns to see if the interface speciﬁc resistances could be estimated

from interfacial alloying of N1 and N2. In this chapter, we extend these studies to

67

determine the spin-diffusion lengths in sputtered Pd, Pt and Au and spin-memory-
losses at Pd / Cu, Pt/ Cu and Au/ Cu interfaces at 4.2K using the same method given
in Ref. [93].

Special interest in the transport pr0perties of Pt and Pd layers, and Pt' and Pd-
based multilayers, derives from several sources: the high polarizability of Pd, the large
atomic number of Pt, and that both produce magnetic anisotropies perpendicular to
the layer planes when very thin layers of these metals and Co are combined in mul-
tilayers [94]. To evaluate whether either metal could be used in magnetic multilayers
displaying GMR, it is necessary to know their transport properties.

So far, studies of CPP-MR in F / N multilayers involving noble metals have focused
upon the N-metals Cu or Ag [61]. Studies of the “inverse” phenomena, current-
induced moment switching or magnon generation in F / N trilayers or multilayers, have
also concentrated upon Cu [37, 95, 96]. For uses such as electron-beam lithographic
fabrication of nanowires, Au has the advantage over Cu and Ag that it does not
oxidize. If Au as the N-metal has small enough resistivity, and long enough spin-
ﬂipping length (spin-diffusion length), and if it produces CPP-MR5 comparable to
those for Cu and Ag, it could be advantageous for both scientiﬁc studies and devices.

In this chapter, we ﬁrst describe the spin-memory-loss detection technique and
then the experiments. Later, we divide the studies into two sections. The ﬁrst section
involves the spin-memory-loss in Pd and Pt, and at Pd / Cu and Pt / Cu interfaces. To
test the effect of spin-orbit scattering, we compare the spin-memory-losses at Pd / Cu
and Pt/ Cu with Nb/Cu and W/ Cu, respectively, as Pd and Nb are in the 5th, and

Pt and W are in the 6th row of the periodic table. The second section involves the

68

spin—memory-loss in Au and at Au/ Cu interfaces. To check whether Au as a spacer
layer gives comparable CPP-MR to Cu and Ag, we compare the MRs of Co—based
exchange-biased spin valves with Au,Cu and Ag Spacer layers. We also compare the
spin-memory-losses at Au/Cu and Pt/Cu interfaces to test whether the dominant
source of spin-flipping is the spin-orbit scattering. If so, then the heavy metals like
Pt and Au should give large spin-ﬂipping and therefore short 13,. In addition, we also
include a comparison between our results and later theoretical calculations of interface
speciﬁc resistances by Xia et al.[97, 98]. Most parts of this chapter are published in

Ref. [99] and [100].
4.2 Spin-memory-loss detection technique

We obtain the interface resistances and spin-memory-loss results using a CPP “spin-
memory-loss detector”[101, 93, 61], which is a Py-based exchange-biased spin-valve

(EBSV) sandwiched between two Nb strips.

The structure of the CPP spin-memory-loss detector is:
Nb(200)/Cu(10)/FeMn(8)/Py(24)/Cu(10)/X/Cu(10)/Py(24)/Cu(10)/Nb(200),

where all the thicknesses are in nanometers. At 4.2K the Nb superconducting leads
provide a uniform current through a well deﬁned area A. The magnetization of
one 24 nm thick Py layer is pinned by an 8 nm thick adjacent antiferromagnetic
FeMn layer, and the magnetization of the other 24 nm thick Py layer is free to
rotate from parallel (P) to antiparallel (AP) to that of the pinned layer in a small
magnetic ﬁeld (~20 0e) [101, 93]. X is either a single layer N-metal or a multilayer

of alternating N1 / N2 metals to be studied. We measure the speciﬁc resistances (area

69

of CPP current ﬂow A times resistance R) AR(AP) and AR(P), as well as their
difference, AAR = AR(AP) — AR(P). By varying the thickness of a single layer of
N-metal (tN) or varying the number of bilayers (N) of N1 / N2, one can make a plot of
AAR vs. tN or N. Valet-Fert [12] analysis of these plots give the spin—memory-loss in
bulk N-metal or at N1 / N2 interfaces. For the spin-memory-loss detector AAR decays

exponentially due to the spin-ﬂipping as:

o If X is a single layer N-metal with thickness tN, then AAR oc exp(—tN/l§f),

where 1?, is the spin diffusion length in N-metal.

o If X is an alternating multilayer of N1 /N2 with N bilayers, then AAR oc
exp(—2N5N1/N2), where 6N1 /N2 measures the spin-memory-loss at an N1/N2

interface [93, 61].

The interface speciﬁc resistance ARM/cu is derived from the slope of AR(AP) vs.

N graph for X=[N / Cu] N inserts, which is given by
2ARN/Cu + pNtN + pCutCu, (41)

where the factor of 2 comes from 2 interfaces for one layer of N-metal.

For X=N(t) inserts, AAR ﬁrst drops fast due to the formation of two N/Cu
interfaces then continues to decrease more slowly as tN increases. To derive 1:} we
analyze the data for thicknesses larger than the total thickness of the two N /Cu
interfaces (2tN/Cu), which include a constant contribution from two N/ Cu interfaces
in addition to the contribution from bulk N-layers. The VP equation for an EBSV

with a single N-metal insert with thickness tN reduces to

70

exp ("11131- — 26N/Cu)
3f

AARK
ARactz've (AP)

(4.2)

For the Py-based EBSV, since 13” << tpy the ARactwe(AP) includes only the “active”
region in the EBSV that is bounded by the amount of 13" in the 24 nm thick Py-layers
and the inserts between them (see Chapter 2). Therefore the denominator in Eqn.
4.2 includes an additional speciﬁc resistance pNtN + 2ARN/Cu.

Similarly, the VP equation for the EBSVs with X=[N / Cu] N inserts reduces to

exp [- (26N/Cu + 551' + :-8&) N]
sf sf

AAR oc
ARactz've (AP)

 

(4.3)

The denominator of Eqn. 4.3 includes an additional speciﬁc resistance N (pNtN +
poutcu + 2ARN/Cu) because the X=[N/ Cu] N insert is in the active region. Since 1:}
and (SN/Cu appear simultaneously in Eqns. 4.2 and 4.3, we ﬁnd their values by self-
consistent numerical ﬁts to our data. For the numerical VF analysis, the parameters
for the other constituents of the Py-based EBSV are: ppy = 123 an, ﬂpy = 0.73,
p32,, = 263.3 nom, (5.” = 5.5 nm, way/C, = 0.7, 2AR;,y/Cu = 1.00 i: 0.08 mm2,
ppeMn = 875 j: 50 an and ARFeMn/py = 1.0 :l: 0.4 me2 taken from Ref. [102],
whereas pm is measured using the van der Pauw technique (see Appendix A). This
method can be used to determine the spin-memory-loss at any N1 / N2 interface, but
in that case the interface speciﬁc resistances and spin-memory-loss of both N1 / Cu

and N2/ Cu must be known.

71

4.3 Experiment

The samples are sputter deposited onto (100) Si substrates. The details of our sample
fabrication and measurement techniques are described in Section 3.1. To study the
spin-memory-loss within the Pd, Pt and Au layers, we inserted certain thicknesses
(1-20 nm) of each in the spin-memory-loss detector as described in Section 4.2. To
study the spin-memory-loss at the Pd / Cu, Pt / Cu and Au / Cu interfaces we inserted
[Pd(3nm)/Cu(3nm)]~, [Pt(3nm)/Cu(3nm)]N and [Pt(3nm)/Cu(3nm)]N bilayers in
the spin-memory-loss detector. We chose to make 3 nm thick layers of Pd, Pt, Au
and Cu to eliminate the possibility of pin holes within the sputtered layers of these
metals. The measuring uncertainties in individual values of AR(AP) (and in the larger
values of AAR) are dominated by a i570 uncertainty in our ability to determine the
area A. Uncertainties in AAR much larger than this value are indicated by error bars.
Data ﬂuctuations illustrate the reproducibilities of AR(AP) or AAR for independent
samples. We also made 300 nm thick sputtered thin ﬁlms of Pd and Pt and 200
nm thick Au ﬁlms to measure their resistivities. The van der Pauw measurements
of sputtered Pd, Pt and Au ﬁlms at 4.2K gave residual resistivities of PM = 40 :l: 3
an, ppt = 42 :l: 6 an, and pm, 2 19 i 6 an (see Appendix A).

The typical sputtering rates for the constituents of the spin-memory-loss detector
are given in Table 4.1. The sputtering rates for the metals given in Table 4.1 are
not exactly the same for every sample but within 10% of those listed in Table 4.1.
The main reason for these variations is the gradual reduction in the volume of the

sputtering targets. All metals except FeMn are dc-triode sputtered using 2” diameter

72

targets. FeMn is dc-magnetron sputtered using a 1” diameter target, therefore the

deposition rate of FeMn is much smaller.

 

Metal Nb Py FeMn Cu Pd Pt Au
Deposition Rate (A/s) 5.5 5.5 0.5 10 4.5 7 5

 

 

 

 

 

 

 

 

 

 

 

Table 4.1: Typical sputter deposition rates of Nb, Py, FeMn, Cu, Pd, Pt and Au.

4.3.1 Examples of AR(H) with X inserts

As examples of our data, Fig. 4.1 shows the resistive “hysteresis curves”, AR(H)S,
for EBSVs with no insert and 10 nm thick inserts of Pd, Pt and Au. Fig. 4.1 (a), (b)
and ((1) show both minor loops, where the pinned layer remains pinned, and major
loops, when the ﬁeld H is raised until the pinned layer unpins. Notice that both
major and minor loops show the same constant values of AR(AP) as H increases.
Also, the values of AR(P) are the same for large positive and negative ﬁelds. The
switching ﬁelds of the pinned Py-layer are different for magnetic ﬁelds applied in
different directions due to the shift in the hysteresis curve of the pinned Py-layer.
The decreases in AAR from (a) to (b), (c), or (d) are due to the spin-ﬂipping in the
bulk and the formation of two Pd/ Cu, Pt/Cu or Au/ Cu interfaces. AAR for the
sample in Fig. 4.1 (c) is very small due to the strong spin—ﬂipping in both the 10
nm thick Pt insert and the two Pt/ Cu interfaces. To measure such small changes,
we measure the resistance at a ﬁxed ﬁeld 100 times and take the average. Since
the EBSVs are in the AP state at 50 Oe and in P state at -200 and 400 Oe (see
Fig. 4.1 (a), (b) and (d)), we measured the resistance of this sample ﬁrst at -200
Oe, where the magnetizations of Py-layers are aligned parallel, then at 50 Oe, where

the magnetization of the free Py-layer is ﬂipped causing an anti-parallel alignment

73

and ﬁnally at 400 Oe, where both magnetizations are again aligned parallel. We also

repeated these measurements one more time in the same order as before.

 

 

 

 

 

 

 

 

 

 

 

. a noX . b - 10
28.01() AP 21.8 ”W ) AP
27.6« -. ‘ Fe
1 q \7 21.6- A]
27.21 1 \, 1 [
1 1 1 1 1 l .
NA 26.8 '. [ “A 21.41 1 ‘1 o
1 l 1 , l 1 l
5 26-4- ‘ 1 CF? ‘ 1 l 1
b 1“] ] 2: 21.2. 1 ]
< \11 < b l.
256 .2 ‘_\ 21.0‘ p \1.“ X ,
252’.“ - 1M" '\\:>~<>.
1' 6 20-8 ' I V I ' I IIIIIII
-300 -200 -1160 200 300 400 -300 -200 -100 o 100 200 300 400
H(Oe) H(Oe)
(c) X:Pt(10) . (d) X=Au(10)
26.52- iAP 224‘ AP
‘ 222
26.48 l \. l
“A “A 220- [
5 26.44 5 21.8. l
03 ‘ E 21.6:
< 26.40- [ < . [
‘ P P 210%»
1
26.36 .............. 21.2 ..............
-200 -100 o 100 200 300 400 -300 -200 100 o 100 200 300 400
H(Oe) H(Oe)

Figure 4.1: Examples of AR vs. H for Py-based
X=Pd(10), (c) X=Pt(10) and (d) X=Au(10)

EBSVs (a) with no insert X, (b)

Fig. 4.2 shows examples of variations in AR(H) vs. N of EBSVs with [Pd(3)/Cu(3)] N

inserts. Fig. 4.3 shows examples of variations in AR(H) vs. N of EBSVS with

[Pt(3) /Cu(3)] N inserts.

74

(a) no x , (b) X=[Pd(3)lCu(3)]N=2

 

 

 

 

 

 

 

 

 

 

 

 

‘ l ]
21.54 \ | '
”A "A 21 s ] | [
E 21.0 E l
53 . $3 I
(I 20.5] 1‘ (I 21.6 i [
4 , \ < 1
200 ‘\ 214 )@ \N o\
' ‘An‘ W? ' - ;\\\:_:V’ W
195 l
’-300 -200 450 0 100 200 300 400 -300 .200 -100 0 100 200 300 400
H (06) H (00)
. (c) X-.-[Pd(3)/Cu(3)]N=5 . (d) X=[Pd(3)lCu(3)]N=10
262“: i 30.404-
2022« ,
20.20; 30.400
“E 20.10; N
g . :5; 30.396
cc 20.101 0:
< 1 <
20.14. 30.392-
20.12. g 3 ‘
.............. 30388.-.-.......-.-
300 -200 .100 0 100 200 300 300 -200 .100 0 100 200 300
H(013) H (09)

Figure 4.2: ABS vs. H for some EBSV samples (a) with no insert X (b) with
X=[Pd(3) /Cu(3)] 10:2 (c) with X=[Pd(3)/Cu(3)]N=5 (d) with X=[Pd(3)/Cu(3)]N=10.
Each data point in (c) and ((1) represents an average of 100 measurements, and each
error bar represents the standard deviation of the mean.

4.4 Spin-Memory-Loss at 4.2K in Sputtered Pd,
Pt, and at Pd/ Cu and Pt / Cu interfaces

We start with inserts of just Pd or Pt. Fig. 4.4 shows AAR vs. thickness, tpd or tpt,
of Pd (ﬁlled triangles and solid curve) or Pt (ﬁlled diamonds and solid curve). For
comparison, we show also previously published data for inserts of Nb (open triangles
and dashed curve) and W (open diamonds and dashed curve) [93]. In all four cases,
AAR behaves as expected, ﬁrst decreasing rapidly as the two interfaces with Cu form

(due both to the additional resistance of these interfaces and to spin-memory loss at

75

 

 

 

 

 

 

 

 

 

 

23.51 (‘l "0" 30.10. (b) X=[Pt(3)/CU(3)] N=1
1
23°] 30.05. i
22.5. . j 1
1 30.00.
«A 22.09 ...A
a 1 E ‘
1:, 21.51 [ g 29.951
0: ‘ . ‘ c: ‘
<1: 21.0] K [ < 29.904 i [
20.51 k ‘
x W 29.85~ l
20.0-.-,..-.......
-300 -200 -100 0 100 200 300 400 -200 -100 o 100 200 300
H (08) H (09)
(e) X=[Pt(3)/CU(3)] N=2 . (d) X=[Pt(3)/Cu(3)] N=3
‘ 30.4821
30.84 1
30.450- l
“7‘ 30.82~ NA 30,450:
E E
53 l “q 30 4441 "
I: 30.80« E ‘ A
< < ‘ 0
30.4381 ]_
30.78- ‘
[ 30.432. 1
-260f-160ﬁ 6 - T P 260 ' 36? .2'00 ' 4'00 ' 6 P 160 ' 260T 360 j
H (09) H (08)

Figure 4.3: ABS vs. H for some EBSV samples (a) with no insert X (b) with
X=[Pt(3)/Cu(3)]~=1 (c) with X=[Pt(3)/Cu(3)]~=2 (d) with X=[Pt(3)/Cu(3)]N=3.
Each data point in (b), (c) and (d) represents an average of 100 measurements, and
each error bar represents the standard deviation of the mean.

the interfaces), and then more slowly due to spin-memory loss (ﬁnite spin-diffusion
length) in the bulk metal. The ﬁts of the Pd and Pt data in the large t regimes give
l5d=25f§° nm and If} =14 j: 6 nm, where the larger uncertainty on the “up side”
for Pd arises because the technique is less reliable when 1,; is larger than the largest

layer thickness (20 nm) used. These values are comparable to those for Nb and W,

19]" 2:251?0 nm and Ir; = 4.8 :t 1 nm [93]. They are also comparable to the mean-free

76

paths Apd = 21 :l: 5 nm and Apt = 20 :t 6 nm estimated from the equation

2
Mum) = (700 —> 12%)me ’ (4.4)

 

where the residual resistivity p is in an [103]. Comparable values of [sf and /\
can be understood if the defects producing the residual resistivities are much lighter
than Pd and Pt, because the large atomic number difference between the host and
scatterer should produce strong spin-orbit scattering, but the likely small differences
(almost unity) in the number of tranSport electrons / atom should produce only modest
transport scattering. For example, the mirror image case of Pd or Pt impurities in Cu
gives only moderate residual resistivities per atomic percent impurity (~ 1chm/at.%
for Pd and ~ 2uQcm/at.% for Pt), [103] but relatively large spin-orbit cross sections,
0,0 ~ 2.5 x 10‘18 cm2 for Pd and ~ 4 x 10‘17 cm2 for Pt [104]. Putting appropriate

values from Refs. [103] and [104] into the Valet-Fert [12] equation

M30
6 ’

 

ls; = (4.5)

which relates lsf to the product of the transport mean-free path A and the spin-orbit
mean-free path A30, gives estimates of [sf/)1 ~ 1 for Pt in Cu and ~ 2.5 for Pd in Cu.

The curves shown in Fig. 4.4 are ﬁts to all of the data by means of numerical
solutions of the appropriate equations of Valet and Fert [12] using the independently
determined values of the properties of the EBSV from Ref. [102], the van der Pauw
layer resistivities for Pd and Pt noted above, and the interface speciﬁc resistances
from the data shown in Fig. 4.5, plus values of 1?, and (SN/Cu found self-consistently

from the data of Fig. 4.4 and the data for multilayer inserts in Fig. 4.6. As we will

77

10 r’ﬁ"’_‘.'"—'T— V—"T "T’w'T"‘-V WfﬂTV‘T'M'w‘TF—i+ “H U—T"? "_”T"'-' _"Y'——"_ 1'. If?

4-4.L1LITLJR 4. A .»_. l_'.,

AAR (mm?)

 

l
A
l__ l_._.LE1 14412;.w4. LAA. ;_.l4 4

5 10 15 20
t (nm)

N
01

Figure 4.4: AAR vs. tpd or tpy for Pd (ﬁlled triangles and solid curve), Pt (ﬁlled
diamonds and solid curve), Nb (open triangles and dashed curve), and W (open
diamonds and dashed curve). The dotted curves represent the behaviors of AAR
expected for no spin-memory-loss in either the bulk Pd or Pt or at the Pd/Cu or
Pt / Cu interfaces.

describe below, the self-consistent solution for Pd yields a slightly different value of
If,“ = 27 nm, which lies within the uncertainty of the simpler ﬁt described above.
Fig. 4.5 shows how AR(AP) increases with an increase in N for inserts of
[Pt(3)/Cu(3)]N and [Pd(3)/Cu(3)]N, as well as for [Nb(3)/Cu(3)]N and [W(3)/Cu(3)]N.
Least-squares straight line ﬁts to the data give (after correcting for the contributions
of the 3 nm thick Pd, Pt, and Cu layers) 2Ade/Cu = 0.9:l:0.1 me2 and QARpt/Cu =
1.5:l:0.1 film”. These values are smaller than those for 2ARNb/0u = 2.2:t0.3 me2 and
2AI'1’w/Cu = 3.1 :t 0.2 f9m2. If these ARs are due simply to interfacial alloying, then
they are consistent with the smaller resistivities per atomic percent impurity given in
Ref. [103] for Cu with Pd (~ luQcm/at.% impurity) and Pt (~ 2uQcm/at.% impu-

rity) than those for Cu with Mo (we use Mo here because no value is given in [103]

78

100 r—_—Y—' ‘T 'MT‘ﬁTra‘ ”7" " f“ ’Y i’Y'* ’77 a- " fﬁiTi‘ﬁ R‘Y"__H‘—_——“T—_ 7*"? ‘T'Uj

    

 

, ,o[W/Cu]

80 ’
NE I I .
“f: 6° , ’ ale/Cul [Pt/Cu]
O. ’ i 2 ’ .
s ’
“<5 40 ~ d/Cu]

20

0 L_._L_.A'_h_x_m_r_1_ ;_ .'_ ;_ .._ i.__x_._L_‘_ .l_._1__ ._____ 2 , J
0 5 1 0 1 5 20 25
N

Figure 4.5: AR(AP) vs. N for [Pd/Cu]N and [Pt/Cu]~ inserts compared with
[Nb/Cu]~ and [W/Cu]~. The symbols are the same as in Fig. 4.4. The lines are
least square ﬁts.

for Nb in Cu) and W (both ~ 4chm/at.% impurity). But such a simple comparison
should be treated cautiously, both because the thicknesses of the interfacial alloys
might differ for the different metal pairs, and because band structure effects could be
important [105, 97].

Last, Fig. 4.6 compares the decreases in AAR vs. N for inserts of [Pd(3)/Cu(3)]~
or [Pt(3)/Cu(3)]~ with similar data for [Nb(3)/Cu(3)]N and [W(3)/Cu(3)]N. Simple
linear ﬁts to the data in the form ln(AAR) vs. N determine values (after correcting for
spin-ﬂipping in the Pt or Pd layers) of (Sm/Cu = 0.24i3;3§ and 5Pt/Cu = 0.9i0.1, where
we’ve allotted greater uncertainty toward large 6Pd/Cu (i.e., we have determined this
uncertainty by underweighting the data points for large N). These values of 6Pd/Cu
and 6Pt/Cu are comparable to those of (Sm/Cu 2' 0.19 i 0.05 and (SW/Cu = 0.96 :t 0.1,

respectively. If spin-orbit scattering in alloyed interfaces is the dominant source of

79

 

    
 

 

 

 

10 7 l I r 1 r
1 . -:
“E
Nb/Cu
g 0.1 : I l 1
(I : :
g 1
j
A .
0.01 5- 3
= 1
0001 . 7 . . . . L I l l . L
0 2 4 6 8 10 12

Figure 4.6: AAR vs. N for [Pd/Cu]~, [Pt/Cu]N, [Nb/Cu]N, and [W/Cu]~ inserts.
Note the similarity of behaviors of [Pd/Cub; and [Nb/Cu]N and of [Pt/Cu]N and
[W/ Cu] IV- The symbols are the same as in Fig. 4.4. The lines are least square ﬁts.

6, then we would expect 6Pd/Cu z 6Nb/Cu [and (Sm/Cu z (SW/Cu, since Pd and Pt are,
respectively, in the same chemical rows in the periodic table as Nb and W. Indeed,
these pairs of data sets in Fig. 4.6 are very similar. Again, however, band structure
details may complicate the analysis.

If we simply combine the best ﬁt values of 6Pd/Cu and (Sm/Cu with the best ﬁt values
of If!" 2 25 nm and If; 2 14 nm from an equivalent ﬁt to the data in Fig. 4.4, then
the resulting Valet-Fert analysis of the data in Fig. 4.4 falls above the data. That is,
these best ﬁt values of 6 are slightly too small to account for the reduction in AAR
in Fig. 4.4 due to the formation of Pd/ Cu and Pt/ Cu interfaces. The curves shown
in Fig. 4.4 use the alternative parameters 15" = 27 nm, lff‘ = 14 nm, (5193/00 = 0.30,
and 6Pt/Cu = 0.99. Since all four values lie within the uncertainties in the original

values, we retain the original values and their uncertainties as our best estimates.

80

 

AAR (tom?)

.0
d

0.01 L 1 1 L i l 1 1 #141 . L r 1 l l #1 l 4; l l

0 5 10 15 20 25

1 (nm)

Figure 4.7: AAR vs. tN for N 2 Au (ﬁlled circles and thick solid curve), V (dotted
curve), Nb (dashed curve), Pd (broken curve), and Pt (thin solid curve). The datum
at t Au = 15 nm is only a lower bound, as its AR(H) curve did not give a ﬂat AR(AP)
region. That point is not included in the ﬁt to the Au data.

4.5 Spin-Memory—Loss and CPP-MR at 4.2K in
Sputtered Multilayers with Au

Fig. 4.7 shows AAR vs. Au thickness, t Au. For comparison, we show also the best
ﬁt curves to previously published data for inserts of V (dotted curve), Nb (dashed
curve) [93], Pd (broken curve) and Pt (thin solid curve). In all cases, AAR behaves as
expected, ﬁrst decreasing rapidly as the two interfaces with Cu form (due both to the
additional resistance of these interfaces and to spin-memory loss at the interfaces),
and then more slowly due to spin-memory loss (ﬁnite spin-diffusion length) in the
bulk metal. The slope of the Au data in the large t M regime gives [f]? = 351?5 nm,

where the uncertainty on the high side is large because the technique is less sensitive

to values of lff“ much larger than the largest layer thickness used (here t M = 20 nm).

81

7o ﬁ‘ﬂT-“Wﬁ‘fkff 1 1 T Y ‘7 f V‘ I f

[ [V/Cu]

L”. [Nib/Cu] ‘ [Pt/CU]
50 [‘ , ' .
1 / [Pd/Cu] -‘

AR(AP) (mez)

[Au/Cu]

 

25

Figure 4.8: AR(AP) vs N for [Au/Cub; inserts compared with [V/Cu]N , [Nb/ Cu] N,
[Pd /Cul1v, and [Pt/ Cu] N. The symbols and lines are the same as in Fig. 4.7.

Such a range of values is comparable to those for V, Nb, and Pd, but smaller than
for Pt [93, 99].

The curves in Fig. 4.7 are ﬁts to all of the data by means of numerical solutions of
the appropriate equations of Valet and Fert [12] using the independently determined
values of the properties of the EBSV given in Ref. [102], the van der Pauw layer
resistivities (for Au, the value given above), and the interface speciﬁc resistances and
values of 6 from the data in Figs. 4.8 and 4.9. To obtain the Au ﬁt shown, we used
IQ“ = 35 nm and (SM/Cu = 0.20, the latter near the upper range of uncertainty of
6Au/Cu (see later). The curve using the best ﬁt to Fig. 4.9 alone, 6Au/Cu = 0.13, would
have essentially the same “slope”, but would fall above the data.

Fig. 4.8 shows how AR(AP) increases with increasing N for inserts of [Au(3)/Cu(3)] N
as well as for [V(3)/Cu(3)]~, [Nb(3)/Cu(3)]~, [Pd(3)/Cu(3)]~, and [Pt(3)/Cu(3)]~.

A least-squares straight line ﬁt to the Au data gives (after correcting for the con-

82

 

    

 

 

 

‘ 1
N5 ,.
=3 0.1 : -:
3 = :
0.01 : _:
= [Pd/Cu] ‘ . -
[Pt/Cu] [NB/Cu] [Au/Cu]
0.001 1 I 1 1 _l 1 1 1 1 l 1 1 1 1 i 1 L 1 J j 1 1 ‘_ 1
0 5 10 15 20 25

N

Figure 4.9: AAR vs. N for [Au/Cu]N, [V/Cu]N, [Nb/Cu]N, [Pd/Cu]N, and [Pt/Cu]~
inserts. The symbols and lines are the same as in Fig. 4.7. AAR for the N = 15
datum dropped by a factor of ten over three successive measurements. We show
the average and ﬂuctuations, but neglect this point in our ﬁt. In contrast, repeated
measurements of the N = 10 and 20 data remained stable to within the uncertainty
bars shown.

tributions of the 3 nm thick Au and Cu layers), 2ARAu/Cu = 0.35:3:3g fflm2, which
overlaps within mutual uncertainties of the value 2ARAu/Cu = 0.30 i 0.01 ern2 pre-
viously derived using a completely different method [106]. ARM/Cu is smaller than
ARV/Cu: ARNb/Cua [7], ARPd/Cu and ARPt/Cu [99]-

Fig. 4.9 compares AAR vs. N for [Au(3)/Cu(3)]~ [V(3)/Cu(3)]N, [Nb(3)/Cu(3)]~,
[Pd(3)/Cu(3)]N, and [Pt(3)/Cu(3)]~ inserts. Self-consistent VF analyses of the data
for [Au(3)/Cu(3)]~ and the data in Figs. 4.7 and 4.8 give (SM/Cu = 0.13f3j33, larger
than for [V/ Cu] N, but smaller than for [Nb/ Cu] N, [Pd/ Cu] N, and [Pt/ Cu] N [93, 99].
The curve shown for Au/ Cu is for (SM/Cu = 0.13, the best ﬁt to the data of Fig. 4.9

alone. The value of 5.40/00 = 0.20 used in Fig. 4.7 would correspond to a line going

just above the datum at N = 10. Given the small values of AAR for N = 20, such a

83

 

 

AR (152 m2)

 

 

13.5 ‘A ., ’
13.3
1;
Co/Cu
12.9 ‘2 ... m - , ‘°' ’
12.7 . 1 1 1 % 1
- 1500 -1000 -500 0 500 1000 1500

H (08)
Figure 4.10: AR(H) vs. H for representative Co/Ag, Co/Au, and Co/Cu EBSVs
with too = 6 nm and tN = 10 nm for N = Ag, Au, or Cu. The data around H = 0
overlap for minor and major loops; the former extend up to about + 250 Oe, the
latter to + 1500 Oe.
line cannot be ruled out.

Finally, to illustrate the size of the CPP-MR for Co/Au, Fig. 4.10 compares
representative hysteresis curves of AAR vs. H for Co-based EBSVs with Co thickness,
tCo = 6 nm, and ﬁxed thickness tN = 10 nm for N=Cu, Ag, and Au. The values of
AR(P), averaged over two or more samples each, increase from about 14 to 15 to 16
fme2 from Cu to Au to Ag, but the data sets overlap within mutual ﬂuctuations of

i1.5 — 2 mez. The equivalent averages for AAR increase from about 0.28 to 0.35 to

0.48 film2 from Cu to Au to Ag, with uncertainties of $0.03 — 0.05 fflm2 each [69].

84

 

 

 

 

N pN (an) 1?; (nm) (SN/Cu 2ARN/Cu (meQ)
Pd 40 :t 3 251;” 0.241383 0.9 i 0.1
Ft 4210 14i6 0940.1 1540.1
Au 19 i 6 35125 0.13333 0.35f3;3,5

 

 

 

 

 

 

 

Table 4.2: A table of resistivities and spin-diffusion lengths of Pd, Pt and Au, and
the spin-memory-losses at Pd/ Cu, Pt/ Cu and Au / Cu and the speciﬁc resistances of
these interfaces at 4.2K.

4.6 Summary

We measured the spin-diffusion lengths of our sputtered Pd, Pt and Au layers as well
as the spin-memory-losses at Pd/ Cu, Pt / Cu and Au/ Cu interfaces at 4.2K using the
spin-memory-loss detector [93]. Table 4.2 shows the measured values of resistivities
and spin-diffusion lengths of sputtered Pd, Pt and Au, the spin-memory-losses (6) at
Pd/ Cu, Pt/ Cu and Au / Cu interfaces and interface speciﬁc resistances of 2Ade/Cu,
2ARpt/Cu and 2ARAu/Cu. The spin-diffusion lengths in sputtered Pd and Pt are long
enough so that nanometer thick layers of either could be used in magnetic multilayers
without excessive spin—ﬂipping (spin-memory-loss) in the layers themselves. The re-
sistivity of sputtered gold, p Au z 192t6 an, is low enough, the spin-diffusion length
in sputtered Au, If,“ 2 35:?5 nm, long enough, and the CPP- AAR of Co/ Au EBSVs
with t Au = 10 nm is large enough, that 10 nm thick Au layers can be used in magnetic
multilayers both for studies of basic physics and for devices. Later, J i et al. measured
the spin diffusion length of sputtered Au in lateral spin valves using a completely dif-
ferent technique [107] and found If!" = 63 :1: 15 nm at 10K, which is in agreement
with our result at 4.2K. The interface speciﬁc resistance, 2ARAu/Cu = 0.354135%: fﬂmz,
overlaps an earlier value [106].

The interface speciﬁc resistances, Ade/Cu and ARpt/Cu, are smaller than those

85

for ARM/Cu and ARw/Cu. In contrast, the spin-ﬂipping probabilities at Pd/ Cu and
Pt/ Cu interfaces are, respectively, very similar to those at Nb/ Cu and W/ Cu inter-
faces. These last similarities are consistent with spin-orbit scattering in interfacial
alloys being the source of those probabilities. However, although Au is next to Pt
in the same row of periodic table the spin-memory-loss at Au/ Cu interface is much
smaller than that for Pt / Cu and W/ Cu, which may be due to the low resistivity and

and the absence of d-electron states at the Au Fermi surface.
4.7 Comparison of measured ABS with theory

We compare the interface speciﬁc resistances measured with superconducting leads in
the CPP direction with later no—free—parameter calculations with real Fermi surfaces
by Xia et al.[97, 98] assuming perfect specular interfaces and 50—50 alloy interfaces.
Interestingly, the samples fabricated using sputtering show good agreement with the
calculated values of ARs if the lattice mismatch at the interfaces is very small, i.e
the difference between the lattice parameters of the metals forming the interface is
less than 1 % (see Table 4.3). Even, the calculated ARs for perfect and 50—50 alloy
Pd / Cu, Pt / Cu and Au / Cu interfaces are not bad, differing only about 50% to a factor

of 2 from our experimentally measured ARs (see Table 4.3).

86

 

 

 

 

 

 

 

 

 

 

M1 / M2 2AR (exp) 2AR (perfect) 2AR (50-50) a(1,"”/a{,”2 Aao
(“21112) (mm?) (film?) (A) (‘70)

Ag/Au 0.1[106] 0.09[98] 0.12[98] 4.09/4.08[1] 0.24
Co(fcc)/Cu 1.0[61] 0.9[98] 1.1[98] 3.56/3.61[1] 1.4
Fe/Cr 1.6[108] 1.5[105],1.9[98] 1.6[98] 2.87/2.88[1] 0.35
Pd/Cu 0.9 :5 0.1 1.5[109] 1.6[109] 3.89/3.61[1] 7.8
Pt/Cu 1.5 :1: 0.1 3.92/3.61[1] 8.6
Au/Cu 0.351835 0.45[109] 0.7[109] 4.08/3.61[1] 13

 

Table 4.3: A comparison between the experimental and theoretical values of ARs at
4.2K. Calculations agree with the experimental values if the lattice parameters of the

metals constituting the interface are close.

87

 

Chapter 5

Changes in magnetic scattering
anisotropy at a ferromagnetic /
superconducting interface

5.0.1 Overview

There is now great interest in both static and transport properties of ferromag-
netic/superconducting (F/ S) metallic interfaces. Examples of topics of interest in-
clude: reductions in the superconducting transition temperature Tc upon injection
of a polarized current [110]; propagation of a polarized current through a supercon-
ductor [111]; proximity effects between S and F metals [112, 113, 114, 115]; the F/S
interface resistance [116]; predictions that the current-perpendicular-to—plane (CPP)
magnetoresistance (MR) should be zero when measured with superconducting leads
[117]; and subsequent arguments that it won’t be if strong spin-ﬂipping is present
[118], or if exchange splitting between spin up and spin down electron bands plays an
important role [119, 120].

CPP-MR experiments with superconducting Nb leads in contact with F- metals
Co, Fe, Ni and the F-based alloys Py and Coo_91Feo,09 (Refs. [61, 121, 122]) showed

that F/ S interfaces contribute only a constant term to each current channel of the

88

2CSR model, with interfacial scattering parameter ’yp/S = 0. Moreover, inserting
10 nm of Cu or Ag between the F -metal Co and the S-metal Nb appeared to leave
yp/S z 0 (Refs. [61, 121]) quite different from the values of 700/61,, z 7Co/Ag z 0.8 for
Co/ Cu and Co/ Ag interfaces not in contact with a superconductor [61, 31]. Direct
measurements for simple Nb/F/Nb and Nb/N/F/N/Nb sandwiches with different
thicknesses of F showed that N = Cu or Ag inserts produced no systematic changes
in AR [123].

K. Bid [57] measured the changes in the scattering anisotropy at a C0/ Nb interface
at 4.2K upon inserting Cu, FeMn and Ru and combinations of these metals in between
superconducting Nb and ferromagnetic Co [57]. To detect very small changes in the
CPP-MR, he designed a hybrid spin-valve that would work as a null detector i.e.
when there is no insert between Co and Nb the MR is zero. The use of the null
detector provides a more precise check for the effect of inserts. Bid [57] observed
that the scattering asymmetry changed for inserts of the combinations of Ru/ Cu,
Cu/FeMn but stayed the same for inserts of just Cu and FeMn between Co and Nb.
We shall call such changes in the MR of the null detector (or scattering anisotropy)
‘magnetic activation’ of F/S interfaces. Here, we conﬁrm Eid’s results by repeating
some of his studies and extending them to additional inserts to understand how the

scattering asymmetry is affected. Most of this chapter is published in Ref. [124].

5.0.2 Introduction

We present evidence of a new phenomenon at F/S interfaces at 4.2K, activation of

the magnetic scattering anisotropy at C0/ Nb interfaces when certain non-magnetic

89

metals (or combinations) X are inserted between the Co and Nb as part of a care-
fully designed CPP-MR hybrid spin-valve (SV) [61]. More precisely, as described in
detail below, the SV is designed so that when X is absent, the system SV/ Nb pro-
duces almost zero change in speciﬁc resistance, AAR = AR(AP) - AR(P), between
the parallel (P) and anti-parallel (AP) magnetic orderings of the two F-layers in the
SV. Here A is the cross-sectional area through which the CPP current ﬂows at 4.2K
(see inset in Fig. 5.1). Such a device is, thus, very sensitive to small changes in
the magnetotransport properties caused by inserting X. We show that inserting X
consisting of some metals or combinations of metals leaves AAR unchanged, which
we interpret as leaving the scattering asymmetry at the F/ X S interface, ”yr/XS, ap-
proximately 0. In contrast, insertion of other metals or combinations changes AAR
(i.e., giving non-zero Viv/x3), and whether AAR becomes more positive or negative
is set by the previously known sign of 'yp/X, the anisotropy without S. “Activation”

of such changes in AAR seems to require strong spin ﬂipping between F and S.

5.0.3 Experiment and Results

In this study we focus upon yp/XS. We justify this focus by showing ﬁrst that none
of the inserts X of interest change AR between Co and Nb, and that all show least
a partial proximity eﬂect (i.e., Cooper pairs propagate from S to Co through X).
Details of our sample preparation, and measuring techniques are given in Chapter 3.
Fig. 5.1 shows AR for 20 nm of Co alone, and with inserts of N = 4 nm of Ru, 6 nm
of Au, or 2 nm of FeMn. 20 nm of Co is thin enough so that the data all lie near the

extrapolated value for zero Co thickness. The lack of signiﬁcant increases in AR in

90

Fig. 5.1 shows at least a partial proximity effect in N, especially for FeMn (p ~ 850
an)[93] and Ru (p ~ 100 nflm)[125], where AR should otherwise have increased
by ~ 3.5 and 0.8 me2, respectively, more than twice the actual deviations from the
value for Co(20). FeMn is also unique in the present study, in that a layer only 1
nm thick generates very strong spin-ﬂipping, [93] and that various of our unpublished

studies lead us to believe that F/FeMn interfaces are magnetically inactive (i.e., have

”VF/FeMn = 0)-

ARs for [Nb/”Nb] sandwiches

141

12 ‘ FeMnZ/CoZO/FeMnZ
10 -
1 Ru4/0020/Ru4

8 J I I I
0020 Au6lCo20/Au6

AR (mm’)

 

 

 

Figure 5.1: ABS for Nb/X/Nb sandwiches with X =Co(20), Ru(4)/Co(20)/Ru(4),
Au(6)/Co(20)/Au(6), and FeMn(2)/Co(20)/FeMn(2). Thicknesses are in nanome-
ters. Inset: The circular sample with bottom (V-, I-) and top (V+, I+) crossed Nb
strips. Current (I) and voltage (V) connections are indicated. A is the area through
which the current ﬂows at 4.2K.

To isolate the behavior of a single Co/ X / N b structure, thus focusing upon 7Co/xs,
we constructed a SV for which AAR is near zero when X is absent. We then examined
the effect of inserting X. We achieved the clear AP and P states required for a

hybrid SV by choosing F-metals and thicknesses to give very different saturation ﬁelds,

91

Py(24) (H, N 20 Oe) and Co(2) (H, N 300 Oe). Here, and hereafter, all thicknesses
are in nm. Our SV has the form Nb/Cu(1)/Py(24)/Cu(20)/Ru(2)/Co(2)/X/Nb.
Signiﬁcant spin-ﬂipping in the Py layer [122] and at Cu/Ru [125] and Co/Cu [68]
interfaces means that the speciﬁc resistances of such a spin-valve cannot be described
using a simple two-current series-resistor (2CSR) model. However, a modiﬁed 2CSR
model so greatly simpliﬁes the explanation of our experiment, that we use it, subject
to the caveat that our explanation is only schematic, strengthened by our belief that
the model does not distort the essential features of the argument. If we assume that
the Co/ Nb interface is magnetically inactive (i.e. that it has Woo/N5 = 0), then for no
X, the 2CSR model taking account of strong spin-flipping only within the Py would

give [61, 23, 12] [Chapter 2]:

AAR 0‘ (ﬂPypf’ylffy 'l' 7Py/CuARFDy/Cu)(HCoPE:otCo + ’YCo/RuARE'o/Ru)1 (5-1)

where 6;» is the bulk scattering anisotropy in metal F, p} = (pr + pf?) / 4, and If]? is
the spin-diffusion length in Py [61, 122]. The essential features of Eqn. 5.1 are: (1)
that the contributions from the Py and the Co appear separately as a product; and
(2) because ”Yea/Ru is negative [125], the two terms in the second parentheses have
Opposite signs. Choosing tCo = 2 nm should bring the sum of those two terms close
to zero [61, 125]. Fig. 5.2 shows that the resulting AAR ~ —0.05 fflm2 with no X is
indeed much smaller than the value AAR ~ +0.4 ﬁlm2 that we ﬁnd [57] when there
is no Ru to the left of the Co(2) layer in the SV and also no X.

If, now, we insert into Eqn. 5.1 a layer X as listed above, and if the Co/ X interface

is magnetically active, (Vac/x3 7E 0), then this interface contributes an additional term

92

AR (mmz)

14.2
14.1
14.[-0
13.9
13.8
13.7
13.6
13.5

13.4
-1000

[MM 50lCu1/Py24lCu20/Ru2IC02IﬂNb1 50]

 

W3

I ' I

Y I I

 

L l 1 l

a”

AP

Au2

A: “114 A )-‘~24
wﬁér‘kw
Wu.

noX

‘

 

 

—500 0

H (06)

500

1 000

Figure 5.2: AR vs. H for Nb150/Cu1/Py24/Cu20/Ru2/Co2/X/Nb150 with no X,
X =Ru(2) and X =Au(2). The AP states are labeled for each case.

to Eqn. 5.1, giving

AAR (X (IBPyp‘Pylffy + VPy/CuAREy/Cu) (ﬁCopaotCo + 7Co/RuAR2‘o/Ru + A/Co/XSARE’o/x)-

(5.2)

To focus upon ’YCo/xs. we have assumed that ARGO /XS — _ARCo/X ~ 0.5 fﬂmz,

the already measured and very similar values of AREo/X for X =Cu, Ag, Ru, and Au.

If this additional term is positive (1100/X5 > 0), AAR should become more positive

and if it is negative (VCo/Xs < 0), AAR should become more negative. Fig. 5.2 shows

how AAR became more negative upon insertion of X = 2 nm of Ru, but larger and

positive upon insertion of X = 2 nm of Au.

In Fig. 5.3 we collect together AAR data for a series of different metals and com-

binations of metals. The uncertainty bars for each X were obtained by summing the

93

squares of the individual uncertainties in AAR, taking the square root, and dividing
by the number of samples measured. The values of AAR and the number of samples
used in the statistics are given in Table 5.1. To within measuring uncertainties, X
= Cu, Ag, FeMn, Ag/Cu, and Cu/Ag, all leave AAR unchanged. In contrast, X
= Ru, Ru/Cu, and Ru/FeMn all make AAR more negative, and X = Au, Cu/Ru,
Cu/FeMn, Cu/ Au, and Ag/ Ru all invert the sign of AAR, making AAR positive. To
within experimental uncertainties, the positive values for Au, Cu/ Ru, Cu/Au, and
Ag/Ru, are all the same (with that for Cu/FeMn somewhat larger). These similar
values without FeMn are consistent with similar values of ’YCo/x for X =Cu [61], Ag
[23], and Au. While, strictly, we have not measured 700/411: and ARE/WAu separately,
the similarity of our (mostly unpublished) CPP-MR data for Co or Py with Au to
those for Co or Py with Cu and Ag strongly suggests that the interfacial parameters
for Au are similar to those for Cu and Ag. It is of key importance to note that AAR
becomes more negative whenever Ru is next to the Co, and more positive when Cu
or Ag is next to the Co but separated from the Nb by Ru, Au, or FeMn. These be-
haviors show that the change in AAR is due to a change in the Co/ X interface, since
’YCo/Ru < 0, but yea/Cu a: 700/119 > 0. We consider next why some X leave AAR un-
changed (i.e., leave Woo/x5 = 0 in the presence of superconducting Nb), while others

change it (i.e., give I’YCo/XSI > 0).
5.0.4 Discussion

As noted in the introduction, Taddei et a1. argued [118] that the presence of strong

spin-ﬂipping is necessary for the appearance of a full CPP-MR. While there are ques-

94

5 Nb(150)/Cu(1)/Py(24)/Cu(20)/Ru(2)/Co(2)/ X/Nb(150)

 

 

1 Cu(10)/FeMn(2)
0'10‘ AW) l Ag(10)/Ru(2)
c 10 IR 2
1 *1 1"”
m" Cu(10)lAu(2)
E 000
g ' Ag(10) Ag(10)/Cu(10)
[ES-0.05: i l,: M 2 Ru(2)/Cu(10) [ l
1 00(10) 9 "( ) Cu(10)/Ag(10)
1 noX
0104 f [
[ Ru(2)
-0.15- Ru(2)/FeMn(2)

Figure 5.3: AAR for samples of the form Nb(150)/Cu(1)/Py(24)/Cu(20)/Ru(2)/
Co(2) / X /Nb(150) for a variety of X, including no X. Samples with italic labels
leave ’Yco/xs = 0, whereas samples with bold labels give l7Co/XS] > 0.

tions about the applicability of their argument, no other simple potential explanation
for the different behaviors shown in Figs. 5.2 and 5.3 presents itself to us. We, thus,
ask whether the data in Figs. 5.2 and 5.3 can be understood in terms of different
amounts of spin-ﬂipping between the Co and Nb for different X, with weak Spin—
ﬂipping leaving X inactive (i.e., yea/x5 = 0) but strong spin-ﬂipping making it active
(7Co/XS > 0)1

All of the metallic layers N except FeMn are much thinner than their respective
bulk spin-diffusion lengths. Thus, spin-ﬂipping within those metals cannot be the

source of the differences in their activities. The only potential source is spin-ﬂipping

95

 

 

 

X [# of samples] AAR (film?)

 

 

 

 

 

 

no X [6] —0.057 i 0.008
Cu(10)[7] —0.039 :1: 0.008
Ag(10) )[2] —0.037 i 0.014

FeMn(2 )2[2] —0.029 :1: 0.014

Au(2 ('2) _ 0.060 3: 0.012

Ru(2 ) ’3‘ —0.093 :t 0.009

 

 

 

Ru2( )/Cu(16) 1 —0.080 :t 0.019
ISu(10)/Ru(2 ) :4: 0.044 :1: 0.009
Ru2( )/FeMn(2) [1] —0.099 i 0.023
Cu(10)/FeMn(2 ) [3] 0.101 :1: 0.012
Cu(10) /Au(2) [1] 0.059 3: 0.022
Ag(10)/Cu( 1(0) [2] —0.046 :1: 0.020
Ag(10) /Ru(2 )) [2] 0.066i0.014
Cu(10)/Ag(10) [2] —0.039 :1: 0.013

 

 

 

 

 

 

 

 

 

 

 

 

 

Table 5.1: AARs and the number of samples used to determine the error bars for
inserts X in Fig. 5.3

at the N/ S interface. Unfortunately, we do not have a way to measure spin-memory
loss at the interfaces of these metals with superconducting Nb. We can, however,
measure such loss at interfaces with non-superconducting Nb [93]. Since Co/ Au and
Co/ Ru have their interface magnetic anisotropies activated when the Au or Ru are
in contact with superconducting Nb, whereas Co/Cu does not, we examined spin-
memory loss at the interfaces of these three metals with non-superconducting Nb,
using the method described in Ref. [93]. In that method, a multilayer of the form
[N (3nm) / Nb(3nm)] N, Where N is the number of bilayers, is inserted into the middle
of a Py-based spin-valve, in which one Py-layer is exchange-bias pinned and the other
is left free to reverse in a small magnetic ﬁeld. Insertion of the multilayer should then

cause AAR to decrease exponentially with N as
AAR OC exp(—2NdN/Nb), (5.3)
where (SN/Nb characterizes the probability of spin-ﬂipping at the N / Nb interface. Fig.

96

AAR (me2)
.o

0.01

 

0.001 L. ..-..._.......-;......L..-- ...... L.--...-_--...i._...;--,._,........_. 1.- .. ....-.” .-l...__..:....-i.... .......i... 4...”).-. ...'.. _l

0 2 4 6 8 1 0 12
N
IX(3)/~b(3)1

Figure 5.4: AAR vs. N for inserts X=[Ru(3)/Nb(3)]~, X=[Au(3)/Nb(3)]1v
and X=[Cu(3)/Nb(3)]1v in exchange-biased spin valves of the form
Nb(150)/Cu(10)/FeMn(8)/Py(24)/Cu(10)/X/Cu(10)/Py(24)/Cu(10)/Nb(150).

5.4 shows a plot of log(AAR) vs. N for N = Cu, Au, or Ru. The Au/Nb and Ru/Nb
data decrease much more rapidly with increasing N than do the data for Cu/ Nb,
showing that spin-flipping is several times stronger at Au / Nb and Ru/ Nb interfaces
than at Cu/ Nb interfaces, when the Nb is normal. Since increasing N also increases
the amount of N = Cu, Au, or Ru, as well as of Nb, in the spin-valve, the data of Fig.
5.4 must be corrected for spin-ﬂipping within the “bulk” of these layers to correctly
obtain (SN/Nb. These corrections are, however, modest in size and similar for Cu, Au,
and Ru. Thus, they do not affect the relative sizes of (SN/Nb inferred directly from Fig.
5.4. Assuming that spin-ﬂipping at Au / Nb and Ru / Nb interfaces is also stronger than
at Cu/Nb when the Nb is superconducting, we have a rationale for activation of the

Co/ Au and Co/ Ru interfaces by spin-ﬂipping at the Au/ Nb and Ru/ Nb interfaces.

Turning, lastly, to FeMn, if ”yea/FeMn = 0 as noted above, then strong spin-ﬂipping

97

 

in the FeMn can only activate (30/ X if X = N / F eMn, where FeMn is inserted between
N and the superconducting Nb. Indeed, we see in Fig. 5.3 that the Co/FeMn interface

is not activated, but the Co/Cu/FeMn interface is.

5.0.5 Summary

We constructed a multilayer that let us measure anisotropy in spin-dependent scat-
tering at a single F / XS interface, when the non-ferromagnetic metal (or combination)
X is sandwiched between the ferromagnetic metal F and superconducting (S) Nb. We
assume that the proximity effect allows Cooper pairs to pass through X, so that we
are still studying an effectively F / S interface. Our results conﬁrm earlier conclusions
[61] that such an F / X interface is magnetically inactive (insertion of X produces little
or no change in AAR) when X =Cu, Ag, or FeMn alone, or combinations involving
only Ag and Cu. In contrast, when X =Ru or Au, or combinations involving either of
these two metals or FeMn that itself is not in contact with the Co, the F /X S inter-
face becomes magnetically active, producing signiﬁcant changes in AAR including, in
some cases, changes in sign. For X =Cu, Au, and Ru, we showed that these different
magnetic activities correlate with the strength of spin-ﬂipping at the interfaces of
X with non-superconducting Nb. Qualitatively these results are consistent with the
arguments of Taddei et al. [117, 118] that spin-ﬂipping can affect the CPP-MR when
the sample has superconducting leads. To determine if this argument provides the

correct explanation for our data, further theoretical analysis is needed.

98

 

 

Chapter 6

Controlled normal and inverse
CIMS and MR in magnetic
nanopillars

6.1 Introduction

Although CIMS is expected to result from spin polarization of the current, it has yet
to be shown that CIMS can be manipulated (e.g., inverted) by changing that polar-
ization. Prior studies of CIMS in F1/N/F2 nanopillars, where F1 is a thick (ﬁxed)
ferromagnet, N is a normal metal and F2 is a thin (free) ferromagnet which is free to
switch with a high-density spin polarized dc current, involved only positive scattering
anisotropies in both F-metals and at F / N interfaces, i.e. minority electrons are scat-
tered more strongly in F1, F2 and at F1 / N and N/ F2 interfaces. The current is then
positively spin polarized in the F layers, i.e., carried mainly by majority electrons. In
such “standard” conditions, electrons ﬂowing from F1 to F2 (negative charge current,
I < 0) switch the moment M; of F2 from antiparallel (AP) to M1 (high resistance R)
to parallel (P) to M1 (low R). Conversely, positive I (> 0) switches F2 from P to AP.
We call these behaviors “normal” CIMS and normal current-perpendicular—to—plane

(CPP) magnetoresistance (MR). We present CIMS experiments exploiting the possi-

99

bility of inverting the spin anisotropy by doping F1 or F2 with an impurity (Cr) that
scatters majority spin electrons more strongly [61, 126, 127, 108, 128, 129, 105, 97].
We thus show, for the ﬁrst time, that inversion of the spin anisotropy can invert the
CIMS direction, i.e., invert the signs of I for AP to P and P to AP transitions. We
also ﬁnd inversions of the MR (larger R for the P state) with appropriately doped
samples, as expected from prior CPP-MR results at low temperature [61, 126]. Mea-
surements at 295 and 4.2 K show that the qualitative behaviors of both CIMS and
MR are independent of temperature over this range for our samples.

Analysis of the switching behaviors lets us discriminate between models of CIMS.
We divide the standard models of spin-transfer torque (STT) used to describe CIMS
into two classes, ballistic [33, 34, 130] and diffusive [86, 35, 73, 131, 132, 74]. Both pre-
dict that changing scattering anisotropies can invert the MR and/ or CIMS. However,
their expectations need not agree. In ballistic transport, the spin anisotropy comes
only from reﬂections at the F/N interfaces. Inverting CIMS is predicted to require
negative anisotropy at F1 / N [130]. Inverting the MR should require opposite scatter-
ing anisotropies at F1 / N and N/F2 [61]. In diffusive transport, the spin anisotropy
of scattering within the F layers is also important, so that one must consider the
net anisotropy of each F layer (i.e., the resultant effect of the bulk of F and its F / N
interface). An additional effect, spin accumulation, can either support or compete
with the effect of polarized current [86, 35, 73, 131, 132, 74]. In this Chapter, in
addition to determining the relation between spin anisotropies and CIMS direction,

we answer four questions relevant to understanding CIMS.

100

 

1. Is the CIMS direction set only by interface scattering anisotropy? No.
2. Can impurity scattering within the layers be important? Yes.

3. Do the anisotropies of F1 and F2 play different roles for the CIMS direction of

F2? Yes.
4. Can spin accumulation be important? Yes.

Most of this Chapter is published in Ref. [89].

6.2 Experiment and Results

Sample preparation and measurement techniques are described in detail in Chapter 3.
Here, we brieﬂy note that our F1 / N / F2 nanopillars have approximately 130 x 70 nm
lateral dimensions, are ion-milled only partly through N to minimize dipolar magnetic
coupling between F1 and F2, and I + represents electron ﬂow from F2 to F1.

To determine how changing spin anisotropies changes CIMS directions, we com-
bine in different ways three pairs of materials: Py / Cu (Py= Ni84Fe16), with both bulk
and interface anisotrOpies positive [61]; Fe(Cr) / Cr [Fe(Cr)=Fe95Cr5], with both nega-
tive [126, 127, 108, 129, 105, 97]; and Ni(Cr) / Cu [Ni(Cr)=Ni95Cr5], with thick enough
Ni(Cr) so its negative anisotropy dominates the positive anisotropy of the Ni(Cr) / Cu
interface [133]. The net anisotropy is found using the MR. Here, we present the gen-
eral results of the study but the details of samples with Fe(Cr) will be given in the
thesis of Al-Haj Darwish.

Py and Py-Cu interfaces both have positive scattering anisotropy [61]. In accord

with prior data [48], Fig. 6.1 shows that Py(24)/Cu(10)/Py(6) nanopillars (layer

101

 

 

 

thicknesses in nm) give normal MR and normal CIMS. At both 295 and 4.2K, the
MR transitions from P to AP occur after H passes through zero, consistent with
little or no magnetic coupling. The agreement between minimum and maximum val—
ues of dV/dI for the MR and CIMS curves shows that the switching is complete.
Figures 6.2 and 6.3 also show weak coupling and complete switching. In contrast
to Py, Ni(Cr) has negative scattering anisotropy but Py/ Cu and Ni(Cr)/Cu inter-
faces both have positive scattering anisotropies. Ni(Cr)/Cu systems are special in
the sense that they have opposite bulk and interface anisotropies that allows us
to investigate the ballistic and diffusive models of transport for CIMS. Figure 6.2
shows the case, Ni(Cr)(20)/Cu(20)/Py(10). Combining net negative anisotropy for
Ni(Cr) with net positive anisotropy for Py gives the expected inverse MR, and now
inverse CIMS. Figure 6.3 shows a way to achieve inverse MR with normal CIMS, us-
ing Py(24)/Cu(10)/Ni(Cr)(4). For each sample, the switching directions of MR and
CIMS are the same at 295K and 4.2K. Each switching behavior was independently
reproduced, and no inconsistent switching was seen.

As expected for the MR [61, 126], when the net scattering anisotrOpies for F1 and
F2 are the same (Fig. 6.1), the MR is normal, and when they are opposite (Figs. 6.3
and 6.2), the MR is inverse. New for CIMS, when the net scattering anisotropy for
F1 is positive, CIMS is normal (Figs. 6.1 and 6.3), and when it is negative (Fig. 6.2),
CIMS is inverse. For these samples, the direction of CIMS is set by the net scattering
anisotropy of F1 and is independent of that of F2. Table 9.2 lists the signs for F1, F2
bulk, F1 /N, N/F 2 interfaces, and F1(net), F2(net), as well as for the observed MR

and CIMS. The MR and CIMS ﬁgures of samples involving Fe(Cr) are given in Ref.

102

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

1.65
1.60»
51.55.
'1':
5 1.50—
'5
1.45}
1.40. g . . _ ,295",
1.4. é“. T TAP,
1.3. $1.0. P.
a 1: . > I.
1:; 1,2- 04 0 0.4
S "(“991 AP
'0 1.1_ K J .
1.0 :
P ... 4.2K
-4 -2 o 2 4 6
|(mA)

Figure 6.1: Py(24)/Cu(10)/Py(6) data at 295K (top) and 4.2K (bottom) showing
normal MR (dV/dI vs. H at I = 0) in the insets and normal CIMS for dV/dI vs. I
in the main ﬁgures at H = 0 0e for 295K and at H = 20 Oe for 4.2K. In all ﬁgures,
I > 0 represents electron ﬂow from F2 to F1.

 

 

 

 

 

 

[89].
F1 / N/ F2 F1 F1 / N F1(net) F2 N / F2 F2(net) MR CIMS
PY/CU/Py + + + + + + + +
Ni(Cr)/Cu/Py - + - + + + - -
Py/Cu/Ni(Cr) + + + - + - - +
Fe(Cr)/Cr/Fe(Cr) - - - - - - + - I
Py/Cu/Cr/Fe(Cr) + + + - - - - + i

 

 

 

 

 

 

 

 

 

 

 

Table 6.1: Scattering anisotropies (+=positive, -=negative) of F1, F1 /N, and their
net anisotropy F1(net), those of F2, N/F2, and their net anisotropy F2(net), and the
observed MR and CIMS (+=normal, -=inverse).

Comparing the results of Py/Cu/Cr/Fe(Cr) and Py/Cu/Ni(Cr), which both have

inverse MR and normal CIMS, and Fe(Cr)/Cr/Fe(Cr) with normal MR and inverse

103

 

 

NlCr/Cu/Py

  

 

 

 

 

 

3.14
3.12 .
3.10 -
E". 3.08 l
a
% 3.06 ,
3.04 .
3.02 ..
3.00 A
$2.30
2.45 g
‘ >
82 4o U2.25 '05 0 0.5 4
: ' H (ROB)
g 2.35 \P
2 30 AP\ _\ 4.2K
. ::\-:: ——,—««-’/l
“5 W
I (mA)

Figure 6.2: Ni(Cr)(20)/Cu(20)/Py(10) data at 295K (top) and 4.2K (bottom) show-
ing inverse MR (dV/dI vs. H at I = 0) in the insets and inverse CIMS for dV/dI vs.

I at H = 0 in the main ﬁgures.
Py/Cu/NICr
1.12

 

   

dV/dl (a)

1.11

 
 

dV/dl (0)

1.1 .

 

 

 

l (mA)

Figure 6.3: Py(24)/Cu(10)/Ni(Cr)(4) data at 4.2K showing inverse MR (dV/dI vs.
H at I = 0) in the inset and normal CIMS for dV/dI vs. I at H = 0 in the main

ﬁgure.

104

 

 

CIMS and Ni(Cr)/Cu/Py with inverse MR and inverse CIMS given in table 9.2,
shows that, when bulk scattering predominates, the CIMS direction is independent
of the scattering anisotropy of F1 /N and N/F2. Finally, dominance of the bulk
contribution of scattering anisotropy in either F1 [e.g., Ni(Cr) in Fig. 6.2] or F2
[Ni(Cr) in Fig. 6.3] is inconsistent with ballistic transport through the nanopillar,
where the interfaces must dominate the scattering. While ballistic STT models cannot
describe our data involving Ni(Cr), the CIMS directions in all cases studied accord
with the ballistic prediction of [130] if the scattering anisotropy at the F1 / N interface
is simply replaced by the net anisotropy for F1. For diffusive transport, the current
polarization in N depends upon the net scattering anisotropies of both F1 and F2, and
CIMS depends upon both the spin-polarized charge current and spin-accumulation
effects [86, 35, 73, 131, 74]. Equation (1) reproduces Eq. 5 of Ref. [74] for the torque
I‘P at a small angle from the P state (for I‘AP, replace P by AP). Our notations for

F1 and F2 are reversed from [74].

P P 'P P
_I:_ = [{ Upgl’N + JUSN} (1 _ e—tN/AN) + {UFZI’FI +j7l;,F1} e—tn/AN] (6.1)

XMZ X (M2 X M1)

Equation 6.1 comes from an extension of the Valet-Fert [12] model of CPP-MR
to noncollinear states. up is the Fermi velocity in N, mi, and jﬁN are the spin
accumulation density [73, 131, 132, 74, 12] and spin-current density in N just outside
the F2—N interface calculated for the P state, m’F’l and jg, F1 are the same quantities
in F1 just inside the Fl-N interface, t N and AN are the thickness and mean-free path

of N, and NI is a unit vector in the direction of M. The second set of {} braces

105

dominates the usual case when tN << AN and the ﬁrst dominates if tN >> AN. The
signs of spin current and spin accumulation can support each other or compete.
Using the best parameters from CPP-MR experiments [61, 133], we calculate spin
currents and spin accumulations [61, 12], and insert them into PP or PAP. Except for
Fig. 6.2, the signs of spin current and accumulation always agree and are as expected
from the sign of the net spin anisotropy of F1 seen by MR. The case for Fig. 6.2
is more complex. For P to AP, the spin accumulation dominates PP and gives the
observed inverse CIMS at I < 0; due to the particular parameters of Ni(Cr) and Py,
the spin current alone would predict normal CIMS. For AP to P, the spin current
dominates PAP and gives the observed inverse CIMS at I > 0. Thus, we reproduce

the behaviors in Fig. 6.2.

6.3 Summary

We have shown that judiciously chosen pairs of ferromagnetic metals or alloys can
produce all four combinations of normal and inverse MR and current-induced mag-
netization switching at both 4.2 and 295 K. The MR is normal if the net scattering
anisotropies of F1 and F2 have the same sign, and inverse if they do not. For the
samples studied, the CIMS direction is set solely by the net anisotropy for F1, al-
though in Fig. 6.2 this result requires dominance of spin accumulation for the P to AP
transition. This latter result, as well as the inverted MRs in Figs. 6.2 and 6.3, show
that the interpretation of MR and CIMS must generally take account not only of the
interface scattering assumed in ballistic models but also the scattering (and diffusion)

within the F layers. As the widely accepted mechanism of CIMS is a quasi-interfacial

106

 

 

absorption of the transverse component of the spin current [33, 130, 73, 131, 74], the
importance of scattering within the F layers might seem surprising. However, in a
noncollinear magnetic conﬁguration, the transverse spin current in the frame of F2
is related to the longitudinal one in F1, and a global treatment [73, 131, 74] of the
longitudinal and transverse components of the spin current and spin accumulation

requires the diffusive aspects of the CPP-MR theory [61, 12].

107

Chapter 7

CIMS in permalloy-based
nanopillars with Cu, Ag, and Au
spacers

7. 1 Introduction

Most experimental studies of current-induced magnetization switching (CIMS) in fer-
romagnetic/ nonmagnetic/ ferromagnetic F / N / F trilayer metal nanopillars have used
N:Cu as the Spacer layer [46, 134, 76, 135, 136, 137, 41, 48]. None has yet used
NzAg or Au. Ag has the potential advantage for devices of sometimes giving a
larger Current-Perpendicular-to-Plane (CPP) magnetoresistance (MR) with Permal—
loy [Py=Ni,84Fe,16] [138]. Au has the advantage of being insensitive to atmospheric
contamination. We, thus, decided to compare MR and CIMS data at 295K and 4.2K
for Py-based nanopillars with N=Cu, Ag, and Au. Most of this chapter is published

in Ref. [139].
7 .2 Experiment and Results

We fabricated magnetically uncoupled Py-based nanopillars with a standard structure

of Cu(80)/Py(24)/N(10)/Py(6)/Au(115), where thicknesses are in nm, N=Cu, Ag,

108

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Py24/Cu10/Py6 Py24/Ag 1 0le6 Py24/Au1 O/Py6
1 3.08 ”71
1.48] MR '4: 1"" lg] 3.064 MR [Th] 1.764 MR 1] [l
i l] i] 304. I '[ ~ ”51 [11‘
A 1.464 ] l ] A ' ] ] A 1 [(1.
a l ] - 1 : S 4 g 1.741 l l]
i 1.44] An‘ I ‘ 'ﬁ 302‘ [ ] . i: 1.73] ']]I
5 1 ( 1 5 l ‘1 ' €1.72) ‘11.
1'42] ll [ 1 2981 ] J ] 1.71?1 [[11
140]”,wﬂ' l l =.[ [V - v 2.96 L] \ 17o}———————"" L-—————
' $0 0.5 0.0 0.5 1.0 43 0.5 0.0 0.5 1.0 -13 05 03 03 1.0
H (kOe) H (kOe) H (kOe)
l 3.08. 1'77l ‘ \
L48 ‘ CIMS / rk‘\\_\ 3.06] ClMS (A"\\\\ 176] CIMS [ \]\\
] .. WU 1.754
a 1.43 A 3.04\ ] A 174] ]
c . . .
.1; ] [ a f g 3-02‘] \5. ] [ ‘2’ 173*! )
\ 1.44 - ‘ . \ f‘ x . 4;
> > \ > . '
n ] ] S 1, 3.00] \ ] ] U 172]\
1.421 2.98] I 1 7, j \4.
] l ] l l \
-2 -... _) 2.964 \ ) 1,70]
1 .40 —“V“q“ +::~—r———v-' [ i x v v W ”I ~—v—-v—v— r v r r v *T r
-6 4 -2 o 2 4 6 -6 -4 -2 o 2 4 6 —6 4 -2 o 2 4 6
Hm) I (NW | (MN -1-

Figure 7.1: MR (top) and CIMS (bottom) at 295K for Py/N/Py nanopillars with
N=Cu (left), Ag (middle) and Au (right). Layer thicknesses are in nm.

or Au is the Spacer layer, the thick Cu layer is the bottom contact, and the thick
Au layer is the top contact. I + represents electron ﬂow from Py(6) to Py(24). The
approximate sputter deposition rates for the metals and alloys are given in Table 8.1.

The details of the sample fabrication and measurement are given in Chapter 3.

 

Target Py Cu Ag Au
Deposition rate (A/s) 5 10 8 5

 

 

 

 

 

 

 

 

Table 7.1: Approximate sputter deposition rates of Py, Cu, Ag and Au.

At 4.2K, 10 nm is much less than our best estimates of the spin-diﬂ'usion lengths
in Cu (0.5:3j35pm) and Ag (0.4 :1: 0.2pm) [69] and three times less than our minimum
estimate [100] for Au (35125 nm). The room temperature spin-diffusion lengths are
probably factors of 2 — 3 smaller, [140, 141] still much longer than 10 nm for Cu and
Ag, but no longer so for Au. Thus, 10 nm thick Au as an interlayer is expected to

give smaller MR than Ag and Cu.

109

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Py24/Cu10/Py6 Py24/Ag10/Py6 Py24/Au10/Py6
2.65
1.12. 123.
MR MR ﬂ [7 MR ﬂ P
2.60.
g "08‘ g g 1.24.
- — 2.55. -
U . 2 2
% "°‘ .2 3 ,2,
1.00] 2.50« ' l
0.96. 245‘ 1.16
-1.0 0.5 do 0:5 To 1.0 05 0.0 Tots 1‘0 40 0.5700 T015 ' 1'.o
H (kOe) H (kOe) H (kOe)
‘ CIMS
2.65-
1.124 CIMS ‘ CIMS 1.28‘
‘ 2.601 A 1
:g‘ "08‘ § 21.241
15 104‘ i 2'55" s
g ' > “a
U 2.50. 120‘
1.00. ]
0.96‘ 2'45‘ 1.16
45-4.?0 216 341-2 0 2 46 4'5 332*0'2'4'6
l (MA) l(mA) I (mA)

Figure 7 .2: MR (top) and CIMS (bottom) plots at 4.2K for the samples in Fig. 7.1.
Layer thicknesses are in nm.

 

 

 

 

 

 

 

 

N-# of samples R(P) (Q) AR(Q) MR (%) AI, (mA) H, (kOe)
295 K
Cu-5 1.38 :1: 0.02 0.070 :1: 0.002 5.1 :l: 0.3 5.7 :l: 0.3 0.18 :l: 0.02
Ag-3 2.75 :l: 0.14 0.096 :1: 0.004 3.5 :l: 0.2 3.1 :t 0.4 0.13 :l: 0.03
Au-13 1.78 :l: 0.05 0.061 :1: 0.003 3.4 :l: 0.1 5.6 :l: 0.4 0.14 :l: 0.02
4.2 K
Cu—2 0.94 :l: 0.02 0.144 :1: 0.006 15.3 :1: 0.3 7.6 :l: 0.2 0.32 :l: 0.05
Ag-3 2.26 :l: 0.12 0.192 :1: 0.005 8.5 i 0.4 6.9 :l: 0.2 0.20 :l: 0.05
Au—3 1.31 i 0.07 0.123 :1: 0.007 9.4 :l: 0.06 6.8 :l: 1 0.18 :l: 0.02

 

 

 

Table 7.2: Average values of R, AR, MR(%), AI, and H, at 295K and at 4.2K of
Py/N/Py with N=Cu, Ag and Au.

Fig. 7 .1 compares representative MR and CIMS switching data at 295K for sam-
ples of Py(24)/N(10)/Py(6) nanopillars with N=Cu, Ag, and Au. Fig. 7.2 compares
the same quantities for the same samples at 4.2K. Table 9.2 compares average val-
ues of the total resistance, R(P), when the magnetic moments of the two layers
are oriented parallel (P) to each other, the change in resistance upon switching,

AR = R(AP) — R(P) between the antiparallel (AP) and P magnetic states, the

110

 

magnetoresistance, MR(%): (AR/ R) x 100%, the difference AI, = I: — I 8‘ between
positive (+) and negative (-) switching currents, 1,, and the upper magnetic switching
ﬁeld, H,, over 3 — 13 samples of each type at 295K (t0p three data sets) and over 2 — 3
samples of each type at 4.2K (bottom three data sets). As a measure of the scatter
in our values, we list also the standard deviation of the mean for each quantity.

The average resistances of the three sets of samples differ, due to an unknown mix
of variation in areas (smaller area increases R, does not change MR, and decreases
A1,) and contact resistances (larger contact resistance increases R, decreases MR,
and may leave I, unchanged). Thus, a precise comparison between them cannot be
made. However, at both temperatures, the MRs and AI, are roughly similar for all
three metals. The average total resistances of the samples in P states given in Table
9.2 include the sheet resistances of the bottom Cu lead and top Au lead, which are
the same for all samples. The resistances of the samples with N=Ag(10) interlayer
are almost twice as much as the resistances of the other samples with Au and Cu
interlayers. Since we fabricate all the nanopillars in the same way, a factor of 2
in the total resistance of a nanopillar suggests that the main reason for these high
resistances is due to bad contacts, because it is very unlikely to have sample areas
twice as small. The MRs for the samples with N=Ag(10) give larger AR, which is in
good agreement with the CPP-MR experiments [138] but smaller MR(%), which is
not expected. This also makes us believe that the high resistances of our samples with
N =Ag(10) are due to large contact resistances. Although the samples with Au give
smaller AR and MR, they show switching currents similar to the ones with Ag and

Cu. The switching currents of nanopillars with N=Au(10) are interestingly similar

111

 

 

to the ones with the ones with N=Cu(10) even though the MRs of nanopillars with
NzAu(10) are smaller. We conclude that Ag and Au represent potentially viable

alternatives to Cu for studies of CIMS physics and for CIMS-based devices.

112

Chapter 8

Reduction of spin polarization and
spin transfer by synthetic
antiferromagnets

8. 1 Introduction

Much CIMS research has been done to try to reduce the switching currents of nanopil-
lars so that they could be used in next generation MRAMs. However, there are
possible applications of nanopillars which require higher switching currents like read
heads of hard drives. To use nanopillars in read heads the switching currents must
be so large so that the magnetization of a nanomagnet ﬂips only due to the external
magnetic ﬁelds of the magnetic bits on a hard disk. In this Chapter we present a
method to increase the switching current by using a synthetic antiferromagnet (SAF)
i.e. antiferromagnetically coupled layers as a polarizing layer. Since a similar study
was also done by Emley et a1. around the same time and already published (see Ref.

[142]), we did not publish our results.

113

 

 

8.2 Experiment and Results

When two ferromagnetic layers are separated by a thin enough nonmagnetic layer, the
RKKY exchange interaction between the two F-layers couples them ferromagnetically
or antiferromagnetically depending on the thickness of the separating N-layer. It is
known that 0.6 nm thick Ru between two Co layers gives antiferromagnetic coupling
[13]. To study the effect of a synthetic antiferromagnet (SAF) as a polarizing (ﬁxed)
layer, we fabricated two magnetically uncoupled nanopillars with the structures: (a)
Cu(80)/Co(20)/Ru(0.6)/Co(4)/Cu(10)/Co(4)/Cu(5)/Au(110) (see Fig. 8.1), where
Co(20)/Ru(0.6)/Co(4) is an SAP; (b) Cu(80)/Co(20)/Cu(10)/Co(4)/Cu(5)/ Au(110)
and compare the MR5 and the switching currents of these devices at room tempera-
ture. The approximate sputter deposition rates of the constituents of our samples are
given in Table 8.1. Further details of sample fabrication and measurement are given

in Chapter 3.

 

Target Au Co Cu Ru
Deposition rates (A/s) 6.7 4 10 0.5

 

 

 

 

 

 

 

 

Table 8.1: The approximate sputter deposition rates for Au, Co, Cu and Ru.

Fig. 8.2 shows the MR curve for a sample shown in Fig. 8.1. The highest resistance
state is the AP alignment of the two Co(4) layers. The P alignment of the two Co(4)
layers gives a pseudo P state which is in between the AP state and the total P state
of all Co layers, which is not shown in Fig. 8.2 because the strong antiferromagnetic
coupling between Co(20) and Co(4) layers require ﬁelds higher than 3 kOe to align.
The average MR (P and AP alignments between Co(4) layers) for 5 samples with

an SAP polarizing layer is 2.23 % and the average MR for 4 samples with a Co(20)

114

 

 

Au(110)

  

1' 601(5):.3

 

Figure 8.1: A cross-sectional cartoon of the structure of a magnetically uncoupled
nanopillar with a synthetic antiferromagnet (SAF). All thicknesses are in nm.

 

 

 

 

 

 

 

0.96-
. )Com
095 }SAF
' _ \ —. A /
.. ‘ \
:9} 0.94- 1 [ l 1
2 (- AR=0.015
> ‘ 1 \
u I \
0.93- , n x
0.92-
-3 Y -2 ' -'1 ' 0 ' i ' 2 ' 3
H(kOe)

Figure 8.2: MR curve for a Co(20)/Ru(0.6)/Co(4)/Cu(10)/Co(4) nanopillar at zero
current at room temperature. The highest resistance antiparallel state represents
antiparallel alignments of two Co(4) layers. The parallel alignments of Co(4) layers
represent a pseudo parallel state, which gives a higher resistance than the complete
parallel alignments of all Co layers.

polarizing layer is 5.4%, which shows that an SAF reduces the spin polarization and
therefore MR.

Fig. 8.3 shows CIMS from AP to P state but does not show the P to AP state

115

 

0.96 4

    

 

5‘: 0.95-
1;” 4R=0.015
‘3 0.94. <
] Al>31.5 mA
0.93-

 

 

q
q
q
-

l V l V r V

-20'-1'5'-10 -5 0 5'10‘15'20
|(mA)

Figure 8.3: CIMS curve for the nanopillar in Fig. 8.2 at zero ﬁeld at room temper-
ature. The positive switching current is not seen because it is higher than the limit

(18 mA) of our current supply.

because of the limit (18 mA) of our dc-current supply. The difference in the switching
currents AI = 1+ — I ‘ for the Co(4) nanomagnet with SAF polarizing layer is
greater than 31.5 mA. None of our 5 samples with an SAF polarizer showed a positive
current switching (from P to AP). However, the average 61 for 5 samples is > 31 mA.
Whereas, the average AI for 4 Co(20)/Cu(10)/Co(4) nanopillars is 8.5 mA. This
dramatic increase in the switching currents of a Co(4) nanomagnet is attributed to
the reduction in the spin polarization and thus reduction in the spin transfer to the
Co(4) nanomagnet due to an SAF polarizer. Since only the top Co(4) nanomagnet
switches with current, the change in resistance, AR for CIMS ( see Fig. 8.3) is the

same as the AR for MR (see Fig. 8.2) between the P and AP alignments of the two

Co(4) layers.

116

 

 

8.3 Summary

We fabricated magnetically uncoupled nanopillars with a SAF polarizer and observed
a reduction in spin polarization that results in a decrease in MR as seen in Fig. 8.2.
The reduction in the spin polarization also causes a reduction in the spin-transfer to
the Co(4) nanopillar, which in turn switches at a higher current by a factor of more
than 3.6 compared to our magnetically uncoupled Co(20)/Cu(10)/Co(4) samples.
Our results agree with the results of Emley et al. [142], in which they studied the
same effect in nanopillars with Co(ll.5)/Ru(0.7)/Co(8)/Cu(6)/Co(2)/Cu(2)/Pt(30)
structure, where they patterned all Co layers into a nanopillar shape but adjusted
the thicknesses of Co layers in the SAF such that the dipolar coupling to the free
layer is minimal. Thus, we essentially study similar geometries. They also observed a
reduction in the ARs and also an increase in the switching current densities by a factor
of ~ 2.3 compared to their magnetically uncoupled Co(40)/Cu(10)/Co(3) samples.
Although it is not clear whether the comparisons between the MRs of samples with
and without SAFs are legitimate due to their unequivalent structures, the comparison
between their switching currents is legitimate because only the free Co nanomagnet
switches with current.

We conclude that the spin polarization and spin transfer could be reduced by
using an SAF as a polarizing layer in a nanopillar. Such reductions in the spin
transfer requires higher current densities for CIMS, which may be useful for some

technological applications such as CPP-MR read heads.

117

Chapter 9

Effect of asymmetric leads on the
switching currents of nanopillars

9.1 Introduction

Manschot et al. [58] recently predicted that J: , the positive switching current den-
sity needed to ﬂip the switching layer from parallel (P) to antiparallel (AP) to
the ﬁxed layer, could be reduced by up to a factor of ﬁve in a Cofixed/Cu/Cofree
nanopillar through the use of asymmetric leads having different effective resistances,
r“ = p x l,, = resistivity times spin-diffusion length. Assuming a simple one-
dimensional model with no magnetic coupling in which the bulk scattering is ne-
glected, they argued that combining a large T‘ for the lead B outside of the switching
layer, with a small 1"“ for the lead A outside of the ﬁxed (polarizing) layer, could
reduce J: by up to a factor of ﬁve. They also predicted that reversing such a B and
A pair would invert the CIMS, i.e. that positive current would switch the sample
from AP to P. Manschot et al. assumed that the F- and N-layers within the leads
of primary interest were much thinner than their spin-diﬂ'usion lengths, and that all
of the interfaces were the same and also produced no spin-ﬂipping. While these as-

sumptions are not necessarily valid for our samples, we adopt them for the moment

118

 

 

 

to obtain simple Manschot predictions.
9.2 Experiment

Testing the Manschot prediction requires F-layers with [ff longer than their layer
thicknesses, tp, since otherwise the effects of the leads on both the magnetoresistance
(MR) and CIMS will be minimal. Then one needs non-magnetic leads N, on the two
sides of the nanopillar with very different values of r*, and with lengths thicker than
their respective values of l,;, so that 7“ can have its full effect. As F -metal, Manschot
et al. proposed using Co, for which If!” (4.2K) is probably ~ 60 nm [143] and lff"(RT)
(RT = room temperature) is probably ~ 38 nm [143]. We adopt their proposal of Co
and chose a thickness of 20 nm to be less than the likely room temperature 13". For
the N-lead pair they proposed Pt for the smaller 1“ and Cu for the larger 7“. For our
sputtered metals at 4.2K we estimate 7‘}, = ppt(42 nﬂm) xlff‘(14 nm)as 0.6 ern2
[99], and r2.” = p0,,(5 nﬂm)xlf}“(500 nm)=2.5 ern2 [144]. Cu as top lead outside of
B, however, cannot have full effect in our nanopillars, since its thickness would have
to be << If," before reaching the Au capping layer that we need to obtain clean top
contacts to our nanopillar. For large 7" we have chosen instead the alloy AgSn(5%),
which has both a much shorter lffsn ~ 25 nm and a larger r3193" ~ 5 fﬂm2 [144]. In
this paper we compare switching currents for samples with A = Pt and B = AgSn
and, for comparison, with just Cu outside the ﬁxed layer and Au outside the switching
one (i.e., with A = B = 0). We estimate 72,, = pA.,(20 n9m)xl§‘f"(60 nm)=1.2 f9m2.
Combining these various metals or alloy gives the values of rr“ = rig/r; listed in

the second column of Table 9.2. With increasing temperature, p will increase and l,;

119

 

will decrease, but we assume that their product 1““ should be nearly independent of
temperature.

Our sample preparation and fabrication techniques are described in detail in Chap-
ter 3. As illustrated in Fig. 9.1, to minimize dipolar coupling between the two Co
layers of our nanopillars, we stop our ion-milling within the middle Cu spacer layer, so
that the in-plane length scales of both our ﬁxed Co layer and part of the Cu layer are
several microns. Our samples thus differ from the simple uniform linear assumption
of Manschot et al. [58].

The multilayered nanopillars to be compared had the forms:
1. Cu(80)/Co(20)/Cu(10)/Co(4)/Cu(5)/Au(110) (Fig. 9.1a)

2. Cu(80)/Pt(20)/Co(20)/Cu(10)/Co(4)/AgSn(30)/Au(110) (Fig. 9.1b)
3. Cu(80)/AgSn(30)/Co(20)/Cu(10)/Co(4)/Pt(20)/Au(110) (Fig. 9.1c)

with all thicknesses in nm. The approximate sputter deposition rates for the con-
stituents of our samples are given in Table 9.1. For ease of description, we label the
samples by A/X/B, where X = Co(20)/Cu(10)/Co(4), A is the insert of interest to
the left of X, and B is the insert to the right of X. The Cu(80) layer is the bottom
contact and the Au(110) is composed of ~ 10 nm of the last layer of the nanopillar
plus an additional 100 nm thickness of the top contact. The thicker and much wider
bottom Co layer switches more easily than the thinner and narrower bottom Co layer
upon application of a magnetic ﬁeld, but won’t switch upon application of a current.
The AgSn(5%) alloy was made by inserting Sn plugs into a Ag sputtering target. We

estimated the atomic percent Sn from measurements on separate 200 nm thick test

120

 

ﬁlms using Energy-Dispersive-Spectroscopy (EDS), which gave ~ 5 at.%, and mea-
surements of the residual resistivity, which gave ~ 150 an corresponding to S, 4 at.

%. (103].

 

 

] Target ] Cu [ Pt]Au ] Co ] AgSn(5%) [
] Deposition rate (A/s) ] 10 [ 4 ] 3T4.5] 7 ]

 

Table 9.1: Approximate sputter deposition rates of Cu, Pt, Au, Co and AgSn(5%).

Au(110) AU(110)

 

Figure 9.1: Cross-sectional cartoons of (a) X=Co(20)/Cu(10)/Co(4), (b)
A=Pt(20)/X/B=AgSn(30), and (c) A=AgSn(30)/X/B=Pt(20) magnetically uncou-
pled nanopillars. All thicknesses are in nm.

A problem in determining reliable values of I j is the irreproducibility of coupling,
magnetoresistance, and switching currents, due to uncontrolled variations in nanopil-
lar shape and to incompletely controlled interface resistance with the top contact. In
our analysis, we include data only for samples that showed both nearly single step
magnetoresistance (MR) occurring after the magnetic ﬁeld H passed zero (consistent
with weak or no magnetic coupling, which is also stipulated by Manschot et al.), and
CIMS with values of resistances in the P and AP states that agreed well with those
for the MR. Fig. 9.2 shows three examples of MR and CIMS for samples with A =

Pt(20) and B =AgSn(30). Fig. 9.3 shows three examples for the inverse situation,

121

 

A = AgSn(30) and B = Pt(20). Note that the values of I, vary substantially from
sample to sample. In both cases we use the standard convention that positive current
1+ ﬂows from the thick to the thin Co layer. However for the case of A = AgSn(30)
and B = Pt(20) we did not observe inverse CIMS as predicted by Manschot et al.
[58]. Table 9.2 lists average values (with standard deviations) of: the resistance in
the AP-state, R(AP), magnetoresistance, MR(%), difference in switching currents,
AI = 1+ — I ', switching current for negative, I ‘, and switching current for positive
current, 1+, for samples with Pt and AgSn on opposite sides of the samples and, for
comparison, samples from diﬂ'erent sources with A = B = 0, i.e. with just Cu and
Au contacts. For each sample set we also list the number of samples averaged.

We focus upon the quantity considered by Manschot et al., namely I + (since all
of our samples have the same nominal area, 1+ is prOportional to J +). Manschot
predicted that 1+ would be much smaller for A = Pt, B = AgSn than for a nanopillar
with symmetric contacts, and that 1"” would be ‘inverted’ for A = AgSn, B = Pt. In
contrast, our average values of I + for both orientations of Pt and AgSn are similar
to each other ( to within mutual uncertainties), and are also similar to those for our
‘comparison’ samples with A = B = 0. While the comparison samples are not strictly
‘symmetric’, their values of rr‘ are intermediate between those for A = Pt, B = AgSn
and A = AgSn, B = Pt. As checks upon our values of AI and I + for A = B = 0, we
list values of both obtained independently by a different one of us, and also our best
estimate of AI from data from the Cornell group in ref. [141]. For this comparison,
we took the Cornell value of AJ = 14 x 107 A/cm2 from Fig. 2 in [141] and multiplied

by the nominal area of A = 70 x 130 nm2 to get rounded AI = 12.7 mA. Both of

122

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

2322 1604-3b#3
M
2.31. [’—1
230.
El 229-
a l
S 2281
'U .
2274
2.26] ARIO.064
. MR=2.84%
225_ Av vim .w
-1.0 03 0:0 015 110
H (kOe)
2.0201 1579-2b#5
2.0154 [— F]
g 2.0101
6 .
3 2.005-
'0 1
2.000-
. AR=0.024
1.995- JAR-1.2%
1.990 . . - . - .
-1.0 0.5 0.0 0.5 1.0
H(kOe)
1.45 1604-4642 K N
1.44-
g .
a 1.43-
3 .
'0
1.42-
ARI-0.04
MR-2.84%
1.41- _ - _
-1.0 I 05 0:0 I 0:5 ' 1:0
H(kOe)

dV/dl (a)

dV/dl (a)

dV/dl (a)

2.4012 1

2.36 4
1
1

2.32 1
1
2.28]
1

2.24

4

2.06 ~
2.04 -

2.02 4

 

2.00-
1

604-3b#3

AI=10.2mA
5 '1'0 1‘5

-15’-1'0'-5' 0 '

MMA)

1 579-2b#5

  

AI-21 mA

 

1.981

1.48.
1.47;
1.464
1.45]
1.44.
1.43]
1.42 9
1.41 -

 

-16-645 0 ' 5 [1071’s

1 6044b#2

Al-19.5 mA

 

35V 0 ' 5107115
I(mA)

-15 -10

Figure 9.2: Examples of MR and CIMS curves for A=Pt(20) / X / B=AgSn(30) nanOpil-
lars at room temperature.

these other results are consistent with ours to within mutual uncertainties.

123

2075-J 16042“, 2.201 1604-Za#4

 

 

 

 

 

 

 

 

  

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

._ 2.18.
2.070] *
1 2.16] \
A 2065-; A 2.141.
S 1 $3. 1,
g 2.060] g 2'12"]
U .' '0 2.10]
2.0554 MR 208]
] .1.05°/. 1
2.050 2'06] Al=31.7 mA
1 ' Y ' I ' I ' l 2. i: r ' Y ' T ' f ‘ I l ‘ l ' T E
-1.0 -0.5 0.0 0.5 1.0 -20 -15 -10 -5 0 5 1o 15 20
H(kOe) l(mA)
33515844,,“ 3-8 1584-1b#1
( 3.7
3.34-]
,1 3.6]
§ 332-j g’: l?
5 ' 1 5 3-51
% ‘= 1% "
3.30] 3.4]
[ L MR: 2 % 3'3 J‘ 211.21 mA
3.28] ]
' l ' 1 1 1 = 1 3.2 ' 1 1 1 1 = 1 1 ' 1 - 1 1 1 '
-1.o -o.5 0.0 0.5 1.0 -15 -1o -5 0 5 10 15
H(kOe) |(mA)
1.89. .
' 1584-1b#3 1584-16113
[ ﬂ 2.001
1.88- [[ ‘
1.964
S S .
§ 1.87- g 1-92~
"U 'U 1
]_ MR-1.4% 1.884
1-86— r‘ - ‘ ‘ ._ . Al-18mA
' 1*1' ﬁ'1 1.84171'.11 YﬁY'I'T'I'I
-0.6 03 0.0 0.3 0.6 -20 -15 -1o -5 o 5 1o 15 20
H(kOe) |(mA)

Figure 9.3: Examples of MR and CIMS curves for A=AgSn(30)/X/B=Pt(20) nanOpil-
lars at room temperature.

9.3 Summary

We have tested the Manschot prediction [58] that 1*, the positive current needed

to switch Co/Cu/Co nanopillar from thrla2f4’ state to the AP state could be greatly

 

 

 

 

 

 

 

A/X/B (# of samples) rr’ R(AP) (0) MR (%) A] (mA) 1" (mA) 1+ (mA)
Pt(20)/X/AgSn(30) (13) 8.3 2.2 2.6 15 i 5 —7.5 :t 2.5 7.5 :t 2.5
AgSn(30)/X/Pt(20) (4) 0.12 3.3 1.5 20 d: 7 —9.5 i 3.5 10.3 :h 5
x (4) 0.5 1.4 5.4 8.5 i: 1.6 —2.85 :l: 0.6 5.7 i 1.2
X (AlHajDarwish) (8) 0.5 1.0 5.6 13.5 :l: 2 —5 :t 1 8.5 :l: 2
x (Cornell [141]) (4-10) 0.24 12.7

 

 

 

 

 

 

 

 

Table 9.2: Average values of rr" = rg/rjq, R(AP), MR(%), AI = I+ — I“, I" and 1+
for our samples of the form Cu(80)/A/X/B/Au(110) with X = Co(20 )/Cu(10)/Co(4),
and A, B = Pt(20), AgSn(30), or nothing. The AlHajDarwish samples are similar to
our ones without Pt and AgSn, except that the thickness of the ﬁxed Co layer is 30
nm.

reduced by placing a lead with large 1"“ outside the switching F -layer and one with
much smaller 1"“ outside the ﬁxed layer. We successfully fabricated magnetically
uncoupled nanopillars with equivalent asymmetric leads as proposed by Manschot et
al. [58] yet we did not observe any reduction in the positive switching current. Upon
switching the sides of the asymmetric leads we did not observe the predicted inverse
CIMS behavior either. We observed little difference between 1+ or AI for nanopillars
with A = Pt and B = AgSn and nanopillars with A = AgSn and B = Pt. Either

the Manschot assumption of a one-dimensional geometry, or some other part of the

analysis (e.g. neglecting bulk scattering), is inadequate for describing our nanopillars.

125

Chapter 10

Effect of lower demagnetizing ﬁelds
on the switching currents of
nanopillars

10. 1 Introduction

The high critical switching current densities of nanopillars has been an obstacle for
some technological uses. Katine et al. [46] derived the critical switching currents

from the original theory of Slonczewski [33] as:

ayeS

 

I: = WlHCIIHD +2’ITM] (10.1)
_ _ 0:785 71 _ 7r
1c _ g(7r) [Heff( ) 2 All (102)

where a is the phenomenological Gilbert damping constant, 7 is the gyromagnetic
ratio, M is the magnetization of the nanomagnet, S is the total spin of the nano-
magnet, and He” 2 Han cos(0) + Hm — He“ is the sum of anisotropy, external
and exchange (but not demagnetizing) ﬁelds. 9(0) is a coefﬁcient that depends on
the spin polarization of the thick (ﬁxed) ferromagnet and the relative orientations of
magnetizations.

For Hm = 0 the switching currents become dependent only on the anisotropy,

126

exchange and demagnetizing ﬁelds of which the contribution from the demagnetizing
ﬁeld is much larger compared to the exchange and anisotropy ﬁelds for a C0 thin ﬁlm
that has in—plane anisotropy. The demagnetizing ﬁeld for a ferromagnetic thin ﬁlm is
the ﬁeld required to pull the in-plane magnetic moment of the sample perpendicular
to the plane and is proportional to the magnetization M of the sample. Due to the
large shape anisotropies of thin ﬁlm ferromagnets, they require very high external
ﬁelds (couple of Teslas) to demagnetize.

In this chapter we study the effect of lower demagnetizing ﬁelds on the switch-
ing currents of nanopillars by comparing the MRs and the switching currents of
Au / Co(t) / Au nanomagnets (where for Au/Co(1 nm) / Au nanomagnet the easy axis is
normal to the sample plane) with the MR5 and the switching currents of Cu/ Co(t) / Au

nanomagnets given in Ref. [141].

10.1.1 Perpendicular anisotrOpy in Au/ 00/ Au systems

For a thin ﬁlm ferromagnet the magnetization usually lies in the plane of the sample
due to large shape anisotropy [145]. Previous studies of epitaxially grown very thin
(< 1 11111) CO layer sandwiched between two Au layers gives perpendicular anisotropy
[146] i.e. the magnetization easy axis is perpendicular to the 00/ Au interface. The
possible reasons for perpendicular anisotropy of such ferromagnetic thin ﬁlms are the
large anisotropy of the orbital moment [147] and the strain at a Co/Au interface
resulting from a large (14%) lattice mismatch [148]. The perpendicular anisotropy
decreases as the thickness of Co increases and disappears at about too = 20 A for a

Au/Co/ Au system and at tCo = 10 A for a Cu/Co/Cu system [149]. Although most

127

studies on perpendicular anisotropy are done with epitaxially grown samples, similar
behaviors were also observed in sputtered Au / Co/ Au systems [147] for up to 1.2 nm
thick Co [150].

As the thickness of a C0 layer sandwiched between two Au layers increases above
1 nm the strength of the perpendicular anisotropy decreases with some magnetic
domains still preferring to stay in the perpendicular direction and others preferring to
stay in-plane in a multi-domain sample. The interfacial interaction between Co/Au
causing the perpendicular anisotropy is always present but as the thickness of Co
increases the net anisotrOpy becomes in-plane. Having some magnetic domains lying
perpendicular to the plane of a C0 thin ﬁlm should give demagnetizing ﬁelds smaller

than those for a C0 layer with the same thickness but with in-plane anisotropy.
10.2 Experiment

To determine the easy axes of of Co(t) layers sandwiched between Au layers, we sput-
tered Cu(40)/Au(10)[Co(t)/Au(10)]x5 samples and measured their magnetizations
using a SQUID magnetometer by varying the angular position of the sample with re-
spect to the external magnetic ﬁeld. All thicknesses are given in nm. The Cu(40) is a
buffer layer sputtered on a (100) Si wafer, Au(10) is selected to minimize exchange cou-
pling between the Co layers and the structure is repeated 5 times to improve signals for
a SQUID magnetometer. Fig. 10.1 shows the magnetization vs. angle at room tem-
perature for Cu(40)/Au(10)[Co(1)/Au(10)]x5 and Cu(40)/Au(10)[Co(2)/Au(10)]x5
in an external magnetic ﬁeld of 1000 Oe. The easy axis for Co(l) is in the perpendic-

ular direction, whereas the easy axis for Co(2) is in the plane of the sample. Fig. 10.2

128

shows the magnetization vs. angular position of a Cu(40)/Au(10)[Co(1)/Au(10)]x5
sample at 500, 1000 and 5000 De at room temperature, which also shows that the
magnetization of Co(l) sandwiched between Au layers has a perpendicular easy axis.
The external ﬁeld had to be increased up to 5000 De to eliminate any angular depen-

dence.

Hm=1000G @ 295K

    
  
  

«l.
2.0x10 5,30
1
1.5x10“«

1 .0x104 1

N

M (emu)

5 Au(10)[Co(1)/Au(10)]x5
5.0x10' 1

 

0'0 ' I ' I I I ' I ' I ' I ﬂ I
0 50 100 150 200 250 300 350
angular position a

 

Figure 10.1: The magnetizations of Au[Co(1)/Au(10)]x5 (ﬁlled symbols) and
Au[Co(1)/Au(10)]x5 (open symbols) samples as a function of angular position 0
in a magnetic ﬁeld of H = 1000 De at room temperature. 0 is the angle between H

and the sample plane.

The easy axes of Au / Co(t) / Au samples were also tested by making mm2 Nb cross-
stripped samples in the following forms and measuring AAR at 4.2K where Nb leads

superconduct:
o Nb(150)/Cu(10)/Py(20)/Cu(10)/Co(t)/Cu(20)/Nb(150)

. Nb(150)/Cu(10)/Py(20)/Au(10)/Co(t)/Au(10)/Cu(10)/Nb(150)

129

Au10/(Co1/Au10)x5 @ 295K

2.0x104~bmWWu%ﬁﬁmms Hm=sooos
J.#¢.. f

15x10‘~."° 80° \ 9" WQs

- v f «20%;! f 0%) 0., Hm=1000G

1 933) o
<><> 0 00

 

g 1.0x10‘-§§0 28 (A %
5" , 059 35 Hm=5006
5.0x10'51
0.0 -

 

0 5h 7100'150'200'250'300'350'400
angular position 0

Figure 10.2: The magnetization of a Au[Co(1)/Au(10)]x5 ﬁlm as a function of an-
gular position 19 in external magnetic ﬁelds of H = 500,1000 and 5000 Oe at room
temperature. 0 is the angle between H and the sample plane.

The sample fabrication and measurement techniques for these macroscopic samples
are given in detail in Chapter 3. Fig. 10.3 shows that AAR of Au/Co(t)/Au samples
decreases as too decreases, whereas AAR of Cu / Co(t)/ Cu samples hardly varies from
t = 4 nm to t = 1 nm. The fact that the easy axis for Au/Co(1)/Au layers is normal
to the sample plane and yet we still have non-zero MRs much smaller than those
for Cu/Co(1)/Cu samples indicates that Co(l) layer sandwiched between two Au(10)
layers has partially perpendicular anisotropy.

All measurements described above were done for multi-domain polycrystalline
samples. For a mono-domain nanomagnet the physics may or may not be the same
as there is no means of determining the easy axes of a mono-domain thin Co sand-
wiched between Au or Cu layers. The only way to see whether there is any partial
perpendicular anisotropy is to compare the MRs of nanOpillars where a thin Co is

sandwiched between Au or Cu layers.

130

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N
A I n

10- O o
0
$5 08- 8
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. o

02.: open: Nb/Cu/Py/Cu/Co(t)lCule
(ﬁlled: NblCu/Py/Au/Co(t)lAu/Nb

0.0v1vur1v
012 3 4

tCO (nm)

 

 

 

-

(”J

Figure 10.3: A comparison between AAR/AAR(C0(4)) of mm2 Nb cross-stripped
CPP samples with Co(t) sandwiched between Au and Cu layers at 4.2K. The inset
shows a picture of a CPP-sample, where the CPP current flows through the overlap
area of Nb strips at 4.2K.

We fabricated Cu(80)/Co(20)/Au(10)/Co(t)/Au(115) nanOpillars with t =1, 2, 3
and 4 nm and measured the MR and switching currents. The details of the fabrication
and measurement of our nanopillars are given in detail in Chapter 3. The sputter
deposition rates were ~ 2 A/s for Co, ~ 10 A/s for Cu, and ~ 5 A/s for Au. Fig.
10.4 shows that the MR of these nanopillars decreases with decreasing t, which is also
consistent with the magnetization and CPP-MR measurements given above. The
quantitative studies in Ref. [141] show that the MR of their Co/Cu(10)/Co(t)/Au
samples for t = 1.2, 3.6, and 6 nm hardly vary as t decreases, indicating that the
(30/ Cu and Co/Au interfaces are the dominant factors determining the MR in these
thickness ranges, also consistent with our data for macrosc0pic CPP samples (see Fig.
10.3). These similarities in MRs of Co/Au/Co(t)/Au nanopillars and macrosc0pic

samples makes us believe that the easy axes directions in both cases are the same.

131

To compare the critical current densities of our nanopillars, we need to divide the
currents by the areas of the samples. As it was mentioned in Chapter 3 getting the
exact shape and size of these nanopillars is almost impossible using our fabrication
technique. However, the value of the resistances of these nanopillars show little de-
viation from one another, suggesting that the areas are not grossly different. Our
nanopillars have nominal sizes of ~ 130 x 70 nm and we approximate the effective
areas of our samples as ~ 0.01pm?

Fig. 10.5 shows a comparison of the difference in switching currents A] = 1+ — I ‘
vs. ta“ for our Co/Au/Co(t)/Au nanopillars (ﬁlled circles) and Co/Cu/Co(t)/Au
nanopillars (open circles) of Ref. [141]. Interestingly the critical currents of our
Co/Au/Co(t)/Au nanopillars are very similar to those for Co/ Cu / Co(t) / Au nanopil-
lars given in Ref. [141]. Since the samples in Ref. [141] have a Au capping layer
on top of the free Co(t) layer, the demagnetizing ﬁelds for these samples may also
be smaller. However, the MRs of those samples show little or no variation with the
thickness of Co, making us believe that those samples have in-plane anisotropy. For
an elliptical Co(1) nanomagnet of the size ~ 130 x 70 nm lateral dimensions with
in-plane anisotropy, the approximate demagnetizing ﬁeld is nearly the same as the
demagnetizing ﬁeld of a macroscopic Co(1) ﬁlm, which is in the order of a couple of
Tesla. However, by sandwiching a Co(1) layer between two Au(10) layers, we reduce
the demagnetizing ﬁeld (~ 0.5 Tesla see Fig. 10.2) of the nanomagnet which in turn
should switch at a current smaller than the critical currents of Co/Cu/Co(1)/Au
nanopillars [141]. The similarities in the switching currents of Co/ Au/ Co(t) / Au and

Co/ Cu/ Co(t) / Au nanopillars suggests that the demagnetizing ﬁeld makes little con-

132

 

tribution to the STT, inconsistent with equations 10.1 and 10.2.

 

 

5-
1

4. i

a- 1
§
E2-

1
1.1
0 . . . , . , , 1
0 1 2 3 4 5

tCo (nm)

Figure 10.4: MR vs. t3“ of Co(20)/Au(10)/Co(t)/Au(115) nanopillars at room
temperature. The MR increases as t increases up to 3-4 nm thick Co and then
saturates.

 

 

25-
1
l l
"E 154 l
o
2 1 1
”52 1o~ l
5 l f
5-
gi
O f I I I
0 2 4 6
too(nm)
Figure 10.5: A comparison of AJ = LIE—”4P - J1”) "P vs. ta,“ for

Co(20)/Au(10)/Co(t)/Au(115) (ﬁlled circles) and Co/Cu(10)/Co(t)/Au (open cir-
cles) [141] nanopillars.

133

10.3 Summary

We observed that our sputtered large Au/Co(1)/Au samples have their easy axes
normal to the plane, and therefore they should have lower demagnetizing ﬁelds than
a C0( 1) layer with in-plane anisotropy. The non-zero MR measurements of both mm2
samples and nanopillars indicate that the net anisotropy of an Au/Co(t)/Au system
is not perfectly perpendicular at least for t =1 and 2 nm. The critical switching
currents of Co/ Au/ Co(t) / Au nanOpillars with lower demagnetizing ﬁelds measured
at room temperature surprisingly show almost no difference with the critical switching
currents of Co/ Cu/ Co(t) / Au nanopillars of Ref. [141], inconsistent with equations
10.1 and 10.2. A linear decrease of the switching currents with decreasing thickness

of the Co/Au/Co(t)/Au nanomagnets was also observed within the thickness ranges.

134

Chapter 11

Summary & Conclusions

In this thesis we have completed two studies involving current-perpendicular-to—plane
magnetoresistance (CPP-MR)at 4.2K, and ﬁve studies of current-induced magnetiza-

tion switching (CIMS) at room temperature and 4.2K.

CPP-MR

We have measured the spin diffusion lengths of sputtered N=Pd, Pt and Au.
The interest in these metals arises from the fact that when combined with very thin
layers of Co they give perpendicular anisotropy, which has potential uses in magnetic
storage. To determine the effects of spin-orbit coupling on the spin memory loss
at an interface formed by two nonmagnetic metals, we measured the spin memory
losses at N / Cu interfaces and compared the results with previously published values
of Nb/ Cu and W/ Cu. We found that the spin diffusion lengths of Pd, Pt, and Au
are long enough so that these metals could be used in devices without excessive loss
of spin polarization. Furthermore, the spin memory losses at Pt/ Cu and W/ Cu are
similar, and spin memory losses at Pd / Cu and Nb/ Cu are also similar, both of which
are consistent with Spin-orbit coupling being the major source of such losses at these

N/ N interfaces. However, for Au/ Cu, the spin memory loss is smaller than those

135

for Pt/ Cu and W/ Cu. The sources of this weak spin-memory loss might be the low
resistivity of Au, and the absence of d-electron states at the Au Fermi surface. We also
measured the interface speciﬁc resistances of N / Cu and compared them with the no-
free—parameter calculations for perfect and 50-50 alloy interfaces. For Ag/Au, Co / Cu,
Fe/ Cr, Pd/ Cu, and Au/ Cu, agreement (or disagreement) between the measured and
calculated interface speciﬁc resistances correlates with small (large) lattice mismatch
between the metals forming the interface.

Previous studies of CPP-MR with superconducting leads showed that the scatter-
ing anisotropy at a simple F/ S interface is zero (i.e. no spin dependent scattering).
Similar behavior was found with X 2 Ag or Cu inserts between various F -metals and
S 2 Nb. Eid developed a ‘null-detector’ method for studying scattering anistropy at
F/N/S ‘interfaces’, observing that inserting certain Xs could change (activate) the
anisotropy. We reproduced his results and extended them to other X. We conﬁrmed
that X = Cu, Ag, or Cu/Ag, as well as the antiferromagnet FeMn, all gave no ac-
tivation. In contrast, X = Au or Ru gave activation, even for thicknesses as small
as 2 nm. Combinations of non-activating metals (Cu or Ag) next to F, with strong
scattering metals next to S (An, Ru, or FeMn) gave activation of the Cu or Ag. From
these results, we concluded that activating an F/ S interface seems to require strong

spin-ﬂipping at the N/ S interface.

CIMS
Previous studies had shown that alloying F-metals with Cr produces negative bulk

scattering anisotropies. By combining such alloys with themselves and with other

136

ferromagnets having positive bulk scattering anisotropies, we were able to produce
all four combinations of normal and inverse CPP-MR and normal and inverse CIMS.
In our samples, the direction of CIMS was set only by the net scattering anisotr0py
of the thick (polarizing) layer. The net scattering anisotropy of the thin (switching)
layer was irrelevant. The importance of bulk properties of our F-layers in the observed
behaviors of CIMS, shows that a diffusive treatment of the transport is crucial for
explaining our results.

To test whether Ag and Au present viable alternatives to the commonly used Cu
spacers, we compared CIMS switching currents for 10 nm thick N=Ag, Au and Cu
spacers in a standard Py- based nanopillar structure. Since the Ag and Au spacers
gave comparable switching currents to Cu, we concluded that Ag and Au could be
used as alternative spacers in devices.

We used a synthetic antiferromagnet (SAF) as a polarizer layer in a nanopillar
structure to successfully reduce the Spin polarization and spin transfer into a nano-
magnet, thereby increasing the switching currents by a factor of more than 3.6 relative
to those for nanopillars with a Co(20) polarizer. Such increases in switching currents
by using an SAF as a polarizer could be useful in devices that require higher switching
currents, such as CPP-MR nanopillar read heads.

Using magnetically uncoupled (as stipulated by Manschot et al.) Co/Cu/Co
nanopillar structures, we tested predictions by Manschot et al. that using asym-
metric leads with very different effective resistances (resistivity times spin diffusion
length) could: (a) reduce the positive switching current by up to a factor of 5, and

(b) invert the switching current when these leads are reversed. Combining Pt (small

137

effective resistance) with AgSn (large effectiveresistance) did not signiﬁcantly change
the switching currents nor inverted the CIMS upon lead reversal.

A switching current formula written down by Katine et al. predicts lower switching
currents for nanomagnets with lower demagnetizing ﬁelds. To study the effect of
demagnetizing ﬁelds on the switching currents for CIMS, we fabricated magnetically
uncoupled Co/ Au / Co(t) / Au nanopillars (which should have perpendicular easy axes
and therefore lower demagnetizing ﬁelds at least for t = 1— 2 nm) and compared their
switching currents with those of Co/ Cu / Co(t) / Au nanopillars (which should have in-
plane easy axes) of Albert et al. Interestingly, the difference between positive and
negative switching currents for the two systems are almost the same, independent of t,
suggesting that the demagnetizing ﬁelds do not have a strong effect on the switching

currents of nanopillars.

138

Appendix A

Determining the resistivities of
metals and alloys

A. 1 Overview

To analyze our CPP-MR data we need to know the resistivities of the metals and
alloys we use in our samples. To measure the resistivities of metals we used the van
der Pauw [151] method, which can be used to measure the resistivity of any thin ﬁlm
metal or alloy of arbitrary shape without holes by measuring its sheet resistance.

The van der Pauw method gives us the resistivity in the CIP direction but actually
what we need is the resistivity in the CPP direction. If we had perfect crystal cubic
samples, the resistivities in both directions would be exactly the same but due to the
columnar growth of our sputtered samples the CIP and CPP resistivities need not be
the same.

Our technique developed for measuring the resistance R in the CPP direction in-
volves sandwiching ferromagnetic (F) metals between two superconducting Nb strips.
Such a resistance measurement is limited to temperatures below the superconducting
transition temperature of Nb. A plot of AR (area times resistance) vs. thickness of

F-metal is obtained by measuring the resistances (4.2K) and the areas of samples with

139

various thicknesses of the F-metal sandwiched between the Nb strips. The slope of
this graph gives the resistivity of the F-metal. For normal (N) metals this technique
can be used by putting F-layers next to both Nb leads to eliminate the proximity
effect [59].

Test comparisons of CIP and CPP resistivities at 4.2K gave similar values as

illustrated by the resistivities of Co and CoZr (2.5% Zr) in Table A.1 [57].

 

 

 

Metal CIP madman) CPP p4.2x(uﬂcm)
Co 5.14: 2 4.0 :l: 0.7
CoZr (2.5% Zr) 20.7 :1: 1.0 19 i 1

 

 

 

 

 

Table A.1: A comparison between the CIP and CPP resistivities of Co and CoZr at
4.2 K [57].

Since these van der Pauw CIP resistivities and the trilayer CPP resistivities of
similarly sputtered ﬁlms agree, we choose to measure the CIP resistivities using the

van der Pauw method, which is much easier.
A.2 Technique

Van der Pauw instructs us to measure the resistance of a thin metal ﬁlm as shown in

Fig. A.1 [151]. The resistivity of the metal in Fig. A.1, p, is then given by:

7ft R1+R2
-_- A.1

 

where t is the thickness of the ﬁlm and f is the correction factor to the resistivity,
which is a function of R1 /R2. If R1 = R2 then f = 1. However, if R1 95 R2
(asymmetric cases) one has to determine the value of f from f vs. R1/R2 graph

shown in Fig. A.2. Notice that in Fig. A.1 there is no current contact between

140

voltage leads and no voltage contact between current leads: as such connection would

invalidate Equation. A.1.

 

 

 

 

 

 

(1) (2)

Figure A.1: The van der Pauw geometry for a thin ﬁlm metal: (1) the resistance is
R1 = [Vab/chl , (2) the resistance is R2 = [Vbd/Ical .

 

 

 

 

 

 

 

 

 

Figure A.2: The graph of correction factor f vs. R1 / R2, which is plotted using a
numerical solution to Equation A.2.

(Rl/R2)-1 1112 _l LII—2-
(RI/821+1' f ‘2‘” f (A2)

 

cosh [

141

A.2.1 Samples and Measurement

Our samples are thin ﬁlms of metals or alloys sputtered on Si substrates, and have
close to square or rectangular shapes with 200-300nm thicknesses. The four-probe
measurement was done by putting indium on each corner of our thin ﬁlms and cold
pressing on the leads. It is important to make the contacts on the periphery of the
ﬁlms because the van der Pauw derivation for resistivity assumes the contacts on the
periphery of the ﬁlm [151]. The samples were mounted on a ”quick dipper” which we
used to dip the sample into a liquid helium dewar to make low temperature (4.2K)
measurements. We also used the same dipper to measure at room temperature.
Measurements were made using a dc current source and a nanovoltmeter. We used
very low currents (< 30mA) and immediate voltage reading to prevent excess heating.
To check the linearity in the resistances we used several different currents ranging
between 1-30 mA. These measurements were also repeated with reversed currents to

eliminate any thermoelectric effects and voltage offset of the nanovoltmeter.

A.3 Results

The metals studied in this thesis are Pd, Pt, Au and Nb. We also studied a silver and
tin alloy Ag0_95Sn0,05. Using the van der Pauw technique we measured the resistivities
of these conductors both at room temperature and at 4.2K except Nb, because Nb
superconducts at 4.2K. We measured the transition temperatures of Nb samples and
also measured the resistivities of Nb ﬁlms at 12K. Following tables show the individual

samples’ resistivities calculated from their respective resistances using Equation A.1

and A.2.

142

 

 

 

 

 

 

 

 

 

 

 

 

 

Au Samples p(295K) (an) p(4.2K) (an) Ap(an)
h1387—8a 47.83 25.24 22.59
h1388-8a 32.40 9.16 23.24
h1388-8b 43.99 20.72 23.27
h1397—8a 36.05 12.64 23.41
h1397-8b 60.19 32.90 27.29
h1398-8a 43.46 18.79 24.67
h1398—8b 48.15 22.14 26.01
h1399-8a 38.06 13.80 24.26
Averages 44 :l: 8 19 i 6 24 :t 2

 

 

 

Table A.2: The resistivities of sputtered gold (Au) samples at 295K and 4.2K and
their differences. Ap = p(295K) — p(4.2K).

 

 

 

 

 

 

 

 

 

 

Pt Samples p(295K)(an) p(4.2K)(nflm) Ap(n§2m)
h1363~8a 164.00 37.45 126.55
h1364—8a 168.17 50.79 117.38
h1368-8b 153.94 36.51 117.43
h1368-8a 167.20 47.08 120.12
h1364-8b 153.30 38.21 115.09
Averages 161 :l: 7 42 i 6 119 :l: 4

 

 

 

Table A.3: The resistivities of sputtered platinum (Pt) samples at 295K and 4.2K
and their differences. Ap = p(295K) — p(4.2K).

 

 

 

 

 

 

 

 

Pd Samples p(295K)(nf2m) p(4.2K)(an) Ap(an)
h1344-8b 155.0 42.28 112.72
h1345—8b 150.3 36.76 113.54
h1344-8a 166.7 40.75 125.95
Averages 157 :l: 8 40 :l: 3 117 j: 7

 

 

 

Table AA: The resistivities of sputtered palladium (Pd) samples at 295K and 4.2K

and their differences. Ap = p(295K) — p(4.2K).

 

 

 

 

 

 

 

 

 

 

Nb Samples p(295K)(an) p(12K)(an) TC(K) Ap(an)
h1387-8b 194.55 43.58 8.88 150.97
h1369—8a 231.64 90.38 8.45 141.26
h1345—8a 203.34 46.57 9.00 156.77
Averages 210 :l: 11 60 :l: 15 8.78 :l: 0.17 150 :l: 5

 

 

Table A.5: The resistivities of sputtered niobium (Nb) samples at 295K and 12K
and their differences. T 6 represents superconducting transition temperatures. Ap =
p(295K) — p(12K).

143

 

 

 

 

 

 

 

 

Ag0,958n0,05 Samples p(295K)(an) p(4.2K)(an) Ap(an)
h1579-5a 170 149 21
h1579—5b 172 152 20
Averages 171 i 1 150.5 2!: 1.5 20.5 :l: 0.5

 

 

Table A6: The resistivity of sputtered silver-tin (Ago,95Sn0,05) samples at 295K and
4.2K and their differences. Ap = p(295K) — p(4.2K).

A.3.1 Discussion
Matthiessen’s Rule

The electrical resistivity is a result of scattering of conduction electrons by other

electrons, impurities, dislocations, phonons etc. If these scattering mechanisms are

completely independent of each other, i.e. presence of one mechanism does not alter

the way other mechanisms function, then, the total resistivity is simply the sum of

resistivities due to all scattering mechanisms. This is known as Matthiessen’s rule

[1]. There are two types of scattering mechanisms in a metal:

1. Temperature independent e.g. elastic impurity scattering, (neither the concen-

tration of impurities nor their interaction with electrons is affected by temper-

ature).

2. Temperature dependent e.g. electron-electron and electron-phonon scattering.

According to the Matthiessen’s rule the total resistivity is given by;

p(c, T) = p(T) + 0(6),

(A.3)

where T is temperature and c is the concentration of impurities [1]. According to

Eqn. A.3, the addition of impurities merely adds a constant component to the total

resistivity, leaving the temperature dependent part unchanged. In reality Eqn. A.3 is

144

never exactly valid. There are always deviations from the Matthiessen’s rule, which
exist for several reasons which are explained in Ref. [152], but basically because the
scattering mechanisms are not completely independent. However, in many cases these
deviations are small compared to either p(T) or to p(c) or to both, and Matthiessen’s
rule represents quite a good approximation to the experimental results.

Let’s consider two samples, one pure and the other has impurities. From Eqn. A.3
one can immediately see that the difference in the resistivities ( Ap = p(T1) — p(T2))
at different temperatures depends only on temperature not impurities. Therefore, in
theory both samples should have same Ap because the impurity scattering is temper-

ature independent.

Comparison with other measurements

Table A.7 shows that Ap of our sputtered nominally-pure—metal samples differ from
the Ap of pure metal counterparts by ~ 4.5-12 ‘70, which is in the ballpark of observed
deviations from Matthiessen’s rule [152]. On the other hand, for a well deﬁned alloy of
Ag0,95Sn0,05 the difference between the Ap of pure Ag and Ap of sputtered Ag0_95Sno,o5

is 25% which is larger than those for sputtered pure metals.

 

 

 

 

 

 

 

 

 

 

Metal Apl,u,.e(nﬂm) Apspuueredmﬂm) Difference Ap(%)
All 22 24 :l: 2 10 2t 10
Pt 106.2 119 :l: 4 12 :l: 4
Pd 105.8 117:1:7 11:1:7
Nb 143.5 150 :l: 5 4.5 :l: 3.5
Agoessnoos 16 20 25

 

 

Table A.7: A comparison between the Aps of pure Au, Pt, Pd, Nb and Ag and
their sputtered counterparts. Ap = p(295K) — p(4.2K) except for Nb, for which
Ap = p(295K) — p(12K). The values of resistivities of the pure metals are taken
from Ref. [103].

145

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