CHAPTER 3 EXPERIMENTAL METHODS
3.9 Properties measurements 44
Superconducting Quantum Interference Device magnetometry (SQUID)
A Superconducting Quantum Interference Device (SQUID) uses the properties of electron-pair wave coherence and Josephson Junctions to detect very small magnetic fields.
The MPMS5 system comprises of two main sections: the dewar, probe and SQUID assembly, and the electronic control system. The probe contains a high precision temperature control system, allowing measurements between 2 - 400K and superconducting electromagnet, giving a field of up to 5.5 Tesla.
The dewar consists of an inner liquid helium reservoir and outer liquid nitrogen jacket, to reduce excessive liquid helium boil off. The liquid helium is used for both maintaining the electromagnet in a superconducting state and for cooling the sample space. Samples are mounted within a plastic straw and connected to one end of a sample rod, which is inserted into the dewar/probe. The other end is attached to a stepper motor, which is used to position the sample within the center of the SQUID pickup coils. These measurements have a sensitivity of 5×10-8 E.M.U. To measure the magnetic properties of (Ga, Mn)As epi-layers, the sample is cooled down to 10K under an external field of 1000 Oe applied along the measurement axis. Then field is removed, and projection of the remnant magnetization along the measurement axis is recorded as a function of temperature.
CHAPTER 4
RESULTS AND DISCUSSION
4.1 Effect of substrate orientation on arsenic precipitation in low-temperature-grown GaAs
Arsenic precipitation in “superlattice” structures of alternately undoped and [Si] = 3 × 1018 cm-3 doped GaAs grown at 250 °C on (100), (311)A, and (311)B GaA substrates has been studied. The cross-sectional bright-field TEM images of Samples H1, H2 and H3 are shown in Figs. 4.1(a), (b), and (c). For reference, plan-view TEM images of arsenic precipitations are shown in Figs. 4.2(a), (b), and (c). It is found that upon annealing at 800℃, the precipitates are totally confined in Si-doped regions, forming into two-dimensional arrays located approximately at the center of each Si-doped layer. The precipitates in each LT
“superlattice” structure also differ in density and size. Arsenic precipitates in Sample H3 (on GaAs(311)B substrate) are slightly denser and larger than those in Sample H2 (on GaAs(311)A substrate, and both are remarkably denser and larger than those in Sample H1(on GaAs(001) substrate. This can be explained by the varying excess arsenic incorporations in LT-GaAs grown on differently oriented substrates. The brighter zones observed near the top and bottom of each image are the AlAs layers, which thus provide good markers for the periodic structure. It can be seen that most of the spherical-like arsenic clusters are confined in Si-doped regions. The preferential precipitation can be explained by the so-called strain-driven vacancy-assisted diffusion mechanism [Chang-98-587, Hsieh-96-1790]. This is because Si-doped GaAs contains more Ga vacancies than intrinsic GaAs; moreover, in the former,
7.20, and 7.58 ×1015 cm-3 for Samples H1, H2, and H3, respectively. The number density of arsenic clusters in each sample is determined by assuming that the penetration thickness of electron beam of TEM is 100 nm.
The DXRD patterns of as-grown control Samples H4, H5 and H6 were recorded in the vicinity of the GaAs (004) reflection for the (001) substrate and (311) reflections for (311)A and (311)B substrates, and the resulting rocking curves are shown in Fig. 4.3 Clearly, two peaks due to the Bragg reflections from the LT-GaAs epilayer and the GaAs substrate are present in all control samples. The peak separations of 72, 144, and 162 arcsec, and the corresponding perpendicular lattice mismatches (Δa/a)⊥ of 0.05, 0.138, and 0.155% were observed for control LT-GaAs Samples H4, H5, and H6, respectively. This DXRD result agrees generally with what O’Hagan and Missous have observed. [O’Hagan-97-2400, Missous-97-197]
The lattice expansion should be attributed mainly +to growth conditions such as the growth temperature Tg and the ratio of As to Ga fluxes, JAs/JGa.
The presence of two well-defined peaks indicates the existence of a single-crystal layer containing excess arsenic other than bulk GaAs in antisite and interstitial positions. The lattice expansion of the LT GaAs epilayer is directly proportional to excess arsenic content in the layer by double-crystal X-ray diffraction [Lavrent’eva-02-118]. A linear correlation exists between the lattice mismatch (Δa/a)⊥ and the concentration of excess arsenic [Asex]. This correlation dependence can be used to estimate the total concentration of excess arsenic from the DXRD data. Therefore, the [Asex] values of the LT-GaAs layers grown on (100), (311)A, and (311)B substrates were estimated approximately to be 1.0, 2.7, and 3.1 ×1020 cm-3, respectively.
As summarized [Asex] and (Δa/a)⊥
values are dependent on J
As/JGa and growth temperature. In this study, the (001) and (311) substrates were mounted side by side on therelated to the JAs/JGa ratio and growth temperature but to the surface bonding configuration.
[Wang-02-2965; Wang-85-826]. On a (311)A surface, the single-dangling-bond sites are arsenic sites and the double-dangling-bond sites are Ga sites. The (311)B surface, on the contrary, has an opposite bonding arrangement, as shown in Fig. 4.4 Such bond configurations would cause the incorporation of excess arsenic in Ga sites to form AsGa defects more easily on the (311)B than on the (311)A surfaces. Our DXRD result is consistent with this analysis that the LT-GaAs grown on the (311)B surface has a larger lattice expansion than that grown on the (311)A surface.
(001) (a)
Figure 4.1(a) Cross-sectional TEM bright-field images of Sample H1 (The annealed LT-GaAs on GaAs(001) substrate)
(311)A (b)
Figure 4.1(b) Cross-sectional TEM bright-field images of Sample H2 (The annealed LT-GaAs on GaAs(311)A substrate)
(311)B (c)
Figure 4.1(c) Cross-sectional TEM bright-field images of Sample H3 (The annealed LT-GaAs on GaAs(311)B substrate)
(a)
(001) 50 nm
Figure 4.2(a) Plan-view TEM images of Sample H1 (The annealed LT-GaAs on GaAs(001) substrate)
(311)A
(b)
50 nm
Figure 4.2(b) Plan-view TEM images of Sample H2 (The annealed LT-GaAs on GaAs(311)A substrate)
(c)
(311)B 50 nm
Figure 4.2(c) Plan-view TEM images of Sample H3 (The annealed LT-GaAs on GaAs(311)B substrate)
-1200 -900 -600 -300 0 300 600 900 1200 0 (311)B
(311)A (001)
162 Sec 144 Sec 72 Sec
In te n s it y (a rb . u n its )
Δ2θ ( Arc sec)
Figure 4.3 DXRD rocking curves of as-grown LT-GaAs of Samples H4, H5, and H6: (004) reflection for sample grown on (001), (311) reflections for samples grown on (311)B and (311)A substrates
Ga As
(311)A
(311)B
Figure 4.4 Schematic illustration of ideal atomic structures at (311)A and (311)B surfaces
4.2 Effects of doping type and concentration on precipitation of arsenic clusters in low-temperature-grown GaAs
In this study, the effects of doping type and concentration on arsenic precipitation in low-temperature-grown GaAs upon postgrowth annealing at 600, 700 and 800°C were investigated. Figures 4.5(a), (b), and (c) show the bright-field TEM images of Samples B1, B2 and B3 after 30 s anneals at 600, 700, and 800°C, respectively. The bright stripes seen in each picture are the AlAs layers. It is noted that AlAs layers provide good markers among the (i-n-i) or (i-p-i) regions. Arsenic precipitation depletion was observed in all Be-doped layers for all annealing temperatures; moreover, the arsenic clusters accumulate toward the undoped layers.
The DXRD rocking curves of GaAs (004) reflection for the three control Samples B4, B5 and B6 are shown in Fig. 4.6 Clearly, peak separations of 108 and 72 arc sec, caused by excess arsenic incorporated in LT GaAs, were observed in Be-doped and undoped LT-GaAs epilayers, respectively; however, no obvious peak separation was observed in the Si-doped LT GaAs. This suggests that the Be doping enhanced the incorporation of excess arsenic into the lattice of LT GaAs while the Si doping suppressed this effect. According to a previous report [Liu. 95-279], the concentration of arsenic antisites, [AsGa], is directly proportional to the lattice expansion of the LT-GaAs epilayer. Therefore, we can conclude that [AsGa]p3> [AsGa]i
> [AsGa]n3, where [AsGa]x denotes the arsenic antisite concentration in the x-doped control sample. Moreover, during postgrowth annealing, the reduction of strain caused by the AsGa
defects is also a driving force for precipitation. For the
i-p
3-i region, the
AsGa defects diffused from the p3 to i layers, and for the i-n3-i region from i to n
3layer, which is opposite to
the VGa defects. Our experimental result about i-n3-i region and all i-p-i regions is therefore
consistent with the vacancy-assisted arsenic antisite diffusion mechanism [Melloch-92-3509;Chang-992442; Chang-98-587]. Interestingly, a “dual” arsenic precipitation phenomenon were observed
1018cm-3doped layer for all annealed temperatures; however, arsenic precipitates are depleted away the other two lightly Si-doped regions (n = 2 x 1016and 2 x 1017cm-3) toward the undoped regions. The latter depletion behavior contrasts with the earlier observation.
[Chang-99-2442].In their work, the i-n1
-i-n
2-i-n
3-i multiplayer structure did not consist of any
isolation layer to prevent the inter-diffusion of defects and doping effect from n1 to n2 or n3. As a result, arsenic precipitation behaviors were disturbed and the correlation of doping concentration to arsenic precipitation could not be clearly identified.In order to understand the “dual” arsenic precipitation phenomenon in Si-doped GaAs layers indicates that other effects must be considered in addition to the vacancy-assisted arsenic antisite diffusion mechanism. The Debye length can influence the arsenic precipitation at the i/n interface, and the depletion of the “n” regions to a depth of up to a few Debye lengths would occur at each i/n interface [O’Hagan-96-8384]. Moreover, the Debye lengths at the current annealing temperatures in our structure are estimated to be around 560, 180, and 50 Å for n1,
n
2 and n3 layers, respectively. Therefore, the total depletion depth should be ~1120,~360 and ~100 Å for n1, n2 and n3 layers, respectively, due to the double n/i interfaces in each active region. Therefore, since each Si-doped layer is 350 Å thick, it is likely that complete depletion occurred in the n1 and n2 layers but an incomplete depletion in the n3 layer, the latter was also noted as an “arsenic accumulation” layer.
Figure 4.7 shows the average number density of arsenic clusters in each i-x-i region after annealing at 700 and 800 °C, respectively. The volume is determined by assuming that the penetration thickness of electron beam of TEM is 100 nm. The analysis reveals that the cluster densities of the i-p-i regions are higher than those of the i-n-i regions under the same
Figure 4.5(a) TEM bright field images of Sample B1 showing arsenic precipitates in different active regions after annealing at 600°C
(a)
i
(b)
n
1i i n
2i i n
3i i p
1i i p
2i
100 nm i
p
3i
Figure 4.5(b) TEM bright field images of Sample B2 showing arsenic precipitates in different active regions after annealing at 700°C
i
Figure 4.5(c) TEM bright field images of Sample B3 showing arsenic precipitates in different active regions after annealing at 800°C
i p
100 nm
3
i
-1 000 -500 0 500 1000
72 arc sec
undoped
108 arc sec
Be-doped Si-doped
Norm al iz ed x- ray count s ( a .u .)
Arc Sec
Δ2θ ( Arc sec)
Figure 4.6 DXRD rocking curves of GaAs (004) for 1-μm-thick undoped (Sample B6),
2 4 6 8 10
800
0C 700
0C
Cluster Density (X 10 15 cm -3 )
i-n1-i i-n2-i i-n3-i i-p1-i i-p2-i i-p3-i
Figure 4.7 Arsenic cluster number density in each i-x-i region after annealing at 700°C (Sample B2) and 800°C (Sample B3), where x = n or pi i
4.3 Effects of thickness and post-annealing on the magnetic properties of Ga
0.93Mn
0.07As
In the study focusing on the effects of thickness and post-annealing on the magnetic properties of Ga0.93Mn0.07As, where high concentration Mn content exists. The M-H curve of post-annealed Sample C1 is shown in Fig. 4.8, which indicates a typical ferromagnetic hysteresis behavior and easy axis near [110] direction. All samples show a ferromagnetic state at 10 K as confirmed by the M-H curves. Figure 4.9 shows the M-T curves of as grown and post-annealed Samples C1, C2 and C3, which were measured in a small magnetic field (H = 1 Oe) applied in-plane direction. The Curie temperature (Tc) of post-annealed Samples C1, C2 and C3 are 160, 135 and 70 K, respectively, which are 80, 40 and 10 K higher than that of the as-grown samples. It is noted that the Tc of the annealed Sample C1 is considerably higher (160 K) than that of the Sample C3 (70 K). It is apparent that the Tc, depends on the thickness of Ga0.93Mn0.07As epi-layers, and the Tc value increases as the thickness decreases. Figure 4.10 shows the Tc
values of the as-grown and post-annealed Samples C1, C2 and C3.
The low-temperature growth of (Ga,Mn)As inevitably leads to numerous point defects in (Ga,Mn)As, presumably AsGa, and MnI both resulting in heavy compensation in (Ga,Mn)As due to the donor nature of those defects. This compensation prevents Mn atoms from acting as acceptors. Therefore, the MnI defects and their diffusion during growth and annealing play a key role in determining the magnetic properties of (Ga, Mn)As epilayers. During growth, the MnI atoms tend to diffuse to the surface of the (Ga, Mn)As epilayer, resulting in an increase in the effective hole concentration in the bulk and thus enhancing the exchange
Ga0.93Mn0.07As epi-layer, respectively. The rocking curves of Sample C3 layers clearly reveal an expansion of the lattice constant along the growth direction, indicating that samples are tetragonal and coherently strained to the GaAs substrate. When growing at a temperature of about 230℃, LT-GaAs buffer layer and LT-(Ga, Mn)As epi-layers are non-stoichiometric containing As , Mn and MnGa Ga I defects, resulting in an expansion of the lattice. The Ga0.93Mn0.07As epi-layers peaks are 696 and 636 arcsec for post-annealed and as-grown Sample C3. Annealing decreases the lattice expansion of Ga0.93Mn0.07As epi-layers by about 60 arcsec. But the peak separation of LT-GaAs buffer layers for post-annealed and as-grown Sample C3 shows no obvious difference. It reveals that annealing effect removes not the AsGa
defects, but the Mn defects. A large proportion of donor-like MnI I defects of Ga0.93Mn0.07As epi-layers are removed by annealing. This results in the strengthening of the hole-mediated interaction between substitutions Mn of Sample C1. Sample C3 of Ga0.93Mn0.07As epi-layer was 1000 nm thick and its corresponding out-diffusion is insufficient in comparison with those of thin Samples C1 and C2.
The Curie temperature of DMS can be greatly increased by a decrease in thickness, and via annealing treatment Annealing treatment is essentially to remove excess MnI from the interstitial sites in the lattice to decrease the donor-like defects, which may cause an increase in hole concentration and Tc. In other words, the diffusion path of Mn atoms for the thinner DMS thickness is much shorter, which may result in a more effective removal of excess MnI
from the lattice and a greater increase in Tc.
The magnetic field dependence of magnetization (M-H) curves was measured along [110], , [100] and [010] axis, respectively. M-H curves of as-grown and annealed Sample C1 are shown in Fig. 4.12 (a) and (b). The easy axis of as-grown and annealed Sample C1 can be rotated from [100] direction to [010] direction by low-temperature
] 10 1 [−
-200 0 200 -2
-1 0 1 2
(001)Ga0.93Mn0.07As
] 1 1 0 [
] 1 0 1 [
−10K
M ( x 10 -5 emu)
H (Oe)
H along H along
Figure 4.8 M-H curves of (Ga, Mn)As DMS (Sample C1) measured with in-plane magnetic field applied along [110] or [1−10] directions at 10K
0 20 40 60 80 100 120 140 160 180 0.0
0.5 1.0 1.5 2.0
annealed C1 as-grown C1 annealed C2
M ( x 10 -5 emu)
T (K)
as-grown C2 annealed C3 as-grown C3
Figure 4.9 M-T curves of as-grown and post-annealed (Ga, Mn)As DMS (Samples C1, C2 and C3) measured under 1 Oe in-plane magnetic field
25 nm 100 nm 1000 nm 40
60 80 100 120 140 160 180
Δ T
C
T C (K )
Thickness
as-grown annealed
Figure 4.10 Curie temperature values (TC) vs. thickness of as-grown and post-annealed (Ga, Mn)As DMS (Samples C1, C2 and C3)
-1000 -500 0 500 10
100 1000 10000 100000 1000000
Ga
0.93Mn
0.07As
GaAs sub.
LT-GaAs
Intensity ( Counts )
as grown C3 annealed C3
Δ2θ ( Arc sec)
sec
Figure 4.11 The DXRD rocking curves of as-grown and post-annealed (Ga, Mn)As DMS (Sample C3)
-600 -400 -200 0 200 400 600
Figure 4.12(a) M-H curves of as-grown (Ga, Mn)As DMS (Sample C1) measured with in-plane magnetic field applied along [110], [1−10] , [100], [010]
-600 -400 -200 0 200 400 600
Figure 4.12(b) M-H curves of post-annealed sample C1 measured with in-plane magnetic field applied along [110], [1−10], [100], [010] directions, respectively, at 10K
4.4 Effects of substrates orientation on the magnetic properties of Ga
0.93Mn
0.07As
The effects of substrate orientation on magnetic properties of Ga0.93Mn0.07As were grown on (001), and (311)A GaAs substrates. The M-T curves of as grown and post-annealed Samples D1 and D2 were measured in small magnetic field (H = 1 Oe) applied in-plane directions. The diamagnetic contribution from the substrate has been subtracted. Figure 4.13(a) shows M-T curves of as-grown and post-annealed Samples D1, with Ga0.93Mn0.07As epi-layer on (001) oriented substrate, measured with a small in-plane magnetic field of 1 Oe applied.
Figure 4.13(b) shows M-T curves of as-grown and post-annealed Samples D2, with Ga0.93Mn0.07As epi-layer on (311)A oriented substrate, measured with a small in-plane magnetic field of 1 Oe applied.
Annealed samples showed Tc value of 160K if (001) oriented and 140K if (311)A oriented. It is apparent that in post-annealed samples, Tc is dependent on the orientation of substrates. Compared with the as-grown and post-annealed samples increments in Tc
dependent on the orientation of substrates are exhibited. The increment of Tc value of Ga0.93Mn0.07As was 80K if (001) oriented and 60K if (311)A oriented.
All samples showed a ferromagnetic state at 10 K as confirmed by the M-H curves.
Figure 4.14(a) shows M-H curves of Ga0.93Mn0.07As epi-layer on (001) oriented substrate Sample D1, measured with in-plane magnetic field applied along [110], ] directions at 10K. Figure 4.14(b) shows M-H curves of Ga
] 10 1 [−
0.93Mn0.07As epi-layer on (311)A oriented substrate Sample D2, measured with in-plane magnetic field applied along [011−], [2−33]
Figure 4.15 shows the DXRD rocking curve of (004) as-grown and post-annealed Sample D3 exhibiting three distinct peaks for GaAs substrate, LT-GaAs buffer layer and Ga0.93Mn0.07As epi-layer, respectively. The rocking curves of Sample D3 layers clearly reveal an expansion of the lattice constant along the growth direction, indicating that samples are tetragonal and coherently strained to the GaAs substrate. When growing at a temperature of about 230℃, LT-GaAs buffer layer and LT-(Ga, Mn)As epi-layers are non-stoichiometric containing AsGa, MnGa and MnI defects, resulting in an expansion of the lattice and a high concentration of AsGa and MnI defects with a donor-like character. The Ga0.93Mn0.07As epi-layers peaks are 696 and 636 arcsec for post-annealed and as-grown Sample D3. Annealing decreases the lattice expansion of Ga0.93Mn0.07As epi-layers by about 60 arcsec. But the peak separation of LT-GaAs buffer layers for post-annealed and as-grown Sample D3 shows no obvious difference. It reveals that annealing effect removes not the AsGa defects, but the MnI defects.
A large proportion of donor-like MnI defects of Ga0.93Mn0.07As epi-layers are removed by annealing. The model proposed [Blinowski-03-121204], MnI may form antiferromagnetic coupling pairs with the nearest substitutional MnGa, which would also suppress the total ferromagnetic exchange interaction and further reduces Tc. Therefore, the MnI defects diffusion during growth and low temperature annealing play a key role in determining the magnetic properties of (Ga, Mn)As epilayers.
In the previous reports [Lee-05-6399], the effect of substrate orientation on arsenic precipitation in low-temperature-grown GaAs, had proved that AsGa defects on (311)A were higher than (001) substrates. Similarly in Ga0.93Mn0.07As epilayers, AsGa defects were higher on (311)A than on (001) oriented. AsGa defects cannot be completely removed during low temperature annealing. Disparity of the hole concentration of the Ga0.93Mn0.07As epilayers are compensated by the donor-like As antisite (AsGa) defects, resulting in higher
Δ T
c of-20 0 20 40 60 80 100 120 140 160 180 200 0.0
0.2 0.4 0.6 0.8 1.0 1.2 1.4
1.6 annealed sample
as-grown sample
(001) Ga
0.93
Mn
0.07
As
M ( x 10 -5 emu)
T (K)
(a)
Figure 4.13(a) M-T curves of as-grown and post-annealed (Ga, Mn)As DMS (Samples D1), with Ga0.93Mn0.07As epi-layer on (001) oriented substrate, measured under 1 Oe in-plane
0 50 100 150 200 0.0
0.2 0.4 0.6 0.8 1.0 1.2 1.4
1.6 annealed sample
as-grown sample
(311)A Ga
0.93Mn
0.07As
M ( x 10 -5 emu)
T (K)
(b)
Figure 4.13(b) M-T curves of as-grown and post-annealed (Ga, Mn)As DMS (Samples D2), with Ga0.93Mn0.07As epi-layer on (311)A oriented substrate, measured under 1 Oe in-plane magnetic field
-200 0 200
Figure 4.14(a) M-H curves of Ga0.93Mn0.07As epi-layer on GaAs(001) oriented substrate (Sample D1), measured with in-plane magnetic field applied along [110], [1−10]] directions at
-600 -400 -200 0 200 400 600
Figure 4.14(b) M-H curves of Ga0.93Mn0.07As epi-layer on GaAs(311)A oriented substrate (Sample D2), measured with in-plane magnetic field applied along , directions at 10K
] 1 01
[ − [2−33]
-1000 -500 0 500 10
100 1000 10000 100000 1000000
Ga
0.93Mn
0.07As
GaAs sub.
LT-GaAs
Intensity ( Counts )
sec
as grown D3 annealed D3
Figure 4.15 The DXRD rocking curves of as-grown and post-annealed (Ga, Mn)As DMS (Sample D3)
4.5 Property of (In0.52
Al
0.48 1-x) Mn
xAs / In
0.52Al
0.48As / InP layer structure
A series of diluted magnetic semiconductors, (In0.52Al0.48)1-xMnxAs (0 < x ≤ 0.11), was grown on InP substrate by low-temperature molecular beam epitaxy. The cross-sectional image projected along [110] zone of Sample E4 is shown in Figure 4.16(a) is the bright field TEM image of the interfacial region between (In0.52Al0.48)0.92Mn0.08As epilayer and In0.52Al0.48As buffer layer and (b) is the lattice image with (In0.52Al0.48)0.92Mn0.08As epilayer on the top and In0.52Al0.48As buffer layer at the bottom. Both RHEED patterns and TEM measurements confirmed that the epitaxial (In0.52Al0.48)1-xMnxAs active layer has a zinc-blend structure. Figure 4.17 shows the DXRD rocking curves of (004) for Sample E2, E4, and E5 with x = 0.05, 0.08, and 0.11, respectively. Obvious peak separation due to the incorporation Mn atoms was observed. The lattice expansion is about 0.20, 0.29, and 0.50 % for Samples E2, E4, and E5, respectively, increasing with the increase in Mn content.All samples of magnetization were measured with the direction in-plane magnetic field. The total magnetization of (In
] 10 1 [−
0.52Al0.48)1-xMnxAs/InP samples includes two components, i.e., the magnetization of (In0.52Al0.48)1-xMnxAs epilayer and InP substrate.
Therefore, in order to obtain the “net” magnetization of the active layer, the magnetization of the “bare” InP substrate with same size must be separately measured and carefully subtracted from the total magnetization. Inset of Fig. 4.18 shows the magnetization as a function of applied field of (In0.52Al0.48)0.95Mn0.05As/InP and corresponding “bare” InP substrate with same size. We can find that the M-H curve of InP substrate shows almost a linear relationship, which suggests that the semi-insulating InP substrate exhibits diamagnetic behavior, similar to that of semi-insulating GaAs substrate. Compared with that of the InP substrate, M-H curve of (In0.52Al0.48)0.95Mn0.05As/InP shows a similar linear relationship with a smaller absolute slope.
Figure 4.18 shows the “net” magnetization as a function of applied field of
Therefore, The Sample E2 [(In0.52Al0.48)0.95Mn0.05As] is belived to possess the paramagnetic-like behavior at 5K. It should be noted that similar phenomenon was also observed in Sample E1 with x = 0.03.
Samples with x ≥ 0.06 showed a ferromagnetic state at 5 K as confirmed by the M-H curves. For example, the M-H curves of Sample E4 and E5 are shown in Fig. 4.19 Obviously,
Samples with x ≥ 0.06 showed a ferromagnetic state at 5 K as confirmed by the M-H curves. For example, the M-H curves of Sample E4 and E5 are shown in Fig. 4.19 Obviously,