Multiple stacking of InGaAs/GaAs (731) nanostructures

We studied the multilayering effects of InGaAs quantum dots (QDs) on GaAs(731), a surface lying inside of the stereographic triangle. The surfaces after stacking 16 InGaAs layers were characterized with highly non-uniformity of QD spatial distribution. The bunched step regions driven by strain accumulation are decorated by QDs, therefore GaAs(731) becomes a good candidate substrate for the growth of QD clusters. The unique optical properties of the QD clusters are revealed by photoluminescence measurements. By adjusting the coverage of InGaAs, a bamboo-like nanostructured surface was observed and the quantum dots aligned up in clusters to separate the “bamboo” into sections.

The morphological evolution during InGaAs deposition on GaAs high index surfaces has been a subject of numerous experimental efforts for better understanding of growth physics and potential applications of the resulted nanostructures [1][2][3][4][5][6][7][8].
Most of the researches chose a surface index located between two low index planes out of the three, (100), (110), and (111).
In other words, the popular choice was a surface index on the sides of the stereographic triangle defined by the three low index planes [3]. For example, GaAs(311), one of the most popular choice, locates between the (100) and (111) planes.
Both quantum wires and quantum dots (QDs) were observed during the InGaAs growth on GaAs(311) [1,2,6,8]. Especially during the processes of stacking multiple layers of InGaAs/GaAs(311) QDs, two-dimensional or one-dimensional arrays of QDs were achieved [8]. In the letter, we study the multilayering effects of InGaAs nanostructures on GaAs(731), a surface index lying within the stereographic triangle. The (731) plane became particularly interesting after the dominating facets of InAs/GaAs(100) QDs were identified to this index [9].
Massive amount of investigations have been performed to reveal the growth behavior of InGaAs on GaAs(100) due to the importance of the resulted QDs for optoelectronic applications [10][11][12].
Experiments were carried out in a solid source Molecular Beam Epitaxy (MBE) 32P Riber system. The adopted GaAs(731) substrates were of n-type and epitaxial ready. After oxide desorption, a GaAs buffer layer of 500 nm was grown at the substrate temperature of 580°C. The GaAs growth rate was one monolayer per second (ML/s) under a constant As beam equivalent pressure (BEP) of 1×10 -5 Torr. Sixteen layers of 7 ML In 0.4 Ga 0.6 As nanostructures were subsequently deposited, partitioned by GaAs spacer in different thicknesses, 120 ML, 70 ML, and 50 ML, respectively. The growth of In 0.4 Ga 0. 6 As layer was performed at 540°C and As BEP of 4.4×10 -6 Torr. After the growth of first 5 MLs of GaAs spacer, the growth temperature was rapidly increased up to 580°C and As BEP was turned back to 1×10 -5 Torr. One more sample, grown in addition to the above series, had the GaAs spacer of 70 ML but the In 0.4 Ga 0.6 As coverage of 5.7 ML. Following the growths, samples were  GaAs (11 5 2), just 2.28° away from GaAs(731), has lower surface energy according to theoretical calculation and was observed to be a stable surface by scanning tunneling microscopy (STM) [13]. Therefore, GaAs(731) is better described as a vicinal GaAs(11 5 2) surface, with straight steps along [-113] and ragged steps along . The large terraces that are free of QDs shown in Fig. 1(a) resulted from the strain-driven step bunching. It is well documented that QDs prefer to stepped regions for nucleation and growth, especially ragged steps due to its active reaction. The observation in Fig.   1(a) demonstrates that GaAs(731) is a substrate candidate for the growth of InGaAs QD clusters. The corresponding PL spectrum from this sample is shown in Fig. 2. The strong and sharp peak about 1.38 eV can be assigned to the InGaAs wetting layer (WL) covering the whole surface. The assignment is consistent with the high peak energy and large terraces without QDs. Different from usual QDs reported where the PL peak from WL is weak, the carriers excited in InGaAs/GaAs(731) WL cannot find nearby QDs to relax and therefore recombine in WL giving rise to a strong signal. The relatively weak peak around 1.25 eV is expected from the clustered QDs.
By reducing the GaAs spacer thickness from 120 ML to 70 ML, more strain could transmit from the underneath InGaAs layer to the subsequent layers. As a result of the strain-driven mechanism, Fig. 1(b) reveals that the InGaAs QDs grow bigger in size and the free-of-QDs terraces disappear. Consistently, the WL peak disappears from its associated PL spectrum in Fig. 2 and the QD peak around 1.25 eV becomes dominant. Please notice, this peak is from 15 buried layers of QDs. The luminescence from the exposed layer of surface QDs is usually not notable, as in the PL spectrum form the sample with 120 ML spacer. However, the PL contribution from surface QDs can be enhanced as the spacer becomes thinner, as the consequence of carrier tunneling from the buried QDs to the surface. In line with the literature [6], the new developed PL peak around 1.0 eV is assigned to the layer of surface QDs.
With further reducing the GaAs spacer to 50 ML, more strain is accumulated to the surface and the QDs further ripe and even elongate along [-113], as revealed in Fig. 1(c). Due to the thinner spacer, more carriers are able to tunnel from buried layers to the surface. The PL peak from surface QDs becomes stronger, and eventually dominates the spectrum as shown in Fig. 2.
As noticed, the amount of accumulated strain is the key in determining the nanostructure configuration in Fig. 1 and the optical properties in Fig. 2. Specially, the sample with less accumulated strain shown in Fig. 1(a) demonstrated the potential of GaAs(731) for growing QD clusters. Therefore, to pursue better QD clusters, we introduce a new sample with less InGaAs coverage, 5.7 ML, in order to significantly reduce the  Fig. 3. Comparing Fig. 3(a) to Fig. 1(b), two samples with the same spacer thickness of 70 ML, we are able to rebuild the QD clusters by reducing the accumulated strain. The surface is characterized by steps plus terraces again and QD clusters decorate the ragged steps.
Remarkably, the QDs in clusters are aligned as short chains. The visual appearance of the resulted surface morphology resembles itself as a jungle of bamboo. The bunched steps along [-113] visualize the individual bamboo sticks and aligned QD clusters separate each stick into sections. Very interesting, the QD clusters are nearly equally distributed along [-113]. Its periodicity is better resolved in the two-dimensional image of the autocorrelation functions calculated from the AFM image [14]. Figure 3(c) shows a line profile along [-113] through the center of Fig. 3(b). The separation between satellite peaks reveals that the periodicity is about 565 nm. The separations between "bamboos" are not uniform but still resolved from has a high potential for applications in blue and UV light-emitting and laser diodes as well as in high temperature and high power electronic devices [1,2]. Up to date, the synthetic approaches for 1D GaN nanomaterials have relied on various physical or chemical methods, such as carbon nanotubes (CNTs) [3][4][5] and anodic alumina membrane [6,7] template-originated methods, Fe, Co, Ni and Au metal-catalyst vapor-liquid-solid (VLS) growth [8][9][10], SiO 2 , Fe 2 O 3 , and B 2 O 3 oxide-assisted growth (OAG) [11,12], Precursor-based methods metalorganic chemical vapor deposition (MOCVD) [13], and hydride vapor phase epitaxy (HVPE) [14].
Although the above-mentioned groups are different from each another, all methods are capable of generating 1D GaN nanomaterials with a characteristic single-crystalline structure.
Liliental-Weber [15] and Lan [16] have reported preferentiallyoriented 1D GaN nanomaterials consisting of small GaN platelets. Furthermore, Zhao [17] and Ogi [18] have studied the  nanoribbons are crystallized on the surface of CNTs, and, while a N 2 gas having impurities of O 2 , is introduced into the furnace, the chemical reactions proceed as follows: where a Ga 2 O vapor is generated in line with the reaction described by Han et al [3].
The vapor cannot be oxidized further after it reaches the surface of CNTs, because the CNTs serve as a shelter for the O 2 impurity in N 2 gas; (ii) Ga 2 O 3 nanoribbons are transformed into GaN nanoribbons while NH 3 is introduced into the furnace; the involved reaction may be written as follows, where GaN platelets may initially nucleate on the side-surface of a Ga 2 O 3 nanoribbon and the subsequent growth of GaN platelets propagates along the whole nanoribbon.
To confirm the growth mechanism of the yellow-colored Ga 2 O 3 nanoribbons, we carried out additional experiments.
Before the second step, we examined a thin layer of the white-colored materials on the surface of CNTs. The materials were determined to be Ga 2 O 3 nanoribbons having morphologies depicted in Fig. 1  found to be GaN. There were nanoribbons covered with nano-platelets. A nanoribbon is shown in Fig. 1(c), its axis is parallel to the (0002) plane, as shown in Fig. 1(d), which is the [112 0] zone axis diffraction pattern. The similarity between the morphologies in Fig. 1(a) and Fig. 1(c) suggests that the proposed growth mode for the nanoribbons depicted in Fig. 1(a) is correct. The reasonability of the proposed growth mode may require other evidence additive to the fact of similar morphologies of the two kinds of the GaN nanoribbons.
Detailed HRTEM investigation may become effective.       The formation of GaN nanowires filled in CNTs can be explained as direct ammoniation of Ga nanowires filled in CNTs. However, such explanation may find a better proof if a pure GaN nanorod not covered with a CNT may be found in the ammonia-treated materials. In fact, such material was found, as shown in Fig. 8(a). This rod has an axis perpendicular to the (11 00) plane, as displayed in Fig. 8(b) and (c). The most intriguing fact is that new contrast fringes appear in Fig. 8(a). It is believed that the fringes are originated from the thickness difference [21]. According to this, we can conjecture that the cross-section of the rod may have two kinds of possible shapes shown in Fig. 9. The above-presented analysis made us possible to estimate the thickness and the cross-section shape using the observed contrast fringes. It is noted though that a more definite confirmation using a slice method of cross-section is probably needed [22].
The formation of a pure GaN nanorod can also be explained due to ammoniation of Ga in CNTs. When a metallic Ga with a density of ~5.90 g/cm 3 is gradually transformed to GaN with a density of ~5.90 g/cm 3 , the volume increased by 20% as the molecular weight increased by 20%. This extra volume part gradually grows and leads to a resultant pure GaN nanorod whose cross-section may become non-round as compared to a round shape of the initial Ga nanorod.
To sum up, the presented transformation scheme consists of two steps and relies on conversion of a given 1D material to another 1D material in a given gaseous atmosphere. This approach has been used to synthesize ZnS [23] and GaN nanotubes [24,25]. It is believed here that the process is an important practical method to synthesize novel 1D material and to fabricate complicated arrays of those with nanoparticles.
We fabricated one-dimensional GaN nanoribbons, Over the past decades, single-walled carbon nanotubes (SWNTs) used as building blocks of nanoscale electronic devices have attracted great attention due to their quasi one-dimensional structure and unique properties [1][2][3].
Depending on their diameter and chirality, SWNTs can exhibit either metallic or semiconducting behavior. However, the chirality of nanotubes can not be controlled normally until now [4]. How to obtain purely semiconducting SWNTs (semi-SWNTs) becomes one of the major obstacles to the widespread applications of SWNTs for high performance electronics [5].
Several separation approaches for the enrichment of one certain electronic type of SWNTs have been developed, such as electrical breakdown [6][7][8][9], ultracentrifugation [10], selective plasma etching [5], electrophoresis [12] and selective chemical functionalization [6]. But the problem is that the posttreatments processes are tedious beyond controlled and often have the disadvantage of damaging or contaminating SWNTs samples [7,8]. Recently, preferential growth of Semi-SWNTs by chemical vapor deposition (CVD) methods by boron/nitrogen co-doping was reported to produce high percentage of Semi-SWNTs or a specific chirality distribution SWNTs [10][11][12][13]. These pioneering studies indicate the possibility for selective synthesis of Semi-SWNTs by heterodoping methode.
S contained compound has often been used as an additive in the synthesis of CNTs [14] for increasing the yield of SWNTs or      For better understanding the S-doping effects into metallic SWNT structure, calculations are carried out with the density functional theory implemented provided by Dmol 3 package [19].  Figure 5a demonstrates the source-drain current increase with increasing negative gate voltage. Figure 5b is the transfer characteristic of as-synthesized S-doped SWNTs-FET which shows that the on/off ratio is more than 10 3 by sweeping gate voltage at room temperature. In contrast, similar thin film field-effect configuration based on the undoped SWNTs sample shows the on/off ratio of less than 10, which means that the undoped sample may contain a lot of M-SWNTs.
CNT-FETs device based on the as-synthesized randomly networked S-doped SWNTs film not just a few CNTs showed a typical field effect characteristic feature, in which the drain current increased with increasing negative gate voltage, this reveals that the S-doped SWNTs is a p-type semiconductor [22][23][24][25][26][27][28][29][30][31][32][33][34]. A schematic of the transistor is shown as an inset in Generally, there are exogenous and endogenous poisons overloading in human body [1]. Acute medicament (e.g. soporific) toxicosis is the most common exogenous case. Multi-amine grafted SBA-15 and FSBA-15 was prepared according to published procedures [14,15]. C. Preparation of mesoporous silica KIT-6. The mesoporous silica with cubic Ia3d symmetry (designated as  was prepared according to the published procedure using tri-block copolymer Pluronic P123 as template by adding of n-butanol in an acidic aqueous solution [16]. Typically, 6 g of surfactant P123 and 6 g of n-butanol were dissolved in the solution of 217 g of distilled water and 11.4 g of concentrated

CHARACTERIZATION
Concentrations of all toxins are characterized using UV-Vis absorbance, which was measured using Shimadzu UV-3101PC spectroscope. However, all amine-grafted mesoporous silicas show higher adsorption capacities than their pure parent mesoporous silicas.
The same results can also be obtained from  C. Effect of initial concentrations of bilirubin. Figure 3 shows  F. Adsorption for other toxins. Pure and amine-grafted mesoporous silicas have also been applied to adsorb other toxins, such as cholic acid, uric acid, creatinine. As shown in Table 2, pure mesoporous silicas also show high adsoprtion   [3] structure with common processes due to the dissolution between polymer layers.
Good balance between electron and hole injection is very important for the functioning of both OLEDs and PLEDs. In PLEDs, the most common way to achieve this balance is by blending the electron and hole transporting materials. The electroluminescent (EL) efficiency is higher when the electron and hole transport materials are separated and attached to the cathode and anode layers, respectively. Another way to achieve this balance could be by producing a structure in which the two materials are graded along the film direction of growth. As there is no interface between the electron and hole transport materials in this structure, it is expected that there will be an increase in the lifetime of the devices. For small organic molecules, graded structures can be fabricated by codeposition processes [4][5][6][7] and annealing [8], but for polymers it is difficult to control the gradient profile in the active layer due to dissolution by the solvent. There have been attempts to achieve polymer graded structures, such as molecular-scale interface engineering [9], self-organization [10], and thermal transfer process [11]; however, the processes involved are not simple.
We have developed a vacuum spray (VS) method [12][13][14][15][16] that allows fabrication of polymer films with a controllable dye distribution along the film direction of growth, by which graded and/or bilayer structures can be prepared. This report describes the production of PLEDs based on films with homogenous and graded structures prepared using this VS method. For reference studies, devices based on spin-coated films were also investigated.  Adding an electron transport material to the hole transport material is one way to achieve better electron injection and to balance the carriers, and thus higher luminance can be expected. In this study, the electron transport material (Alq 3 ) was added into the hole transport (P3HT) material and the results are shown in Fig. 2. In the case of PLEDs with uniformly blended layers, both the current and the luminance of PLEDs prepared with the spray layer were higher than those of PLEDs prepared with the spin-coated layer. This may be because solvent residue is left in the spin-coated layer and/or that Alq 3 aggregates easily in layers prepared by the spin coating process due to the solvent residue. With the VS method, where almost no solvent residue is left in the layer, Alq 3 molecules did not migrate after reaching the substrate, which markedly impeded the aggregation.
As a general trend, PLEDs with a blend structure have higher working voltage [17], and in this study the turn-on voltage (V on , at the minimum luminance of 0.1 cd/m 2 ) of the PLEDs with a spin coating layer (see Table 1) confirmed this trend. However, the V on of the PLEDs with a blend layer prepared by VS method was about the same as that of PLEDs composed of pure structures. This low V on is an advantage of the VS method for PLED applications. Unfortunately, the highest measured luminance (L max ) of the PLEDs with a blend layer was lower than the L max of PLEDs with pure structures. In the blend structure case, the efficiency was even lower than that of the PLEDs with a pure P3HT active layer.
To improve the luminance of PLEDs, the blended layer was exchanged for a graded doped layer where the structure consisted of a P3HT matrix in which the concentration of Alq 3 was increased linearly from 0 to 20%. Thus, while there is no Alq 3 in the polymer near the interface with the PEDOT:PSS layer, the concentration of Alq 3 increases linearly in the polymer until about 20% in the region near the interface with the cathode. We have reported the structure of these graded layers prepared by VS method analyzed by scanning transmission electron microscopy (STEM) [14,15]. Although in the present study the polymer film was thinner, the graded structure is expected to be similar to that reported previously.
The graded structure device has a turn-on voltage of 9 V, which is about the same as that of the pure and blend structure devices (Tab. 1), however, it emits much brighter light than the pure and graded structure devices (Fig. 2b). As the heterojunction interfaces are virtually eliminated in the graded structure, this increase in brightness was probably because there is a better balance of charge carrier in such a way that there is a higher recombination probability of holes and electrons [5], leading to a large increase in PLED performance. In the present study, the optimized highest ratio of Alq 3 to P3HT was about 1:4, but suitable ratios have still to be determined for other combinations of electron and hole transport materials.
All PLEDs showed very similar EL spectra to that of the graded structured PLED shown in the insert of Fig. 2b. In addition, all PLEDs emitted yellow-orange light and the preparation method and structure of active layers did not influence the color. Thus, only P3HT emit light, while Alq 3 acted only as an electron transport material without emitting light.
In summary, this report described the successful fabrication of a PLED in which the polymer active layer had graded doping with electron transport molecules by the VS  Since its recently discovered quasi-one-dimensional nanostructures, Cu(TCNQ) has attracted renewed attention because of its large surface to volume ratio and size effects which exhibit perfect crystal structure and opens up prospects for high-density nanoelectronics devices [7,8]. Recently, some researches on the switching properties and prototype device fabrication of the Cu(TCNQ) nanostructures have been reported.
Liu [9] and Xiao [10] reported the directed integration of  nanowire (as shown in Fig. 4A). The sample for measurement here was 400 nm in diameter and the spacing between the two electrodes was 2.1 µm (Fig. 1B). When the voltage applied to the device was increased to the threshold of 34 V, the current rapidly increased, indicating the sample transfer from a high-resistance to a low-resistance state. After the voltage was swept down to 0 V, the sample turned back to high-resistance.   The strategy for the synthesis of the magnetic cores embedded in HMS is designed and presented in Fig. 1 Table 1. Here, similar pore size distributions for all samples, around 2.4 nm, can be found in Table 1 eV is an attractive material due to its weak sensitivity in the visible wavelength range and high sensitivity in the DUV range (i.e., it has solar blindness).
Intrinsic diamond shows a cut-off wavelength of 225 nm beyond the flame UV emission. It is basically difficult to use only thick intrinsic diamond for flame sensing. However, the submicron thin film technology and impurity engineering provide the opportunity to apply diamond to flame sensing. By using submicron thick boron-doped homoepitaxial diamond layers on type-Ib diamond substrates and proper device structures, it is able to tune the overall properties of diamond photodetectors [3]. For example, the dark current can be extremely low due to the depletion of free holes in the submicron epilayer [4]. The threshold wavelength of the diamond detector can be tailored to be larger than 225 nm [4][5][6][7].
The photoconductivity gain can be as larger as 100 upon DUV light illumination. Thus, the diamond photosensor developed is able to detect the CO and OH emission bands in the hydrocarbon flame.
In this work, we demonstrate the solid-state flame sensor by using the high-sensitivity diamond DUV photosensor. In addition, we show the flame sensing performance of the diamond photosensor.   nm light, which indicates the photoconductivity gain property as reported previously [4,5]. The gain mechanism is explained by existence of electron trap with high capture rate and low emission rate, which provides the significant increment of the hole lifetime and the hole concentration [3,6]. This model is supported by the predominant current transport due to hole.
Since we also observed the gain property of the MSM photoconductor at an applied voltage lower than 1 V [7], the interaction between the epilayer and the Ib-substrate is believed to be responsible for the kinetic gain mechanism.   [8,9], studied the DW motion based on the s-d electron model and collective coordinates of the DW. They found that the DW motion will be dominated by the effective force on the DW, which is rooted from the reflection of the electrons incident on the DW. On the other hand, S. Zhang and Z. Li [10,11] pointed out that the out-of-plane STT was crucial for the stead current induced DW motion and proposed a non-adiabatic STT due to spinrelaxation mechanism. Another non-adiabatic STT was proposed by J. Xiao [12], which is due to the misalignment of non-equilibrium spin density and local effective field. The out-plane spin transfer torques can stem from either spin relaxation or non-local non-adiabatic effects [9]. However, the link between the microscopic STT and the effective force on the DW is missing.
In this paper, we start from the Landau-Lifshitz-Gilbert

EFFECTIVE FORCE AND OUT-OF-PLANE STT
The current-driven domain wall dynamics is studied based on the Landau-Lifshitz-Gilbert equation, the connection between effective force and out-of-plane STT on each site is obtained. The LLG equation reads, where the magnetization M can be written as the function of polar angles as indicated in Fig. 1: We assume that the magnitude of the magnetization will not change when the DW moves. We can rewrite the on sites LLG equations Eq. (1) with sphere polar angle coordinates (θ i , φ i ) and then apply the Walker ansatz so that Euler angles are treated as functions of position and time.
The DW is parallel to the z direction, θ and ϕ are the function of position and time. θ =θ(z −X(t)), ϕ =ϕ(z, t), X is defined as the wall center's position within the "rigid wall" approximation, and ϕ(z, t) is simply reduced to ϕ(t) in the model used in following derivation. The variable X and ϕ can serve as collective coordinates to describe the wall motion.
Namely, we can get the dynamic information of the domain wall if we know the time evolution of (X(t), ϕ(t)). (In the simplest picture, we may let ϕ constant so that the localized spins change their orientations within a plain). The effective field on each local magnetization sin : The above two sets of equations can be solved numerically for arbitrary STT form and effective field. However, when current and applied magnetic fields are not very large, rigid wall is a good approximation to study the DW motion. Applying Walker's ansatz [16], the localized spins in Neel domain wall structure (Shown in Fig. 2 . We can immediately come to some interesting results after comparing Eq. (8) (9) with effective forces and effective torques [9]. It turns out that the out-plane spin torque , which appears in Eq. (9) corresponds to the force on domain wall, while the in-plane spin torque contributes to the z component of torque on the wall as a whole: Although it is widely accepted that different components of STTs can be written in the form of Eq. (4), the origin of out-of-plane STT could be stem from different physics. In our calculation, we focused on the nonadiabatic term pointed out by J. Xiao and M. Stiles [12].

NUMERICAL CALCULATION AND DISCUSSION
In the following, numerical calculation is used to DOI: 10.5101/nml.v1i1.p34-39 http://www.nmletters.org investigate the torque as the function of domain wall width λ.
After introducing the calculation methods we used, we subsequently investigate the in-plane and out-plane STTs. The obtained calculations are in accordance with our model and can be helpful to the discussion about effective force.

A. Calculation of STT in Domain Wall
The structure used in the calculation is current in plane (CIP) domain walls in ferromagnetic materials. We apply free electron model in our calculations where the current is along (110) direction in fcc structure using the lattice constant of Cobalt a Co = 3.549 Å. And we set the energy split between majority and minority spins in free electron energy band as 1.69eV, which equals the value of exchange energy split of bulk Cobalt. Our numerical approach is based on the Tight-binding linear muffin-tin orbital formulism [13,14]. Scattering wave function is obtained by the wave function matching method [15]. The rigid potential approximation is employed to simulate the DW. Here STTs in our study can be defined as the difference between the incoming and outgoing spin current at R site.

B. Out-plane STT and Effective Force
Let us focus on the out-plane component of STTs. From our numerical calculations, we find a decaying oscillation for out-of-plane STT as shown in Fig. 4. From the inset in can see that c j obtained from our model are position dependent rather than constant. This is reasonable because the non-adiabatic torques is nonlocal and has vibrating behaviors.
Combining the calculated results, we will discuss the DW velocities expressed in Eq. (11). When the rigid wall still holds, it is suitable to express the terminal velocities v(t→∞)=−c j /α.
Considering c j is not constant but oscillating, we may use the average to express the terminal velocity. When the wall is thicker, the increase of inertia or wall mass makes it harder to move, thus the velocities get smaller and consequently larger driving currents are required. In our model, the velocities will indeed drop significantly when the out plane torques (c j terms) vanish in thicker walls, as depicted in Fig. 6. However, we need to keep in mind that the Walker's model we use to describe rigid walls will break down under high external field [16,17] or high current density [9]. In that case, the domain wall would be oscillating so the definition of DW velocities is no longer available.
The forces at different wall structures are calculated from integration within the narrowed region of c j (z) using ) ( ) (  [4][5][6]. In last decade the fast photo-induced charge transfer from a conjugated polymer to fullerene showed to be very efficient and this system considerably speed up polymer based solar cell development [7,8]. Nowadays best polymer based solar cells having a electrical conversion efficiency in the range of 3-8% [9][10][11][12].
However, still vigorous period of research is needed to refine the structure processing and to develop the solar cell fabrication techniques.
One-way to improve the efficiency is blending of polymer with a second inorganic nano-composite. On the other hand, it is accepted in photovoltaic community that to achieve a considerable improvement in solar cell efficiency a new physics as quantum-confined structures have to be applied [13][14][15]. One of such approach involves so called "hybrid" photovoltaic structures, consisting of blending of organic polymers with inorganic semiconductor nanocrystals (quantum dots), possessing of quantum confinement size effect (e.g. CdS, CdSe, and CdTe) [16]. Better yet, similar to fullerene, the nanocrystal can also serves as an acceptor for electrons [17] or potentially as a donor when it is doped.
Silicon nanocrystals (Si-ncs) are one potential low dimensional nanocomposite that fulfills all criteria for the next generation of solar cell development [13]. It has been demonstrated that it can be also efficiently used for bulk-heterojunction formation when blended with conjugated polymers [17,18]. Natural micro/nano phase-separation between the polymer and the solid state Si-nc get the morphology/structure to obtain photogenerated carrier separation [17,18]. In addition, it is expected that Si-ncs with quantum confinement size effect may provide a significant boost in carrier generation efficiency from a phenomenon so called carrier multiplication [19][20][21][22]. The principle of the phenomenon is sketched in Fig. 1 responds to an absorbed photon by producing multiple electron-hole pairs. So far in many types of nanocrystals the phenomenon has been widely reported [20][21][22]. It has to be stressed higher multiplication rates were reported for low band gap nanocrystals (e.g. CdSe). However, since the Si-ncs possess non-toxicity for the environment, abundance and established silicon based photovoltaic technologies Si-ncs even lower carrier multiplication rate could be more favorable material for solar cell production process at lower cost.
Processing  Si-ncs bright room temperature PL feature is used to witness them. When the laser beam is shined through the solution red PL is observed ( Fig. 2(b)). This proofs that we deal with the Si-ncs with quite good quality. Low concentration defects (surface or bulk) do not hinder emission significantly. Further structural analysis [25] witnessed the presence Si-ncs with diamond-like structure with average size about 3 nm.
Corresponding PL spectra is shown in inset of the Fig. 2(b). The PL spectrum is rather broad with maximum centered at 2 eV. As reported elsewhere the PL in Si-ncs with grain size of less than 5 nm originate mostly from quantum confinement effects [25][26][27]. Similar arguments could be used to describe luminescence properties of our sample. Therefore one can expected that a significant spectrum broadening could be ascribed to large nanocrystal size distribution.
Recently it has been demonstrated that such a way prepared Si-ncs blended with conjugated polymer form a bulk-heterojunction [17,28]. As the freestanding Si-ncs concentration within the polymer can be easily tuned, the photoconductivity response of the blend at different nanocrystalites concentration has been optimized. It has been observed that the photo-and dark-conductivity ratio varied as a function of Si-ncs concentration (Fig. 3). As the Si-ncs concentration is increased also the ratio between photo and dark increases and reaching maximum at around 40 wt.%. After reaching the peak the ratio again decreased. It has to be noted that the photoexcitation in silicon nanocrystals is fundamentally different from those in polymers or organics nanoparticles.

Particular, whereas light absorption in Si-ncs (quantum dot)
results to the direct generation of mobile charged carriers, however, the higher (more than 10 times) light absorption in polymer leads to the generation of excitons [17]. The exciton in polymer is associated with binding energy in excess of 0.2 eV.
Therefore the exciton and the charged carrier transport in higher absorbing polymer [28] is the heart of the photocarrier generation of the blend. As a result an excitionic energy transfer might be an efficient mechanism to pump Si-ncs through the excitation of conjugated polymers [17]. We stress that the exciton is charge-neutral and is transported due to the diffusion process only. Contrary to that a charged carrier is transported by both diffusion and/or drift in the built-in electric fields. In our case field results from difference in work function of P3HT polymer and ionization potential of the nanocrystal that dissociates photogenerated excitons [17]. When the concentration is further augmented, most likely the P3HT lamella-like structure become more perturbed, which results in a hole mobility decrease [29]. Those perturbations then limit carrier transport and photoconductivity response of the blend at higher Si-ncs concentrations as 40 wt%.   Since V OC is relatively high the generated photocurrent is weak.
This is due to the multiple factors that have to be taken into account. In order to improve the photocurrent generation the oxidation and formation of oxide shell has to be avoided during  This work was also partially supported by a NEDO project.
scanning electron microscope, and the optical characteristics were analyzed by photoluminescence spectra at roomtemperature. The photoluminescence mechanism for both pure ZnO nanocrystals and Mn doped ZnO are discussed as well. ZnO nanocrystals were fabricated by a simple vapor phase transport (VPT) process, as shown in Fig. 1, where the system is consisted of a large horizontal quartz tube furnace, a vacuum system, a gas meter, and a temperature controller. Pure zinc acetate dihydrate as the source material was placed at the bottom of a one-endsealed slender quartz tube that was positioned at the center position of the furnace. For uniform growth, cleaned Si (111) and quartz wafers as the substrates were placed at end of the large tube; in front of the slender tube but far away from it. Prior to the fabrication, the furnace was pumped at vacuum and heated to 100℃ and kept for 2 hours to remove the water moisture in the zinc acetate dehydrate. Then the furnace was heated to the growing temperature of 500℃. At the same time, a mixture gas of O 2 /Ar was loaded through the furnace and kept a constant flowing during the growing process. The growing process was carried out for 30 min, and the samples were taken out when the system cooled down to room temperature.
For Mn doped ZnO Ncs, the synthesis process is identical with above except that the source material is the mixture of zinc acetate dihydrate and manganese acetate tetrahydrate with different ratios. After the mixture decomposing at a high temperature of 500 o C, the mixed vapor of Zn/Mn was drawn out of the small tube by the flowing gas and condensed on the substrate as well. When the mixed Zn/Mn liquid droplets deposited and crystallized, ZnO Ncs were uniformly doped with Mn.
The morphology of the products was investigated by a Shimadzu SS-550 super scanning electron microscope (SEM). And the composition was analyzed by X-ray diffraction (XRD) on a RINT2000 vertical goniometer with Cu Κа radiation (λ=0.1541 nm). The photoluminescence (PL) measurements were performed by using a 310 nm excitation source at room temperature. Figure 2(a) illustrates the surface morphology of ZnO Ncs on Si substrate investigated by SEM. Clearly, ZnO nanoparticles with an average size of 18 nm are uniformly dispersed on Si substrate, and a few of them aggregate together, which because new small ZnO nanoparticles were nucleated continually in the growth process. At the same time, small ZnO Ncs has large surface effect than that of the bulk material, leading to the nanocrystals readily get together. Figure 2(b) is a SEM picture taken from the quartz sample, where the morphology of ZnO Ncs is similar to that of Fig. 2(a) except that the density of ZnO Ncs is lower, indicating that ZnO Ncs prefer nucleate on Si substrate to quartz. The X-ray diffraction pattern of the ZnO Ncs on Si is shown in the upper panel in Fig. 3, while the standard diffraction spectrum of the bulk wurtzite ZnO is in the down panel. We can see that the diffraction peaks and the interplane spacing are well matched to the standard diffraction pattern of wurtzite ZnO, demonstrating that our products have a distinct formation of wurtzite ZnO nanosrystals. The average diameter of ZnO DQs is 19 nm estimated from Scherrer's equation. The size is very accordance with the surface SEM image. These results indicate that the ZnO Ncs obtained in our experiment are of high quality with standard crystal shape and acceptable size.   shows an absorption spectrum of the pure ZnO Ncs on Si substrate at room temperature. It can be seen that the absorption intensity decreases sharply as the wavelength is over 355 nm, which can be defined as the absorption edge, corresponding the absorption of the intrinsic bandgap of ZnO. From the value of the absorption edge, we can estimate the bandgap of energy of ZnO Ncs to be about 3.54 eV. Figure 4(b) shows the photoluminescence (PL) spectrum of ZnO Ncs measured at room temperature. In the PL spectrum, ZnO Ncs exhibit a strong and predominant UV emission peaked at 377 nm, originating from the band to band emission of ZnO Ncs. And a much lower blue-green peak is positioned at 435 nm that is attributed to the surface defect from oxygen vacancies or zinc interstitials. For comparison, the PL spectrum of bulk ZnO is also given in a red line, where the date is magnified of 100 times. The UV and the blue emissions is absent in the bulk ZnO at room temperature, while a weak large band peaked at 500 nm (Kelly color) is appeared, attributing to the emission from the surface energy levels. Therefore, the predominant strong UV emission in ZnO Ncs results from the quantum-confined band-edge emission and the quantum size effects. Additionally, In the conventional growth methods, in order to obtain a prominent UV emission, the as grown ZnO Ncs usually require some accessional treatments, such as surface modification and annealing. In our experiment, ZnO Ncs sample achieves a very strong UV emission through a simple VPT process without any additional treatment in the experiment, showing that the as grown ZnO Ncs are of high quality. Figure 5 shows the absorption spectrum of the Mn doped ZnO Ncs sample measured at room temperature. There are two absorption bands, one is located at 208 nm, and the other is a large band with a sharp absorption edge at 357 nm. The former corresponds to the absorption of MnO that has a larger band gap (4.2 eV), which makes the first absorption edge taking a blue shift. The latter is a combined absorption of ZnO:Mn Ncs, pure ZnO Ncs, and the surface defect states. Figure 6 shows the photoluminescence spectra of ZnO:Mn Ncs sample. For comparison, the PL spectrum of the pure ZnO Ncs is shown in panel (a). Figure 6(b) shows the PL spectrum of ZnO: Mn Ncs when the material ratio of Zn/Mn at 95:5. We can see that the PL property has taken a significant change by doping Mn impurity. The UV emission from the band edge of ZnO is almost quenching, replaced by two emissions with almost identical intensity in the spectrum, one is located at 409 nm, and the other is placed at 435 nm. The former can be considered as the UV emission shifting to longer wavelength of 408 nm because of Mn ions introducing an impurity in the energy bandgap of ZnO. The latter originates from Mn relevant compounds, such as Mn 3 O 4 etc. With the ratio of Zn/Mn increase to 95:10, the blue emission of 435 nm becomes stronger, while the UV emission decreases much a lot along with shifting from 408 to 390 nm, that is, the major UV emission shifts to short wavelength again.
From the PL spectra, we know that Mn impurity is triumphantly doped into ZnO Ncs. The doping way of Mn impurity in ZnO Ncs either is an interstitial doping or a substitutional doping. If the substitutional doping occurs, Mn ion will substitute Zn ion of ZnO to form MnO. MnO has a larger band gap (4.2 eV) that may lead to the UV emission taking a blue shift in the PL spectra. However, in our experiment the UV emission shifts to longer wavelength after doping Mn, hence we deduce that the doping is more an interstitial doping than a substitutional doping. In addition, the UV emission shifts to longer wavelength initially, then back to short wavelength again. The similar phenomena also happened in other's researches on Mn doped ZnO [13,14], CdS and ZnSe. According to Bylsma's [15] second-order perturbation theory on the similar phenomena in Mn doped ZnSe bulk materials, we can give a explanation. If the doping concentration of Mn is low enough (<2%), the d orbit of Mn has strong exchange interactions with the s and p orbits of ZnO, which can be considered as a short range disorder spin system. The first interaction decreases the energy of conduction band bottom, and the second one increases the energy of valence band top, so that the band gap of the product become narrower than before, resulting in the red-shift of UV emission. With the concentration of Mn increasing (>2%), the band gap is broadening, and the emission moves back to the UV field again.
Moreover, the relative intensity ratio of the UV and the blue (435 nm) emissions changes with the ratio of Zn/Mn in the source material. With the concentration of Mn increasing, the intensity of the blue emission enhanced much a lot, which comes from surface defect levels associated with oxygen vacancies, Mn 3 O 4 etc. The enhanced blue emission is mainly attributed to the increasing of Mn impurities on the surface of ZnO Ncs. Because the decomposing temperature of Mn acetate is higher than that of Zn acetate, it decomposes later than Zn acetate in the growth process. Therefore, more Mn impurities are covered on ZnO Ncs than inter-doped them, on which Mn further combines with oxygen to form Mn 3 O 4 , a more stable structure, resulting in more oxygen vacancies and enhancing the green emission.
ZnO Ncs have been grown and doped with Mn by a simple technology of a vapor phase transport (VPT) process. The as grown ZnO Ncs in size of 19 nm have a distinct of wurtzite structure. Without any additional treatment, ZnO Ncs sample exhibit a strong and predominant UV emission in the PL spectra at room temperature. For Mn doped ZnO Ncs, a UV emission and a blue emission are observed, and the position of the UV emission shifts to longer wavelength direction because of Mn introduced an impurity level in the bandgap. With the concentration of Mn increasing, the blue emission enhanced much a lot due to the strong exchange interaction in the short range spin system and the excess Mn 3 O 4 on the surface, respectively. Therefore, it is imperative to develop simple and template-free methods for fabrication 1D magnetic materials with expected advantages such as low cost, friendly environment, high purity and large-scale production prospection.
Preparation of nanostructures in the aqueous phase may be a preferred approach due to its significant advantages: nonflammable and nontoxic, environment friendly, safety and feasible for a large scale production. However, some limitation of the aqueous approach also should be considered. For example, the as-synthesized products obtained at relatively low temperature are often amorphous, and additional treatment protocols are required to achieve good crystalline of the resulting materials. Additionally, the as-synthesized nanometer-sized Ni particles tend to be oxidized in aqueous solutions.
In this study, we present an approach for fabrication Ni nanowires with good crystalline in an aqueous phase solution under normal pressure in absence of any inorganic or organic templates. The influences of magnetic field strength, concentration of Ni ions and pH values on the formation and morphology of Ni nanowires were investigated. The magnetic properties of Ni nanostructures with different morphology were also evaluated.
All chemicals were of analytical reagent grade and used without any further purification. In a typical process, nickel chloride (NiCl 2 ·6H 2 O) was disolved into the 100 mL mixture of deionized water and ethanol with a volume ratio of 5:4 to form a transparent absinthe-green solution. An appropriate amount of hydrazine hydrate (N 2 H 4 ·H 2 O, 85 wt%) solution as a reducing agent was added until the absinthe-green sedimentation turned to light blueviolet colour. NaOH solution (5 mol/L) was used to adjust the pH value of the mixture. The solution was then placed in a 60°C water bath and the bath was fixed in a magnetic field. The static magnetic field with interval of 30 cm and magnetic pole area of 1200 cm 2 was generated by direct current flow, and the strength can be controlled from 0.005 T to 0.5 T by adjusting the current intensity. After about 30 minutes, loose black floccules was formed and floated on the solution surface. The products were filtered and washed repeatedly with distilled water and ethanol by using magnetic field, and then dried at 60°C for 12 hours.
The size and morphology analyses were performed using field emission scanning electron microscope (SEM, Ultrazeiss, Zeiss). The phase structure was characterized by X-ray polycrystalline diffractometer (XRD, D8 Advance, Bruker) using Cu Ka radation with graphite monochromator. The hysteresis properties were measured on a vibration sample magnetometer (VSM, Lake Shore 7400). Figure 1 shows the SEM images of resulting products with Ni + concentration of 0.01 M under different applied magnetic fields. It can be seen that the magnetic field has strong influences on the morphology of products. In the absence of the magnetic field, only some bulky particles were formed and no 1D material was observed (see Fig. 1a), whereas, under a low magnetic field of 0.05 T, some short, thick and flexural wires were observed (see Fig. 1b). When the strength of the magnetic field was increased to 0.2 T, the longer and thinner nanowires can be formed and no separated particles were observed. The wires have aspect ratio of about 1000 with average diameter of 200 nm and length of 200 μm (see Fig. 1c). There is little change observed on morphology of nanowires when magnetic fields increased from 0.2 T to 0.5 T. The influence of concentration of Ni ions on morphology of Ni nanowires has also been observed. Figure 2 shows the SEM images of products with different Ni ion concentration under the same magnetic field of 0.2 T. It can be seen that the dimeter of the nanowires increased with increasing Ni + concentration. It is about 200 nm for concentration of 0.01 M, and about 600 nm for 0.05 M. When the concentration increased to 0.1 M, there is no wires observed, whereas, the clusters with diameter from 200 to 600 nm were produced.
The reaction time has strong influence on the morphology of the products. There are no products observed less than 10 min. Some loose black floccules were formed and floated on the solution surface after 10 min later and for about 30 min the solution became transparent, clarifying and colorless. All the products were floated on the solution surface. However, as the reaction time elongated to 2 h, the nanowires will join together to form silvery-white flakes. For about 3 h later, the grey-colored nanowires disappeared gradually, and bright silvery-white films were formed on the beaker wall. The products dependence of reaction time was displayed in Fig. 3.
The XRD pattern of the Ni nanowires prepared at 0.01 M concentration under 0.2 T magnetic field is shown Fig. 4. Three diffraction peaks corresponding to different crystallization directions of [111], [200] and [220], respectively, indicate the nanowire has a cubic crystal structure which is in agreement well with the Nickel standard card (no. 65-0380). There is no other impurity observed, suggesting the prepared Ni nanowires have high purity. Figure 5 shows the hysteresis loops of the Ni nanowires which prepared at magnetic field of 0.2 T. The magnetic property of commercial Ni powder which has particle size of about 1 um was also investiged. The nanowires have coercivity of 260 Oe which is much larger than 31 Oe of the powder. However, the saturation magnetization of the Ni nanowires is 41 emu/g, which is much less than the 55.5 emu/g of the powder sample. The large coercivity of Ni nanowire attributes to the shape anisotropy due to its large aspect ratio. This phenomenon has been observed on other Ni nanowires which prepared by templated-assited methods. The reduced magnetization of Ni nanowires may result from the surface effect of Ni nanowire structures, in which lots of Ni atoms at the surface are not aligned along its magnetic anisotropy direction in order to reduce the the surface energy. Therefore, the magnetic moments of these Ni atoms cannot be aligned along the magnetic field due to the strong exchange interaction.
The possible mechanism for the formation of Ni nanowires in alkali solution under magnetic assistant may be understood as following: nickel ions were firstly reducted by strong reduction agent of hydrazine hydrate and turned to tiny spherical particles. Then the magnetic Ni particles aligned along the magnetic field direction under the magnetic driving force. The nickel nanowires retained their linear structure after kept in ultrasonic bath for 10 minutes, which proved that the nanowires showed a good mechanical strength. It was found that the pH value of the reaction solution is sensitive to the reduction process of nickel ion by hydrazine hydrate. When the pH value was higher than 13.70, the nickel ion can be completely reduced within 30 min, whereas, when the pH value was adjusted to 13.00, it will need 2 hours to finish the reaction. If the pH value was lower than 13.00, the reaction would not take place and there was no any Ni nanowires or particles produced. This is attributed to the different oxidation reduction potentials φ 0 of the reaction for different pH values. For a low pH condition, the reaction takes place as follows and the reduction potential can be calculated: N 2 H 5 + -4e=5H + +N 2 , φ 0 = -0.23 V (1) Ni 2+ +2e=Ni, φ 0 = -0.25 V (2) whereas, for a high pH condition, they will be: Ni(OH) 2 +2e=Ni, φ 0 = -0.72 V (3) N 2 H 4 +4OH --4e=N 2 +4H 2 O, φ 0 = -1.15 V (4) The oxidation reduction potentials at low pH condition are higher than those for higher pH values. This will lead to the reaction much easier for higher pH conditions.
In this method all the reaction agents are common and low-cost. The reaction process is simple and fast. It provides a new strategy to prepare other magnetic nanowires. In addition, the amount of products can be up to 100 g/day in our laboratory. So it is also a feasible and potential approach for large scale synthesis of Ni nanowires. Although the diameter of the nanowires can be tuned by changing magnetic field and the concentration of nickel ion, the diameter less than 200 nm and more smooth surface nanowires with single crystal structure which synthesized in an ambient aqueous condition is still a challenge. It may be obtained under a much higher magnetic field or to find more suitable experimental conditions. The reaction mechanism should be understood much more. The more detailed research is ongoing.
A simple, low-cost, environment-friendly approach of preparation magnetic Ni nanowires was developed. In this method, the nanowires were fabricated in an ambient aqueous solution at normal pressure by assistant of magnetic field without any inorganic or organic templates or any other surfactants. The prompt wires have mean diameter of 200 nm and length up to 200 μm. It was found that the diameter of wires increases with increasing strength of magnetic field, and there is no wires formed in absence of magnetic field. The pH value and the concentration of nickel ion has also strong influences on formation and morphology of nanowires. The wires formed only below 0.1 M of Ni concentration and pH value should be up to 13 at 60°C. This method provides a new approach to fabricate magnetic nanowires in an ambient condition and may be the most promising candidate to produce large-scale magnetic nanowires.