JP5238189B2 - Method for producing stretch-strained germanium thin film, stretch-strained germanium thin film, and multilayer structure - Google Patents

Description

  The present invention relates to a method for producing a stretched strain germanium film, a stretched strain germanium film, and a multilayer structure that can be suitably used in the fabrication of a semiconductor device, particularly a field effect transistor having a strained germanium channel.

  The performance improvement of the silicon ultra-large scale integrated circuit (Si-ULSI), which forms the foundation of today's advanced information processing industry, is the Si-MOSFET (Metal-Oxide-Field Conductor Effect Transistor), which is its basic element. ) Has been supported by improved performance and higher integration. Further, the performance of each element has been improved by proportional reduction (scaling) of the transistor structure. The scaling rule is to improve the current driving capability by reducing the channel length or reducing the thickness of the gate insulating film while keeping the electric field applied to the gate insulating film constant. However, improvement in device performance according to this scaling law has become difficult due to the short channel effect and the increase in tunnel leakage current accompanying the reduction in the thickness of the gate insulating film.

Under such circumstances, in order to achieve higher speed and higher functionality of the device, attention has been focused on a technology that does not depend on element miniaturization. One of them is a technique using a high mobility non-Si material for the channel portion of the MOSFET. Among these, germanium (Ge), which is the same group IV element as Si, has attracted attention. The reason is that Ge has a higher mobility than Si in both electron and hole mobility. In recent years, reports on MOSFETs having Ge as a channel have increased (for example, Non-Patent Documents 1, 2, and 3). Regarding Ge-pMOSFETs, high-current drive Ge-pMOSFETs (hafnium oxide films (HfO ₂ ) used as gate insulating films) with gate lengths reduced to 125 nm from IMEC at International Electron Devices Meeting (IEDM) in 2006 (For example, Non-Patent Document 4), the hole mobility of a long channel device is about three times that of bulk Si-pMOSFET (HfO ₂ is used for the gate insulating film), and the gate length. It has been reported that it has a similar mobility at 125 nm. On the other hand, many reports have been made on Ge-nMOSFET (for example, Non-Patent Documents 5 and 6). It has been reported that it has twice the mobility (Non-patent Document 7).

  On the other hand, attention is also paid to a technique for improving carrier mobility by applying a strain to a crystal in a channel to change a band structure. Various studies have been conducted on strained Si channels in which in-plane biaxial strain is added to Si, and electron and hole mobility have been verified theoretically and experimentally (for example, Non-patent documents 8 to 12). In 2003 IEDM, IBM is an electron transfer in a strained (elongated strain ε˜1.45%) Si-MOSFET fabricated on a silicon-on-insulator (SOI) substrate compared to a bulk Si-MOSFET. It has been reported that a performance improvement of about 2 times in terms of degree and about 1.4 times in terms of hole mobility has been obtained (Non-patent Document 13).

  On the other hand, Fischetti et al. Pointed out by theoretical calculation that even when strain is applied to Ge, improvement in carrier mobility can be expected by changing the band structure, as in Si (Non-patent Document 14). When strain is applied in the Ge (001) plane, the electron mobility rapidly decreases with a strain amount of 0.986 with respect to compressive strain. In addition, the electron mobility with respect to the tensile strain increases slowly when the strain amount is 1.013 or less, and rapidly increases when the strain amount is 1.013 or more. On the other hand, hole mobility increases with compression and extension strain. The rate of increase is greater for compressive strain.

Experimentally, a conventional Si channel p-type MOSFET is a Ge channel p-type MOSFET having compressive strain (compression strain ε˜1.20%) in a plane formed on a strain relaxation Si _0.3 Ge _0.7 buffer layer. There is a report that the hole mobility has increased by about 20 times (Non-patent Document 15).

  However, there are almost no reports on the electron and hole mobility of stretch-strained Ge-nMOSFETs. One reason for this is that a method for forming a Ge layer having a large tensile strain has not been established.

  As described above, a Ge channel device having an extension strain in the (001) plane is a very attractive device because high carrier mobility for both electrons and holes can be expected. Realization of stretch-strained Ge channels is a very important challenge towards a new generation of ultrafast devices that replace Si and strained Si devices.

<Conventional stretch-strained germanium film forming method>
If the target is not limited to an electronic device, there are several reports regarding the formation of a Ge layer having an in-plane tensile strain. A group of Kimerling et al., Massachusetts Institute of Technology, USA has reported the growth of a Ge layer to which an extension strain of about 0.20% is applied for the purpose of optical device application (Non-Patent Documents 16 to 19). They developed a seed-Ge layer on a Si (001) substrate at 350 ° C. by 30 nm by UHV-CVD (Ultra High Vacuum Chemical Deposition) method at 350 ° C., then increased the growth temperature to 600 ° C. and grown a Ge layer by 1 μm. Let Thereafter, in order to reduce threading dislocations in the Ge layer, cyclic heat treatment at 700 ° C. to 900 ° C. is performed 10 times to return to room temperature. Since Ge has a larger thermal expansion coefficient than Si, the Ge layer is subjected to tensile strain when the temperature is lowered to room temperature. The amount of tensile strain applied to the Ge layer can be controlled by the difference between the temperature lowering start temperature and room temperature. From the evaluation by the X-ray diffraction method, the control range is from 0.15% to a maximum of 0.21%. .

They have also reported that elongation strain can be applied to the Ge layer up to 0.25% by forming C54-TiSi ₂ on the back surface of the Si substrate (Non-patent Document 20).

The result of Kimerling et al. Is very useful in realizing a group IV semiconductor optical device using a stretched strained Ge layer on a Si substrate. However, when considering application to ultrafast devices, considering the theoretical calculation of Fischetti et al., The amount of strain is 1.02 times the mobility of bulk Ge for electrons and 1. It can be said that it is still small enough to make it 05 times. In order to obtain this amount of strain, a thick Ge layer of 1 μm and a thick C54-TiSi ₂ layer of 3.8 μm are required on the back surface of the Si substrate, which is unrealistic from the viewpoint of the MOSFET fabrication process. There is also a problem in terms of cost. In order to form a very thin Ge layer having a large tensile strain, a strain relaxation buffer layer having a lattice constant larger than that of bulk Ge is required. Furthermore, in order to apply a desired strain to the tensile strained Ge layer, it is very important to control the lattice constant of the buffer layer serving as a base.

<Conventional Germanium Tin Layer Formation Method>
Therefore, we focused on Ge _1-x Sn _x layer to which tin (Sn) having a larger lattice constant than Ge was added. Since Sn is the same group IV semiconductor as Si and Ge, and the α-Sn phase has the same diamond structure as Si and Ge, formation of a Ge _1-x Sn _x layer in which Sn is dissolved in Ge is formed. Thus, it is possible to realize a buffer layer having a lattice constant larger than that of Ge and for applying an in-plane stretching strain to Ge. In particular, the present invention has developed a technique for heteroepitaxial growth of a high-quality Ge _1-x Sn _x layer on a Si substrate and realization of in-plane stretch strain Ge thereon.

However, the following difficulties are expected for heteroepitaxial growth of a high-quality Ge _1-x Sn _x layer on a Si substrate. The lattice constant of the α-Sn phase is 0.64892 nm, and the difference from the lattice constant of Ge, 0.56579 nm, is very large at 14.7%. Sn has a melting point of 232 ° C., and a phase transition from the α-Sn phase to the β-Sn phase occurs at 13.2 ° C. Furthermore, in the thermal equilibrium state, the solid solubility of Sn in Ge is about 1.0 at. % (The amount of tensile strain that can be applied to the Ge layer is 0.15%) and is very low. Therefore, when an attempt is made to grow a strain-relaxed Ge _1-x Sn _x layer having a high Sn composition on the Si substrate, Sn precipitates in the Ge _1-x Sn _x layer without being replaced by crystal lattice positions. It is easily predicted that problems such as inability to maintain surface and interface flatness due to the low melting point of the Ge _1-x Sn _x layer will occur. That is, it is considered that there are many technical problems in the epitaxial growth of the high Sn composition strain relaxation Ge _1-x Sn _x layer on the Si substrate.

Here, reports on Ge _1-x Sn _x so far will be described. From the latter half of the 1980s, application of Ge _1-x Sn _x layers on Si (001) and Ge (001) substrates to optical devices began to be studied, and various growth methods enable high Sn composition (0.10 <x < 0.40) single crystal Ge _1-x Sn _x has been tried (for example, Non-Patent Documents 21 to 23). However, in any growth method, polycrystallization, amorphization, and β-Sn phase precipitation are observed during the growth of the Ge _1-x Sn _x layer.

Since late 1990, a group of Kouvekis et al. At Arizona State University has reported the growth of Ge _1-x Sn _x layers on Si (001) substrates by UHV-CVD (Non-patent Documents 24 and 25). Their group is characterized by the use of (Ph) SnD ₃ and SnD ₄ gas as the CVD material for Sn. SnH ₄ has a low bond energy of Sn—H and is very unstable at room temperature, whereas (Ph) SnD ₃ and SnD ₄ gases using deuterium have Sn—D bonds as Sn—H bonds. In comparison, since it is thermally stable (Non-patent Document 26), it is not decomposed at room temperature, but is decomposed at 250 ° C. to 350 ° C. near the growth temperature. A Ge _1-x Sn _x layer on a Si (001) substrate grown by using these source gases is observed by Ge _1-x Sn _x / Si (001) by TEM (Transmission Electron Microscopy) observation. It has been confirmed that edge dislocations are formed at the substrate interface. In addition, in the sample having a Sn composition of 2.0% and a film thickness of 500 nm, many stacking faults and threading dislocations are observed, but there is no precipitation of Sn. However, the detailed dislocation structure and strain relaxation mechanism have not been elucidated, and no detailed discussion has been made on samples having a high Sn composition.

They also reported that in 2007, the strain-strained Ge was realized using the Ge _1-x Sn _x layer directly formed on the Si substrate as the strain applying layer. However, the Sn composition in the Ge _1-x Sn _x layer in this report is as low as 2.5%. Therefore, in their method, a high-quality strain relaxation Ge _1-x Sn _x layer having a sufficiently large lattice constant cannot be formed, and the strain amount of the strained Ge layer formed thereon is 0.25%. There was a problem of being small (Non-patent Document 27).

The group of Greene et al. At the University of Illinois has reported the growth of a strained Ge _1-x Sn _x layer on a Ge (001) substrate by a low temperature MBE method (Non-patent Documents 28 and 29). They report that the film thickness at which the strained Ge _1-x Sn _x layer can be grown epitaxially changes for a certain growth temperature and Sn composition. However, they do not mention the formation of a strained Ge layer on a Ge _1-x Sn _x layer.
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<Substrate technology suitable for germanium tin film formation>
In order to form a material having a lattice constant much larger than that of Si, such as a strained Ge _1-x Sn _x layer, as a buffer layer having high crystallinity, it is necessary to make a crystal as a buffer layer as much as possible. Therefore, it is considered important to use a substrate having a small lattice constant difference. From this point of view, it is promising to use a Ge substrate instead of the Si substrate. However, among the substances present on the earth surface, Si is the second most abundant after oxygen, and the ratio thereof is about 28%, whereas Ge is very little as 1.5 ppm. Furthermore, Ge has a high density and is easily broken, so that it needs to be handled more carefully as a substrate than Si that has been widely used in the past. Further, Ge has a lower thermal conductivity than Si, and thus is not suitable as a substrate material for electronic devices. For the above reasons, it is predicted that using a Ge substrate instead of a conventional Si substrate in the LSI industry will be a very big problem from the viewpoint of product cost and process transport system. Therefore, there is a demand for a technique for manufacturing an ultrahigh-speed device using Ge on a Si substrate that is familiar to LSI manufacturing. In the future, in order to realize ultrahigh-speed devices using high mobility materials such as tensile strained Ge based on mature Si technology, the lattice constant of the underlying buffer layer will be precisely controlled and high quality will be achieved. Formation of a heteroepitaxial film is the biggest problem.

In addition, the technique for realizing such a heteroepitaxial layer having a large lattice constant on the Si substrate is attractive from the viewpoint of growing other III-V group compound semiconductor materials. The group III-V compound semiconductor has an extremely high electron mobility as compared with many other semiconductors. In particular, in the case of InSb, the electron mobility is about 80000 cm ² V ⁻¹ sec ^{−1 in} a bulk material. It has special characteristics. Therefore, in recent years, attention has been increasing as a channel material for next-generation MOSFETs like Ge. However, their lattice constant is considerably larger than that of Si, and the lattice mismatch strain with Si reaches 4.2%, which is about the same as Ge in GaAs, 11.6% in InAs, and 19.3% in InSb. . For the formation of crystallographically high-quality III-V compound semiconductor layers, in order to overcome the enormous lattice mismatch between the compound semiconductor material and Si, a relatively large lattice constant is formed on the Si substrate. Forming a heteroepitaxial layer is considered an effective means.

Conventional reports have been made by using a Ge layer heteroepitaxially grown on a Si substrate (which is called a “virtual Ge substrate”) and forming a Ge _1-x Sn _x layer and Ge layer on the Ge layer. It was possible to form an extension strained Ge layer having a large strain amount exceeding. Specifically, Si (001) after forming the epitaxial Ge layer that strain relaxation on the substrate, the upper layer by forming an epitaxial Ge _1-x Sn _x layer that strain relaxation, further Ge _1-x Sn _x layer on the upper layer An epitaxial strained Ge layer having the same in-plane lattice constant was formed. The use of the virtual Ge substrate promotes effective strain relaxation due to threading dislocations in the Ge layer, so that a Ge _1-x Sn _x layer having a large Sn composition and a high strain relaxation rate is formed. Became possible.

By adopting a virtual Ge substrate, strain relaxation of the Ge _1-x Sn _x layer was promoted, and a Ge _1-x Sn _x layer having a larger in-plane extension strain could be formed. As a result, the present inventors succeeded in forming an extension strain Ge having an in-plane lattice constant of 0.4015 nm in the (110) direction and a strain amount of 0.35%, which has a strain amount larger than that of any conventional report.

Here, a result of an attempt to grow a strain relaxation Ge _1-x Sn _x layer on a virtual Ge substrate as an underlying substrate is shown. The virtual Ge substrate is a Ge thin film formed on a Si substrate and having a very flat surface and 100% strain relaxation. The advantages of using a virtual Ge substrate are listed below. As described above, in the LSI industry, the use of Ge as a bulk substrate causes significant problems in terms of cost and handling. On the other hand, the virtual Ge substrate is a virtual substrate obtained by heteroepitaxially growing a thin Ge layer on the Si substrate, which is the core of the existing LSI industry. It is thought that it becomes advantageous to.

Below, the substrate cleaning method, the virtual Ge substrate fabrication method, the Ge _1-x Sn _x layer growth on the virtual Ge substrate, and the strain relaxation Ge _1-x Sn _x layer dislocation structure and strain relaxation mechanism are described in detail. The results of the investigation are shown.

  1 and 2 are views for explaining a method for producing a Ge / GeSn / Ge / Si structure according to the present invention.

First, a manufacturing method of a strained Ge / strain-relieved Ge _1-x Sn _x layer multilayer structure on a virtual Ge substrate will be described with reference to FIG.

<Production of virtual Ge substrate>
A Si substrate 11 made of a Si (001) substrate, which is a Si single crystal substrate, was prepared as a single crystal substrate. After this Si substrate is chemically cleaned, it is introduced into an ultra-high vacuum apparatus (not shown), and the substrate is used in a vacuum of 1 × 10 ⁻⁶ Pa or less using a general vapor deposition apparatus such as a Knudsen cell (K cell). A Ge layer having a thickness of 40 nm was epitaxially grown while maintaining the temperature at 200 ° C. Thereafter, the sample was taken out into the air and subjected to a heat treatment at 700 ° C. for 1 minute in a nitrogen atmosphere, thereby relaxing the strain in the Ge layer and forming the strain-relieved Ge layer 12. The sample at this stage is called a virtual Ge substrate.

<Preparation of strain relaxation GeSn layer>
After the virtual Ge substrate was chemically cleaned, it was introduced into the ultrahigh vacuum apparatus, and the surface was cleaned by heat treatment at 450 ° C. for 30 minutes in a vacuum of 1 × 10 ⁻⁶ Pa or less. Thereafter, the substrate temperature was kept at 200 ° C., and Ge and Sn were simultaneously deposited by a K cell and an arc plasma gun (APG), respectively, thereby epitaxially growing a Ge _1-x Sn _x layer. At this time, the Sn composition x was, for example, 8.0%. Thereafter, the samples were removed into the atmosphere, 600 ° C. in a nitrogen atmosphere, formed by heat treatment of 10 minutes, to relieve the strain of Ge _1-x Sn _x layer, the strain relaxation Ge _1-x Sn _x layer 13 did.

<Preparation of stretch-strained Ge layer>
After the chemical cleaning of the sample, it was introduced into the ultra-high vacuum apparatus, and the surface was cleaned by heat treatment at 450 ° C. for 30 minutes in a vacuum of 1 × 10 ⁻⁶ Pa or less. Thereafter, the substrate temperature was kept at 300 ° C., and Ge was vapor-deposited by a K cell to epitaxially grow the Ge layer. As a result, the extension strained Ge layer 14 was formed, and the multilayer structure 10 was produced.

Next, a method for producing a multilayer structure composed of a strained Ge / Sn composition gradient multi-stage Ge _1-x Sn _x layer on a virtual Ge substrate will be described with reference to FIG.

<Production of virtual Ge substrate>
A Si substrate 21 made of a Si (001) substrate, which is a Si single crystal substrate, was prepared as a single crystal substrate. After this Si substrate is chemically cleaned, it is introduced into an ultra-high vacuum apparatus (not shown), and the substrate is used in a vacuum of 1 × 10 ⁻⁶ Pa or less using a general vapor deposition apparatus such as a Knudsen cell (K cell). A Ge layer having a thickness of 40 nm was epitaxially grown while maintaining the temperature at 200 ° C. Thereafter, the sample was taken out into the air and subjected to a heat treatment at 700 ° C. for 1 minute in a nitrogen atmosphere to relieve the strain in the Ge layer, thereby forming the strain-relieved Ge layer 22 as described above. The sample at this stage is called a virtual Ge substrate.

<Preparation of Sn composition gradient multi-stage Ge _1-x Sn _x layer>
After the virtual Ge substrate was chemically cleaned, it was introduced into the ultrahigh vacuum apparatus, and the surface was cleaned by heat treatment at 450 ° C. for 30 minutes in a vacuum of 1 × 10 ⁻⁶ Pa or less. Thereafter, the substrate temperature is kept at 200 ° C., and Ge and Sn are simultaneously deposited in the K cell and the APG, respectively, so that a Ge _1-x Sn _x layer of the first stage has a film thickness of 60 nm at a substrate temperature of 200 ° C. A Ge _0.990 Sn _0.010 layer was formed. Thereafter, the samples were removed into the atmosphere, 600 ° C. in a nitrogen atmosphere, formed by heat treatment of 10 minutes, to relieve the strain of Ge _1-x Sn _x layer, the strain relaxation Ge _1-x Sn _x layer 23 did.

Thereafter, the sample was taken out into the air and subjected to a post heat treatment (PDA: Post Deposition Annealing) at 700 ° C. for 10 minutes in a nitrogen atmosphere. The sample was again chemically cleaned and then introduced into the ultra-high vacuum apparatus, and the surface was cleaned by heat treatment at 450 ° C. for 30 minutes in a vacuum of 1 × 10 ⁻⁶ Pa or less. As the second stage Ge _1-x Sn _x layer, a substrate temperature was maintained at 200 ° C., and a Ge _0.970 Sn _0.030 layer having a thickness of 70 nm was grown. Thereafter, the sample is taken out into the air and subjected to a heat treatment at 700 ° C. for 10 minutes in a nitrogen atmosphere, thereby relaxing the strain in the Ge _1-x Sn _x layer, and the second stage strain relaxation Ge _1-x Sn _x. Layer 24 was formed.

Thereafter, the sample was taken out into the air and subjected to PDA treatment at 700 ° C. for 10 minutes in a nitrogen atmosphere. The sample was again chemically cleaned and then introduced into the ultra-high vacuum apparatus, and the surface was cleaned by heat treatment at 450 ° C. for 30 minutes in a vacuum of 1 × 10 ⁻⁶ Pa or less. A Ge _0.933 Sn _0.067 layer having a thickness of 50 nm was grown as a third stage Ge _1-x Sn _x layer at a substrate temperature of 200 ° C. Thereafter, the sample is taken out into the air and subjected to a heat treatment in a nitrogen atmosphere at 700 ° C. for 10 minutes to relax the strain in the Ge _1-x Sn _x layer, and the third-stage strain relaxation Ge _1-x Sn _x. Layer 25 was formed.

<Preparation of stretch-strained Ge layer>
After the chemical cleaning of the sample, it was introduced into the ultra-high vacuum apparatus, and the surface was cleaned by heat treatment at 450 ° C. for 30 minutes in a vacuum of 1 × 10 ⁻⁶ Pa or less. Thereafter, the Ge temperature was epitaxially grown by keeping the substrate temperature at 250 ° C. and depositing Ge by K cell. As a result, the extension strained Ge layer 26 was formed, and the multilayer structure 20 was produced.

Here, the Sn composition graded multi-stage Ge _1-x Sn _x layer has three stages, but in order to improve the Sn composition and to generate appropriate strain relaxation, the graded multi-stage Ge ₁ -stage Ge ₁ or more is provided. It is of course possible to form a _-x Sn _x layer.

<How to clean the sample>
In order to epitaxially grow a desired crystal layer on Si, Ge and virtual Ge substrates, it is necessary to clean the substrate surface in advance. The substrate cleaning method performed in this example will be described in detail below.

First, a method for cleaning the substrate will be described. The Si substrate was washed by boiling in an ammonia solution (NH ₄ OH: H ₂ O ₂ : H ₂ O = 1: 6: 20) heated to 40 ° C. for 10 minutes, and then overflowed in ultrapure water for 10 minutes. Dry the substrate by blowing. Thereafter, the substrate is introduced into the sample exchange chamber of the ultra-high vacuum apparatus and further transferred to the growth chamber in the ultra-high vacuum state. In the growth chamber, gas such as moisture contained in the substrate and the substrate holder is released by pre-annealing at 600 ° C. for 10 minutes in advance. The temperature is lowered to room temperature, and an amorphous Si layer of about 1 nm is deposited on the substrate surface. Thereafter, the substrate is heated at 850 ° C. for 15 minutes. This creates a clean surface. At that time, a 2 × 1 surface reconstruction structure peculiar to a Si (001) clean surface was confirmed by a reflection high energy electron diffraction (RHEED) method.

Next, a method for cleaning Ge and virtual Ge substrates will be described. After overflowing in ultrapure water for 10 minutes, washing is performed for 5 minutes in an ammonia solution (NH ₄ OH: H ₂ O = 1: 4) at room temperature. This etches the native oxide film on the Ge and virtual Ge substrates and removes heavy metals present on the surface. After overflowing in ultrapure water for 10 minutes, washing is performed for 2 minutes in a sulfuric acid solution (H ₂ SO ₄ : H ₂ O = 1: 7) at room temperature. This removes contamination such as metal and carbon on the surface. Next, after overflowing in ultrapure water for 10 minutes, the Ge substrate is immersed in an H ₂ O ₂ solution for 30 seconds and the virtual Ge substrate for 5 seconds at room temperature. Thereby, a protective oxide film is formed on the Ge and virtual Ge substrates. Immediately thereafter, the substrate is dried by nitrogen blowing. In the case of a virtual Ge substrate, care must be taken because the Ge layer is entirely oxidized by the H ₂ O ₂ solution unless this step is performed quickly. Thereafter, it is introduced into the sample exchange chamber and transferred to the growth chamber. In the growth chamber, heat treatment was performed at 450 ° C. for 30 minutes in an ultrahigh vacuum. This creates a clean surface. At that time, a 2 × 1 surface reconstruction structure peculiar to the Ge (001) surface was confirmed by RHEED. Note that the above-described cleaning method is an example, and a method of forming a clean surface by a general cleaning method such as another chemical cleaning method or a cleaning method by physical etching may be used.

(Comparative example)
<Ge _1-x Sn _x layer formed on Si and Ge substrates>
The Si or Ge substrate was cleaned according to the method described above to clean the surface. After the surface cleaning, a Si buffer layer having a film thickness of 50 nm is deposited on the bulk Si (001) substrate by electron beam deposition at a substrate temperature of 600 ° C. and a deposition rate of 0.1 Å / s, and on the bulk Ge (001) substrate A Ge buffer layer having a thickness of 100 nm was grown by a Knudsen cell (cell temperature 1200 ° C.) at a substrate temperature of 400 ° C. at a deposition rate of 0.28 Å / s. After surface cleaning and buffer layer growth, a 2 × 1 surface reconstructed structure was confirmed by RHEED. A Ge _1-x Sn _x layer was epitaxially grown on each substrate at a substrate temperature of 200 ° C. The reason why the substrate temperature is as low as 200 ° C. is to suppress the precipitation of β-Sn during the growth of the Ge _1-x Sn _x layer. A Knudsen cell and an arc plasma gun were used for the deposition of Ge and Sn, respectively. After the growth of the Ge _1-x Sn _x layer, some samples were taken out into the atmosphere and subjected to PDA at 600 ° C. for 10 minutes in a nitrogen atmosphere.

<Growth and Crystal Structure Evaluation of Ge _1-x Sn _x Layer on Si (001) Substrate>
First, the crystallinity and dislocation structure of the Ge _0.98 Sn _0.02 layer on the bulk Si (001) substrate will be described. Hereinafter, the Sn composition x when described as a Ge _1-x Sn _x layer is a value evaluated by Rutherford Backscattering Spectrometry (RBS). FIG. 3A is a cross-sectional TEM dark field image (diffracted wave vector g ₀₀₄ ) of a sample obtained by growing a Ge _0.98 Sn _0.02 layer by 240 nm on a bulk Si (001) substrate at a substrate temperature of 200 ° C. Many threading dislocations were observed in the Ge _0.98 Sn _0.02 layer, and the threading dislocation density was estimated to be 10 ¹² cm ⁻² or more. Further, as shown by the squares, it can be seen that there is a region where almost no diffraction contrast is seen in the vicinity of the surface. FIG. 3B and FIG. 3C show a high-resolution TEM image and a limited field diffraction pattern of a portion surrounded by a square in FIG. It can be seen that an amorphous Ge _1-x Sn _x phase is formed on the epitaxially grown Ge _1-x Sn _x layer. In addition, a halo pattern is also observed in the limited field diffraction pattern, and the formation of an amorphous Ge _1-x Sn _x phase can be confirmed.

In the growth by the low temperature MBE method, the film thickness has a transition from epitaxial growth to amorphous growth in the growth process regardless of homo and hetero epitaxial growth (Non-patent Document 31). Gurdal et al. Report that when Ge _1-x Sn _x is grown on a Ge (001) substrate at a growth temperature of 100 ° C., the reduction of h _epi due to strain becomes significant when the Sn composition is 9% or more (non-conversion). Patent Document 33). In our experiment, it is considered that the Ge _0.98 Sn _0.02 layer transitioned from epitaxial growth to amorphous growth during growth because of low temperature growth and a large lattice mismatch of 4.5%. In this case, the stacking fault could not be confirmed in FIG. 3A, but it is considered that a large number of threading dislocation groups cause amorphization.

FIG. 4 shows the results of X-ray diffraction two-dimensional reciprocal lattice mapping (XRD-2DRSM) measurement of this sample in the vicinity of the Ge _1-x Sn _x (224) peak. The diffraction peak of the Ge _1-x Sn _x layer was clearly observed, and the Sn composition was estimated to be about 1.0% from the Vegard rule. This is about half of the Sn composition estimated from RBS, and about half of the originally introduced Sn atoms are not replaced by the lattice positions of the Ge _1-x Sn _x crystal, and are included in the amorphous phase. It is suggested that.

In FIG. 4, a very broad Ge _1-x Sn _x diffraction profile along the [110] direction was observed. The in-plane crystal domain size L can be expressed by the following formula from the full width half maximum (FWHM) Δθ (rad) in the [110] direction of this profile.
L = λ / 2Δθsinθ _B (1)
Here, λ is the wavelength of the X-ray, and θ _B is the Bragg angle in the (004) plane reflection of the substrate. From the formula (1), the domain size of the crystal was estimated to be about 30 nm. This minute crystal domain formation is presumed to be due to threading dislocations in the film. These results show that the direct growth of Ge _1-x Sn _x layer on bulk Si (001) substrate is difficult even with Sn composition as low as 2%.

<Growth of Ge _1-x Sn _x layer on Ge (001) substrate and evaluation of dislocation structure>
In this section, we discuss the growth and crystallinity of Ge _1-x Sn _x layers on bulk Ge (001) substrates grown under three conditions. FIG. 5 shows a sample structure of the Ge _1-x Sn _x layer. Samples were prepared under the growth conditions of (0.03, 40 nm), (0.025, 220 nm), and (0.081, 30 nm) (Sn composition, Ge _1-x Sn _x layer thickness). FIG. _{6A shows} a cross-sectional TEM Akeno image (diffracted wave vector g ₀₀₄ ) when a Ge _0.977 Sn _0.023 layer is grown on a bulk Ge (001) substrate by 40 nm. Pseudomorphic Ge _0.977 Sn _0.023 layer is grown without forming dislocations. Moreover, the amorphous phase as observed on the Si substrate was not observed. FIG. 6B shows a cross-sectional TEM Akeno image (diffracted wave vector g ₀₀₄ ) of the Ge _0.977 Sn _0.023 layer after the sample was subjected to PDA treatment at 600 ° C. for 10 minutes. No dislocation was observed even after PDA treatment.

Next, FIG. 7A shows a cross-sectional TEM Akeno image (diffracted wave vector g ₀₀₄ ) immediately after a Ge _0.975 Sn _0.025 layer is grown on a bulk Ge (001) substrate by 220 nm. A Ge _0.975 Sn _0.025 layer is grown on _{pseudomorphic} . The reason why dislocations are not introduced even when the film thickness is increased is that the film grows at a low temperature, so that the film is in a non-equilibrium state and grows while containing strain. The reason why the amorphous phase is not observed in the Ge _0.975 Sn _0.025 / Ge system compared to the Ge _0.98 Sn _0.02 / Si system is that the growth temperature is the same as that at 200 ° C. It can be said that the cause is large. Further, since there are no threading dislocations, it is considered that amorphization near the surface induced by defects does not occur.

7B shows a cross-sectional dark field image (diffracted wave vector g ₀₀₄ ) of a Ge _0.975 Sn _0.025 layer obtained by _subjecting the sample to 600 (PDA treatment at C for 10 minutes). Misfit dislocation is Ge _0.975 Sn _0.025 / bulk Ge (001) introduced at the substrate interface, which can be observed with diffraction wave vectors g ₀₀₄ and g ₂₂₀ , indicating that these dislocations are 60 ^° dislocations, and some of them are indicated by arrows. Propagation towards the bulk Ge (001) substrate was also observed.

FIG. 8 shows a planar TEM image of the sample shown in FIG. Two 60 ^° dislocations propagating in two <110> directions were observed at the Ge _0.975 Sn _0.025 / bulk Ge (001) substrate interface. Dislocations are introduced from the film surface in a half loop shape, and after reaching the substrate interface, misfit components remain at the interface, and penetrating components are considered to have escaped from the film. The threading dislocation density was estimated to be 10 ⁶ cm ^-2 or less. Further, as indicated by the circles, a structure that appears to correspond to the dislocation form in which a part of dislocations propagates toward the substrate in FIG. 7B was also observed in the planar TEM image. An enlarged image of the circled portion in FIG. 8 is shown in FIG. It can be confirmed that dislocation lines of 60 ^o dislocation do not intersect.

Such a dislocation form is thought to be formed as a result of the reaction that occurs when the dislocation lines of two intersecting 60 ^° dislocations have the same size and parallel Burgers vectors. At the point where each dislocation crosses, a dislocation splitting reaction (annihilation reaction) occurs, and a new dislocation line is created. This splitting reaction is a reaction in which two dislocations with parallel Burgers vectors avoid an interaction between dislocations and split up and down at the intersection to reduce the length of the dislocation line, resulting in an energetically stable dislocation structure. (Non-Patent Document 34).

Next, FIG. 10A shows a cross-sectional TEM Akeno image of a sample obtained by growing a Ge _0.919 Sn _0.081 layer on a bulk Ge (001) substrate to a thickness of 30 nm. In spite of the high Sn composition of 8.1%, Sn does not precipitate and a Pseudomorphic Ge _0.919 Sn _0.081 layer grows. FIG. 10B shows a bright cross-sectional image of the Ge _0.919 Sn _0.081 layer obtained by _{subjecting the} same sample to PDA treatment at 600 ° C. for 10 minutes. A large number of dot-like precipitates can be observed in the vicinity of the Ge _0.919 Sn _0.081 / bulk Ge (001) substrate interface. A lattice image and a diffraction pattern obtained by observing the precipitate with a high-resolution TEM are shown in FIGS. Moire fringes are observed from the lattice image. From the fact that the moire fringe spacing is 2.70 nm and the diffraction pattern, this moire fringe is formed by the interference of the Ge 111 and β-Sn 200 planes, and the Ge _0.919 Sn _0.081 / bulk Ge (001) substrate interface It can be seen that a β-Sn phase is precipitated in the vicinity.

These results indicate the following: In the case of a low Sn composition, dislocations are not introduced into the Ge _1-x Sn _x layer even when PDA treatment is performed below a certain film thickness, and after a PDA treatment, Ge _1-x Sn _x is introduced after PDA treatment. Misfit dislocations are introduced at the / Ge substrate interface. However, even when the film thickness is more than a certain level, in the case of a high Sn composition, it is considered that after the PDA treatment, the β-Sn phase precipitates near the interface and the strain in the film is relaxed. It was found that a Ge _1-x Sn _x layer can be grown pseudomorphically on a bulk Ge substrate up to an Sn composition of 8.1% by a low temperature MBE method at a substrate temperature of 200 ° C.

<Measurement result of strain relaxation rate of Ge _1-x Sn _x layer on Ge (001) substrate>
FIG. 12 shows the X-ray diffraction two-dimensional reciprocal lattice mapping (XRD-2DRSM) measurement results in the vicinity of the GeSn (224) diffraction peak immediately after the growth of the Ge _0.975 Sn _0.025 layer on the bulk Ge (001) substrate. Since a GeSn (224) diffraction peak is observed immediately below the Ge (224) diffraction peak, it can be seen that a Ge _0.975 Sn _0.025 layer has grown on _{pseudomorphic} as observed by the cross-sectional TEM.

Next, the strain relaxation rate in the [110] direction of the Ge _0.977 Sn _0.023 layer, Ge _0.975 Sn _0.025 layer and Ge _0.919 Sn _0.081 layer on the bulk Ge (001) substrate from the XRD-2DRSM near the GeSn (224) diffraction peak is determined. evaluated. FIG. 13 shows the reciprocal lattice points of the Ge _1-x Sn _x (224) diffraction peak obtained from XRD-2DRSM. The dotted line in the figure is an ideal line when the crystal is considered as an isotropic elastic body and deformed according to the Poisson's ratio. FIG. 14 shows the Sn composition in the Ge _1-x Sn _x layer immediately after the growth and after the PDA treatment obtained from the strain relaxation rate in the [110] direction and the VegarD rule. Regarding the strain relaxation rate in the [110] direction, the strain relaxation rate in the [110] direction does not change before and after the PDA treatment in a sample having a low Sn composition and a critical film thickness or less. In the sample having a low Sn composition and a critical film thickness or more, a change is observed in the strain relaxation rate in the [110] direction before and after the PDA treatment, but it is a very small value (1% to 5%). This is thought to be due to the limited number of misfit dislocations that occur and proliferate. Further, in the sample having a low Sn composition, there is no change in the Sn composition before and after the PDA treatment, which indicates that there is no precipitation of β-Sn phase in the Ge _1-x Sn _x layer. However, in the sample having a high Sn composition, the Sn composition was changed from 8.1% (immediately after growth) to 2.4% (after PDA treatment). This is considered to be because the β-Sn phase was precipitated at the interface in relaxing the strain in the film as observed in the TEM image.

<Production of virtual Ge substrate>
From here, a method for producing a virtual Ge substrate will be described. After the Si (001) substrate was cleaned by the above-described method, a Ge layer was grown to 40 nm by molecular beam epitaxy using a Knudsen cell while keeping the substrate temperature at 200 ° C. in the ultrahigh vacuum apparatus. Thereafter, the sample was heat-treated at 700 ° C. for 10 minutes in the same vacuum apparatus. Some samples were taken out into the atmosphere as they were after the Ge layer was grown, and heat treatment at 700 ° C. for 1 minute was performed in a nitrogen atmosphere.

15 and 16 show that the substrate temperature is 200 on the Si (001) substrate. A TEM image of a sample obtained by growing a Ge layer at 40 ° C. at 40 ° C. and then performing heat treatment is shown. FIGS. 15A and 15B are a cross-sectional dark field TEM image (diffracted wave vector g ₂₂₀ ) and a planar TEM image when heat treatment is performed at 700 ° C. for 10 minutes in an ultrahigh vacuum apparatus, respectively. is there. 16 (a) and 16 (b) are cross-sectional dark field TEM images (diffracted wave vector g) when a Ge layer is grown and taken out to the atmosphere and subjected to rapid heat treatment at 700 ° C. for 1 minute in a nitrogen atmosphere. ₂₂₀ ) and a planar TEM image.

  In both samples, dislocations periodically arranged at the Ge / Si interface are observed from the cross-sectional TEM image. Furthermore, a dislocation network structure extending in two <110> directions can be clearly observed from a planar TEM image. This corresponds to a dislocation at the Ge / Si interface observed in the cross-sectional TEM image, and the dislocation interval was estimated to be about 10 nm. The theoretically calculated dislocation interval is 9.5 nm, assuming that the lattice mismatch between Si and Ge is 4.2%, and that the strain is completely relaxed using the Burgers vector of edge dislocations. . Therefore, it can be estimated that the strain relaxation of Ge / Si is predominantly performed by edge dislocations. The strain relaxation rate in the [110] direction and [001] direction of the Ge layer after the heat treatment estimated by the XRD-2DRSM measurement is 100% and 100% in both the ultra-high vacuum heat treatment and the rapid heat treatment in a nitrogen atmosphere. there were. Considering the observation result of the TEM image, it was confirmed that a Ge layer having a Ge lattice constant in both the [110] and [001] directions was formed on the bulk Si (001) substrate.

When the two samples are compared with respect to crystallinity, the threading dislocation densities estimated from the planar TEM images of FIGS. 15B and 16B are 10 ⁸ to 10 ⁹ cm ⁻² and 10 ¹⁰ to 10 ^11. cm ^-2 . A sample heat-treated in an ultra-high vacuum has a long heat treatment time of 10 minutes, and therefore has fewer threading dislocations than a sample heat-treated for 1 minute in a nitrogen atmosphere. However, from the cross-sectional TEM image, FIG. 15A shows a roughened surface after the ultra-high vacuum heat treatment, whereas FIG. 16A shows a flat surface after the rapid heat treatment in a nitrogen atmosphere.

  FIGS. 17A and 17B show atomic force microscopic (AFM) images of the sample surface after the ultra-high vacuum heat treatment and after the rapid heat treatment in a nitrogen atmosphere. The mean square roughness (RMS) of the surface shows a large value of 9.0 nm in the case of the ultra-vacuum heat treatment, but a very small value of about 1.0 nm in the case of the rapid heat treatment in a nitrogen atmosphere. . This result shows that roughing due to migration of Ge surface atoms can be suppressed by performing rapid thermal processing in an atmospheric pressure nitrogen atmosphere. Based on the above results, it can be seen that the method of rapid heat treatment in a nitrogen atmosphere is suitable for the production of the strain relaxation Ge layer toward the virtual Ge substrate. In the following examples, a virtual Ge substrate manufactured using a method of relaxing strain by a rapid heat treatment method was used.

<Growth and dislocation structure evaluation of Ge _1-x Sn x layer to the virtual Ge substrate>
The growth of the Ge _1-x Sn _x layer on the virtual Ge substrate will be described below. The virtual Ge substrate was cleaned according to the method described above to clean the surface. After the surface cleaning, a 2 × 1 surface reconstructed structure was confirmed by RHEED. A Ge _1-x Sn _x layer was epitaxially grown on the virtual Ge substrate at a substrate temperature of 170 to 200 ° C. A Knudsen cell was used for the vapor deposition of Ge, and an arc plasma gun was used for the vapor deposition of Sn. After the growth of the Ge _1-x Sn _x layer, some samples were taken out into the atmosphere and subjected to PDA treatment at 400 to 600 ° C. for 1 to 10 minutes in a nitrogen atmosphere. The structure of the manufactured sample is shown in FIG. Further, FIG. 19 shows the growth conditions of the manufactured sample.

<Dislocation structure and surface morphology of low Sn composition (3.0% or less) Ge _1-x Sn _x layer>
FIG. 20A to FIG. 20H show RHEED patterns in the process of growing a Ge _0.99 Sn _0.01 layer at 1300 nm at a growth temperature of 170 ° C. on a temporary Ge-conceived Ge substrate. It can be seen that the RHEED pattern changes from streak to spot as the Ge _0.99 Sn _0.01 layer starts to grow. Further, at the time of the film thickness of 400 nm shown in FIG. 20 (e), a polycrystalline ring pattern began to be observed as indicated by an arrow. This indicates a transition from single crystal growth to polycrystalline growth. As the growth continued further, the polycrystalline ring pattern was more clearly observed.

FIG. 21A shows a cross-sectional TEM dark field image (diffracted wave vector g ₀₀₄ ) immediately after the growth of the sample shown in FIG. _21H , and FIG. Indicates a pattern. A ring pattern due to the polycrystalline growth was also observed from the limited field diffraction pattern. An amorphous phase having a uniform contrast observed even in the sample on the Si substrate as shown by a white arrow in the cross-sectional TEM image of FIG. It was also observed that stacking faults exist as indicated by black arrows around threading dislocations. However, the result of RHEED observation shows epitaxial growth at a film thickness of 200 nm, and there is a difference between the observation results of TEM and RHEED. This suggests that the transition from the epitaxial layer to the polycrystalline and amorphous phases that occurred in the vicinity of the surface of the Ge _0.99 Sn _0.01 layer during the film thickness growth of 400 nm extends to the deep part of the Ge _0.99 Sn _0.01 layer through defects. In any case, it can be seen that the Ge _1-x Sn _x / virtual Ge substrate system also has a critical film thickness that allows epitaxial growth in the same manner as on the Si substrate. After this, experiments were conducted with a thin film thickness so that polycrystalline and amorphous growth did not occur.

22 (a) and 22 (b), a cross-sectional TEM of a sample subjected to PDA treatment at 600 ° C. for 10 minutes immediately after growing a Ge _0.99 Sn _0.01 layer on a virtual Ge substrate by 40 nm at a growth temperature of 200 ° C. A dark field image (diffracted wave vector g ₀₀₄ ) is shown. FIG. 23A and FIG. _{23B show} a cross-sectional TEM dark of a sample subjected to PDA treatment for 10 minutes at 600 ° C. immediately after the Ge _0.973 Sn _0.027 layer was grown to 40 nm on the virtual Ge substrate at a growth temperature of 200 ° C. Each field image (diffracted wave vector g ₀₀₄ ) is shown. 24 (a) and 24 (b) show cross-sectional TEM dark field images of samples subjected to PDA treatment at 600 ° C. for 10 minutes immediately after a Ge _0.978 Sn _0.022 layer was grown to 210 nm on a virtual Ge substrate ( The diffraction wave vectors g ₀₀₄ ) are shown respectively. From FIG. 22A, FIG. 23A, and FIG. 24A, in the Ge _1-x Sn _x layer immediately after growth, an epitaxial layer is formed up to the surface, and is observed in the sample on the Si (001) bulk substrate. It was observed that the amorphous phase as formed was not formed. Although a number of threading dislocations in the Ge _1-x Sn _x layer is observed, these, Ge ₁ threading dislocations present in the virtual Ge substrate through a Ge _1-x Sn _x / virtual Ge substrate interface _-x Sn _x is inherited in the _x layer.

Figure 25 shows a cross-sectional TEM images of as-grown observed at diffraction vector g ₂₂₀ of the same region of the same sample shown in FIG. 24 (a). Also in the observation of the diffraction vector g _220, as in the case of observing the diffraction vector g _004, threading dislocations are imaged. This result shows that the threading dislocation existing in the virtual Ge substrate and the Ge _1-x Sn _x layer has the feature of 60 ^° dislocation based on the relationship between the diffraction vector and the Burgers vector.

FIG. 26A shows a planar TEM image of the same sample as that shown in FIG. 24A immediately after the growth of the Ge _0.978 Sn _0.022 layer. _Threading dislocations extending from Ge layer to Ge _0.978 Sn _0.022 layer as indicated by arrows, with two edge dislocation networks extending linearly in the [110] direction formed at the Ge layer / bulk Si (001) substrate interface A short dislocation was observed indicating These threading dislocations were also observed in samples with an Sn composition of 1.0% and 2.7%.

In the Ge _0.99 Sn _0.01 sample after PDA at 600 ° C. for 10 minutes shown in FIG. 22B, no significant change was observed in the dislocation behavior, but FIG. 23A and FIG. The threading dislocations observed in Fig. 3 change dramatically after PDA treatment. As indicated by arrows in the figure, misfit dislocations observed at the Ge _1-x Sn _x / virtual Ge substrate interface and dislocation groups bent in the Ge _1-x Sn _x layer are observed.

FIG. 26 (b) shows a planar TEM image of a sample that was subjected to PDA treatment for 10 minutes at 600 ° C. immediately after the Ge _0.978 Sn _0.022 layer was grown to 210 nm on the virtual Ge substrate. 27 (a) and 27 (b), a Ge _0.99 Sn _0.01 layer having a thickness of 40 nm is grown on a virtual Ge substrate, and a Ge _0.973 Sn _0.027 layer having a thickness of 40 nm is grown at 600 ° C. The plane TEM image of the sample which performed the PDA process for 10 minutes is shown, respectively. From these planar TEM images after PDA treatment, it can be seen that short dislocations indicating threading dislocations observed immediately after growth are reduced. The threading dislocation density was estimated to be about 10 ¹⁰ cm ^−2, and decreased by an order of magnitude from immediately after the growth. These results indicate that PDA treatment promotes the lateral propagation of threading dislocations at the Ge _1-x Sn _x / virtual Ge substrate interface, leading to a decrease in threading dislocation density in the film.

Next, the results of observing the surface morphology of these three types of Ge _1-x Sn _x samples with AFM are shown. FIGS. _28A and _{28B show} an AFM image of a sample surface obtained by growing a Ge _0.99 Sn _0.01 layer and a Ge _0.973 Sn _0.027 layer on a virtual Ge substrate by 40 nm and _performing PDA treatment at 600 ° C. for 10 minutes. Show. FIG. _28C shows an AFM image of a sample obtained by growing a Ge _0.978 Sn _0.022 layer on a virtual Ge substrate by 210 nm and performing PDA treatment at 600 ° C. for 10 minutes. These RMS values were 1.9 nm, 0.6 nm, and 1.0 nm, respectively. In all samples, the flatness of the underlying virtual Ge substrate was maintained after the PDA treatment.

<Dislocation structure and surface morphology of high Sn composition (5.0% or more) Ge _1-x Sn _x layer>
Next, the results of TEM observation of samples having Sn compositions of 5.8% and 8.5% are shown. 29A and _29B , at a growth temperature of 200 ° C., a Ge _0.942 Sn _0.058 layer is _grown on a virtual Ge substrate to a thickness of 40 nm and a Ge _0.915 Sn _0.085 layer is grown to a thickness of 15 nm, and the temperature is 600 ° C. for 10 minutes. A cross-sectional TEM dark field image (diffracted wave vector g ₀₀₄ ) of the PDA-treated sample is shown. 30 (a) and 30 (b) show a cross-sectional TEM dark field image (diffracted wave vector g ₂₂₀ ) of the same sample as in FIGS. 29 (a) and 29 (b). The observation result of the cross-sectional TEM immediately after the growth was the same as the dislocation structure of the sample having the Sn composition of 2.5%. From FIG. 29A and FIG. 29B, threading dislocations propagate in the lateral direction at the Ge _1-x Sn _x / virtual Ge substrate interface, and a contrast of 60 ° dislocation can be observed at the interface. In addition, from FIG. 30A and FIG. 30B, spherical precipitates having a diameter of about 10 nm were observed at the Ge _1-x Sn _x / virtual Ge substrate interface. These are considered to be β-Sn phase precipitates as observed at the Ge _0.919 Sn _0.081 / bulk Ge substrate interface after PDA treatment. From the cross-sectional TEM images, the number density was estimated to be 3.0 × 10 ⁹ cm ⁻² and 2.0 × 10 ¹⁰ cm ⁻² when the Sn composition Sn was 5.8% and 8.5%, respectively. It can be said that the higher the Sn composition, the more likely Sn precipitation occurs.

Next, the result of observing the surface morphology of these two types of Ge _1-x Sn _x layers with AFM is shown. _{31A and 31B show} an AFM image of a sample obtained by growing a Ge _0.942 Sn _0.058 layer on a virtual Ge substrate by 40 nm and a Ge _0.915 Sn _0.085 layer by 15 nm and performing PDA treatment at 600 ° C. for 10 minutes. Indicates. These RMS values were 0.6 nm and 1.2 nm, respectively. In all samples, the flatness of the underlying virtual Ge substrate was maintained after the PDA treatment.

Next, the result about a sample with a high Sn composition and a large film thickness is shown. FIGS. 32 (a) and 32 (b) are cross-sections of samples subjected to PDA treatment for 10 minutes immediately after growing a Ge _0.937 Sn _0.063 layer on a virtual Ge substrate at a growth temperature of 170 ° C. and 300 nm. A TEM dark field image (diffracted wave vector g ₀₀₄ ) is shown. Lateral propagation of threading dislocations was observed after PDA treatment. On the other hand, the film surface immediately after the growth was very flat, but it was observed that the surface roughening occurred remarkably after the PDA treatment.

FIG. 33 (a) and FIG. 33 (b) _show the surface after PDA treatment of a sample obtained by growing a Ge _0.937 Sn _0.063 layer on a virtual Ge substrate by 300 nm using an AFM and a scanning electron microscope (SEM). The results observed using each are shown. 33 (a) and 33 (b), surface roughening was observed, and the RMS estimated from the AFM image of FIG. 33 (a) was a very large value of 22 nm. A precipitate having a number density of 6.4 × 10 ⁷ cm ⁻² was observed on the surface. The reason for determining the precipitate is the difference in contrast of the SEM image in FIG. These precipitates were also observed after PDA treatment at 400 ° C and 500 ° C.

FIGS. 34 (a), 34 (b), and 34 (c) _show an SEM image and a cross-sectional TEM dark field image (diffraction) after a Ge _0.937 Sn _0.063 layer on a virtual Ge substrate was subjected to PDA treatment at 500 ° C. for 10 minutes. The wave vector g ₀₀₄ ) and the limited field diffraction pattern are shown respectively. From FIG. 34 (a), it can be confirmed that the same surface morphology as that obtained after the PDA treatment at 600 ° C. for 10 minutes is obtained. In addition, precipitates were observed on the surface of the Ge _0.937 Sn _0.063 layer from the SEM image. A precipitate is observed from the cross-sectional TEM image of FIG. 34 (b), and {200}, {004} and {301} plane diffraction spots due to the β-Sn phase are obtained from the diffraction pattern of FIG. 34 (c). It was. This revealed that the precipitate on the surface was β-Sn. Regarding threading dislocations, it was observed from cross-sectional TEM observation that even in heat treatment at 400 ° C. and 500 ° C., a 60 ^° dislocation was left at the interface in the lateral direction at the Ge _1-x Sn _x / virtual Ge substrate interface.

When the Sn film is thin with a high Sn composition, a spherical β-Sn phase having a diameter of about 10 nm is deposited at the Ge _1-x Sn _x / virtual Ge substrate interface after PDA treatment, but the flatness of the surface is maintained. However, it was found that when the film thickness is thick, β-Sn phase precipitates of several hundred nm are formed after the PDA treatment, and the flatness of the Ge _1-x Sn _x layer surface is extremely deteriorated. These results show that, in the case of a high Sn composition Ge _1-x Sn _x layer, if an appropriate PDA treatment condition is selected by reducing the film thickness, only the dislocation is moved while suppressing the precipitation of Sn, thereby reducing the strain. It suggests that it can be promoted.

<Evaluation of crystallinity of Ge _1-x Sn _x layer on virtual Ge substrate>
The crystallinity of the Ge _0.978 Sn _0.022 layer on the virtual Ge substrate was evaluated from the shape of XRD-2DRSM near the GeSn (224) diffraction peak. FIGS. 35 (a) and 35 (b) show the XRD-2DRSM measurement results immediately after growth and after PDA treatment at 600 ° C. for 10 minutes, respectively. From FIG. 35A, a very broad diffraction profile extending in the [110] direction and a tail extending in the [001] direction as indicated by the arrows were obtained. Such a diffraction profile was changed to a very sharp and symmetrical diffraction profile by performing PDA as shown in FIG. From the FWHM in the [110] direction, it was found that the in-plane crystal domain size was changed from 50 nm to 73 nm before and after the PDA treatment.

FIG. 36 shows changes in domain size evaluated from the results of XRD-2DRSM for samples immediately after Ge _1-x Sn _x growth and after PDA treatment. In samples where Sn precipitation was not observed significantly, the domain size increased after PDA treatment. This indicates that threading dislocations in the Ge _1-x Sn _x layer were reduced by the PDA treatment, which is consistent with the TEM observation result. In the result of FIG. 35B, the peak due to the virtual Ge substrate disappears after the PDA process. This is because, as seen in FIG. 24F, half-loop dislocations are introduced into the Ge layer by the reaction between threading dislocations after the PDA treatment.

<Measurement result of strain relaxation rate of Ge _1-x Sn _x layer on virtual Ge substrate>
FIG. 37 shows GeSn (224) diffraction peak positions immediately after growth and after PDA treatment of all Ge _1-x Sn _x samples epitaxially grown on the virtual Ge substrate, evaluated by the XRD-2DRSM method. The dotted line in the figure is an ideal line when the crystal is considered as an isotropic elastic body and deformed according to the Poisson's ratio. It can be seen that all the samples have film strain relaxation. However, from FIG. 37, the peak position was observed on the inner side of the theoretical line where the strain was relaxed according to Poisson's ratio after PDA treatment for the samples with Sn composition of 2.7%, 5.8% and 8.5%. This is considered to be caused by precipitation of Sn.

FIG. 38 shows the strain relaxation rate in the [110] direction of the Ge _1-x Sn _x layer before and after the PDA treatment estimated from the change in the GeSn (224) diffraction peak position shown in FIG. For samples with Sn compositions of 2.7%, 5.8%, and 8.5% in which the Sn composition immediately after growth was observed on the Sn composition side lower than the theoretical line after PDA treatment, the peak position after PDA treatment The strain relaxation rate is obtained from the Sn composition.

For comparison, the [110] direction strain relaxation rate of a Ge _1-x Sn _x layer formed on a conventional bulk Ge substrate is also shown. It can be seen that when a virtual Ge substrate is used, a certain amount of strain relaxation occurs in any sample. In addition, the [110] direction strain relaxation rate of the sample on the bulk Ge substrate is very small, but in the case of the sample on the virtual Ge substrate, the [110] direction strain relaxation rate is a very high value. This result is also evident from the presence of misfit dislocations at the Ge _1-x Sn _x / Ge substrate interface seen when TEM observation is performed. Compared with the sample on the bulk Ge substrate, the heat treatment conditions are the same, but the strain relaxation rate depends decisively on the difference in the substrate. These results indicate that threading dislocations existing in the Ge _1-x Sn _x layer grown on the virtual Ge substrate effectively contribute to strain relaxation. FIG. 39 shows the (110) plane spacing after PDA processing and the amount of strain that can be applied to Ge. All of the (110) plane spacing after the PDA treatment shows a value greater than or equal to the (110) plane spacing of Ge, and has a possibility of becoming a buffer layer for growing an extension strained Ge layer.

<Growth of strained Ge layer>
Hereinafter, the results of actually growing an extension strained Ge layer on the strain relaxation Ge _1-x Sn _x layer will be shown. FIG. 40 shows a cross-sectional structure of the manufactured sample. The strain-relieved Ge _1-x Sn _x layer was cleaned in the same manner as the virtual Ge substrate, transferred again to the growth chamber, and then subjected to a cleaning heat treatment similar to that in the virtual Ge substrate. The Ge layer to be the strained Ge layer was grown at a growth temperature of 300 ° C. The designed value of the Sn composition immediately after the growth was 8.0%.

FIG. 41A shows a cross-sectional TEM dark field image (diffracted wave vector g ₂₂₀ ) of Ge (24 nm) / Ge _0.922 Sn _0.078 (24 nm) / virtual Ge substrate. _{Precipitation} of β-Sn phase was observed at the Ge _0.922 Sn _0.078 (24 nm) / virtual Ge substrate interface. FIG. 41B shows a cross-sectional TEM dark field image (diffracted wave vector g ₀₀₄ ) in the same observation region of the same sample. FIG. 41C shows a cross-sectional TEM dark field image (diffracted wave vector g ₀₀₄ ) in another region. As observed in FIG. 41 (b) and FIG. 41 (c), contrast due to 60 ^° dislocation is observed at the interface of Ge _0.922 Sn _0.078 (24 nm) / virtual Ge substrate, and strain relaxation in the film due to dislocation is observed. You can see what is happening. On the other hand, no contrast of dislocations slid laterally is observed at the Ge / Ge _0.922 Sn _0.078 interface. This suggests that the Ge layer is grown pseudomorphically with respect to the underlying Ge _1-x Sn _x layer (because the β-Sn phase is precipitated, the Sn composition is x). That is, it shows that a Ge layer having a tensile strain having a larger (110) plane spacing than that of bulk Ge is formed.

FIG. 42 shows XRD-2DRSM measurement results in the vicinity of the GeSn (224) peak and the strained Ge (224) peak of the same sample as that shown in FIG. Although it overlaps with the virtual Ge peak, a peak due to strain Ge can be observed. When the epitaxial Ge layer is grown _{pseudomorphically on the} strain relaxed Ge _0.955 Sn _0.045 buffer layer, the (110) plane spacing matches the strain relaxed germanium tin (Ge _0.955 Sn _0.045 ) buffer layer, and the (001) plane spacing is in accordance with the Poisson ratio. Get smaller. Therefore, on the two-dimensional reciprocal lattice map, it is expected that the diffraction peak of the strained Ge layer appears on the straight line according to the Poisson ratio from the virtual Ge substrate peak and directly above the strain relaxation Ge _0.955 Sn _0.045 buffer layer. . On the other hand, as described above with reference to FIG. 35 (b), the peak due to the virtual Ge substrate has a half-loop dislocation introduced into the virtual Ge substrate after the PD ₁ treatment of the Ge _1-x Sn _x layer. As a result, the peak intensity decreases. However, the Ge (224) peak after fabrication of the virtual Ge substrate and the virtual Ge substrate (224) peak position immediately after the growth of the Ge _1-x Sn _x layer are always measured at the Ge (224) peak position obtained from theoretical calculation. It was. That is, the lattice constant of the virtual Ge substrate does not change at all immediately after the Ge _1-x Sn _x layer growth. Therefore, when obtaining the strained Ge (224) peak, the evaluation was performed on the assumption that the virtual Ge (224) peak position is at the Ge (224) peak position obtained from theoretical calculation. From the results of XRD-2DRSM shown in FIG. 42, the Sn composition of the Ge _0.922 Sn _0.078 buffer layer, the peak position of the strained Ge layer, and the amount of strain were evaluated. From the GeSn (224) peak position, the Sn composition could be estimated to be 4.5%. The [110] direction strain relaxation rate was 58%.

FIG. 43 shows a diffraction profile at the position indicated by the dotted line in FIG. FIG. 43 also shows the profiles of the virtual Ge substrate and the strained Ge layer obtained by performing peak separation on the obtained diffraction profile. It was confirmed that a strained Ge layer having a tensile strain in the [110] direction from the virtual Ge substrate and the strained Ge layer peak position was formed on the strain-relaxed Ge _0.955 Sn _0.045 buffer layer.

From FIG. 42 and FIG. 43, the (110) plane spacing of the strain-relaxed Ge _0.955 Sn _0.045 buffer layer and the strained Ge layer was calculated. The (110) plane spacing values of the strain relaxed Ge _0.955 Sn _0.045 buffer layer and the strained Ge layer were 0.4016 nm and 0.4015 nm, respectively. This shows that the strained Ge layer is grown by _{pseudomorphic with} respect to the strain-relaxed Ge _0.955 Sn _0.045 buffer layer. This agrees with the result of the cross-sectional TEM observation shown in FIG. Further, the strained Ge layer has an elongation strain of 0.35% with respect to the bulk Ge.

<Formation of Sn Composition Gradient Multi-stage Ge _1-x Sn _x Layer and Strained Ge Layer>
Next, for the purpose of increasing the amount of strain applied to the stretch-strained Ge layer, a method of growing the Ge _1-x Sn _x layer in multiple stages by inclining the Sn composition was studied. At this time, in order to increase the Sn composition, it is necessary to suppress the precipitation of Sn from the Ge _1-x Sn _x layer. As shown in FIG. 37, when the Ge _1-x Sn _x layer having an Sn composition of 2.7% is heat-treated, Sn precipitation occurs and the Sn composition is reduced to 2.5%. Became clear. On the other hand, when the Sn composition is 2.5% or less, it can be said that Sn does not precipitate. This result suggests that Sn precipitation is promoted when the amount of misfit strain of the Ge _1-x Sn _x layer relative to the underlying Ge layer increases due to an increase in Sn composition. When the misfit amount with respect to the Ge layer in the Ge _1-x Sn _x layer in the case of the Sn composition of 2.5% is obtained, the lattice constant ratio is 0.37%. Accordingly, each misfit amount of the growing Ge _1-x Sn _x layer is smaller than 0.37% in terms of the lattice constant ratio with respect to the virtual Ge layer or Ge _1-y Sn _y layer as a base. The Sn composition of the Ge _1-x Sn _x layer in the stage was set.

Hereinafter, a sample preparation method will be described. FIG. 44 shows a cross-sectional structure of the manufactured sample. After the surface of the virtual Ge substrate produced by the above-described method is washed with the above-described ammonia solution, a Ge _0.990 Sn _0.010 layer having a thickness of 60 nm is grown at a substrate temperature of 200 ° C. as a Ge _1-x Sn _x layer in the first stage. did. Thereafter, the sample was taken out into the air and subjected to PDA treatment at 700 ° C. for 10 minutes in a nitrogen atmosphere. The sample was again washed with an ammonia solution and then introduced into an ultra-high vacuum apparatus, and subjected to a heat treatment at 450 ° C. for 30 minutes to clean the surface. A Ge _0.970 Sn _0.030 layer having a thickness of 70 nm was grown at a substrate temperature of 200 ° C. as the second stage Ge _1-x Sn _x layer. Thereafter, the sample was taken out into the air and subjected to PDA treatment at 700 ° C. for 10 minutes in a nitrogen atmosphere. The sample was again washed with an ammonia solution and then introduced into an ultra-high vacuum apparatus, and subjected to a heat treatment at 450 ° C. for 30 minutes to clean the surface. A Ge _0.933 Sn _0.067 layer having a thickness of 50 nm was grown as a third stage Ge _1-x Sn _x layer at a substrate temperature of 200 ° C. Thereafter, the sample was taken out into the air and subjected to PDA treatment at 700 ° C. for 10 minutes in a nitrogen atmosphere. The sample was again washed with an ammonia solution and then introduced into an ultra-high vacuum apparatus, and subjected to a heat treatment at 450 ° C. for 30 minutes to clean the surface. Finally, a 20 nm-thick Ge layer was grown at a substrate temperature of 250 ° C. as a strained Ge layer.

45 (a) and 45 (b), the prepared Sn composition graded multi-stage Ge _1-x Sn _x layer and the cross-sectional TEM image of the strained Ge layer sample, and the atomic resolution in which the strained Ge / GeSn interface is enlarged. Each TEM image is shown. Linear contrast due to misfit dislocations can be observed at the interface between the virtual Ge substrate and the Ge _1-x Sn _x layer at each stage. This indicates that the Ge _1-x Sn _x layer is relaxed by propagation of misfit dislocations. In addition, threading dislocations remaining in the Ge _1-x Sn _x layer can also be observed.

On the other hand, such misfit dislocations are not observed at the strained Ge / Ge _1-x Sn _x layer interface on the surface as shown in FIG. Even in the observation of atomic resolution as shown in FIG. 45 (b), the appearance of dislocations introduced at the interface of the strained Ge / Ge _1-x Sn _x layer indicated by the arrow is not observed, but in the [110] direction. You can observe how the atoms are aligned. This indicates that the Ge layer is grown pseudomorphically on the Ge _1-x Sn _x layer, suggesting that strain is applied to the Ge layer.

46 (a) and 46 (b) show a sample immediately after growth of the third stage Ge _1-x Sn _x layer and a sample in which a Ge layer is grown on the Sn composition-graded multi-stage Ge _1-x Sn _x layer. The XRD-2DRSM observation result in the vicinity of the Ge (224) reciprocal lattice point is shown. It can be seen from FIG. 46A that the virtual Ge substrate and each Ge _1-x Sn _x layer are formed without being mixed. Further, from the comparison between FIG. 46A and FIG. 46B, the strain relaxation of the uppermost Ge _1-x Sn _x layer progresses by the heat treatment after the third stage Ge _1-x Sn _x layer growth. You can see that In FIG. 46B, a Ge (224) diffraction peak is observed at a position different from the virtual Ge substrate, and it can be seen that a strained Ge layer is formed.

47 (a) and 47 (b), each stage of a strained Ge / Sn composition graded multi-stage Ge _1-x Sn _x sample evaluated by the XRD-2DRSM method, (224) diffraction peak positions and symbols of each layer And the strain relaxation rate of each Ge _1-x Sn _x layer at each stage evaluated and evaluated. As can be seen from FIG. 47A, the strain relaxation of the Ge _1-x Sn _x layer is promoted as the process proceeds. Further, it can be seen that the Sn composition is maintained even after strained Ge growth in the Ge _1-x Sn _x layers in the first stage and the second stage, and the precipitation of Sn is suppressed. In the Ge _1-x Sn _x layer in the third stage, the Sn composition slightly decreases even after strained Ge growth, but it can be seen that the Sn composition maintains a high value of 5.5%. It can also be seen from FIG. 47 (b) that a large strain relaxation rate exceeding 80% has been achieved after the growth of the strained Ge layer. From the position of the strained Ge (224) peak shown in FIG. 47 (a), the (110) plane spacing and strain amount of the strained Ge layer are evaluated to be 0.4028 nm and 0.68%, respectively, and are produced by this method. It can be seen that the strained Ge layer has a large amount of strain that has not existed in the past.

According to the theoretical calculation of Fischetti et al., These strained Ge layers prepared above have a carrier mobility of 1.1 times or more for electrons and 1.3 times or more for holes compared to the carrier mobility of bulk Ge. Expected to have a degree. In addition, the maximum strain amount of the strained Ge layer reported by Kimerling et al. And the Arizona State University group described above is 0.25%. It can be said that the strain amount of the strained Ge layer is a large value.
The additional notes are described below.
(Appendix 1)
A method for forming a multilayer structure for a semiconductor device, comprising: a step of forming a germanium layer 12 above a silicon substrate 11; a step of forming a germanium tin mixed crystal layer 13 thereon; and a tensile strain above the step. A method of forming a multilayer structure 10 including a step of forming a germanium layer 14.
(Appendix 2)
The method for forming a multilayer structure 10 according to appendix 1, wherein the germanium-tin mixed crystal layer 13 has a tin composition of 4.5% or more.
(Appendix 3)
3. The method for forming a multilayer structure 10 according to appendix 1 or 2, wherein the lattice constant of the (110) plane in the substrate in-plane direction of the stretched strain germanium layer 14 has a value of 0.4015 nm or more.
(Appendix 4)
The threading dislocations that continue from the germanium layer 12 to the germanium tin mixed crystal layer 13 are propagated at their interfaces to form misfit dislocations, thereby promoting strain relaxation in the germanium tin mixed crystal layer 13. The method for forming a multilayer film structure 10 according to any one of appendices 1 to 3, wherein:
(Appendix 5)
A method for forming a multilayer structure for a semiconductor device, comprising a step of forming a germanium layer 22 above a silicon substrate 21, and a multilayer comprising a plurality of germanium tin mixed crystal layers 23 to 25 having different tin compositions above the step. A method of forming a multilayer structure 20 comprising: sequentially forming a structure; and forming an extension strained germanium layer 26 thereon.
(Appendix 6)
The multilayer according to appendix 5, wherein a misfit amount of crystal lattice generated between the layers in the multilayer structure including the germanium layer 22 and the germanium tin mixed crystal layers 23 to 25 is smaller than 0.37%. A method for forming the film structure 20.
(Appendix 7)
The method for forming a multilayer structure 20 according to appendix 5 or 6, wherein the tin composition of the germanium-tin mixed crystal layer 25 is 4.5% or more.
(Appendix 8)
The multilayer film structure 20 according to any one of appendices 5 to 7, wherein the lattice constant of the (110) plane in the in-plane direction of the stretched strain germanium layer 26 has a value of 0.4015 nm or more. Forming method.
(Appendix 9)
The threading dislocations that continue from the germanium layer 22 to the germanium-tin mixed crystal layers 23 to 25 are propagated at their interfaces to form misfit dislocations, so that the strain in the germanium-tin mixed crystal layers 23 to 25 is increased. The method for forming a multilayer structure 20 according to any one of appendices 5 to 8, wherein relaxation is promoted.
(Appendix 10)
A multilayer structure for a semiconductor device, which is a germanium layer 12 formed above a silicon substrate 11, a germanium tin mixed crystal layer 13 formed above the germanium layer, and an extension strained germanium layer formed thereabove. 14. A multilayer structure 10 including 14.
(Appendix 11)
11. The multilayer film structure 10 according to appendix 10, wherein a tin composition in the germanium tin mixed crystal layer 13 is 4.5% or more.
(Appendix 12)
12. The multilayer structure 10 according to appendix 10 or 11, wherein the lattice constant of the (110) plane in the substrate in-plane direction of the stretched strain germanium layer 14 has a value of 0.4015 nm or more.
(Appendix 13)
A multilayer structure for a semiconductor device, which is composed of a germanium layer 22 formed above a silicon substrate 21 and a plurality of germanium-tin mixed crystal layers 23 to 25 having different tin compositions formed thereon. And a multilayer film structure 20 comprising an extension strained germanium layer 26 formed thereon.
(Appendix 14)
14. The multilayer according to appendix 13, wherein a misfit amount of a crystal lattice generated between the layers in the multilayer structure including the germanium layer 22 and the germanium tin mixed crystal layers 23 to 25 is smaller than 0.37%. Membrane structure 20.
(Appendix 15)
15. The multilayer film structure 20 according to appendix 13 or 14, wherein the germanium tin mixed crystal layer 25 has a tin composition of 4.5% or more.
(Appendix 16)
The multilayer film structure 20 according to any one of appendices 13 to 15, wherein the lattice constant of the (110) plane in the substrate in-plane direction of the stretched strain germanium layer 26 has a value of 0.4015 nm or more. .

The figure for demonstrating the multilayer-film structure which has the strained Ge layer of this invention. The figure for demonstrating the multilayer-film structure which has the strained Ge layer of this invention. (A) Cross-sectional TEM image immediately after growth of Ge _0.98 Sn _0.02 layer on Si (001) substrate. (B) A high-resolution TEM image near the surface. (C) Limited field diffraction pattern. XRD-2DRSM measurement result in the vicinity of the Ge _0.98 Sn _0.02 layer Ge _1-x Sn _x (224) peak grown on the Si (001) substrate. Sample structure of a Ge _1-x Sn _x layer grown on a Ge (001) substrate. Ge _0.977 Sn _0.023 layer on a bulk Ge (001) substrate (a) Sample immediately after growth, and (b) Cross-sectional TEM image after PDA treatment at 600 ° C. for 10 minutes. Ge _0.975 Sn _0.025 layer on bulk Ge (001) substrate (a) Sample immediately after growth, and (b) Cross-sectional TEM image after PDA treatment at 600 ° C. for 10 minutes. Planar TEM image of Ge _0.975 Sn _0.025 layer grown on bulk Ge (001) substrate. 9 is an enlarged image of a circled portion in FIG. (A) Immediately after growth of a Ge _0.919 Sn _0.081 layer on a bulk Ge (001) substrate, and (b) Cross-sectional TEM image after PDA treatment at 600 ° C. for 10 minutes. (A) Lattice image and (b) diffraction pattern obtained by observing the precipitate in FIG. 10 (b) with a high-resolution TEM. XRD-2DRSM measurement result in the vicinity of the GeSn (224) diffraction peak of a sample immediately after the growth of a Ge _0.975 Sn _0.025 layer on a bulk Ge (001) substrate. Summary of reciprocal lattice point positions of GeSn (224) diffraction peak evaluated by XRD-2DRSM in a Ge _1-x Sn _x layer growth sample on a bulk Ge (001) substrate. Summary of strain relaxation rate in the [110] direction obtained from the reciprocal lattice point of the GeSn (224) diffraction peak in the Ge _1-x Sn _x layer growth sample on the bulk Ge (001) substrate. (A) Cross-sectional dark field TEM image and (b) Planar TEM image of a sample heat-treated at 700 ° C. for 10 minutes in an ultra-high vacuum apparatus after growing a Ge layer on a Si (001) substrate. (A) Cross-sectional dark field TEM image and (b) Planar TEM image of a sample subjected to rapid heat treatment at 700 ° C. for 1 minute in a nitrogen atmosphere after growing a Ge layer on a Si (001) substrate. AFM images of sample surfaces grown on a Si (001) substrate, (a) after ultra-high vacuum heat treatment, and (b) rapid heat treatment in nitrogen atmosphere. Sample cross-sectional structure of a Ge _1-x Sn _x layer grown on a virtual Ge substrate. Growth conditions for a Ge _1-x Sn _x layer grown on a virtual Ge substrate. RHEED pattern in the process of growing a Ge _0.99 Sn _0.01 layer on a virtual Ge substrate to 1300 nm. (A) Si substrate clean surface, (b) Immediately after Ge layer growth, (c) Virtual Ge substrate clean surface, (D) Ge _0.99 Sn _0.01 layer 200 nm, (e) 400 nm, (f) 600 nm, (g) 800 nm, (H) After 1300 nm growth. (A) Cross-sectional TEM dark field image immediately after growth of Ge _0.99 Sn _0.01 layer 1300 nm on virtual Ge substrate, and (b) Restricted field diffraction pattern in circles in FIG. Cross-sectional TEM dark field images of samples subjected to PDA treatment immediately after (a) and (b) 600 ° C. for 10 minutes after growing a Ge _0.99 Sn _0.01 layer having a thickness of 40 nm on a virtual Ge substrate. Cross-sectional TEM dark field images of samples subjected to PDA treatment immediately after (a) and (b) 600 ° C. for 10 minutes after a 40 nm- _thick Ge _0.973 Sn _0.027 layer was grown on a virtual Ge substrate. Sectional TEM dark field images of samples subjected to PDA treatment for 10 minutes at (a) and (b) 600 ° C. after a Ge _0.978 Sn _0.022 layer having a thickness of 210 nm was grown on a virtual Ge substrate. A cross-sectional TEM image obtained by observing the same region as shown in FIG. 24A of the sample immediately after the Ge _0.978 Sn _0.022 layer was grown by 210 nm with the diffraction vector g ₂₂₀ . Plane TEM images of samples immediately after (a) and after (b) PDA treatment in which a Ge _0.978 Sn _0.022 layer having a thickness of 210 nm was grown on a virtual Ge substrate. (A) A sample obtained by growing a Ge _0.99 Sn _0.01 layer having a thickness of 40 nm on a virtual Ge substrate and then subjected to PDA treatment at 600 ° C. for 10 minutes, and (b) Ge _0.973 Sn having a thickness of 40 nm on the virtual Ge substrate. A planar TEM image of a sample subjected to PDA treatment at 600 ° C. for 10 minutes after growing a _0.027 layer. After growing a 40 nm-thick (a) Ge _0.99 Sn _0.01 layer and (b) a Ge _0.973 Sn _0.027 layer and (c) a 210 nm-thick Ge _0.978 Sn _0.022 layer on a virtual Ge substrate, AFM image of sample surface after PDA treatment A cross-sectional TEM of a sample subjected to PDA treatment at 600 ° C. for 10 minutes after growing (a) a Ge _0.942 Sn _0.058 layer _having a thickness of 40 nm and (b) a Ge _0.915 Sn _0.085 layer having a thickness of 15 nm on a virtual Ge substrate. Dark field image (diffracted wave vector g ₀₀₄ ). A cross-sectional TEM of a sample subjected to PDA treatment at 600 ° C. for 10 minutes after growing (a) a Ge _0.942 Sn _0.058 layer _having a thickness of 40 nm and (b) a Ge _0.915 Sn _0.085 layer having a thickness of 15 nm on a virtual Ge substrate. Dark field image (diffracted wave vector g ₂₂₀ ). An AFM image of a sample subjected to PDA treatment at 600 ° C. for 10 minutes after growing (a) a Ge _0.942 Sn _0.058 layer _having a thickness of 40 nm and (b) a Ge _0.915 Sn _0.085 layer having a thickness of 15 nm on a virtual Ge substrate. . A cross-sectional TEM dark field image of a sample subjected to PDA treatment at 600 ° C. for 10 minutes after a 300 nm-thick Ge _0.937 Sn _0.063 layer was grown to 300 nm on a virtual Ge substrate (a) and (b) 600 ° C. for 10 minutes. (A) AFM image and (b) SEM image of the surface of a sample obtained by growing a Ge _0.937 Sn _0.063 layer having a thickness of 300 nm on a virtual Ge substrate to 300 nm and then subjecting it to PDA treatment at 600 ° C. for 10 minutes. (A) SEM image, (b) cross-sectional TEM dark-field image, and (c) limited-field diffraction of a sample subjected to PDA treatment at 500 ° C. for 10 minutes after growing a Ge _0.937 Sn _0.063 layer having a thickness of 300 nm on a virtual Ge substrate pattern. XRD-2DRSM measurement results after (a) immediately after growth of the Ge _0.978 Sn _0.022 layer on the virtual Ge substrate and (b) after PDA treatment at 600 ° C. for 10 minutes. Domain size change immediately after growth of the Ge _1-x Sn _x layer grown on the virtual Ge substrate and after PDA treatment. GeSn (224) diffraction peak position immediately after growth and after PDA treatment of a Ge _1-x Sn _x sample epitaxially grown on a virtual Ge substrate, evaluated by XRD-2DRSM method. Virtual Ge substrate and a bulk Ge Ge _1-x Sn _x layer grown on the substrate, the distortion relaxation ratio for [110] direction of Ge _1-x Sn _x layer before and after the growth immediately and PDA process. The Ge _1-x Sn _x layer grown on the virtual Ge substrate has a (110) spacing after PDA treatment at 600 ° C. for 10 minutes and the amount of strain that can be applied to Ge. Sample structure of a strained Ge / strain relaxed Ge _1-x Sn _x layer fabricated on a virtual Ge substrate. (A) Ge (24 nm) / Ge _0.922 Sn _0.078 (24 nm) / cross-sectional TEM dark field image (diffraction wave vector g ₂₂₀ ) of virtual Ge substrate sample. (B) Cross-sectional TEM dark field image (diffracted wave vector g ₀₀₄ ) in the same observation region of the same sample. (C) Cross-sectional TEM dark field image (diffracted wave vector g ₀₀₄ ) of another region of the sample. XRD-2DRSM measurement result of Ge (24 nm) / Ge _0.922 Sn _0.078 (24 nm) / virtual Ge substrate sample. The diffraction profile on the dotted line in FIG. The dotted line in the figure shows the profile resulting from the virtual Ge substrate region and strained Ge layer evaluated by peak separation. Sectional structure of Sn composition graded multi-stage Ge _1-x Sn _x layer and strained Ge layer sample formed on a virtual Ge substrate. (A) Cross-sectional TEM images of the prepared Sn composition gradient multi-stage Ge _1-x Sn _x layer and strained Ge layer sample, and (b) an atomic resolution TEM image in which the strained Ge / GeSn interface is enlarged. Ge (224) reciprocal lattice of (a) sample immediately after growth of the Ge _1-x Sn _x layer in the third stage, and (b) sample having a Ge layer grown on the Sn composition graded multi-stage Ge _1-x Sn _x layer XRD-2DRSM observation results near the point. (A) (224) diffraction peak position of each stage of strained Ge / Sn composition graded multi-stage Ge _1-x Sn _x sample evaluated by XRD-2DRSM method, and (b) explanation of symbols and each Ge _1-x Strain relaxation rate at each stage of Sn _x .

Explanation of symbols

10: Multilayer structure 11: Si substrate 12: Strain relaxation Ge layer 13: Strain relaxation Ge _1-x Sn _x layer 14: Stretch strain Ge layer 20: Multilayer structure 21: Si substrate 22: Strain relaxation Ge layer 23 : Strain relaxation Ge _1-x Sn _x layer (Sn composition 1.0%)
24: strain relaxation Ge _1-x Sn _x layer (Sn composition 3.0%)
25: Strain relaxation Ge _1-x Sn _x layer (Sn composition 6.7%)
26: Stretch-strained Ge layer

Claims (16)

A method for forming a multilayer structure for a semiconductor device, comprising:
Forming a strain-relieved germanium layer on the silicon substrate ;
Forming a strained germanium tin mixed crystal layer on the germanium layer;
A method of forming a multilayer structure comprising a step of forming an extension strained germanium layer on the germanium tin mixed crystal layer .   The method for forming a multilayer structure according to claim 1, wherein a tin composition of the germanium-tin mixed crystal layer is 4.5% or more.   3. The method of forming a multilayer structure according to claim 1, wherein a lattice constant of a (110) plane in a substrate in-plane direction of the extension strained germanium layer has a value of 0.4015 nm or more.   Propagating strain relaxation in the germanium-tin mixed crystal layer by propagating threading dislocations from the germanium layer to the germanium-tin mixed crystal layer at their interface to form misfit dislocations. The method for forming a multilayer structure according to any one of claims 1 to 3. A method for forming a multilayer structure for a semiconductor device, comprising:
Forming a strain-relieved germanium layer on the silicon substrate ;
Sequentially forming a multilayer structure comprising a plurality of strain-relieved germanium-tin mixed crystal layers having different tin compositions on the germanium layer;
A method of forming a multilayer film structure comprising the step of forming an extension strained germanium layer on the multilayer structure .   6. The multilayer film structure according to claim 5, wherein a misfit amount of crystal lattice generated between the layers in the multilayer structure including the germanium layer and the germanium tin mixed crystal layer is smaller than 0.37%. Forming method.   The method for forming a multilayer structure according to claim 5 or 6, wherein a tin composition of the germanium-tin mixed crystal layer is 4.5% or more.   8. The multilayer structure according to claim 5, wherein a lattice constant of a (110) plane in a substrate in-plane direction of the stretched strain germanium layer has a value of 0.4015 nm or more. Forming method.   Propagating strain relaxation in the germanium-tin mixed crystal layer by propagating threading dislocations from the germanium layer to the germanium-tin mixed crystal layer at their interface to form misfit dislocations. A method for forming a multilayer structure according to any one of claims 5 to 8. A multilayer structure for a semiconductor element,
Germanium layer which is strain-relaxed formed on a silicon base plate,
A strain-relaxed germanium tin mixed crystal layer formed on the germanium layer;
A multilayer film structure comprising an extension strained germanium layer formed on the germanium tin mixed crystal layer .   The multilayer film structure according to claim 10, wherein a tin composition in the germanium-tin mixed crystal layer is 4.5% or more.   12. The multilayer structure according to claim 10, wherein a lattice constant of a (110) plane in a substrate in-plane direction of the tensile strain germanium layer has a value of 0.4015 nm or more. A multilayer structure for a semiconductor element,
Germanium layer which is strain-relaxed formed on a silicon base plate,
A multilayer structure composed of a plurality of strain-relaxed germanium tin mixed crystal layers having different tin compositions formed on the germanium layer;
A multilayer film structure comprising an extension strained germanium layer formed on the multilayer structure.   The multilayer film structure according to claim 13, wherein a misfit amount of a crystal lattice generated between the layers in the multilayer structure including the germanium layer and the germanium tin mixed crystal layer is smaller than 0.37%. .   The multilayer film structure according to claim 13 or 14, wherein a tin composition of the germanium-tin mixed crystal layer is 4.5% or more.   The multilayer film structure according to any one of claims 13 to 15, wherein a lattice constant of a (110) plane in a substrate in-plane direction of the stretch-strained germanium layer has a value of 0.4015 nm or more.