Growth of single-crystal metastable Ge1-xSnx alloys on Ge(100) and GaAs(100) substrates

Journal of Crystal Growth 83 (1987) 3—10 North-Holland, Amsterdam

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GROWTh OF SINGLE-CRYSTAL METASTABLE Ge1~Sn1ALLOYS ON Ge(100) AND GaAs(100) SUBSTRATES

S.I. SHAH and J.E. GREENE Department of Materials Science, the Coordinated Science Laboratory, and the Materials Research Laboratory, University of Illinois at Urbana-Champaign, 1101 West Springfield Avenue, Urbana, Illinois 61801, USA

and L.L. ABELS, Qi YAO and P.M. RACCAH Department of Physics, University of Illinois at Chicago, Chicago, Illinois 60680, USA Received 7 October 1986; manuscript received in final form 3 February 1987

Single-phase metastable Ge1 _~Sn~ alloys with a diamond-cubic structure have been grown with Sn concentrations up to 15% (the maximum equilibrium solid solubility of Sn in Ge is —1%) by bias-sputter deposition. Films with x ~ 0.08 grown on Ge(100) and GaAs(100) substrates were single crystals while films deposited on amorphous glass were polycrystalline with a (220) preferred orientation. In both cases, obtaining single-phase crystalline alloys required the use of low-energy (~180 eV) ion bombardment of the growing film in order to provide collisional mixing of the near-surface region during deposition. The allowable range in growth temperature 1 for obtaining single-phase alloys was a function of x and the ion acceleration potential ~a• The maximum in 7’, varied from 150°C at x = 0.2 to 90°C at x = 0.15. A growth phase map plotted as a function of x, T,, and V~,was determined. Raman spectra of Ge1 _~Sn~ alloys exhibit a two-peak density of states behavior which can be explained by assuming that Sn substitution introduces disorder in the bonding direction and removes the degeneracy of the optical branches.

1. Introduction Vapor-phase deposition techniques have recently been developed for the growth of a variety of new single-crystal metastable semiconductor alloys such as InSb1 _~Bi~ [1], (GaAs)1 _~(Ge2)~ [2], and (GaSb)1_~(Ge2)~[3—5].In each case, the pseudobinary constituents exhibit very limited equilibrium solid solubilities. Nevertheless, in at least the latter two systems, alloys with good thermal and temporal stability have been grown at compositions across the entire pseudobinary phase diagram. Studies of optical [2,6], electrical [5], lattice dynamic [7,8], structural [9,10], and thermodynamic properties [4,11] have been reported and reviews of this new class of semiconducting materials have appeared in the literature [12,13]. A metastable alloy system which may offer promise for infrared applications as well as for

serving as an important test system for investigating the growth of non-isostructural substitutional crystalline metastable alloys is Ge1 ~ Amorphous Ge/Sn mixtures possessing shortrange tetrahedral coordination have been deposited by sputtering [14,15] while Farrow et al. [16] have reported using molecular beam epitaxy to grow Ge1 _~Sn~ films with x ~ 0.99 (essentially the solid solubility limit of Ge in diamond-structure Sn) on InSb(100). However, attempts at growing extended single-phase crystalline solid solutions have up to now been unsuccessful [17,18]. Tin is an allotropic element which transforms from semiconducting diamond-structure a-Sn to the metallic body-centered-tetragonal structure /3-Sn at 13.2°C.The system is a simple eutectic in which the invariant point occurs at a composition close to that of pure Sn and at a temperature 1°C below the melting point of /3-Sn, 232°C —

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0022-0248/87/$03.50 © Elsevier Science Publishers B.V. (North-Holland Physics Publishing Division)

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SI. Shah et al.

/

Single-crystal metastable Ge

[19]. The maximum solubilities of Sn in Ge and Ge in Sn are 1 at% [20] and 0.6 at% [21], respectively, In this paper, we report the first results for the growth of metastable Ge1_~Sn~alloys with Sn concentrations up to 15%. Films with x ~ 0.08 grown on Ge(100) and GaAs(100) substrates were not only single-phase but also single crystals with a diamond-cubic structure while films deposited on amorphous glass substrates were polycrystalline with a (220) preferred orientation. The key feature in the growth of these alloys, as was the case for the (III—V)1 ~(TV2 )~ systems, was the use of low-energy (~180 eV) ion bombardment to provide collisional mixing of the upper one or two monolayers during deposition. The maximum allowable growth temperature was much less than for the (III—V)1_~(IV2)~alloys and ranged from 150°C at x = 0.02 to 90°Cat x = 0.15. A growth phase map plotted as a function of deposition temperature, ion energy, and film composition was determined. Raman measurements of as-deposited films showed that the Ge optical phonon mode at 300 cm~ split into two peaks with the addition of Sn and that the frequencies of both peaks decreased with increasing x. —

2. Experimental procedure All samples were grown by DC-diode sputter deposition in a system which allowed independent control of the target and substrate potentials with respect to the positive space-charge region in the discharge. The system base pressure during these experiments was iO~ Torr, achieved using a 260 1 s~ turbomolecular pump, and sputtering was carried out in ultra-pure Ar (99.999%) which was further purified by passing it through a Ti sponge getter maintained at 90°C prior to introduction into the growth chamber. The sputtering pressure was 25 mTorr. Two-phase, water cooled, targets were fabricated from undoped Ge wafers with room temperature resistivities > 40 Si cm and high-purity (> 99.999%) Sn. New targets were always sputter etched for > 2 h prior to initiating film growth experiments. The target voltage during sputtering —

—

1— ~

alloys on Ge(I00) and GaAs(l00)

was 3 kV resulting in a deposition rate of 0.2 nm s’, in the absence of a substrate bias, at a target-to-substrate separation of 4 cm. The total film thickness was typically 1 ~m. Applied negative substrate biases V~,ranged from 0 to 300 V. The plasma potential V~,which is additive to Va, was measured to be 5 V. Polished undoped Ge(100), Cr-doped semi-insulating GaAs(100), and Corning 7059 amorphous glass wafers were used as substrates in different growth runs. Immediately prior to inserting the substrates into the vacuum chamber, they were vapor degreased by successive rinses in trichloroethylene, acetone, and deionized water and then blown dry in dry N2. The semiconductor substrates were also etched in either dilute HF (Ge) or a 7 : 1 solution of 2 SO4 : H 202 (GaAs) and, without air exposure, rinsed in deionized water and again blown dry in dry N2. The substrates were then thermally contacted to a Ge-coated Mo platten using liquid In. Final substrate preparation for the Ge and GaAs wafers consisted of sputter etching for 30 mm at 600°Cin a 250 V, 15 mTorr, Ar discharge in order to remove the passivating oxide layer. Film growth temperatures 7~were monitored using a chromel/alumel thermocouple soldered to the front side of a Ge-coated substrate with liquid In. The reported values for 1 include the contributions from both the substrate heater and energetic particle bombardment. 1~was maintained constant during deposition using an automatic temperature controller and values quoted are accurate to within ±5°C. Film compositions were determined by energydispersive X-ray analysis (EDAX) in an electron microprobe. Pure Ge and /3-Sn were used as reference standards and matrix corrections for X-ray fluorescence, absorption, and atomic number were carried out using the Magic V computer program [22]. Reported compositions are accurate to within ±0.5%. The structure and crystalline perfection of as-deposited films were investigated by X-ray diifraction (XRD), electron channeling in a scanning electron microscope, and transmission electron microscopy (TEM). Raman spectra from single-crystal alloy films on GaAs(100) substrates were recorded using 100 —

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SI. Shah et al.

/ Single-crystal metastable Ge1

mW, 514.5 nm radiation from an Ar~ ion laser beam incident at near the Brewster angle. The scattered light was collected in a near-backscattered geometry, dispersed by a double monochromator with a 7 cm’ spectral resolution, and detected by photon counting techniques. Optical frequencies were determined by least-square fits of theoretical line shapes to be observed spectra.

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aSn~alloys on Ge(I00) and GaAs(I00)

5

a

12C

Ge095 Sn005 ~

icx

Ge(11l)

Substrate~7O59Glass T~rI5O°C

/3-Sn(200)

-

V0~0

80

-

Ge(220)

60

.~‘ -~

C C

40

—

3. Experimental results

-

-~

All films with x > 0.02 grown at temperatures T1 ~ 75°C with no applied substrate bias were two-phase while films grown at lower temperature were amorphous. However, single-phase crystalline films with x up to 0.15 were obtained over a range in 7~by applying a negative substrate bias Va. Fig. la shows a typical XRD pattern from a 1 ~tm thick Ge0 95Sn005 film deposited on Corning 7059 glass at 7 = 150°C and V~,= 0. Diamondstructure Ge (111) and (220) peaks as well as tetragonal /3-Sn (200), (101), and (220) reflections were observed, Ge095Sn005 films deposited under the same conditions except for the addition of an applied substrate bias ~ 150 V( but less than 180 V as we will show later) were single phase. An only diffraction example is given observed fig. lb showing were anV.XRD and pattern from a peaks filmingrown with Va =the150(220) The (111) alloy reflections. Comparing the relative peak intensities with those obtained from a Ge powder diffraction pattern indicate that the film has a (220) preferred orientation. Using the full peak widths at half-maximum intensity from XRD patterns such as shown in fig. ib, the average grain size was estimated from Scherrer’s equation [23] to be 20 nm. Grain sizes obtained from X-ray peak broadening are typically much smaller than those observed directly in the TEM [24]. This is due to the fact that dislocations and non-uniform strain also contribute to peak broadening. (Better agreement between XRD and TEM results can, in principle, be obtained using the Warren—Averbach method [25] to analyze the Fourier components of the broadened peak and subtract out the effects of in-plane strain. However, the higher-order reflections which

20 0 __________________________________ 24 28 32 36 40 44 48

Bragg Angle, 29 12C

b I

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Ge095 Sn005 Substrate = 7059 Glass

-~

b~

(degree)

-

T1r150°C

>~

Var 150V 80

~~

-

Alloy (111)

-~

~

Alloy (220)

-

40

24

—k—-—- - ~- _ _- ~_. . . ~ 28

32

36

40

Bragg Angle, 29

44

(degree)

48

I

52

Fig. X-ray diffraction spectra from (a) a two-phase film 1.with average composition Ge095Sn05 depositedGe/Sn on a Corning 7059 glass substrate at 150°Cwith no applied bias and (b) a single-phase Ge095Sn05 alloy deposited on Corning 7059 glass at 7, = 150°C and Va = 150 V.

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are required to use this technique typically have very low intensities in films with small grain sizes.) A few thin, 50 nm thick, alloy films were also grown on NaC1(100) substrates (lattice constant a0 = 0.564 nm) and then floated off in deionized water for TEM analysis. Fig. 2 shows a typical diffraction pattern and bright field image from a Ge095 Sn 0.05 film grown at T~= 150°C and Va = 160 V. The average grain size was 80 nm and the arcs in the diffraction rings were evidence of —

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SI. Shah eta!.

/ Single-crystal metastable Ge,

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aSn~alloys on Ge(100) and GaAs(100)

0.580

— 10gm

::::HIIIlI

0

Fig. 3.

Ge1~Sn~

-

0.04 0.08 0.12 Atomic Fraction Sn, x

The lattice constant a0 of Ge1 ~Sna x.

0.18

as a function of

preferred orientation. The positions of the intensity maxima in the (111) and (220) rings indicated that the preferred orientation was (100) in registry with the substrate. Fig. 3 shows the lattice constant a0 of Ge1~Sn~obtained from the XRD spectra from single-phase films grown on glass substrates, as a function of x. Both the (220) and (111) reflections gave identical results. The lattice constant a0 was found to obey Vegard’s law and vary linearly from the lattice constant of pure Ge (a0 0.5657 nm) towards that of diamond-structure a-Sn (a0 0.06489 nm). Lattice constants obtained from films grown on Ge(100) and GaAs(100) substrates also varied linearly with x but were displaced somewhat from the Vegard’s law line due to compressive stress in the films resulting from the increasing film/substrate lattice mismatch as the Sn concentration was increased. As noted above, low-energy ion bombardment of the growing film was essential for obtaining single-phase alloys. This is illustrated in fig. 4 showing a phase map, plotted as a function of T~ and Va~ for Ge095Sn005 alloys grown on Ge(100) substrates. All films deposited with Va less than 50 V were either amorphous or two-phase de=

=

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Fig. 2. Transmission electron microscopy: (a) bright field image and (b) corresponding diffraction pattern from a single-phase Ge09, ~, alloy deposited on a Corning 7059 glass substrate atT,=150°Cand Va=160V.

SI. Shah et al.

/ Single-crystal metastable Ge1

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~Snr alloys on Ge(100) and GaAs(I00)

7

180

Geo.95 Sno.05 160

-

50

~ Single Phase Single Crystal ~-J Polycrysfal ~

75

100

125

-

150

Applied Substrate Bias, V0

175

(V)

200

225

Fig. 4. Growth phase map plotted as a function of the growth temperature 1 and applied substrate bias deposited on Ge(100) substrates.

pending upon the growth temperature. Increasing Va above 50 V allowed the growth of single-phase crystalline films. The maximum growth temperature T~’for obtaining single-phase films increased with increasing Va until T* itself reached a maximum value at Va*. Increasing Va> Va* resulted in a decrease in 2’~’.The minimum “epitaxial temperature” for this alloy was found to be 145°C; thus single crystals could only be obtained over a rather narrow window in growth parameters: 7 between 145 and 155°Cand Va between 140 and 180 V. Similar results were found for growth on GaAs(100) substrates. Fig. 5 is an electron channeling pattern (ECP) from a Ge059Sn005 film on ECP GaAs(100) at 16°grown V. The was ob7=150°C and Va= tamed using a 25 kV beam with a 10 ~tm aperture. The higher order lines are sharp indicating that the film is of good crystalline quality and no change in the ECP was observed while scanning the probe beam over the entire samples demonstrating that it was a complete single crystal. Epitaxial layers were obtained at Sn compositions up to 8%. At higher Sn concentrations, the epitaxial temperature was less than T,. Structural data obtained from X-ray and elec-

Va

for Ge0 95Sn05 films

summarized in fig. 6 as a three-dimensional plot of the maximum growth temperature 7~ for obtaming single-phase films versus both the substrate bias during deposition and film composition. The maximum in the 7*(x, Va) surface de-

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tron diffraction analyses of more than 60 films are

-

Fig. 5. Electron channeling pattern from a 1 ~sm thick Ge 095Sn0, film grown on GaAs(100) at 150°Cwith an applied substrate bias of Va = 160 V.

8

S.f. Shah et aL ~-‘

/

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Single-crystal metastable Ge, ~Snr alloys on Ge(100) and GaAs(100)

200 Ge

1..,, Sn,

~ 150

-

-

~ ioo

-

-

so

creased and moved to lower Va values with increasing x. I~’ranged from 150 to 90°Cas x was increased from 0.02 to 0.15 while Va* decreased from 180 to 100 V over the same composition range. Low-frequency Raman spectra from singlecrystal Ge1_~Sn~ films exhibited a weak diffuse structure with a Stokes shift centered near 270 cm Samples with smaller x values scattered a greater portion of radiation at larger frequency shifts. Smoothed spectral traces for 0.01 0.05. ~.

0.05

/ ‘9 0~10

015

0

—50

—100

—150

—200

—250 lVl

Applied Substrate Bias, Va

—300

.

Fig. 6. Growth phase map for Ge, - Sn . The three-dimensional surface corresponds to the transition from a single-phase metastable alloy to a two-phase mixture as a function of the Sn fraction x, the growth temperature 3’,,, and the applied substrate bias Va.

I

300

I

I

4. Discussion

I

I

I

Ge15 Sflx

-

-

°

\

.

280

-

-

260

-

-

S

Single-phase polycrystalline diamond-structure Ge1_~Sn~alloys with Sn concentrations up to 15%, more than 15 times the equilibrium solid solubility, were grown on amorphous glass substrates while epitaxial metastable layers with Sn concentrations of up to 8% were grown on Ge(100) and GaAs(100) substrates. The essential feature allowing the growth of these unusual alloys was the use of low-energy ion bombardment of the growing film. This can be understood based upon ion/surface interactions. The primary ion species incident at the growing film was2 Ar~ theV.current density was width 0.1 at Vaand150 The plasma sheath mAthe cmsubstrate was estimated from Child’s law at [26] to be <0.8 cm while the mean free path for charge exchange collisions at 25 mTorr was > 1.2 cm. Thus most of the ions incident at the growing =

240

-

I

-

220o

I

0.02

I

I

0.04 0.06 Atomic Fraction Sn, x

I

0.08

o.io

—

Fig. 7. Frequencies, measured using Raman spectroscopy, of single-crystal metastable Ge,~,Sn~alloys as a function of composition.

SI. Shah et al.

/ Single-crystal melastable Ge,

—

~

alloys on Ge(100) and GaAs(100)

9

film experience the full sheath potential given by the sum Va + V~ Va. The Ar~ion energies used

heavier Sn atoms, the frequencies of both optic modes in Ge1 _~Sn~ decrease rapidly from the 300

in the present experiments correspond to a projected ion range in the growing film of the order of one or two monolayers. Ion irradiation during deposition provided collisional mixing of the near-surface region without affecting the “bulk” lattice and 1” initially increased with increasing Va. The collisionally mixed layer was then buried by subsequently deposited material. During deposition, the mixing and overlaying were, of course, happening simultaneously. The alloys remains in solution provided that the growth temperature is sufficiently low to avoid bulk diffusion driven by the free energy difference between the metastable and equilibrium states. However, increasing Va results in an increase in the ion range and eventually gives rise to enhanced diffusion, and hence phase separation, in the bulk film. Thus, there was an optimum value of Va above which 7~*decreased. It should be noted in the above discussion that the amount of ion-bombardment-induced mixing depends not only upon Va but also upon the flux of incident particles. At Va 150 V, the ratio of accelerated ion to neutral vapor flux was of the order of 0.3. In addition, there was a substantial flux of incident energetic neutral particles (primarily Ar) due to accelerated ions which were neutralized and backscattered from the target.

cm’ optic mode of pure Ge to 274 and 237 cm respectively, at x ~ 0.05. The LO and TO branches of Ge at the L point have been reported at 280 and 246 cm’ respectively [27]. From fig. 7 the extrapolations of our data agree well with these reported values. This indicates that Sn substitution affects primarily the bonding direction and, by introducing disorder, removes the LO/TO degeneracy. —

~,

Acknowledgements The authors gratefully acknowledge the financial support of the Materials Science Division of the Department of Energy (S.I.S. and J.E.G.), the Joint Services Electronics Program (J.E.G.), and the Office of Naval Research (L.L.A., Q.Y., and P.M.R.) during the course of this research.

References

=

The window in growth parameter space for obtaining single-phase films decreased at higher Sn concentrations due to the increase in the thermodynamic driving force for phase transformation to the equilibrium two-phase state. The relatively high diffusivity of Sn in Ge and the low melting point of Sn resulted in much lower values of 7~’ for Ge1_~Sn~ than, for example, (GaAs)1_,(Ge2)~ 500—550°C) [3] and (GaSb)1_~(Ge2)5 (450—500°C)[5]. Nevertheless it was still possible to grow single-crystal layers with Sn concentrations up to 8%. The range in x values attainable for single crystals was less than for polycrystalline films, 15%, due to the additional constraint that the epitaxial temperature be less than T* Raman results suggest that single-crystal metastable ~ alloys have a two-peak density of states spectrum. With the substitution of the (—

J.E.

[1] J.L. Zilko and Greene, 1. Appi. Phys. 51(1980) 1549; 51 (1980) 1560. [2] S.A. Barnett, MA. Ray, A. Lastras, B. Kramer, J.E. Greene and P.M. Raccah, Electron. Letters 18 (1982) 891. [3] K.C. Cadien, A.H. Eltoukhy and J.E. Greene, AppI. Phys. Letters 38 (1981) 773. [4] K.C. Cadien, A.H. Eltoukhy and J.E. Greene, Vacuum 31 (1981) 253. [5] J.E. Greene, S.A. Barnett, K.C. Cadien and Ray, J. Crystal Growth 56 (1982) 389. [6] K.E. Newman, A. Lastras, B. Kramer, S.A. Barnett, MA. Ray, J.D. Dow, J.E. Greene and P.M. Raccah, Phys. Rev. Letters 50 (1983) 1466. [7] TN. Krabach, N. Wada, M.V. Klein, K.C. Cadien and J.E. Greene, Solid State Commun. 45 (1983) 895. [8] R. Besserman, i.E. Greene, MV. Klein, T.N. Krabach, T.C. McGlinn, L.T. Romano and SI. Shah, in: Proc. 17th Intern. Conf. on Physics of Semiconductors, San Francisco, 1984, Eds. J.D. Chadi and WA. Harrison

MA.

(Springer, Berlin, 1985). [9] E.A. Stern, F. Ellis, K. Kim, L. Romano, SI. Shah and i.E. Greene, Phys. Rev. Letters 54 (1985) 905. [10] SI. Shah, B. Kramer, S.A. Barnett and J.E. Greene, J. AppI. Phys. 59 (1986) 1482. [11] K.C. Cadien, B.C. Muddle and J.E. Greene, J. Appl Phys. 55 (1984) 4177. [12] i.E. Greene, I. Vacuum Sci. Technol. Bi (1983) 229.

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/ Single-crystal metastable Ge,

[13] S.A. Barnett, B. Kramer, K.T. Romano, SI. Shah, MA. Ray, S. Fang and i.E. Greene, A Review of Recent Results on Single Crystal Metastable Semiconductors: Crystal Growth, Phase Stability and Physical Properties, in: Layered Structures, Epitaxy, and Interfaces, Eds. J.M. Gibson and L.R. Dawson (North-Holland, Amsterdam, 1985) p. 285. [14] R.J. Temkin, G.A.N. Connell and W. Paul, Solid State Commun. 11(1972)1591. [15] Ri. Temkin and W. Paul, in: Proc. 5th Intern. Conf. on Amorphous and Liquid Semiconductors, GarmischPartenkirchen, 1973. [16] R.F.C. Farrow, D.S. Robertson, G.M. Williams, AG. Cullis, G.R. Jones, I.M. Young and P.N.J. Dennis, i. Crystal Growth 54 (1981) 507. [17] G.M. Kuznetsov, L.S. Tsurgan, F.A. Grimel’farb and V.A. Rotenberg, Inorg. Mater. 41 (1975) 1015. [18] S. Oguz, W. Paul, T.F. Deutch, B.Y. Tsaur and D.V.

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,,Sn~alloys on Ge(100) and GaAs(IOO)

Murphy, Appi. Phys. Letters 43 (1983) 848. [19] W. Klemm and H. Stöhr, Z. Anorg. Chem. 241 (1939) 305. [201 CD. Thurmond, F.A. Trumbore and M. Kowalchik, J. Chem. Phys. 24 (1956) 799. [21] CD. Thurmond, i. Phys. Chem. 57 (1953) 827. [22] i. Colby, in: Advances in X-Ray Analysis, Ed. G. Mallett, M. Fay and M.M. Mueller (Plenum, New York, NY, 1968). [23] B.D. Cullity, in: Elements of X-Ray Diffraction (Addison-Wesley, Reading, MA, 1978). [24] J.-E. Sundgren, A. Rockett, i.E. Greene and U. Helmers. son, i. Vacuum Sci. Technol. A4 (1986) 2770. [25] See, for example, H.P. Kiug and L.-E. Alexander, X-Ray Diffraction Procedures (Wiley, New York, 1974) ch. 9. [26] B. Chapman, Glow Discharge Processes (Wiley, New York, 1980). [27] B.N. Brockhouse and P.K. Iyengar, Phys. Rev. 111 (1958) 747.