I. Phys. Chum. Solids. 1978. Vol. 39. pp. 529-538. Pergamon Press. Printed in Great Britain.

EPITAXIAL GROWTH OF Nb3Ge ON NbJr AND NbdW

A.H. DAYEM
BellTelephone Laboratories, Holmdel, NJ 07733, U.S.A.

and A.B. HALLAK~U.S.A.

T.H. GEBALLES,R.B.ZUBECK§
Stanford University, Stanford, CA 94305,U.S.A.

and

G. W. HULL,Jr.
Bell Telephone Laboratories, Murray Hill, NJ 07974,U.S.A.

(Receioed 25 August 1977; accepted in revised form 12 October 1977)

Abstract-This is a report on a cooperative research carried out in Stanford University to investigate the possibility
of using epitaxy to prepare the high T, superconductor NbJGe in an Al5 crystalline structure at the 3: 1
stoichiometry.
NbrIr polycrystalline Mms with the Al5 structure deposited on sapphire were used as substrates for the
epitaxial growth of NbrGe because of the favorable lattice parameter match. The experimental results clearly show
that epitaxial growth indeed occurs and helps to extend the range of homogeneity of the A15 phase up to 26.3 at.%
Ge as compared with the thermodynamic equilibrium boundary at 19 at.% Ge. We also used Nb3Rhfilms as
substratesand found them inferior to Nb& becauseof the multiphasenature of the films.
In addition to extendingthe Al5 phase boundary epitaxyresults in a considerable rise in the superconducting
transition temperature for Ge-rich samples together with a reduction in the transition width. The work suggests that
polycrystalline epitaxy can be an important tool in the synthesis of thin-film intermetallic compounds.

INTRODUCTION
on a substrate held at an optimized temperature to obtain
In thermodynamic equilibrium the Al5 phase in the
a metastable A15 structure with a Ge concentration
Nb-Ge binary system occurs in the composition range
larger than the equilibrium value of 19at.% Ge. But since
17-19at.%Ge[l]. For higher Ge concentrations a two-
the Al5 phase is competing with the stable compound
phase region extends up to 37.5at.%Ge where the Al5
NbsGe3, the structure obtained is extremely sensitive to
phase coexists with the stable compound NbsGes. The
the nucleation conditions at the substrate, intentional or
compound NbXGe does not exist in stable equilibrium.
accidental contaminations, the deposition rate and the
However, recent experiments have shown that a
substrate temperature. Experience shows that the film
metastable A15 phase can be prepared with a Ge
crystalline structure, as well as its T, and transition
concentration > 19at.%, and with a T, onset approach-
width, vary considerably among samples prepared under
ing 23°K. The methods of preparation include splat-
apparently identical conditions.
cooling of a Ge-rich melt [2], controlled-rate
A logical alternative is to grow the Nb3Ge epitaxially
evaporation[3], sputtering from a NbGe target of given
on a substrate having an A15 structure with a lattice
composition [4], and chemical vapor deposition (51.
constant which matches that of the Nb3Ge.[7] Under
In a recent study Hallak et al. [6], have determined the
such conditions it is expected that the Gibbs free energy
range of existence of the metastable A15 phase in elec-
is a minimum if the overgrowth has a lattice which
tron-beam codeposited NbGe films prepared in vacuum
matches that of the substrate. This provides the optimum
( - 10m6Torr) or in the presence of oxygen. They find
conditions for the Nb3Ge to grow in-an Al5 structure,
that the A15 phase boundary can be made to extend to
23 at.% Ge at an oxygen partial pressure - 2 x 10m6Torr since the nucleation and growth of the competing
with the substrate held at 775°C during film deposition. NbsGe3 phase has energetically become less favourable.
In all the above methods one is rapidly quenching a If the free energy minima are sufficiently altered by the
mixture of Nb and Ge vapors with a given composition crystalline substrate the A15 phase boundary may extend
well beyond the 25at.% Ge and hence allowing the
formation of Nb3Ge in a “pure” Al5 phase.
tResearch at Stanford was supported in part by the Air Force Ideally, one would like to have an Al5 single crystal
Office of Scientific Research, Air Force Systems Command,
USAF, under Grant No. AFOSR 73-2435, and by the National
substrate for this epitaxial growth. Lacking this, an
Science Foundation’s Materials Research Laboratory Program alternative is to prepare on the sapphire substrate a
through the center for Materials Research at Stanford. polycrystalline thin film of a compound having a ther-
SDepartments of Applied Physics and Materials Science; also modynamically stable Al5 phase with the proper lattice
at Bell Laboratories, Murray Hill, NJ 07974,U.S.A. constant, and then grow the NbsGe “epitaxially” on top
IPresent address: IBM Research Laboratory, San Jose, CA
95193,U.S.A. of this film. It was found that NbJr, and to a lesser
YPresent address: Department of Physics, University of degree also Nb3Rh, are suitable candidates. In the
Jordan, Amman, Jordan. following sections we will first discuss the properties of

529
530 A. H. DAYEM et al.

evaporated Nb31r and NbsRh thin films. We will next Nb,Ge on sapphire in the composition range from - 16
show that in spite of the polycrystalline nature of these to - 26 at.% Ge. We determined the lattice constant (Fig.
films the Nb3Ge grows epitaxially on these substrates in 1) and T, (Fig. 2) vs composition. We find that the Al5
an Al5 structure over a wide composition range from 17 boundary in these two samples extends to 22 at.% Ge and
to - 26 at.% Ge. Although the highest T, obtained in the lattice constants are considerably smaller than the
these films did not exceed 23”K, both the crystalline equilibrium values depicted by the top solid line in Fig. 1.
structure of the films and the transition width have been At 18at.% Ge, e.g. the film lattice constant is - .36%
greatly improved and the variation from sample to smaller than the equilibrium value. The reason for this
sample greatly reduced. relatively large compression is not known. One can
easily calculate that neither the differential thermal
EXPERIMENTAL expansion between the film and the sapphire substrate,
The samples were prepared on highly polished, nor the lattice mismatch will result in the observed
randomly oriented sapphire single crystal substrates of compression. Two other factors could influence the
dimensions 0.635 x 0.635 x 0.0635 cm3. The geometry of lattice constant of the film. In a Nb-rich sample one
the apparatus permitted placing 10 such substrates in a expects that some excess Nb atoms will occupy Ge-sites
direction parallel to the line joining the Nb and Ge
sources. In this way the composition of the deposited
films varied from one substrate to the next by about
l-l.2 at.% Ge. During evaporation the substrate
temperature was held at 875 i 10°C using a temperature- EQUILIBRIUM
controlled oven.
Two separately controlable electron guns were used, a
one to activate a fixed Nb source while the other
activates the Ge, Ir or Rh sources which were placed in
separate crucibles in a rotating table.
Electron microprobe analysis was used to determine
the composition of the film to an accuracy of - t 1
at.%. The lattice constant was determined from
X-ray diffractometer traces. The diffraction patterns 5.12 1
16 18 20 22 24 26 28 30 32
were further recorded on photographic films using
a Reed diffraction camera. These photographs were at.%Ge
used to determine any preferred orientation as well as Fig. 1. Al5 lattice constant vs composition for two Nb,Ge
the number of crystalline phases present. The supercon- samples deposited on sapphire at T, = 875°C and r = 66 &sec.
Film lattice is under considerable compression relative to its
ducting transition temperature was measured on all state in thermodynamic equilibrium.
samples both inductively and resistively at the same time

r
to an accuracy of ~0.2”K. For most of the samples the 22
inductive T, was about 1K below the resistive one, and
of the same width. 20 -
Nb,Ge FILMS ON SAPPHIRE
Before discussing the effects of epitaxy we first sum- 18 -
marize the properties of the Al5 phase found in bulk
material in thermodynamic equilibrium and in thin films
16 --
deposited on sapphire substrates. The equilibrium phase
diagram of the NbGe system was described by Milller[ll.
As mentioned in the introduction, the region of homo- 14 a-
geneity of the A 15 phase extends from 17 to 19at.% Ge
TEK
where T, rises from 4.8 to 5.9”K and hence stays 12 .B
constant for larger Ge concentrations while the lattice
constant drops from 5.176 to 5.167A. The Al5 phase
10 --
exists in equilibrium with NbsGe3 in a mixed phase
region extending from 19 to 37.5 at.% Ge, while on the
Nb-rich side, the Al5 coexists with a solid solution of Ge f33-

in Nb. The high stability of the NbzGe3 compound pro-

hibits the Al5 phase from extending up to 25 at.% Ge
and hence the stoichiometric Nb3Ge compound does not
exist under thermodynamic equilibrium conditions.
Hallak[8] has studied the effect of the substrate 16 18 20 22 24 26 28 30
temperature T, and deposition rate r on the formation of at. % Ge
the A 15 phase over the range 600 < T, < 900°C and 66 < Fig. 2. Transition into superconducting state vs composition for
r < 130&sec. Using his optimum values of T, = 875°C the two Nb,Ge samples in Fig. 1. The dashed line is T, for bulk
and r = 66 &sec we prepared two sets of samples of material.
 Epitaxial growthof NbjGe on &Ir and Nb,Rh 531

and that some Ge vacancies or impurities on Ge sites will

be incorporated in the film during deposition. Such Ge-
site substitution can result in a film lattice constant which
is smaller than that of the bulk. At the same time it can
account for the higher T, of the films as depicted in Fig.
2. The other factor is the lattice constants of the tetra-
gonal compound NbJGe3: a = b = 10.163 and c =
5.130 .&.In the mixed phase region the A 15 phase grows
in close association with the compound NbxGe3 and
hence its lattice constant may be forced to conform with
that of the NbsGeT to achieve a minimum in the energy
of the solid deposit at T = Z’,.In the single phase region,
the Reed photographs do not show the presence of any
diffraction Iines belonging to Nb5Ge3. Nonetheless, one
may assume that the A15 phase grows around small
cores of the compound NbsGes and hence is subjected to
a similar compression. If these cores are of sufficiently
small dimensions, X-ray diffraction lines originating from
them will be broad and of small enough intensity that
they will not be visible in the Reed photographs.
The transitions into the superconducting state as
measured by a It-point resistance probe are shown in Fig.
2. The transition widths, sensitive to the degree of in-
homogeneity of the film, are rather large especially in the
single phase Nb-rich region. However, the onset T,,
ranging from 17.5 to 20.5”K is much higher than the
T, = 6°K of the bulk material. This high T,, when linked
with the observed reduction in the Al.5 lattice constant
gave an added incentive to the initiation of the present
study. It was argued that, if through epitaxy one could
further compress the lattice constant, higher T,‘s will be
achieved. However, as will be seen later, Nb,Ge films
(0.5 pm thick) grown epitaxially on polyctystalline Nb,Ir
films, tend to acquire their own “natural” lattice
constant. It is quite possible that if the Nb,Ir substrate Fig. 3. X-ray photographusing Reed diffraction camera for a
were in the form of a perfect A 15 single crystal, epitaxial Nb,Ge tiirn (24at.% Ge) deposited on sapphire at T, = 875°Cand
growth of Nb,Ge may behave quite differently and may r = 13.9A/set. The two smallarcs on the extremeleft are the Kp
reflections of the (200) Al5 line and the mystery line; they are
permit moderate compression of the Nb,Ge lattice and
immediately follo&d by their respective Ku retlections. The
possibly also a rise in T,. small arcs off the equator are due to the NbsGes phase growing in
Since epitaxial growth at a given substrate temperature a preferred orientation.
generally improves at reduced deposition rates, we pre-
pared two sets of Nb,Ge samples at rates below those
investigated by Hallak[6], [7]: one in the composition 5.18 t-
range 16-26 at.% Ge at a rate 16.5&sec, the other in the
range 22-29.5 at.% Ge and 13.9 &sec. From Reed pho- 5.47 -
tographs we found that the boundary of the A 15 phase is
now coincident with its equilibrium value of 19at.% Ge. 5.16 -I
Both Reed photographs (see Fig. 3) and diffractometer 8.
traces show that, in addition to the Al5 and NbsGe:, 5.15 -
diffraction lines, there are two “mystery lines” at 28 =
33.84”and 213= 71” which have been previously observed 5.*4 -
by other authors[9]. The measured lattice constants
shown in Fig. 4 fluctuate erratically from sample to 5.13 -
sample but remain confined between the equilibrium
value of 5.167A and the c = 5.13 A of the tetragonal 5s2 ’ ‘ I I I I I

16 18 20 22 24 26 28 30
NbTGel. Comparing Fig. 2 with Fig. 5 one finds that the at.%Gc
reduced rate results generally in a reduction in T,. More
importantly, however, is the huge differences in T, in the Fig. 4. AI5 lattice constant vs composition for two Nb,Cie
sampiesdepositedon sapphireat T, = 875°Cand r = 16.5(0) and
composition range - 25 at.% Ge. 13.9Alsec (0). The phase boundary in the A15 is at 19at.% Ge.
From the above discussion one is lead to the Notice the large scatter in lattice constants obtained in the
conclusion that the crystalline structure and the super- two-phase region.
532 A. H. DAYEM et al.

22 conducting transition of Nb,Ge films vary widely among

r
samples prepared under apparently identical conditions
20 because the relaxation of the metastable Nb3Ge can be
suppressed in a variety of subtle ways.
18
II THE Nb# AND &Rh FILMS ON SAPPHIRE
I The range of homogeneity of the Al5 structure in the
16 Nb-Ir binary system[lO] extends over the wide range
: * t k
21.5-28.5 at.% Ir. Codeposition of Nb and Ir on a
1 sapphire substrate held at 875°C invariably resulted in
14
polycrystalline films with a pure A 15 structure. The Reed
TiK
photograph in Fig. 6 represents the typical diffraction
12
arcs obtained from all films in the proper composition
i 1 I"x: P
6 range. Using diffractometer traces to determine the
10 -

8-
1 lattice constant vs composition we obtained the results
shown by the solid line in Fig. 7. The dashed line
represents the thermodynamic equilibrium values for
bulk samples[lO]. The general agreement between film
and bulk behavior is quite good both for the magnitude
6-
of the lattice constant as well as the position of Al5
phase bondaries. However, the film lattice is slightly
I I I I
16 18 20 22 24 26 28 30
expanded relative to the bulk lattice with the amount of
expansion increasing with decreasing Ir concentration.
at.% Ge
Fig. 5. T, vs composition of samples shown in Fig. 4. The origin of this expansion is not known.
It is obvious that the Al5 Nb,Ir is ideally suited for
studying the epitaxial growth of Nb,Ge. By proper
selection of the composition of Nb,Ir its lattice constant
could be chosen to lie anywhere between 5.125 and
5.169 8. This will enable the experimental determination
of substrate lattice constant best suited for epitaxially
growing Nb3Ge.
We turn now to the Nb-Rh system. It was found that
Al5 Nb,Rh forms peritectoidally at - 122O”C[ll] with
an extremely slow reaction rate. A metallographic esti-
mate puts its homogeneity range at < 1 at.% around
25 at.% Rh. Long time anneals of bulk samples result in a
mixture of three phases: cubic a Nb t A 15+ tetragonal
V. This behavior was reflected in the Nb3Rh films
deposited on sapphire. Films prepared at T, = 875°C give
the X-ray diffraction pattern shown in the Reed pho-
tographs in Fig. 8. Some faint arcs are discernible in the

5.q8r

5.17 -5
I
5.161 -l
5.15
i
5.i4
I
5.13

5.12t

96 18 20 22 24 26 28 30 32 34
at%Ir

Fig. 7. Al5 lattice constant vs composition for two Nb,Ir

samples deposited on sapphire at T, = 87s”C and r = 16Alsec.
Fig. 6. Reed photograph of Nb$ (24.9at.%Ir) deposited on There is a good agreement with the thermodynamic equilibrium
sapphire at T, = 875°C and r = 16A/set. This is a typical pure values (shown by solid circles on the dashed line), but the film
A 15 diffraction pattern. lattice is slightly expanded.
 Epitaxialgrowthof NbrGe on Nb& and NbrRh

Fig. 9. Reed photographs of Nb,Rh sample (22.7at.%Rh)

Fig. 8. Reed photograph of Nb,Rh sample (24at.% Rh) deposited deposited on sapphire at T, = 400°Cand r = 13&sec. Notice the
on sapphireat T, = 875°Cand r = 13bi/sec. The film has a strong disappearance of the preferred orientation shown in Fig. 8.
preferred orientation in the (200)direction. Notice the inverted
arcs whichrender lattice analysis extremelydifficult. 5.14 -

familiar Al5 pattern. The extremely strong spots on the 5.13 -

equator show that the films have a highly preferred
orientation in the (200) direction. However, similar spots 532 - BULK
off the equator are in the form of streaks which join
5.11 -
together at the equator with the (200) spot forming a a
continuous arc of the wrong curvature. This implies that 530 -
different points on the arc have a different diffraction
angle and hence must have originated from crystallites of
different lattice constants. This pattern is indeed ex-
tremely difficult to understand.
oNbtl51 Al5tu 1 o
To get rid of the highly preferred orientation and the I I I tt, ,I I I
1

associating peculiar crystalline structure, we prepared 18 20 22 24 26 28 30 32 34 36 38

some Nb,Rh films at T, = 400°C. The Reed photographs at %Rh
of Fig. 9 show that the films are now polycrystalline with Fig. 10.Lattice constant a0 of A IS or c, or tetragonal u-phase vs
composition for a Nb,Rh sample. Values were determined from
little or no preferred orientation. But the A 15 pattern has one single diffraction line: the (400) of Al5=(004) of o using
also disappeared. This does not imply, however, the both Cu Ka and Cu Kp. Bulk values of the (r phase are shown
complete absence of an A 15 phase. As mentioned earlier by the dashed line; isolated point is for bulk AH. The arrows
the A 15 Nb3Rh forms usually in a mixture with aNb and indicate phase boundaries.
u phases and hence may give its own peculiar X-ray
diffraction pattern. the different phases occuring in a film of a given
Considerable time and effort were hence spent on composition. This lead us to the conclusion that indeed
analyzing the X-ray diffractometer traces to determine all three phases must in general be present in order to
534 A. H. DAYEMet a&.

account for all the observed diffraction lines. However,

since the c lattice constant of the a phase is almost equal
to the lattice constant of the AU, the (200), (210) and
(400) lines of the A15 had to be considered as the (002),
(202) and (004) of the u-phase. Hence, computer evalua-
tion of a, of A 15 and of a, and co of u gave results with
large standard error. The values obtained were scattered
in such a way that no consistent dependence on
composition could be established.
The only consistent set of lattice constant data was
obtained by choosing the (400) line and determining its
position (28) for both the CuKcv and CuK@ radiation
with the largest accuracy provided by the diffractometer.
This gave two estimates of the lattice constant at each
composition. The lattice constants thus determined can
be either the a, of the Al5 structure or the c, of the
tetragonal u phase. Their variation with composition is
shown in Fig. 10 together with the relevant phase boun-
daries (vertical arrows) which occur in equilibria in the
bulk. One finds that around the phase boundary between
the (A I5 + u) and the o phases there is a region where
the lattice constant varies slowly with composition. The
equilibrium values of the c lattice constant of the D
phase are shown by the dashed line while the isolated
point is that of the A 15. In summ~y, we find that Nb,Rh
films contain in general three phases including an A15
structure of a lattice constant - 5.11 iO.02 A in the
composition range 24-32 at.% Rh.
Using SEM many attempts have been made to deter-
mine the grain size of the Nb,Ir and Nb,Rh films either
as evaporated or folIowing various polishing and
decorating procedures. No well-defined grain structure
was resolved up to a magnification of 58,000x. A rough
guess is to place the grain size between 200 and 1500A.
Fig. 11. The pure Al5 structure of epitaxially grown NbJGe
Epitaxial growth of Nb,Ge on Nb,Ir (25at.% Ge) deposited on Nb,Ir at T, = 87s”Cand r = 13.9A/WC.
The term “epitaxy” usually refers to the growth of an
oriented film on a single-crystal substrate where both the extend the boundary of the Al5 phase in the Nb-Ge
substrate and film have compatible crystalline structures system to include the 3 : 1 stoichiometry. Our experi-
with nearly equal lattice constants. The idea of growing a mental results fully confirm this. The Reed photograph in
polycrystalline film “epitaxially” on a polycrystalline Fig. 11 contains the diffraction pattern of a Nb,Ge
substrate is novel and requires some elucidation, especi- (25 at.% Ge) film grown epitaxially on a polycrystalline
ally in view of the fact that all evaporated metallic films NbJIr substrate. The Nb3Ge has grown in pure Al5
usually have a poly~rystalline structure independently of structure. There is no sign of any diffraction arcs
the nature of the substrate. Obviously epitaxial growth of belonging to the Nb5Ge3 or to the “mystery” phase.
a polycrystalline film becomes meaningful, only if one is These arcs are usually encountered in Nb,Ge films
dealing with the formation of a binary alloy of a given deposited directly on sapphire as depicted, for
crystalline structure by the codeposition of two elements comparison, in the Reed photograph Fig. 3. Thus, epi-
on the substrate. The required binary phase has a crys- taxial growth of the A 15 Nb& is indeed feasible with the
talline structure which matches that of the substrate; and boundary of the A 15 now extending up to - 26 at.% Ge,
at the required composition it would form in association well beyond its equilib~um value of 19at.%.
with a second phase of a different crystalline structure. The next question to consider is the effect of Nb,Ir
Since the deposit is initially forced to be related epitaxi- lattice constant on that of the epitaxially grown Nb,Ge.
ally to the substrate so as to achieve a minimum in the By properly choosing the evaporation rates of the Nb
free energy, growth conditions of the required phase are and the Ir we prepared two sets of samples where in one
more favourable than those of the competing second set the intent is to expand, and in the other to contract the
phase. If the presence of the matching substrate provides lattice of the Nb,Ge as depicted by the two dashed
a sufficient reduction in the energy of formation of the lines in Fig. 12. The Nb,Ir substrate under the Nb,Ge
required phase, then growth of the competing phase may was not analyzed directly however the microprobe
be totally prohibited. Thus, the first question encoun- analysis of Nb,Ir films alone always was within 1% of
tered in the present study is whether epitaxy could the value predicted from the previously calibrated
 Epitaxial growth of Nb3Ge on Nb& and NbSRh 535

evaporation. Measurement of the lattice constants of the 24 i-

epitaxially-grown Nb,Ge films indicated, as shown in
Fig. 12 that the Nb,Ge lattice neither expanded nor 22
contracted but rather within the experimental scatter
grew in both cases with its “natural” lattice constant.
Since the only “natural” lattice constant of a given phase 20
is that obtained under thermodynamic equilibrium
conditions, it is worth checking whether our measured 18
values lie sufficiently close to a straight line extrapolation
of the bulk values determined by MiillerlS]. As shown in 16
Fig. 12 this fit is extremely good. Furthermore, we see that YK
the boundary of the Al5 has certainly increased from
44 -
19at.% to - 26 at.% Ge. That the lattice constants of
these Nb,Ge films lie on an extrapolation of the equili-
brium values is expected on general grounds, if epitaxy is 12 -
indeed taking place in the manner discussed in the in- “I
troduction to this section. The presence of the Nb,Ir io -
%
films in a pure Al5 phase have provided the proper
environment in which the minimum in the free energy is
8-
achieved when the Nb,Ge forms also in a pure A15
phase. As usually encountered iu the equilibrium range
of homogeneity of a given phase, the lattice constant 6-
varies linearly with composition. I b .l ”

The effect of epitaxy on enhancing the growth of the 41 I I I I I I I

46 18 20 22 24 26 28 30
A 15 phase is further demonstrated by the resistive tran-
sition into the superconducting state. This is shown in at. % Ge
Fig. 13 for two sets of samples, one directly deposited on Fig. 13. Effect of epitaxy on T,: the lower values are from
Nb,Ge deposited directly on sapphire, the higher values from
sapphire, while the other is epitaxially grown on Nb,Ir Nb,Ge grown epitaxially on Nb,Ir; both under same vacuum
under otherwise identical vacuum conditions, substrate conditions and at T, = 875°C and r = 13.9Alsec. At 25at.%Ge,
temperature and deposition rates. The enhancement of epitaxy causes T, to rise by -9°K and the transition width to
the growth of the Al5 produced by epitaxy is accom- decrease by a factor of a 3.
panied, at 25 at.% Ge, by - 9°K rise in onset temperature 5.18 t-
and a factor of four reduction in the transition width. T,
for epitaxialiy-grown films peaks around 25at.%Ge 5.47
where a minimum in the transition width of -0.YK is
also attained.
5.16
A second attempt to compress the Nb,Ge lattice was
carried out for samples having Nb-rich composition. The
Nb,Ir lattice constants together with those of the epi-
taxially grown Nb,Ge are shown in Fig. 14. Again we

5.18 -
‘\
547 - THERMODYNAMIC
EQUILIBRIUM
&I,
5.42 ! I t I I I I I
5.16 - 24 26 28 30
46 16 20 22

5.15 - at.% Ge
A Fig. 14. Lattice constants of a Nb-rich sample epitaxially grown
5.l4 - on NbJr of lattice constant lying on the dashed line. The
epitaxial film lattice constant lies & an ex~poiation of the
5~3 . equilibria values (compare with Fig. 1).

find that compression could not be achieved and that the

I I t I 1 f Nb,Ge lattice constants fit the thermodynamic equili-
16 18 20 22 24 26 28 30 32 34 brium values. A summary of the dependence of the
at.% Ge
epitaxial Nb,Ge lattice constant on composition is
Fig. 12. Attempt to expand or to contract the lattice of Nb,Ge by shown in Fig. 15, while the corresponding T, onset is
growing it on a NbJr substrate having a lattice constant which shown in Fig. 16. One sees that, within the limits in-
lies on either of the dashed lines. Neither expansion not contrac-
tion takes place. Tbe Nb,Ge relaxes to its natural lattice dicated in Figs. 12 and 14, neither the lattice constant nor
constant: in the single phase region the epitaxial film lattice the T, is particularly sensitive to Nb,Ir substrate lattice
constant lies on an extrapolation of the equilibrium A 15 values. constant. Evidently, the Nb,Ir forces the Nb,Ge to form
 536 A. H. DAYEM et al.

5.16 -

,i , , , , , , ,( ,

16 18 20 22 24 26 28 30 32
L I I I I I
at%Ge I

7c ) 60 50 40 30
Fig. 15. Behavior of lattice constant of Nb,Ge on Nb,Ir vs
ANGLE 28’
at.%Ge for the epitaxially grown films previously discussed.
Single Al5 phase extends from 17 to 26.3at.% Ge. Fig. 17. Diffractometer traces of two Nb,Ge samples epitaxially
grown on NbJr. Both samples have 24 at.% Ge. The Nb;Ir la&
22 - constant is 5.1318, in (a) and 5.165A in (b). The eoitaxial NbqGe
grows in a single Al5 phase with 5.143A‘lktice constant whilk in
(b) the growth is in two distinct A 15phases with lattice constants
20 - 5.143and 5.1778, respectively.

5.18
18 - 0 0
x 5.17
-\
16 - ’ I
5.16 - ‘4, i
T;K
x
14 - 5.15 -
8 f
5.14 -
12 -
x
5.13 -
x
10 - 1
5.12 I ’ ’ ’ ’
x
16 18 20 22 24 26 28 30 32 34
8r at%Ge
Fig. 18. Effect of a large Nb,Ir lattice constant on the epitaxial
growth of Nb,Ge. The substrate (Nb,Ir) lattice constant lies on
61 I I I I I I I the dashed line. The epitaxially grown Nb,Ge has two Al5
16 18 20 22 24 26 28 30 phases, one phase has a lattice constant approximately equal to
al. % Ge that of the substrate lattice constant, while the other A 15 phase
has a fixed lattice constant approximately equal to that of Nb,Ge
Fig. 16. T, onset for Nb,Ge epitaxially grown on Nb,Ir. The
(see Fig. 15).Note that the A 15 boundary for these samples is at
maximum occurs at 25 at.% Ge.
25 at.% Ge.

in an A15 structure with the first monolayers having a constant of Nb3Ge obtained from Fig. 15. Also shown in
lattice constant identical to that of the substrate. In Fig. 18 is a dashed line indicating the estimated lattice
subsequent growth the Nb,Ge lattice gradually relaxes to constant of the Nb,Ir substrate. Since the shift in Nb
its natural state. and Ir evaporation rates required to produce this varia-
tion of the Nb,Ir lattice constant was rather large
EF’FECT OF LARGE Nb,Ir LAlTKE compared to the shifts previously used and since we did
We prepared a set of samples so that the stoi- not make a separate set of Nb,Ir samples with the same
chiometric Nb3Ge grows on a Nb,Ir with a lattice rates adopted here to check its actual composition, we
constant = 5.165& calling for an expansion of the will assume that the variable set of lattice constants
NbsGe lattice by -0.52% which is about five times mentioned above corresponds to the lattice constant of
larger than that selected for one of the samples in Fig. Nb,Ir substrates. Thus, a possible interpretation of the
12. Reed photographs showed that the Al5 phase boun- results shown in Fig. 18 is as follows. Comparison with
dary is at 25 at.% Ge, while the diffractometer traces (see Fig. 15 shows that the Nb,Ge composition which
Fig. 17) showed that almost each Al5 line is split into matches the substrate lattice constant is considerably
two lines of comparable intensity indicating the presence richer in Nb than the composition of the incident vapor.
of two A15 phases. The evaluated lattice constants are It seems likely, therefore, that some Nb-rich regions
shown in Fig. 18. There are two sets of lattice constants, grow in an Al5 structure which matches the substrate
the larger values vary with composition while the smaller while the balance of the material deposited forms as
values remain qonstant approximately equal to the lattice another distinct Al5 NbsGe which is almost
 Epitaxial growth of Nb,Ge on NbJr and NblRh 537

stoichiometric.The validity of this interpretation is rein-

forced by the uniformly high ‘I, onset (> 20K) for the
composition range 23-27 at.% Ge. The determination of
the extent to which the lattice of the Nb,Ge can be
expanded or contracted should await further experiments
in which the Nb,Ge is prepared on a single crystal Nb, Ir.

EPITAXIAL GROWTH OF Nb& ON Nb,Rh

We first prepared one set of samples with both Nb,Rh
and Nb,Ge deposited at T, = 875°C. We found that the
A 15 boundary was at - 22 at.% Ge with a maximum in
I’,. = 21.7” and a width of 2.2”. The rest of samples, at 44
higher Ge concentration, were in the mixed phase region
and hence the evaluated lattice constant showed the t
12
usual large spread and large errors. I-
Two more samples were prepared on Nb,Rh deposited 0
at T, = 400°C. The substrate temperature was then raised 10
to 875°C prior to the deposition of the Nb,Ge. The lattice t
I 0
constant vs composition for these sets is shown in Fig. 6
19. The boundary of the Al5 extends now up to -
t
25 at.% Ge. The values of the lattice constants are about 61 I I I I I I I
0.0035 A smaller than the extrapolated equilibrium 16 18 20 22 24 26 28 30
values. The Reed photographs show that the A 15 Nb,Ge at.% Go
phase is in no way inferior to that obtained on Nb,Ir Fig. 20. Onset 7’, of three Nb,Ge samples grown on Nb,Rh
substrates. The same can be said about the T, of these substrates. The peak appears to be in the vicinity of 24 at.% Ge.
samples shown in Fig. 20, except that the transition
widths, around the 25 at.% Ge composition, are about
three times larger than for samples prepared on Nb,Ir. CONCLUSIONS
As mentioned earlier, the A15 NbzRh occurs always Nb,Ir films deposited on sapphire at 875°C form in a
mixed with the D or with both the o and a-Nb phases. pure A 15 structure in the composition range 20-
The idea of epitaxial growth of A15 Nb,Ge on such a 28.3 at.% Ir with a lattice constant between 5.168 and
substrate is rather hard to accept. But the fact remains 5.125 A. This range of homogeneity of the A I5 phase and
that the lattice constant of the Nb3Rh and the co of the u the film lattice constants agree well with bulk values.
phase are practically equal and hence both phases may Similar agreement between bulk and film behaviors was
participate in encouraging the formation of Al5 Nb,Ge also encountered for Nb,Rh films. These Nb,Rh films,
phase as found experimentally. however, are not suitable for epitaxial growth studies
As far as a study of epitaxial growth is concerned, it is since the films contain 2 or 3 phases including cubic
recommended to use a Nb,Ir substrate rather than a-Nb, Al5 and tetragonal u phases.
NblRh which is extremely difficult to characterize both Nb,Ge films grown epitaxially on Nb,Ir have a pure
in bulk and film forms. A I5 structure in the composition range 17-26.3 at.% Ge
with a lattice constant varying between 5.176 and
5.135A. Thus epitaxy enhances the growth of the Al5
5.16 phase and extends its range to include the 3: I stoi-
chiometry.
NbJGe deposited directly on sapphire has a lattice
constant (Fig. I) considerably smaller than the “natural”
5.16 - P’D lattice constant established by epitaxy (Fig. 15). The
\
f\ presence of vacancies and/or other defects could ac-
5.15 - i count for the shrinkage of the lattice in the absence of
I\D
i \f I epitaxy and might be necessary to stabilize the A15
5.14 - f f
f # EPITAXIAL Nb,Ge phase [ 121.
Epitaxial Nb,Ge films show reproducible properties,

;;I[ ) , , , , , 1,, high T, with a narrow transition width and are practic-
ally insensitive to the ambient vacuum conditions. The
highest T, and minimum transition width occur at
25 at.%Ge. The highest TV observed in our samples is
16 16 20 22 24 26 26 30 32 22YK which by no means exceeds previously observed
at. 3: Go
values. However, no attempt was made to optimize the
Fig. 19. Nb,Ge grown on Nb,Rh. The lattice constants lie on a conditions which may lead to higher T,‘s, viz. substrate
stlzight line at values lower than the equilibrium values. The A15 temperature, growth rate, and low temperature anneal-
phase boundary is at 25.2at.%Ge. ing, all of which could enhance the long range order.
538 A. H. DAYEMet al.

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