An in situ neutron diffraction study of the thermal disproportionation of the Zr2FeD5 system

Article abstract of DOI:10.1016/j.jallcom.2011.02.129

The Zr₂FeD₅ system has been annealed to 680 °C under ultra high vacuum, and studied in situ by neutron diffraction. The system disproportionates through three distinct regions in temperature. Initially, the tetragonal Zr₂FeD₅ (P4/ncc) is retained up to 330 °C, while steadily depleted of D. From 330 °C to 530 °C, a complex multi-phase disproportionation occurs, with the production of cubic ZrD ₂, tetragonal ZrD₂, tetragonal Zr₂FeD ₅ (I4/mcm), and growth of the intermetallic ZrFe₂. At the beginning of the 330-530 °C period, the total atom count from quantitative phase analysis (QPA) indicates the formation of amorphous (a-) Zr ₅₆Fe₄₄. By 530 °C, QPA and peak breadth analysis indicate that ca. 2/3rd of the sample is consumed as very small nanocrystals (<150 coherence length) of strained ZrD₂. From 530 °C to 680 °C, the cubic ZrD₂ is almost entirely consumed and depleted of D to form the final mixture of the intermetallic phases Zr₃Fe and ZrFe₂. QPA of the final intermetallic mixture yields a Zr:Fe ratio greater than that observed in either the arc melted alloy or the initial Zr ₂FeD₅ deuteride, indicating that a ca. Zr ₇₁Fe₂₉ amorphous component was present in the initial arc melted alloy. According to the total atom count by QPA, crystallisation of the Fe richer amorphous Zr₅₆Fe₄₄ phase formed at 330 °C begins at ca. 530 °C, and later by 680 °C, all amorphous phases have completely crystallised to yield a 70.77:26.75:2.47 mol.% mixture of Zr ₃Fe:ZrFe₂:ZrD_2-x.

Full text of DOI:10.1016/j.jallcom.2011.02.129

Journal of Alloys and Compounds 509 (2011) 5515–5524
Contents lists available at ScienceDirect
Journal of Alloys and Compounds
journal homepage: www.elsevier.com/locate/jallcom
An in situ neutron diffraction study of the thermal disproportionation of the
Zr FeD system
2
5
a,b,∗
, L.K.W. Pittb,c, H. Fjellvåg , B.C. Haubackb
c
M.P. Pitt
a
Queensland Micro and Nanotechnology Centre, Griffith University, Brisbane 4111, Australia
Institute for Energy Technology, P.O. Box 40, 2027 Kjeller, Norway
Centre for Materials Science and Nanotechnology, Department of Chemistry, University of Oslo, P.O. Box 1126, 0318 Blindern, Norway
b
c
a r t i c l e i n f o
a b s t r a c t
◦
2 5
The Zr FeD system has been annealed to 680 C under ultra high vacuum, and studied in situ by neutron
Article history:
Received 11 February 2011
Received in revised form 21 February 2011
Accepted 23 February 2011
Available online 1 March 2011
diffraction. The system disproportionates through three distinct regions in temperature. Initially, the
◦
◦
◦
tetragonal Zr2FeD5 (P4/ncc) is retained up to 330 C, while steadily depleted of D. From 330 C to 530 C,
a complex multi-phase disproportionation occurs, with the production of cubic ZrD2, tetragonal ZrD2,
◦
tetragonal Zr2FeD5 (I4/mcm), and growth of the intermetallic ZrFe2. At the beginning of the 330–530 C
period, the total atom count from quantitative phase analysis (QPA) indicates the formation of amorphous
Keywords:
Intermetallics
Nuclear reactor materials
Metal hydrides
Phase transitions
Neutron diffraction
◦
(
a-) Zr56Fe44. By 530 C, QPA and peak breadth analysis indicate that ca. 2/3rd of the sample is consumed as
very small nanocrystals (<150 A˚ coherence length) of strained ZrD2. From 530 C to 680 C, the cubic ZrD2
◦
◦
is almost entirely consumed and depleted of D to form the final mixture of the intermetallic phases Zr3Fe
and ZrFe2. QPA of the final intermetallic mixture yields a Zr:Fe ratio greater than that observed in either the
arc melted alloy or the initial Zr2FeD5 deuteride, indicating that a ca. Zr71Fe29 amorphous component was
present in the initial arc melted alloy. According to the total atom count by QPA, crystallisation of the Fe
◦
◦
◦
richer amorphous Zr56Fe44 phase formed at 330 C begins at ca. 530 C, and later by 680 C, all amorphous
phases have completely crystallised to yield a 70.77:26.75:2.47 mol.% mixture of Zr3Fe:ZrFe2:ZrD2−x.
©
2011 Elsevier B.V. All rights reserved.
1
. Introduction
The Zr1 Fex system is an important alloy system for the nuclear
by the absorption of hydrogen. The Zr Fe intermetallic is typi-
2
cally metastabilised at room temperature by rapid quenching from
the melt. There exist many quenching techniques, from extremely
−x
6
industry, based on the tritium gettering properties of the Zr Fe
rapid at 10 K/s, such as melt spinning, enabling the production of
2
intermetallic [1–6]. The commercial SAES St 198 alloy typically used
for non-evaporable hydrogen isotope gettering is predominantly
crystalline Zr Fe, which forms a very stable hydride, Zr FeH5, with
predominantly bulk amorphous alloys [11], to slower techniques,
such as arc melting, typically 10 K/s, producing predominantly
1
crystalline alloys. Due to a large gradient of cooling rate between
2
2
1
6
very low plateau pressure at ambient temperature, ca. 1 Pa [7].
free sample surface (10 K/s) and the Cu hearth (up to 10 K/s), arc
melted samples are typically inhomogeneous, multi-phase [1], and
may retain a minor amorphous component [15], which is why arc
melted samples are typically flipped and re-melted several times
[16–18], or post annealed to ensure homogeneity. Amorphous
Zr Fe has also been proposed as an outer safety layer in high pres-
2
sure hydrogen storage tanks to avoid hydrogen embrittlement and
micro cracking of the tank [8]. Amorphous (a-) phases can also be
formed from the Zr1 Fex system over a very wide Fe composition
−x
range (ca. 0.27 < xFe < 0.79) (mechanically milled alloys) [9,10], and
Zr Fe is quite stable, displaying a crystallisation temperature (Tcryst)
2
◦
>90 at.% Fe in melt spins [11] and thin films [12,13], and are of great
>700 C for mechanically alloyed material [10], and an onset of crys-
◦
interest for their unique magnetic properties.
tallisation of ca. 400–430 C for melt spun material [11,19], defined
◦
The crystalline Zr Fe intermetallic exists in a small high tem-
by a rapid increase in resistivity which peaks at ca. 630 C, indicating
2
◦
perature range in the binary phase diagram, 780–951 C, with a
crystallisation and subsequent growth of Zr Fe crystallites evolves
2
◦
very narrow range of Zr composition, ca. 0.5 at.% Zr [14], being
metastable at room temperature, and only able to be stabilised
over a 200 C span in temperature. Zr Fe crystallises with the CuAl
2
2
structure type, and initially forms very fine and well dispersed
–3 nm nanocrystals [19]. It is reported in [20] that amorphous
2
Zr Fe can absorb a small amount of hydrogen, 0.2–0.7 wt.%H, which
2
increases Tcryst by ca. 40 K. Crystalline Zr Fe can absorb deuterium
2
∗ Corresponding author at: Queensland Micro and Nanotechnology Centre, Grif-
fith University, Brisbane 4111, Australia.
up to a composition of Zr FeD5, which possesses a tetragonal unit
2
cell in space group P4/ncc [18]. The crystalline Zr FeH5 hydride
E-mail address: mark.pitt@gmail.com (M.P. Pitt).
2
0
925-8388/$ – see front matter © 2011 Elsevier B.V. All rights reserved.
doi:10.1016/j.jallcom.2011.02.129

5
516
M.P. Pitt et al. / Journal of Alloys and Compounds 509 (2011) 5515–5524
◦
◦
Fig. 1. Raw neutron diffraction data series from 20 to 680 C, zoomed from 53 to 73 2ꢀ. Reflections from Zr2FeD5, ZrD2, Zr3Fe and ZrFe2 are visible, indicated by the dashed
lines.
forms part of the greater Zr MHx family (M = Ni, Cu, Co, Fe, etc.), all
of which display unique crystallographic properties upon hydrid-
in a shielded furnace at the diffractometer sample position. The silica tube is con-
nected to a Sieverts apparatus, which was used to deuteride the sample at room
temperature under 1 bar, to a composition of Zr2FeD4.7 by manometry. In situ pow-
der neutron diffraction experiments were carried out using the PUS diffractometer
2
ing, such as the dislocated monoclinic Zr CuD
system [21], or
2
4.71
the unique triclinic Zr NiD structure [22].
2
4.8
[28] at the JEEP II reactor at Kjeller, Norway. A total of 15 patterns were collected
Hydrogenation of Zr Fe can produce several phases depen-
from room temperature up to 680 ◦C. Monochromated neutrons with ꢁ = 1.5550 A˚
2
3
4
dent on the temperature used. When charged at 10 to 8 × 10 Pa
were obtained from a Ge (5 1 1) focussing monochromator. The detector unit con-
3
◦
◦
sists of two banks of seven position-sensitive He detectors, each covering 20 in 2ꢀ
(
H2 pressure at <400 C, it is reported in [16] that a single Zr–Fe
◦ ◦
binned in 0.05 steps). Data was collected in the 2ꢀ range 10–130 . Neutron diffrac-
◦
hydride is formed, but that above 400 C, ZrH2 and ZrFe2 are
tion patterns were analysed by the Rietveld method using RIETICA [29]. Diffraction
lineshape profiles were fitted with a full Voigt function, with the instrumental shape
determined by a NIST Al2O3 standard. Quantitative phase analysis (QPA) has been
performed on the diffraction data collected at each temperature.
produced. A Zr FeH4.5 composition is reported in [17] after hydrid-
2
ing at 0.5–5 atm H2 at 250 K. Hydrogenation of the SAES St 198
◦
getter at temperatures >350 C report decomposition of Zr FeDx
2
into ZrD2 and possibly ZrFe2 [23]. It is also suggested in [23]
that formation of ZrD2 may render the getter irreversible. It was
3. Results and discussion
demonstrated in [24,25] for both Zr Fe and Zr Fe respectively,
2
3
that if the intermetallics are deuterided above the critical tem-
perature for disproportionation, and significant quantities of ZrD2
are formed, the material can be recovered at <973 K by releas-
ing D from ZrD2 under ultra high vacuum. In both cases, the
original Zr Fe and Zr Fe compositions are completely re-formed
By visual inspection of our raw neutron diffraction data series
zoomed over 53–73 2ꢀ in Fig. 1, we can observe most of the phases
◦
involved in the Zr FeD5 disproportionation phase transition. The
2
characteristic Zr FeD
58.00 and 60.24 2ꢀ can be observed to shift to higher 2ꢀ as a
5−x
reflections (1 4 1)/(1 2 3) and (2 4 0) at ca.
2
◦
2
3
at high temperature. Such reformation represents an example
of the so called Hydrogenation-Disproportionation-Desorption-
Recombination (HDDR) process [26], otherwise referred to as
function of temperature, indicating unit cell contraction with tem-
◦
perature, suggesting the release of deuterium up to at least 330 C,
◦
after which the phase is no longer easily observed. Past 330 C,
vacuum annealing recovery. As such, any capacity decay in Zr Fe
◦
2
we can observe the formation of cubic ZrD2 ((2 2 0) at 54.71 2ꢀ),
◦
based tritium getters induced by operating them above 350–400 C
which is initially very broad, suggesting growth starts from small
and producing ZrT2 can be recovered by this annealing process.
◦
nanocrystals. By 530 C, the formation of Zr Fe ((2 4 1)/(1 7 1) at
3
Recent diffraction studies of the disproportionation of Zr FeH5 have
◦
◦
2
66.49 2ꢀ) and ZrFe2 ((1 1 5) at 69.27 2ꢀ) can be observed. Fur-
◦
identified ZrFe2 and ZrH1.8 after absorption at 500 C, and Zr Fe,
3
ther Rietveld refinement reveals the presence of tetragonal ZrD .
2
◦
ZrFe2 and ZrH1.8 after desorption to 800 C [27]. The desorbed
◦
◦
◦
Fig. 2(a)–(c) shows the Rietveld fits at 20 C, 380 C, and 680 C. Fit
sample could not be reproportionated. In [4], it was noted that dis-
2
quality varies from ꢂ = 1.96–2.96.
proportionation of Zr Fe alloys (including SAES St 198) occurred
2
Figs. 3 and 4 give an overview of the complete Zr FeD5 dis-
2
◦
during hydriding above 400 C, and that ZrHx formation could be
proportionation phase transition, with Fig. 3 showing the global
crystalline D/Zr ratio (by Rietveld QPA) of the sample as a function of
temperature, and Fig. 4 showing the crystalline phase proportions
as a function of temperature. As indicated in each figure, the phase
transition can be broken into three distinct regions of tempera-
observed.
In this paper, we follow the disproportionation mechanism of a
fully deuterated sample of Zr FeD5 under ultra high vacuum as a
2
function of temperature by in situ neutron diffraction.
ture, characterised by (i) the depletion of D from tetragonal Zr FeD5
2
◦
2.
Experimental procedure
from 20 to 330 C, (ii) a complex multiple phase region from 330 to
◦
5
30 C, where P4/ncc Zr FeD5−x transforms to I4/mcm Zr FeD5−x,
nanocrystalline cubic ZrD2 appears and consumes ca. 2/3rd of the
sample, and ZrFe2 alloy growth begins, and (iii) an alloy growth
2
2
Zr lumps (99.8% purity) and Fe lumps (99.95% purity) were obtained from Good-
fellow. The elements were Ar arc melted in the stoichiometric ratio 2:1 in a water
cooled Cu hearth. The button was flipped and re-melted six times to obtain homoge-
◦
◦
period from 530 to 680 C, where significant Zr Fe growth occurs
nous Zr2Fe. The alloy was crushed under Ar, and activated at 450 C under ultra
3
high vacuum for 1 h. The alloy was packed in a 6 mm silica-glass tube, and placed
by ZrD consumption, yielding the final products Zr Fe and ZrFe . In
2
3
2

M.P. Pitt et al. / Journal of Alloys and Compounds 509 (2011) 5515–5524
5517
◦
◦
◦
Fig. 2. Rietveld fits from top to bottom, 20 C, 380 C and 680 C. Open circles represent the raw data, with the solid line through them representing the calculation. The
◦
◦
difference profile is the solid line below the pattern. Reflection markers, top to bottom 20 C: P4/ncc Zr2FeD5−x, ZrFe2, and I4/mmm ZrD2; 380 C: P4/ncc Zr2FeD5−x, ZrFe2,
I4/mcm Zr2FeD5−x, Fm3m ZrD2−x and I4/mmm ZrD2; 680 C: ZrFe2, Fm3m ZrD2−x and Zr3Fe.
◦
the following discussion, we will study the phase transition in detail
in each of the three characteristic regions of temperature defined
above, and follow the behaviour of the unit cell and D occupancy of
each phase, and describe how the various phases react with each
other.
distances. In [30], it is observed that at lower D content, the P4/ncc
Zr FeD5−x phase transforms to the tetragonal I4/mcm Zr FeD5−x
phase, and the D ordering is removed, and partial occupancy of
the 32m and 16l (Zr4 tetrahedra) sites occurs. In this sense, there
are four sites available to D, 32m = 16g + 16g, 4b, 16l = 16g, and 16k
2
2
(
Zr Fe tetrahedra). For example, all four of these sites are par-
2 2
3
.1. Deuterium depletion from tetragonal P4/ncc Zr FeD5 from 20
tially occupied with D in the I4/mcm Zr2NiD4.4 structure type [22].
The 16k position has not typically been observed to take D for the
2
◦
to 330 C
I4/mcm Zr FeD5−x phase [18], and as such, we utilise four poten-
2
In this region of temperature, the starting P4/ncc Zr FeD5 phase
has a molar proportion of ca. 91.4%, with minor phases, tetrago-
tial sites in the D depleted P4/ncc Zr2FeD5−x structure, referred to as
D1 (4b Zr4), D2 (16g = 32m/2 Zr3Fe), D3 (16g = 32m/2 Zr3Fe), and D4
(16g = 16l Zr4). To avoid short D–D distances between face sharing
tetrahedra of the split 16g sites (D2 and D3), any occupancy (occ)
of the typically empty D3 site must be less than any occupancy
decrease of the typically full D2 site, i.e. occD3 ≤ 1-occD2.
2
nal I4/mmm ZrD (6.1 mol.%) and cubic Fd3m ZrFe (2.5 mol.%). The
2
2
phase proportion of Zr FeD5 does not change significantly until
2
◦
after 280 C. Initially, the room temperature P4/ncc Zr FeD5 phase
2
possesses a completely ordered D sublattice, with full occupation
of 4b (Zr tetrahedra) and 16g (Zr Fe tetrahedra) sites. The 16g site
As can be noted from Fig. 3, there is a significant loss of D from
4
3
◦
is one split half of the 32m site in the parent Zr Fe structure, and
the sample of ca. 1.6 D/Zr ratio up to 330 C. For the P4/ncc Zr2FeD5−x
2
as discussed in [18], only one half of the original 32m site is occu-
pied with D, which avoids face sharing tetrahedra and short D–D
phase, this is equivalent to a change in composition of Zr2FeD5
to ca. Zr2FeD1.8. No D release occurs from the minor tetragonal

5
518
M.P. Pitt et al. / Journal of Alloys and Compounds 509 (2011) 5515–5524
◦
Fig. 3. The crystalline D/Zr ratio (by Rietveld QPA) of the sample as a function of temperature from 20 to 680 C.
I4/mmm ZrD phase during this period, so we can expect significant
Zr Fe) sites of the P4/ncc Zr FeD5−x structure are significantly
2
3
2
reductions in D site occupancies in the P4/ncc Zr FeD5−x structure,
reduced to ca. 40% and 8% respectively from their originally full
2
◦
consistent with the shift to lower d-spacing of the Zr FeD5−x reflec-
occupancies. From ca. 150 C onward, the D3 (16g = 32m/2 Zr Fe)
2
3
tions in Fig. 1. Fig. 5 shows the D occupancies of every phase during
and D4 (16g = 16l Zr ) sites become partially occupied to ca. 20% of
4
◦
the complete phase transition. The D1 (4b Zr ) and D2 (16g = 32m/2
capacity, with the D3 site falling back to 8% occupancy by 330 C,
4
◦
Fig. 4. Crystalline molar phase proportions of all known crystalline phases as a function of temperature from 20 to 680 C.

M.P. Pitt et al. / Journal of Alloys and Compounds 509 (2011) 5515–5524
5519
◦
Fig. 5. Deuterium occupancies of all known D containing phases as a function of temperature from 20 to 680 C.
and the D4 site continuing to increase into the complex multi-phase
period. From 280 C to 330 C, the molar proportion of the P4/ncc
and an increase in the tetragonal ZrD proportion to 13.4 mol.%, and
ZrFe2 remaining at ca. 1.8 mol.%. Fig. 6 shows the variation of the
2
◦
◦
Zr FeD5−x phase is reduced from 89.2% to 76.6%, commensurate
P4/ncc Zr FeD5−x lattice parameters. Both the basal and the pris-
2
2
◦
◦
with the formation of nanoscopic cubic ZrD2 by 330 C (8.2 mol.%),
matic unit cell parameters show a smooth decrease up to 330 C,
◦
Fig. 6. Variation of the P4/ncc Zr2FeD5−x lattice parameters from 20 to 480 C.

5
520
M.P. Pitt et al. / Journal of Alloys and Compounds 509 (2011) 5515–5524
consistent with the continuous loss of D from the cell as a function
of temperature. According to QPA of all observed crystalline phases,
the temperature range from 20 to 330 C can be characterised by
phase (I), and increasing quantities of both the cubic and the tetrag-
onal ZrD2 phases can be described quantitatively as:
◦
◦
(
330–380 C) Zr FeD
(P) → 0.17Zr FeD
2
1.76
2
1.78
two reaction types, the first describing the simple loss of D from
◦
Zr FeD5 from room temperature up to 280 C:
+ (P) 0.32Zr FeD
(I) + 0.38ZrD (c) + 0.13ZrD2
2
2
0.60
2
◦
(
20–280 C) Zr FeD5 → Zr FeD
+ 1.33D (g)
(1)
+ (t) 0.12D2(g)
(3)
2
2
2.35
2
It can be noted that the discrepancy of Zr and Fe atoms in the
The second reaction describes the formation of nanocrystals of
total atom count is larger again than reaction (2), with only 75% and
cubic (c) ZrD2 and an increase in the tetragonal (t) ZrD2 proportion
◦
◦
50% of the Zr and Fe accounted for. From 380 to 530 C, the I4/mcm
from 280 to 330 C, by consumption of a small amount of the P4/ncc
◦
Zr FeD5−x phase is depleted, and is no longer visible after 530 C.
Zr FeD5−x phase:
2
2
◦
After 480 C, the P4/ncc Zr FeD5−x phase is no longer visible, and
2
◦
(
280–330 C) Zr FeD
→ 0.85Zr FeD
+ 0.1ZrD2
the proportion of the ZrFe alloy phase can be observed to increase
2
2.35
2
1.76
2
at the expense of the I4/mcm Zr FeD5−x phase. The consumption of
2
+
(c) 0.05ZrD (t) + 0.29D (g)
(2)
2
2
◦
the majority of the I4/mcm Zr FeD5−x phase from 430 to 530 C can
2
be described as:
It can be noted from reaction (2) that only ca. 93% of the Zr
and 85% of the Fe is accounted for. This implies the dispersion of
a minor fraction of the sample to an amorphous Zr1 Fex phase.
◦
(
430–530 C) Zr FeD
(I) + 0.21Zr FeD1.78(P) + 0.43ZrD2
2
0.62
2
−x
→
(t) 0.20Zr FeD0.62(I) + 0.22ZrD1.45(c) + 0.79ZrFe
2 2
The presence of amorphous Zr1 Fex phases in the sample is dis-
−x
cussed below (in Section 3.4) with regard to the total crystalline
atom count by QPA.
+
0.42Zr Fe + 0.71D (g)
(4)
3
2
On balance, 95% of the Zr atoms are accounted for, and the Fe
◦
atom count has increased by ca. 44%, indicating that crystallisation
3
.2. Complex multi-phase period from 330 to 530 C
◦
of an amorphous Zr1 Fex phase has occurred by 530 C, releas-
−x
ing Fe to assist in the growth of the alloy phases ZrFe2 and Zr Fe.
3
Inspection of Fig. 4 shows that this period in temperature
is strongly multi-phase, and dominated by the appearance and
growth of cubic ZrD , which at its maximum consumes ca. 2/3rd
of the sample. Growth of the alloy phase ZrFe2 also occurs dur-
Inspection of Fig. 6 shows that the smooth unit cell contraction of
◦
the P4/ncc Zr FeD5−x phase up to 380 C is significantly affected by
2
2
the presence of other phases in the complex multi-phase period
from 330 to 530 C. While the lattice parameters of the I4/mcm
◦
◦
ing this period, with the appearance of Zr FeD
by 530 C. The
3
6.7−x
◦
Zr2FeD5−x phase can be observed to evolve smoothly from the
P4/ncc Zr FeD5−x phase disappears by 480 C, and formation of the
2
◦
◦
P4/ncc Zr2FeD5−x phase after 380 C, the basal axis of the P4/ncc
I4/mcm Zr FeD5−x phase occurs at 380 C, which is also subsumed
2
◦
Zr2FeD5−x phase is significantly stretched to a magnitude similar to
by 530 C. Inspection of Fig. 3 suggests that only a minor amount of
the ZrFe2 unit cell dimension, consistent with a strong increase of
D comes out of the crystalline phases in the sample between 380
and 480 C, which is inconsistent with previously reported thermal
◦
◦
the ZrFe2 proportion after 430 C. Both of the tetragonal Zr2FeD5−x
◦
exist up to ca. 500 C with quite a low D concentration, and the
I4/mcm Zr FeD5−x phase is slightly more stable by ca. 50 C.
desorption spectra showing continuous hydrogen release occur-
ring during this period [30]. To follow the behaviour of each phase
in this complex region of temperature, we break the discussion
into three sections, concerning (i) the tetragonal deuterides P4/ncc
Zr FeD5−x and I4/mcm Zr FeD5−x, and their distribution of D sites
and unit cell variation, (ii) the formation of significant quantities
of the mono deuterides Fm3m ZrD2 and I4/mmm ZrD , and the
very broad peak shape of cubic ZrD , and (iii) the appearance and
and Fd3m ZrFe , and
phase at high temper-
ature. Concluding these sections, we will discuss the presence of
amorphous Zr1 Fex phases in the sample, as indicated by the total
◦
2
3.2.2. Mono deuterides Fm3m ZrD2 and I4/mmm ZrD2
The tetragonal I4/mmm ZrD2 phase is initially present as a
minority phase in the sample after deuteration, at ca. 9 mol.%,
and does not show any significant change in proportion until past
2
2
−x
2
2
◦
2
80 C, commensurate with a decrease in the P4/ncc Zr FeD5−x pro-
2
portion to 76.5 mol.%. The I4/mmm ZrD phase reaches a maximum
of 23.6 mol.% at 380 C, and its growth occurs at the expense of the
P4/ncc Zr FeD5−x phase between 330 and 380 C according to reac-
tion (3). From 380 to 430 C, the I4/mmm ZrD proportion is reduced
2
growth of the alloy phases Cmcm Zr FeD
the weak occupancy of D in the Zr FeD
3
6.7−x
2
◦
3
6.7−x
◦
2
◦
2
−x
by 56%, according to:
crystalline atom count.
◦
(
380–430 C) ZrD (t) + 0.55Zr FeD
(P) → 0.22Zr FeD
2
2
1.78
2
1.78
3
.2.1. Tetragonal deuterides P4/ncc Zr FeD5−x and I4/mcm
2
+
(P) 0.87ZrD1.75(c) + 0.44ZrD (t)
(5)
2
Zr FeD5−x
2
◦
From 330 to 380 C, the occupancies of the D1 (4b Zr ), D2
This equation is not balanced, as reactant and product terms
4
(
16g = 32m/2 Zr Fe) and D3 (16g = 32m/2 Zr Fe) sites of the P4/ncc
have been removed to show only the phase formation contributing
3
3
Zr FeD5−x structure are reduced to zero, while the D4 (16g = 16l Zr )
to tetragonal I4/mmm ZrD2 depletion. The fully balanced 10 term
2
4
occupancy rises to a maximum of ca. 46% occupancy, commensu-
equation shows that 94% of the Zr and 82% of the Fe is accounted
◦
rate with a very large reduction in proportion of P4/ncc Zr FeD5−x
for at 430 C, again suggesting the formation of amorphous Zr
Fex
2
1−x
◦
from 76.5 to 13.1 mol.%. After 480 C, the P4/ncc Zr FeD5−x phase is
phases. The tetragonal I4/mmm ZrD2 phase does not lose any D at
2
no longer visible. The large reduction in P4/ncc Zr FeD5−x content
any temperature. The cubic Fm3m ZrD
between 380 and 430 C, and is formed directly from depletion of
phase grows by 20 mol.%
2
2−x
◦
◦
from 330 to 380 C corresponds to the formation of the I4/mcm
◦
Zr FeD5−x phase, and significant rises in proportion of both the
tetragonal I4/mmm ZrD2 and P4/ncc Zr FeD5−x. After 430 C, the
2
2
cubic and the tetragonal ZrD2 phases. Fig. 5 shows that the D1 (4b
tetragonal I4/mmm ZrD phase is no longer observed. Inspection of
2
◦
Zr ) site in the I4/mcm Zr FeD5−x structure shows the highest occu-
Fig. 4 shows that from 430 to 530 C, the cubic Fm3m ZrD
phase
4
2
2−x
pation of ca. 35%, while the only other occupied sites, D2 and D4,
reaches maximum content where it consumes ca. 2/3rd of the sam-
◦
◦
are both only weakly occupied at <10%. After 530 C, the I4/mcm
ple at 480 C. During this period, the P4/ncc Zr FeD5−x and I4/mcm
2
Zr FeD5−x phase is no longer visible. The consumption of P4/ncc
Zr FeD5−x phases are depleted, and ZrFe alloy growth starts past
2
2
2
◦
◦ ◦
430 C, and Zr Fe can be observed at 530 C. Growth of the cubic
3
Zr FeD5−x (P) from 330 to 380 C to form the I4/mcm Zr FeD5−x
2
2

M.P. Pitt et al. / Journal of Alloys and Compounds 509 (2011) 5515–5524
5521
◦
compression and prismatic expansion from 330 to 430 C, with ꢃa
ca. −0.06 A˚ , and ꢃc ca. +0.18 A˚ , an order of magnitude difference
compared to the 0.01 A˚ change in the cube side. As such, the cubic
Fm3m ZrD2 phase is likelymechanicallystifferthanthe tetragonal
−x
I4/mmm ZrD2 phase, also indicated by the c-axis of the tetragonal
I4/mmm ZrD phase being stretched toward the cube side, which is
2
the only unit cell dimension (c = 4.5730 A˚ ) that is coherently close
in dimension to the cubic Fm3m ZrD2 unit cell (a = 4.7843 A˚ ) in
−x
◦
the 330–430 C range. A mechanical couple between the tetrago-
nal and cubic ZrD2 phases is also implied by the growth of cubic
ZrD2 from tetragonal ZrD according to reactions (4) and (5). The
−x
2
contribution from the initially very broad reflections of the cubic
Fm3m ZrD2 phase are shown in a zoomed section of the neutron
◦
◦
◦
diffraction pattern at 380 C, over 40–90 2ꢀ, in Fig. 8. By 530 C,
the cubic Fm3m ZrD2 reflections have sharpened considerably.
−x
Lineshape analysis suggests the cubic Fm3m ZrD initially forms on
2
◦
the 15 nm length scale, which by 580 C has increased to >100 nm.
After 530 C, the cubic Fm3m ZrD2 phase proportion begins to
◦
−x
rapidly decrease.
3
.2.3. Alloy phases Cmcm Zr Fe and Fm3m ZrFe
3
2
The cubic alloy ZrFe is present as a minority phase of ca. 2 mol.%
2
in the original arc melted sample, and is present in the sample over
◦
the entire temperature interval of 20–680 C. Fig. 4 shows ZrFe
2
◦
growth begins after 430 C, and the growth of ZrFe from 430 to
30 C follows reaction (4). As the tetragonal I4/mmm ZrD2 phase
2
◦
5
Fig. 7. Variation of the lattice parameters of cubic Fm3m ZrD2−x and tetragonal
I4/mmm ZrD2 phases from 20 to 530 C.
is likely mechanically coupled and directly consumed by the cubic
◦
Fm3m ZrD2 phase, growth of ZrFe2 occurs directly by depletion
−x
of the P4/ncc Zr FeD5−x and the I4/mcm Zr FeD5−x phases, mostly
2
2
the I4/mcm Zr FeD5−x phase, as the P4/ncc Zr FeD5−x proportion
2
2
◦
◦
remains below ca. 5 mol.% from 430 to 480 C. From 480 to 530 C,
Fm3m ZrD2 phase can be described by reaction (4), where three
phases are depleted to form two. From 380 to 530 C, the D content
growth of ZrFe2 is mostly from I4/mcm Zr FeD5−x depletion. From
−x
2
◦
the atom balance of reaction (4), it is also likely that an amorphous
◦
of the cubic Fm3m ZrD2 phase is reduced to 72.6% (see Fig. 5).
Zr Fex phase is crystallising up to 530 C, and an excess of Fe from
−x
1−x
Inspection of Fig. 7 shows that the cubic Fm3m ZrD2 unit cell
this amorphous phase may also contribute to the growth of ZrFe .
From 480 to 530 C, the proportion of the cubic Fm3m ZrD
−x
2
◦
dimension appears little affected by this D loss (only a very small
phase
2−x
◦
◦
compression can be observed from 430 C onward, consistent with
decreases by ca. 4 mol.%, and by 530 C, the orthorhombic Cmcm
D loss), and the cell may be dilated in the presence of other phases.
Zr Fe phase can be observed in the neutron diffraction pattern at a
proportion of 10.4 mol.%. From 480 to 530 C, the growth of ZrFe2
3
◦
In contrast, the tetragonal I4/mmm ZrD phase shows distinct basal
2
◦
◦
Fig. 8. Neutron diffraction pattern at 380 C zoomed from 53 to 73 2ꢀ, showing the very broad reflections from cubic Fm3m ZrD2−x. Reflection markers from top to bottom
are P4/ncc Zr2FeD5−x, I4/mcm Zr2FeD5−x, cubic Fm3m ZrD2−x, tetragonal I4/mmm ZrD2 and ZrFe2.

5
522
M.P. Pitt et al. / Journal of Alloys and Compounds 509 (2011) 5515–5524
and Zr Fe alloy phases can be described as:
mal rise, consistent with typical thermal expansion in the absence
of D.
3
◦
(
480–530 C) Zr FeD
(I) + 0.17Zr FeD
2
0.62
2
1.78
+
(P) 4.29ZrD1.66(c) → 0.32Zr FeD0.62(I) + 4.02ZrD1.45
2
3
.4. Amorphous Zr1 Fex phases during thermal
−x
+
(c) 0.46ZrFe + 0.42Zr Fe
(6)
2
3
disproportionation of Zr FeD
5−x
2
Reactant terms have been removed to show only the phase for-
For reactions (2)–(8) described above, we have consistently
observed a discrepancy in the balance of Zr and Fe atoms. For reac-
tions (2), (3) and (5) covering 280–430 C, we observe a loss of Zr
mation contributing to the growth of ZrFe2 and Zr Fe, the majority
3
of which occurs by depletion of the I4/mcm Zr FeD5−x phase, with
2
◦
a minor amount of Zr being sourced from cubic ZrD2 . The atom
−x
and Fe atoms from the observed crystalline phases in the sample.
balance indicates a ca. 6% increase in Zr and a 47% increase in Fe,
again, consistent with the ongoing crystallisation of an amorphous
◦
For reactions (4) and (6)–(8) covering 430–680 C, we observe a gain
in Zr and Fe atoms. Fig. 10 shows the total atom count of Zr, Fe and D
across the entire temperature range of Zr2FeD5 disproportionation.
Table 1 also summarises the change in composition of all crystalline
phases, the temperature range in which they are stable, and what
their unit cell dimensions range over. Inspection of Fig. 10 shows
the temperature range over which Fe recovery occurs is very large,
Zr1 Fex phase. The significant proportion of ca. 10 mol.% Zr Fe
suggests it has nucleated well before 530 C.
−x
3
◦
◦
.3. Zr Fe alloy growth period from 530 to 680 C
3
3
This high temperature period is well characterised by the dom-
inant growth of the alloy phase Zr Fe (see Fig. 4), a moderate rise
◦
◦
covering 200 C from 430 to 630 C. The large temperature range of
3
recovery suggests that several amorphous (a-) Zr1 Fex composi-
−x
in ZrFe2 proportion, and strong depletion of cubic Fm3m ZrD
.
2−x
tions may be present in the sample, and indeed, when summing the
Inspection of Fig. 5 shows that the Zr Fe is weakly occupied by
3
◦
discrepancy from the total (fully crystallised) atom count at 680 C,
D, which can only be sourced from the cubic Fm3m ZrD2 phase,
−x
we can infer that a-Zr71Fe29 must be present in the originally arc
suggesting a morphology of intimate mechanical contact between
the phases, allowing D diffusion that is not blocked by the pres-
ence of ZrFe . There are four D sites available in the Zr FeD
melted sample, and that a-Zr56Fe44 formation occurs from 280 to
◦
4
30 C. It is not unreasonable to expect an amorphous composi-
2
3
6.7
tion in an inhomogenous arc melt close to the Zr Fe composition.
2
structure [30], D1 (4c Zr Fe ), D2 (8f Zr Fe), D3 (8f Zr ), and D4
3
2
3
4
Nor is it surprising to expect the formation of a-Zr1 Fex phases,
−x
(
5
8f Zr ). The D3 and D4 sites show weak occupancy <12% from
4
particularly when the crystalline phases in the phase transition all
have such dissimilar lattice parameters. The most plausible expla-
◦ ◦
80 to 630 C and <5% occupancy of the D2 site at 630 C. This is
consistent with [30], where typically only the Zr4 tetrahedra are
occupied at weak D concentrations. The trigonal bypyramid Zr Fe
nation is that the domination of the cubic Fm3m ZrD2 phase in the
−x
3
2
◦
multi-phase region from 330 to 530 C is the origin of considerable
◦
is not occupied. The Zr Fe structure does not hold any D at 680 C,
3
lattice parameter misfit with other phases, resulting in the spread
and by this temperature, the cubic Fm3m ZrD2
reduced to ca. 2.5 mol.%. From 530 to 580 C, ZrFe2 and Zr FeD
growth occurs mostly from the cubic Fm3m ZrD2
according to:
proportion is
−x
of dislocations and interfacial growth of a-Zr Fex phases. There is
◦
1−x
3
6.7−x
phase
a significant resolvable strain component in the cubic Fm3m ZrD2
−x
−x
lineshape when it is initially forming, and dislocation loop punch-
ing is likely, as is observed for needle like ZrHx precipitates [31].
As such, we would expect considerable microstructure commensu-
◦
(
530–580 C) ZrD1.44(c) + 0.08Zr FeD0.62(I) → 0.86ZrD1.45
2
rate with ZrD2 nucleation, and the formation of a-Zr
Fex phases
−x
1−x
+
(c) 0.10ZrFe + 0.12Zr FeD
+ 0.38D (g)
(7)
2
3
0.34
2
along the grain boundaries. Such features obviously suggest the use
of transmission electron microscopy (TEM) to be vindicated.
The total atom count of D also suggests the presence of D in
This equation has reactant and product terms removed to show
only those phases contributing to ZrFe2 and Zr FeD growth
3
6.7−x
non-crystalline phases. Fig. 3 shows the release of D from crys-
(
the D content in the equation has been balanced to emphasise D
◦
talline Zr FexDy phases is minimal in the 380–480 C temperature
release from the sample). On balance, there is a 9% and 21% increase
1−x
◦
range. The thermal desorption data in [30] clearly show deuterium
in Zr and Fe respectively. This indicates that even at 580 C, there
◦
evolution from Zr FeD5 disproportionation in this temperature
is still amorphous Zr1 Fex crystallising. After 580 C, there is little
2
−x
range. An explanation can be found in [20], where it has been
change in the ZrFe2 proportion, and the growth of Zr FeD
can
3
6.7−x
observed that 0.2–0.7 wt.%H can be solved into a-Zr Fe33. The
be directly attributed to cubic Fm3m ZrD2 depletion as:
67
−x
◦
presence of H in a-Zr67Fe33Hx increases Tcryst by ca. 40 C. Accord-
◦
(
580–680 C) ZrD1.44(c) → 0.05ZrD1.34(c) + 0.97Zr3
ing to the concentration dependent crystallisation temperatures
◦
reported in [10], a-Zr71Fe29 will begin crystallising at ca. 365 C, or
+
Fe 0.69D (g)
2
(8)
◦
ca. 405 C for a-Zr71Fe29Dx, assuming no significant isotope effect.
◦
Reactant and product terms are removed to show the phases
These crystallisation temperatures fall in the 380–480 C range, and
contributing to Zr Fe growth. Here we have the strongest indica-
the obvious inference is that D released from the sample in this tem-
perature range can come from a-Zr71Fe29Dx present in the initially
deuterated arc melt. Crystallisation of a-Zr71Fe29Dx beginning at
3
tion yet of non-crystalline phases in the sample crystallising and
releasing atoms to contribute to the growth of existing crystalline
phases, with a 44% and 41% increase in the Zr and Fe atom count
◦
ca. 405 C is also consistent with the Fe recovery observed from ca.
◦
◦
respectively. At 680 C, we do not yet observe the formation of
430 C onward. The crystallisation of a-Zr71Fe29 and a-Zr71Fe29Dx
Zr Fe, which is below the temperature at which it is observed in
the binary phase diagram, 780 C [14]. Fig. 9 shows the variation
is shown by the dashed lines in Fig. 10. As such, the complex
2
◦
◦
crystalline multi-phase period from 330 to 530 C is further com-
◦
in unit cell parameters of ZrFe2 and Zr FeD
temperature from 530 to 680 C. The absolute change in magni-
tude of these lattice parameters is minor for example compared to
as a function of
plicated by the formation of a-Zr56Fe44 from 280 to 430 C, and the
3
6.7−x
◦
◦
crystallisation of a-Zr71Fe Dx and release of D from ca. 405 C. Zr
29
◦
and Fe recovery after 530 C is also consistent with the crystallisa-
tion of a-Zr56Fe44 starting at ca. 535 C according to [10]. It is also
◦
the variation of P4/ncc Zr FeD5−x lattice parameters in Fig. 6. How-
2
ever, all metrics of the Zr FeD
and fall, commensurate with the weak rise and fall in D occupancy
phase can be observed to rise
clear that the boundaries for crystallisation of a-Zr71Fe29 and a-
3
6.7−x
◦
Zr56Fe44 are very close in temperature to the beginning (330 C) and
◦
across the D2–D4 sites. In contrast, the ZrFe unit cell displays a nor-
end (530 C) respectively of the complex multi-phase period, and
2

M.P. Pitt et al. / Journal of Alloys and Compounds 509 (2011) 5515–5524
5523
◦
Fig. 9. Variation of the lattice parameters of Zr3FeD6.7−x and ZrFe2 from 530 to 680 C.
the presence of a-Zr1 Fex phases may have significant influence
on the temperature dependence of the crystalline disproportiona-
is critical in this regard, and should be carefully assessed for the
presence of a-Zr1 Fex phases by TEM. Regardless of the initial
−x
−x
tion of Zr FeD5. We have not found reports describing a-Zr
Fex
sample state, it appears that a-Zr Fex phase formation occurs
2
1−x
1−x
◦
phases in the literature concerning Zr FeD5 disproportionation,
naturally after 330 C as part of the Zr FeD5 disproportionation
2
2
however, this is the first accurate quantitative description of the
process. This mechanism should be verified and studied further by
TEM.
phase transition. Clearly the initial state of the prepared Zr Fe alloy
2
◦
Fig. 10. Total atom count from all known crystalline phases as a function of temperature from 20 to 680 C. Tcryst of a-Zr71Fe29, a-Zr71Fe29Dx and a-Zr56Fe44 are shown for
reference.

5
524
M.P. Pitt et al. / Journal of Alloys and Compounds 509 (2011) 5515–5524
Table 1
A summary of all crystalline phases occurring during the disproportionation of Zr2FeD5, showing from top to bottom, changes in deuterium composition, the temperature
range in which they are stable, and what their unit cell dimensions range over.
Zr2FeD5−x
P4/ncc
Zr2FeD5−x
I4/mcm
ZrD2
I4/mmm
ZrD2−x
Fm3m
ZrFe2
Fd3m
Zr3FeD6.7−x
Cmcm
Zr2FeD5 → Zr2FeD1.76
Zr2FeD0.62 → Zr2FeD0.60
ZrD2
Minor 20–330
Major 330–430
ZrD2 → ZrD0.91
ZrFe2
Minor 20–430
Major 430–680
Zr3FeD0.34 → Zr3Fe
◦
◦
◦
◦
◦
◦
2
0–480
C
380–530
C
C
C
330–680
C
C
C
530–680
C
◦
◦
a = 6.617(2)–6.95(2) A˚
–
c = 5.39(3)–5.6232(2) A˚
a = 6.512(2)–6.586(1) A˚
–
c = 5.451(1)–5.469(3) A˚
a = 3.436(2)–3.517(4) A˚
–
c = 4.452(6)–4.670(6) A˚
a = 4.755(5)–4.784(5) A˚
–
–
a = 7.087(1)–7.114(3) A˚
–
–
a = 3.354(1)–3.3618(4) A˚
b = 11.035(1)–11.089(2) A˚
c = 8.8325(8)–8.8753(9) A˚
4
. Conclusion
[7] S. Fukada, M. Nishikawa, J. Alloys Compd. 234 (1996) L7.
[
[
8] R. Janot, M. Latroche, A. Percheron-Guegan, Mater. Sci. Eng. B 123 (2005) 187.
9] D. Mishra, A. Perumal, A. Srinivasan, J. Phys. D Appl. Phys. 41 (2008) 215003.
The disproportionation of Zr FeD5 is a considerably com-
2
[10] A.W. Weeber, H. Bakker, Physica B 153 (1988) 93.
plex phase transition, characterised as a function of temperature
[11] X.D. Liu, X.B. Liu, Z. Altounian, J. Non-Cryst. Solids 351 (2005) 604.
[
12] P.T. Korelis, A. Liebig, M. Bjorck, B. Hjorvarsson, H. Lidbaum, K. Leifer, A.R.
Wildes, Thin Solid Films 519 (2010) 404.
by three distinct events, the initial depletion of D from P4/ncc
◦
Zr FeD5−x to 330 C, a complex multi-phase period from 330 to
2
[
13] P. Sharma, H. Kimura, A. Inoue, Phys. Rev. B 78 (2008) 134414.
◦
5
30 C dominated by the growth of cubic Fm3m ZrD , and the
subsequent growth of Zr FeD
[14] F. Stein, G. Sauthoff, M. Palm, J. Phase Equilib. 23 (2002) 480.
[15] A. Inoue, Acta Mater. 48 (2000) 279.
2
−x
◦
from 530 to 680 C. Accompa-
3
6.7−x
[
[
16] F. Aubertin, U. Gosner, S.J. Campbell, J. Phys. F Met. Phys. 14 (1984) 2213.
17] P. Raj, P. Suryanarayana, A. Sathyamoorthy, K. Shashikala, R.M. Iyer, J. Alloys
Compd. 178 (1992) 393.
nying these crystalline transitions, our quantitative analysis infers
two distinct crystallisation events, from a-Zr71Fe29 and a-Zr56Fe44
which correspond closely to the low and high temperature bound-
aries of the complex multi-phase region. Our QPA indicates that
a-Zr56Fe44 has formed naturally as part of the phase transition,
likely a consequence of incoherent matching of lattice parame-
[18] V.A. Yartys, H. Fjellvag, B.C. Hauback, A.B. Riabov, J. Alloys Compd. 274 (1998)
17.
2
[
19] Z. Altounian, E. Batalla, J.O. Strom-Olsen, J.L. Walter, J. Appl. Phys. 61 (1986)
149.
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[
[
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22] M.H. Sorby, H. Fjellvag, B.C. Hauback, J. Alloys Compd. 394 (2005) 107.
ters to the dominant cubic Fm3m ZrD2 phase. TEM is the obvious
−x
choice to verify the a-Zr1 Fex formation at high temperature, by
−x
[23] A. Nobile, W.C. Mosley, J.S. Holder, K.N. Brooks, J. Alloys Compd. 206 (1994)
83.
◦
studying samples quenched from above ca. 330 C. Arc melted
[
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2
presence of amorphous Zr1 Fex phases.
−x
[
[
[
[
2
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