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Mechanical alloying and microstructure of a

Nb–20% V–15% Al alloy

S. Dymek a,*, A. Lorent a , M. Wro bela , A. Dollar b

a Department of Metallurgy and Materials Engineering, University of Mining and Metallurgy,

al. Mickiewicza 30, 30-059 Cracow, Poland bSchool of Engineering and Applied Science, Miami University, Kreger Hall, Oxford, OH 45056, USA

Received 7 November 2001; received in revised form 10 December 2001; accepted 16 December 2001

Abstract

A niobium-based alloy with 20% V–15% Al (at.%) was synthesized by mechanical alloying of elemental

powders. During milling, the splitting of Nb X-ray peaks into two components was observed. Each component

was found to correspond to a niobium solid solution (NbI and NbII) with a different lattice parameter. The

intensities of NbI peaks on X-ray diffraction patterns decreased with the milling time and disappeared completely

after 180 h of milling while the intensities of NbII peaks gradually increased. The powders were hot pressed andmicrostructural and phase analyses of the consolidated material were carried out. The microstructure consisted of

Nb solid solution, Nb3Al-base intermetallic with the A15 crystal structure and dispersoid Al2O3. Also an

unexpected, detrimental, Nb2Al-base s phase was found. The volume fraction of the s phase depended on the

temperature of consolidation. D 2001 Published by Elsevier Science Inc.

Keywords: Mechanical alloying; Niobium aluminides; Intermetallics

1. Introduction

The niobium aluminides are candidate materials

for high-temperature structural applications because

of their high melting points and moderate density [1].

However, despite good high temperature strength, the

severe brittleness of these alloys at temperatures

below 900 C hamper their applications. It was shown

that the ductility of Nb3Al-base alloys can be

improved by reducing the average gain size as well

as by incorporating a ductile phase into the micro-

structure [2]. This ductile phase may be disordered or

ordered solution of Al and ternary elements in niobium (A2 or B2 structures). Also, as was recently

postulated by Cahn [3], it is possible that for strongly

ordered compounds such as Nb3Al, very slight dis-

ordering might have a substantial effect on mechanical

properties. One way of achieving this might be by

ternary ordering of the phase itself. One of the ternary

additions that is considered for this purpose is vana-

dium [4,5]. This research is a continuation of our

study on binary Nb3Al-based alloys [2,6] and deals

with an Nb– Al alloy modified by an addition of

vanadium as a ternary element. The goal of this work is to produce and make a preliminary characterization

of this unique alloy with an emphasis on processing

and phase composition in milled powders as well as in

consolidated material.

1044-5803/01/$ – see front matter D 2001 Published by Elsevier Science Inc.

PII: S 1 0 4 4 - 5 8 0 3 ( 0 2 ) 0 0 1 8 4 - 5

* Corresponding author. Tel.: +48-12-617-2696;

fax: +48-12-617-3190.

E-mail address: [email protected]

(S. Dymek).

Materials Characterization 47 (2001) 375 – 381



2. Material and process selection

The addition of vanadium to the binary Nb–Al

alloy was inspired by the recent works of Horspool et

al. [4] and Tappin et al. [5] who examined a series of Nb – Al – V alloys with different compositions. They

found that Nb–Al–V alloys prepared by transferred-

arc plasma melting consisted of a mixture of two

phases: one, brittle (d), with the A15 structure and

the other, ductile, with the B2 crystal structure. The

volume fraction of a particular phase varied linearly

with the ratio of Al to V content in the alloy; alloys

with Al/V 0.5 exhibited only B2 phase, whereas

those with Al/V 1 were almost completely A15. The

alloys produced by Horspool et al. [4] and Tappin et

al. [5] were ductile in the as-cast condition (single- phase B2 alloys) and also exhibited limited ductility

(about 3% in compression) after homogenization at

1500 C when the d phase nucleated from the B2

alloy. The latter alloy with an Al/V ratio of 0.75

contained 55% vol. of B2 phase and 45% vol. of

A15 phase. In the present study, the Al/V ratio was

also kept at 0.75 and the target composition of the

processed alloy was Nb – 20% V – 15% Al (all com-

positions in this article are expressed in atomic per-

centage unless mentioned otherwise).

The processing route was mechanical alloying(MA) of elemental powders followed by consolida-

tion by hot pressing. It was shown previously [2,6–9]

that such a technique is a feasible route for process-

ing niobium aluminides. The powder metallurgical

technique is one of the best methods for synthesis

Nb – Al alloys because it is very difficult to prevent

Al from evaporation during melting [10]. MA offers

a unique solid-state processing route that does not

require melting of the constituents. Furthermore, the

process, by its nature, introduces insoluble disper-

soids that can substantially alter mechanical, espe-

cially creep, properties. Much of the current work on

MA of intermetallics has been done in shaker mills or

similar laboratory equipment. These kinds of mill

produce a few grams of material and necessarily lead

to fairly restricted characterization of the final prod-

uct. In this research, the Szegvari-type attritor was

used for milling. Such an attritor is capable of pro-

ducing significant quantities of material, sufficient for

extensive characterization.

3. Experimental procedure

Pure elemental powders of constituents (Nb: 99.8

wt.%, Al: 99.5 wt.%, V: 99.5 wt.% pure) were used

as starting materials. The powder diameter was less

than 45 mm ( 325 mesh). The powders were

blended in appropriate amounts to give the nominal

composition of 65% Nb, 15% Al, and 20% V. MA

was carried out in a Szegvari laboratory attritor mill

in an argon atmosphere with the controlled oxygen

level reduced to less than 10 ppm. The milling was

carried out in a sealed stainless steel tank with a3.63-kg (8 lb) charge of 4.76 mm (3/16 in.) stainless

steel balls and a 12:1 ball-to-powder ratio (by

weight). The initial first hour of milling was carried

out at cryogenic temperatures, using liquid nitrogen

as the coolant. During the milling, powder samples

were taken out after 1, 5, 15, 25, 40, 80, and 180 h

(the total milling time), allowing for characterization

of the powder morphology and the progress of MA

process. The powders collected at the end of milling

were sieved through a 45-mm mesh and consolidated

by hot pressing in the argon atmosphere at a pressure of 25 MPa and temperatures of 1300 and

1500 C. The resulting material was fully dense and

free from cracks.

The changes in particle morphology during mill-

ing were examined by scanning electron microscopy.

The phase analysis was performed with CoK a radi-

ation. The X-ray spectra were used for evaluation of

lattice parameters, lattice strains, and X-ray crystallite

size. As a standard, the initial Nb powder annealed at

1000 C for 5 h in vacuum was used. The micro-

structure of the consolidated material was character-ized by light and scanning electron microscopy.

Backscattered electron images utilizing Z -contrast

were used to reveal constituent phases. The phases

were also identified by X-ray diffractometry. The

scanning electron microscopy was supplemented by

the investigation of thin foils in a transmission elec-

tron microscope (TEM). The chemical composition

measurements were performed by energy dispersive

spectroscopy (EDS) on an Oxford-Link system

attached to the microscope. For the quantitative

analysis, a Cliff-Lorimer standardless method for thinsections was used.

Preliminary mechanical testing comprised hard-

ness measurements by a Vickers indenter. The eval-

uation of the fracture toughness K Ic was performed

on indented specimens (with microcracks at indenta-

tion corners) following a procedure described by

Peters [11].

4. Results and discussion

4.1. Structural evaluation of MA powders

Already after 1 h of milling (at cryogenic temper-

ature), particles had more rounded shapes compared

to the initial blend. After 5 h of milling, the particles’

shapes and sizes were characteristic for milled ductile

powders [12] — most particles were round and flat

S. Dymek et al. / Materials Characterization 47 (2001) 375–381376



with particle diameters between 100 and 200 mm

(Fig. 1a). The increase in particle size from the initial

blend was due to the predominance of welding over

fracturing occurring in the relatively soft particles.

Between 5 and 15 h of milling, both particle fracture

and welding occurred; however, the presence of bigger

particles indicates that welding occurred more often

than fracture. The largest particles looked like thosefound after 5 h of milling, but in addition to them,

smaller particles were produced as well. There was a

significant scatter in particle size. After 25 and 40 h of

milling, the particle morphologies were similar. Awide

distribution in particle size was observed as well, and

the shape of particles changed to a more spheroidal

type. However, flat particles were also observed indic-

ating that welding still was at work. After 80 h of

milling, particles became much finer indicating a

predominance of fracturing but the nonuniform par-

ticle surface shows that welding still operated. The 180h of milling produces uniform powder with almost

perfectly spherical particles with smooth surfaces

(Fig. 1b). The average size of powder particle (after

sieving through a 45-mm mesh) was about 6 mm. The

sieved powder was used for consolidation.

4.2. X-ray diffraction analysis of MA powders

The X-ray diffractograms exhibited dependence

on the milling time. The composed figure of diffrac-

tograms after 0, 1, 5, 15, 25, 40, 80, and 180 h isshown in Fig. 2. Peaks from Al, though much weaker

than in the starting powder, were still present till 5 h

of milling and disappeared after 15 h indicating

dissolution of Al in Nb (or V). Peaks from V, on

the other hand, became weaker and weaker with the

milling time but were still present after 40 h of

milling. In general, with the milling time, peaks

became broader and their intensities decreased. The

peak broadening was due to refinement of powder

and due to strains associated with intense plastic

deformation of powder particles. The positions of

peaks also changed with milling time. All peaks

moved right toward larger values of 2q indicating a

change in lattice parameter of the alloy with milling

time. Such behavior was also observed for other mechanically alloyed Nb– Al alloys [7– 9] and is

Fig. 1. Characteristic scanning electron microscope (SEM) images of milled powders: (a) after 5 h, (b) after 180 h.

Fig. 2. The composed figure of X-ray spectra after particular

milling times.

S. Dymek et al. / Materials Characterization 47 (2001) 375–381 377



considered typical for the MA process [12]. In

addition, close to the primary Nb peaks, new satellite

peaks appeared always on the right side of the each

primary peak. The intensity of these secondary peaksincreased with milling time, whereas the intensity of

primary ones disappeared. Such behavior is unique

and was not reported earlier. This is evidence that

during MA, two different Nb-base phases with dif-

ferent lattice parameters coexist in the powder mix-

ture. The phases are labeled NbI and NbII. For both

phases, calculation of lattice parameters, volume frac-

tions, average X-ray crystallite size, and value of

lattice strains was performed. The results are shown

in Table 1. Small differences in lattice parameters for

the NbI phase and continuous decrease in lattice parameters for the NbII phase are apparently associ-

ated with the presence of vanadium. Since atomic

diameters of Nb and Al are almost identical (2.86 and

2.864 A, respectively) and the atomic diameter of V is

much less (2.621 A), one can assume that the NbII

phase is enriched with vanadium while the NbI phase

contains much less vanadium. It appears from the

crystallite size analysis that progress of alloying in the

NbI phase is less advanced (larger crystallites) than in

the NbII phase. Also, the values of lattice strains

demonstrate that the NbII phase corresponds to

smaller particles that are more deformed and more

saturated with vanadium. Such behavior may be

associated with the different diffusion coefficients of

Al and V in niobium. Assuming that the diffusivity of a particular element is proportional to its melting

point, Al diffuses faster and deeper into Nb particles

than V does. Since V diffuses smaller distances, the

average crystallite size of the phase enriched with V

is smaller.

Comparing the behavior of the present alloy and

the alloy Nb–18% Al examined previously [2,6], we

can point out significant differences. The alloy Nb–

18% Al behaved in a ‘‘traditional’’ manner, that is,

the peaks broadened and their intensities decreased.

No satellite peaks were observed. Also, the shifting of peaks from their initial positions occurred much later

(it was observed after 86 h of milling) than in the

presence of vanadium. Though Al dissolves in Nb

relatively fast, the alloy with V has to be milled for a

longer time in order to completely dissolve V in the

Nb-base solid solution.

4.3. Consolidated material

The optical micrographs revealed little about the

microstructure. More information was provided by

Table 1

The dependence of lattice parameters, crystallite size, lattice strains, and volume fraction of NbI and NbII phases on milling time

Milling time (h) 0 1 5 15 25 40 80 180

Lattice parameter NbI (A) 3.304 3.304 3.308 3.308 3.307 3.308 3.303 –

Lattice parameter NbII (A) – 3.293 3.281 3.272 3.262 3.249 3.247 3.247Crystallite size NbI (A) 1091 188 162 142 153 91 69 –

Crystallite size NbII (A) – – 33 37 39 54 56 54

Lattice strain NbI (%) 0.110 0.396 0.451 0.503 0.471 0.741 0.956 –

Lattice strain NbII (%) – – 1.927 1.720 1.615 1.180 1.149 1.180

Volume fraction NbI (%) 100 81 63 39 43 38 21 0

Volume fraction NbII (%) 0 19 37 61 57 62 79 100

Fig. 3. Backscattered electron images of the consolidated material at temperature: (a) 1300 C, (b) 1500 C.




images formed by backscattered electrons in a SEM

(Fig. 3). Utilization of Z -contrast allowed one to

reveal phases with different Nb concentrations, that

is, Nbss (Nb solid solution) and an intermetallic

phase. Since the backscattered electron yield is higher for higher Z (atomic number), the Nbss grains are

brighter. The darker phase corresponds to an inter-

metallic phase (the black particles are oxides).

The X-ray phase analysis showed that the material

consolidated at 1500 C consists mainly of Nbss and d(A15 crystal structure) phases while the material

consolidated at 1300 C contained additionally a s(Nb2Al-based) phase (Fig. 4). The presence of the

Nb2Al phase was unexpected. It is noteworthy, how-

ever, that this phase was also found in the binary

Nb – 18% Al alloy [2]. This finding is only partly con-sistent with work of Tappin et al. [5] who found, in the

alloy with the same composition subjected to anneal-

ing at 1500 C, only two phases: one with the A15 and

the other with the B2 crystal structures. Since no

{100} peaks were observed on X-ray diffractograms,

it appears that the niobium solid solution remains in a

disordered state. However, peaks from Nb were

moved slightly leftward compared to their positions

in the as-milled powder but were still shifted in

relation to the pure Nb position. This shows that

alloying elements, especially vanadium, remain dis-

solved in niobium. This result was confirmed by TEM

studies, which reveal a fine-grained microstructure(about 1 mm) composed of a mixture of Nbss, d and s,

and Al2O3. Small Al2O3 particles embedded in other

phases or appearing on grain boundaries were also

found. The sample consolidated at 1500 C contained

much bigger oxides that can also be seen on scanning

electron micrographs (Fig. 3). The intermetallic

phases could be easily distinguished from the niobium

solid solution: niobium grains contained a high dis-

location density, while in d as well as in s, numerous

stacking faults occurred (Fig. 5). The high density of

dislocations in Nbss grains can be explained by asignificant difference in the coefficient of thermal

expansion between intermetallic and Nbss phases

[13]. The presence of stacking faults may play a role

in ‘‘absorbing deviations from stoichiometry’’ [14].

As was suggested earlier by Smith et al. [14], such

stacking faults may form as a result of the coalescence

of point defects (mainly vacancies) on {100} planes.Fig. 4. X-ray spectra after consolidation.

Fig. 5. TEM micrographs of the material consolidated at

1300 C: (a) small magnification revealing size of grains, (b)

large magnification showing differences between Nbss and

intermetallic phases.




With this regard, the microstructure found in the

present research was similar to that found in the binary

Nb – 18% Al alloy [2,6]. Selected area diffraction

patterns from Nbss did not reveal superlattice reflec-

tions confirming that the solid solution is in a disor-dered state. The intermetallic phases d and s could be

relatively easily recognized by their different crystal

structures (cubic cp8 vs. tetragonal tP30) producing

SAD patterns with different symmetry. The EDS

results for the alloy compacted at 1300 C are shown

in Table 2. The Nb content in Nbss and d phases is

almost identical as well as the content of V in the

intermetallic phases. The atomic ratios in the inter-

metallic phases suggest that V atoms occupy both Nb

and Al sites in the d and s lattices. The vanadium

content in the Nbss phase was relatively high while Alcontent was much lower than that found by Horspool

et al. [4] for an alloy with the same composition. The

lower Al content in the Nbss phase may be the likely

reason that the Nbss phase was in disordered state in

the present research.

The consolidated material was characterized by

hardness measurements. The average Vickers hard-

ness number was 618 and 596 HV for samples

consolidated at 1300 and 1500 C, respectively. For

loads lower than 2 kg, the indentation were free from

cracks while indentations at 5–30 kg loads had well-developed cracks (Fig. 6). These cracks were used for

indentation fracture toughness evaluation followed by

the procedure introduced first by Palmquist [15] and

developed by Peters [11]. In this method, the lengths

of cracks emanating from the impressions are meas-

ured for a range of loads. A plot is constructed of the

applied load and the crack length, and a linear

relationship is approximated. The inverse of the slope

in kilograms per millimeter is a measure of the

fracture toughness (Palmquist indentation toughness

W ). As was shown by Peters [11], the Palmquist indentation toughness may be related to the plain-

strain energy release rate G Ic by a linear relationship

(W = aG Ic, where a was estimated by Peters for

WC– Co hard metals to beabout 4.9 103) and further

to fracture toughness K Ic:

K 2Ic ¼ EG Ic

1 v 2 ð1Þ

where E is Young modulus and n is Poisson ratio

(since the values of E and n were not known for the

examined alloy, the values for pure Nb were used for

the calculations: E = 103 GPa, n = 0.38). The values of

Palmquist indentation toughness W for both samples

(consolidated at 1300 and 1500 C) were similar: W

was about 70 kg/mm. The K Ic calculated from the

above formula was about 4 MPa m1/2. This is only an

approximate value because of errors arising from

crack length measurements and arbitrary assumed

values of a, E , and n. In our experiment, the error may

approach about 30%. However, the obtained value is

similar to that measured by Murugesh et al. [16] for a Nb/Nb3Al composite (about 6 MPa m1/2) and is larger

than the value for a pure Nb3Al compound (about

1 MPa m1/2) [16]. The relatively low fracture tough-

ness is likely due to the presence of the severely brittle

s phase. Elimination of this phase by optimization of

composition and processing parameters is expected to

improve the fracture toughness. Though the Palmquist

test does not have universal validity as a fracture

toughness test, it provides a simple, inexpensive

method of evaluating and comparing brittle materials.

5. Summary and conclusions

Mechanical alloying of elemental powder mix-

tures of 65% Nb, 15% Al, and 20% V leads to the

formation of two niobium solid solution phases (NbI

and NbII) with different lattice parameters. Different

lattice parameters are likely due to different saturation

with alloying elements (Al and V). The lattice para-

meter of the NbI phase is close to the lattice parameter

of pure niobium. The peaks from the NbI phase on theX-ray spectra decreased with milling time and ceased

completely after 180 h of milling. The peaks from the

NbII phase behaved in the opposite manner.

The phase composition in the compacted mater-

ial depended not only on the Al/V ratio but also on

the consolidation temperature: higher temperature

(1500 C) produces less detrimental Nb2Al-based s

Table 2

Chemical compositions of constituent phases (at.%)

Phase Nb V Al

Nbss 70 – 73 20 – 21 6 – 8

d 70 – 71 13 – 14 16 – 17

s 63 – 64 13 – 14 19 – 20

Fig. 6. Example of typical Palmquist cracks.




phase but simultaneously made dispersoid to ripen.

Also, unlike in the works of Tappin et al. [5] and

Horspool et al. [4], the niobium solid solution, in the

present research, remained in the disordered form.

The consolidated material exhibited high hardnessand higher fracture toughness than the pure Nb3Al

compound.

Acknowledgments

This research was sponsored by the Academy of

Mining and Metallurgy, Krakow, Poland; grant no.

10.10.110.256.

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