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The strength of niobium-oxygen solid solutions
THE
STRENGTH
OF NIOBIUM-OXYGEN K. V.
RAVIt
and
R.
SOLID
SOLUTIONS*
GIBALAt
Single crystals of high purity niobium and niobium-oxygen alloys were deformed in tension at temperatures between 4.2”K and 55O’K. At room temperature and above, the yield stress increased monotonically with oxygen content over the entire range of compositions investigated (to -4300 at. ppm). At lower temperatures alloy softening was observed, i.e. the yield stress and its temperature dependence decreased with initial additions of oxygen to approximately 500 at. ppm. With larger solute contents. the yield stress increased monotonically, similar to results at higher temperatures. Correspondmgly, the activation enthalpies and volumes at a constant effective stress first decreased then increased as These results coupled with others in the literature demonstrate a function of oxygen concentration. that dislocation-interstitiel solute interaction is the dominant deformation mechanism of b.c.c. metals et low temperatures. Furthermore, both interstitial and substitutional alloy softening are thr result of solute msocietion (scavenging) processes in solid solution. RESISTANCE MECANIQUE DES SOLUTIONS SOLIDES SIOBIVM-OSYGESE Des monocristaux de niobium de haute pm&C et d’slliages niobium-oxygtine ont &e d&form& en traction entre 4,2OK et 550’K. A temperature ambiant,e et au-dessus, la limite elestique augmente rCguli&ement avec la concentration en oxyg&ne pour tout le domaine de compositions Btudi& (jusqu’it 4300 ppm at.). Aux temp&atures infbrieures, les auteurs ont observe une d&consolidation de l’alliage, c.a.d. que la limite Qlastique et sa variation avec la tempbrature diminuent avec des additions initiales d’oxygkne allant jusqu’b 500 ppm et. environ. Pour des concentrations de solute plus &levees, la limite Blsstique augmente r&uli&ement, comme pour les temp&atures BlevBes. En m6me temps, les enthalpies et les volumes d’activation pour une contrainte effective constante diminuent d’abord puis augmentent en for&ion de la concentration d’oxyde. Ces r&ultats. joints & d’autres donnes par la litterature montrent que l’intersction dislocation-atome de solute est le m&canisme de deformation predominant dens les m&aur C.C.C.aux basses temp&atures. En outre, la d&consolidation it la fois des alliages interstitiels et de substitution est le resultat de processus d’association (raasemblement) du solute dans la solution solide. DIE FESTIGKEIT VOX ?;IOB-SAUERSTOFF-LEGIERUSGES Hochreine Siob-Einkristalle und Siob-Sauerstoff-Legierungen wurden bei Temperaturen zwischen 4,2OK und 550°K im Zugversuch verformt. Bei Zimmertemperat,ur und dariiber nimmt, die Streckgrenze im gesamten untersuchten Zusammensetzungsbereich (bis etwa 4300 ppm) monoton mit dem Sauerstoffgehalt zu. Bei tieferen Temperaturen wurde Legierungsentfestigung beobechtet, d.h. die Streckgrenze und ihre Temperaturabhangigkeit nehmen mit wachsendem Sauerstoffgehelt bis etwa 500 ppm ab. Bei griifieren Sauerstoffgehalten nahm die Streckgrenze iihnlich wie bei den Hochtempereturversuchen monoton zu. Entsprechend nahmen die Aktivierungsenthalpien und -volumina bei konst,anter effektiver Schubspannung als Funktion der Sauerstoffkonzentration zuerst ab und dann zu. Unsere Resultate zusammen mit. anderen veriiffentlichten Ergebnissen zeigen, daR die Wechselwirkung zwischen Versetzungen und Zwischengitter-Ve runreinigungen der entscbeidrnde Verformungsmecha.. nismus kubisch-raumzent,rierter Metalle bei tiefen Tempereturen ist.
experimental
1. INTRODUCTION
The large increase b.c.c.
metals
with
in the yield
decreasing
of the melt,ing temperature of interest rate ture
While
dislocation
at
mechanisms
the yield
possible
have
strength,“)
thought.
in recent
of viewpoints On one
maintained
t,here has been
hand,
many
a majority a
barrier
to
dislocation
mechanism
work intended t’o demoneach
of the Peierls-type (or flow)
of
the
models
of the thermal stress.
viewpoints.
have generally component
the activation
volume (i) to interstitial cont.ent in as evidence that an intrinsic
prevails.
long range, athermal
Solute
atoms
obstacles
are regarded
as
which do not partici-
pate in a direct way in the rate controlling tion
motion.
of
enthalpy,
of experiments
have
of the meta1,(2-4) viz. the Peierls-
potential
of
is an
investigators strength
and theoretical correctness
and t,he activation
been
into two distinct schools of
that the low temperature
property
Nabarro
years
the
cited the insensitivity
a great deal
several
strate
Proponents
~0.2
to account for the origin of the low tempera-
polarization
intrinsic
temperature
has generated
and controversy.
controlling
proposed
and flow &ress of
deforma-
mechanism.
deformation
Advocates of impurity controlled mechanisms have argued that most such
experiments
have been performed on relatively impure
Others assert that the high strength is caused by the
materials and that, in a fc>winstances where high purity
interaction
between
materials
have
interstitial
impurities
mentioned
parameters
There
exists
moving
dislocations
and residual
in solution.(s*6)
in the
literature
an abundance
of
* Received August 22, 1969. t Division of Metallurgy and Materials Science School of Engineering, Case Western Reserve University, Cleveland, Ohio 44106, U.S.A. ACTA
METALLURGICA,
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18, JUXE
1970
623
been
examined,‘6***s)
the
are concentration
above
dependent.
The presence of residual interstitials. of course, is a necessary precursor for strengthening according to these
theories
which
ranged, thermally motion
consider
activatable
interstitials obstacles
as short
to dislocation
in the b.c.c. metals, as for solutes in all other
ACTA
624
Both
classes of materials.
sides have offered
ment between experimental of simplified
and/or
ItfETALLURGICA,
agree-
results and the predictions
flexible
quantitative
theories
as
18. 19iO
VOL.
machined
crystals were chemically
surface damage introduced
These specimens were further purified by outgassing in a valveless and well trapped vacuum system which
support of their positions. The intentional alloying
of b.c.c. metals with either
employed
a diffusion
substitut,ional or interstitial solutes has unfortunately not yet greatly helped resolve this controversy.
of -lO4-1O-g
Yart of the problem
outgassing
arises in that the solutes have
judged
of
~500°K
impurity
hardening
would
consider
impure.
have been found to increase as
well as decrease the low temperature
strength of b.c.c.
metals
such
depending
temperature, strength
upon
variables
and strain rate.‘10-15)
with
initial
addition
The decrease
of
termed alloy or solution softening
as purity,
solute
has
pump.05’
torr
temperature
usually been added to base materials which advocates Also, solute additions
polished to remove
during machining.
of ~2600°K.
the
from the differences for zone-melted
Controlled
base niobium
(2.5 kg/mm2) and outgassed
amounts of oxygen were introduced
the purified crystals by forming
in
different
thicknesses
on the specimen
surfaces
been
diffusing
the oxygen
into the crystals.
The films of
and constitutes
an
various
thickness
were formed
This conclusion
low tempera-
in
b.c.c.
was based on preliminary dependence
in interstitial
(Nb-0)
results which
heated
for 2 hr at ~1300°K.
affected
alloys.
study of the effects of small and controlled of interstitial of high
oxygen
purity
of grain size.
The crystals
center of the stereographic permit comparison
a more
on the mechanical
niobium.
were used in the investigation
reports
Single
crystals
to eliminate the effects were oriented
quenched after.
under
a vacuum
near the
triangle for easy glide to
with the results of other investiga-
experiments
mechanical on
specimens
diffusion
the alloys were homogenized
beam
z3ne
outgassing.
because
it can be purified
melting
and
Oxygen
was selected
solute because
high
by electron
vacuum
oxygen
state
in niobium.
ease with which it can be added in controlled property
solid
as the interstitial
of its high solubility
and the availability
for
t,he
amounts.
of several physical and mechanical
calibrations
of these alloys as a function
of
concentration.
The
compositions
identically room
of
prepared
the niobium-oxygen
polycrystalline
temperature/liquid
the crystals. comparing
Some
pure niobium
external
dia. rods were used to grow oriented
cm
single crystals
in an elect’ron beam zone melter under a vacuum
checks
of
h1O-6
ratios of
were made analysis
Table 1 lists the solut,c
All tensile tests were conducted
on a floor model
Instron machine at a nominal strain rate of 8.3 x lo-* set-r except for strain rate sensitivity which were obtained set-’
and
8.3 x lo-* set-‘.
t,ween 42°K controlled
by cycling
and 550°K temperature
Test
during
the
deformation
8.3 x 1O-5
temperatures
by employing
liquid
Temperatures
baths.
thermocouple All shear stress-
were carried out on a Univac
area being
and
orientation
incorporated
the
computations. proportional
in length with a gage length of 2 cm and a dia. of 2 mm
absent
and as the lower yield
were spark-machined
upper and lower yield points were present.
The
changes into
as the
machine.
be-
were obtained
by an iron-constantan
shear strain calculations 110X computer,
determinations
between
tori-. The crystals were seeded to obtain specimens oriented in the middle of the stereographic triangle for single slip. Tensile specimens 5.5 cm on a ,Servomet
by and
of various specimens.
placed at the center of the specimen.
PROCEDURE
in the form of ~0.3
alloys
wires and by the
helium resistivity
t’hese results with chemical
were measured
2. EXPERIMENTAL
Commercially
for
to check that
were determined from the heights of the Snoek peaks in
concentrations
these studies
annealed
and were not contamin-
weight gain after anodization.
as the base metal
therc-
ated during the diffusion anneal.
tions on such crystals of different residual interstitial was chosen
and were
tests and internal friction
levels.
Niobium
torr
All specimens
into water and tested immediately
Routine
by
acid at
in outgassed
of ~10~*
different time intervals were performed
(Nb-W)
properties
paper
substitutional
complete additions
The present
of
phosphoric
The oxygen was then diffused into
tubes
the extent of alloy and
of 10%
quartz
of the yield strength
high purity niobium and influenced softening
in a solution
metals.
showed that the level of residual interstitials the temperature
anodizing
and
on the specimens
the crystals while they were encapsulated
mechanism
into
anodic oxide films of
different voltages.
strengthening
be
stress at
crystals.
that impurity ture
may
in the athermal
important aspect of the present investigation. In a previous investigation(l*) the authors concluded hardening is the dominant
pressures
for 8 hr at a
The effect of the vacuum
in purifying
(0.5 kg/mm2)
Typically,
were maintained
The critical resolved shear stress (CRSS) was takrn limit
when
yield
points
were
stress when distinct Because
RXVI TABLE
1.
ASD
GIBALA:
STRESGTH
OF
SIOBIUM-OXYGES
Concentrations of niobium 8nd niobium-oxygen”,“’ Oxygen@’ iat. %,
Treatment
M8terial
0.0025 ( 0.0028’~’ 0.017 0.023 0.0432
Siobium
Outgassed
Niobium-oxygen Niobium-oxygen Niobium-oxvnen
Anodized Anodized Anodized
Siobium-oxygen j?iiobium-oxygen Niobium-oxygen
Anodized 50 V Anodized 50 XVtwice” Anodized 50 T- four times’”
Niobium-oxygen
Anodized
5 J* 10 *(’ 20 V
50 V six times”’
0.072 0.141 0.2lW 0.282 10 .-393W’ 0.432
(a) Nitrogen, carbon and hydrogen contents were determined by internal friction (S) a&l chemical analysis (C, H) to be 0.002 & 0.001 at.%, 0.002 + 0.001 at. %, and <0.009 at.%, res ectively. 1!?1Substitutional impurity contents determined by Ledoux and Company to be <450 wt. ppm. Jlajor impurities included tantalum ((200 wt. ppm,), tungsten ( < 100 wt. ppm), h8fnium (
24
SOLID
SOLVTIOSS
of the well defined onset of macroscopic flow for virtually all of the specimens, none of the trends estab~shed in this investigation were itiuenced b\ the particular choice of the strength parameter. 3. EXPERIMENTAL 3.1.
RESULTS
Effects of puri$cation
The stress-strain curves of the purified niobium at various t,emperat~es are presented in Pig. 1. Three stage hardening is observed at temperatures between 250°K and 550°K. Above and below room t,emperature the yield stress changes with an attendant change in the work hardening characteristics. With increasing temperature above 298’K the lengt,h of Stage I decreases, the rate of work hardening in Stage IT increases and the stress at the onset. of Stage 311 increases. At temperatures below -250% the stressstrain curves are parabolic. with yielding followed b> relatively rapid non-linear work hardening. No yield points or twinning were observed down to 42°K;. In Fig. 2 the temperature dependence of the resolved shear stress of the high purity niobium is compared
-I
16
a
I
0. I
i
0.2
(
6
250°K
4-
2
. 298 lK ~~~~
0
FIG.
0
415-K 550’ K 0.4
625
0.8 1.2 SHEAR STRAIN 1. Resolved shrsr stress-shear strain curves at several temper8tures for high purity niobium single crystals.
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CHRlSTl4N AND MASTERS (4)
Q
YlTCnCLL
e
sv, al. (161
DUESBERY f&D H1RSCW I171 PRESENT STUDY
VOL.
18,
1970
level at higher concentrations, a result typically observed in alloy hardening systems. In Fig. 2 there is an indication of such a saturation effect from the similarities in the temperature dependences of zone refined niobium from different laboratories.(**16) It appears that many investigators have mistakenly interpreted interstitial hardening as strictly athermal when in fact they have studied hardening in impure materials, i.e. above a few hundred at. ppm interstitials, for which the changes in temperature dependence had leveled off. 3.2. EflecEsof oxygen additions
100
I
200 TEMPERATURE
--
300 ’ K
FIG. 2. Temperature dependence of the resolvod shear stress for niobium single crystals of different base pu&,ies.“.l@,l7)
with that of niobium of different base purities, as compiled from several investigations. The data of Christian and Masters(*) and Mitchell et oZ.‘16) show the temperature dependence of the yield stress of zone refined single crystals of niobium which we estimate to have ~300-400 at. ppm of residual int,erstitials. Duesbery and Hirsch(17j have purified niobium by outgassing in vacuums of the order of 1O-1o torr. For these crystals the residual interstitial level is probably in the vicinity of 10 at. ppm. The purification clearly reduces the temperat’ure dependence of the shear stress of niobium over the entire temperature range. The differences in the shear stress increase with decreasing temperatures. One indication of the potency of the interstitials to act as hardeners is the observation that the shear stress at, 42°K (29 kg mm-2) of the niobium employed in this work is lower than that of the zone refined crystals(q*l”) at 77’K (~33 kg mms2). Stein and Low’@ have reported a similar dependence of the low temperature strength on interstitial concentration in high purit,y iron. From the compilations of Conrad and Hayes it is known that the temperature dependence of t.he flow stress does not change appreciably with interstitial
Typical stress-strain curves of purified niobium and niobium-oxygen alloys deformed at three of several temperatures used in this investigation are given in Fig. 3. More complete data are presented elsewhere.(16) Yield points were observed only in a few instances in the high concentration niobium-oxygen alloys at low temperatures. At room temperature and above, a monotonic increase in the yield stress with increasing oxygen content occurred and was a~eompanied by an increase in the lengt,h of Stage I, a decrease in the rate of work hardening in Stage II and an increase in the stress at the onset of Stage III. At low temperatures (<298”K) the shear stress of the dilute niobium-oxygen alloys is lower t$han that of the purified niobium. Coupled with this alloy softening the work hardening behavior undergoes a
FIG. 3. Resolved shear stress-shear strain curves for niobium-oxygen single crystals at 77”K, 185°K and 298°K.
RAT-I
\ \ \ \ - \
STREXGTH
GIBALA:
OF
NIOBIUM-OXYGES
\
\
*
aoo2s
0
0.011
‘1
0.141
IO432
\
X
\
EXTRAPOLATED WEAR STRESSES ,FROY FIG.51
\
e >
\
a
\
\ 1:\A i ’
\\ I\\\ 0
\\
v
10 -
\
\
‘A
~
x’, ,p\y \\
\
\x
:\\_
0.
0
I
100
p\ A -&_.-
\ \‘\ -A x\
‘-i_.
zoo
.I.
--_-_*--rC
I
300
SOLID
SOLCTIOA-S
627
of solute concentration at various t,emperatures, Fig. 5. At room temperature the shear stress increases with the square root or cube root of the oxygen conAt t,emperatures below 298°K alloy centration. softening is manifested as an initial decrease in the shear stress with oxygen addition followed by nonlinear hardening at, the higher oxygen levels. The maximum amount of alloy softening is observed to occur at different oxygen concentrations at different temperatures, the minimum generally moving toward lower concentrations with decreasing t.emperature. The rate of hardening following the alloy softening at 77°K reflects the large temperature dependence of the strength of niobium-oxygen alloys at, t~m~r~tures below ~100°K. This was also observed in Fig. 2 for niobium of different. impurity levels controlled in amount but not in type of interstitial.
AT%0
20 N
ASD
400
FIG. 4. Temperature dependence of the resolved shear stress of niobium and niobium-oxygen single crystals. The dotted curve represents extrapolated zero-oxygen stresses taken from Fig. 5.
The stress-strain curves change change with allo$ng. from it parabolic form exhibited by the unalloyed niobium to curves resembling a three stage form with yielding followed by low rates of work hardening in the dilute alloys of niobium. At high oxygen concentrations the shear stress increased with increasing oxygen and the stress-strain curves returned to a parabolic form and occasionally exhibited yield drops. At, 77°K only the first. addition of oxygen (0.017 at,.:,; cent) resuhed in B decrease in the shear stress. At, higher oxygen concentrations twinning was observed along with an increase in the yield stress. The temperature dependence of the shear stress of niobium and three of the niobium-oxygen alloys is shown in Fig. 4. At 298°K and above all the alloys display a higher shear stress than the pure niobium. At temperatures below 298°K the majority of the alloys display a lower shear stress than the unalloyed niobium at, some temperatures. For example, the niobium-O.017 at.YO oxygen alloy has a lower “yield stress than the pure niobium over the entire temperature range below ~250%, whereas the 0.141 at.94 oxygen alloy is stronger than the base niobium above -230°K and below ~1OO’K. This latter behavior is similar t,o behavior observed for Fe-N alloys(1r**9) which display alloy softening. The phenomenon of alloy softening is most strikingly displayed when the shear stress is plotted as a function
3.3. l’hermal activation analysis The activation enthalpy and activation volume for low temperature deformation behavior are often determined by the methods out.lined by Conrad(r) and others. These parameters are given by
and
Al.
X OXYSEN
Fro. 5. The concentration depcudence of the resofved shear stress of niobium-oxygen single crystals in the temperature range 77-3WK
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METALLURGICA,
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AT
4 0 + 0 *
0 0
I
I
100
200 TEMPERATURE
X
0
0.0025 0.017 0.023 0.141 0.282
300
4
0
‘K
FIG. 6. The temperature dependence of the strain rate sensitivity of niobium and niobium-oxygen single crystals.
respectively. Here T* = TV- r,,, T,,is the shear stress, T# is the long range athermal stress, r* is the thermally activatable portion of the total stress, 9 is the shear strain rate and &, = pvAb, where p is the mobile dislocation density, v is the attempt, frequency, A is the area swept out by each successful t’hermal fluctuation and b is the Burgers vector. k and T have their usual meaning, and the subscript p refers to T,, having t,he temperature dependence of the shear modulus p. The partial derivatives in equations (1) and (3) are l.5r
Al
Y 0
4 0.0025 -0 0.017 +
I.0 -
0.072
-& 0.432
(
determined from the strain rate sensitivity of the shear stress at constant temperature and the temperature dependence of the shear stress at constant strain rate. Representative examples of the temperature dependences of the strain rate sensitivit’\r of niobium and niobium-oxygen alloys are given in Fig. 6. Figures 7 and 8 show the activation enthalpy and the activation volume as a function of the effective stress T*. These parameters are functions of oxygen concentrations at low temperatures and at, all effective stress levels. The composition dependence of t,he activation parameters parallels the composition dependence of strength, i.e. the addition of small amounts of oxygen results in alloy softening and a reduction of thtb activation parameters (an increase in the peak temperature in the case of the strain rate sensitivity). whereas at the higher oxygen concentrations an increase in the shear stress is accompanied by a corresponding increase in the activation parameters and a decrease in the strain rate sensitivity peak temperature. These results are quite similar to those obt,aincd in comparable studies of alloy-softened substitutional allovs.(13*20*21) Since the skain rate sensitivitp and t,hcb activation parameters are a function of solute concentration, the rate determining step for the low temperature deformation of niobium is a function of the oxygen concentration of the metal. 4. DISUSSION
FIG.
7. Activation enthalpy vs. effective stress for niobium and niobium-oxygen single crystals.
In studies of the low temperature strength of the b.c.c. metals experimental results are often analyzed in terms of the predictions of various models of lattice
RAY1
ATD GIBALA:
STRESGTH
OF
SIOBIUM-OSTGES
SOLID
+ -0. 0 9*
6%
SOLI-TIOSS
AT YO 00025 0.017 0.0432 0.141 0.282
FIG. 8. Activation volume vs. effective stress for niobium and niobium-oxygen single crystals.
or impurity
hardening.
as evidence
for or against
mechanisms. theories such
These analyses are then used
Apart from objections
and the logic
analyses
investigation distorts
one of these two general
seem
inappropriate
the temperature To
enumerate
present arrived
and established
approach, the
present
of alloy softening
and compositional expected
depend-
for either type of
our discussion
and discuss separately
conclusions
for
because the occurrence
ences from that normally mode1.t
to the individual
of this analytical
succinctly,
we
each of the major
at in this study from our results
information
from the literature.
hardening.
The more potent the hardener,
the concentration be reached. “gradual”
solute interactio?l in
h.c.c. metals c.ol,stitutes n low temperawe
and interstitial
for concentrat,ions
The variations hardened
material
In general. Ar*lA ature.
is given
increasing
schematically
in Fig. 9.
purit,y results in a lower
In 1; and H and a higher v at, constant In
practical
terms.
these
not decrease or increase indefinitely concentration; values
T*, AT/A In 3,
of T(or T*) of an ideal solution
and
give
the
rrppenrance
temper-
paramet#ers as a function
they u-ill reach or approach of
T*, will of
limiting
athermal-type
t Nevertheless, some of these analyses were made as part
of this investigation.‘15’ It was found that the experimental results were in only modest agreement with lattice hardening Dorn-Rajnak)‘3’ and impurity hardening (e.g. (e.g. Fleischer)‘s’ models and then only at higher concentrations (22000 at. ppm), where neither type of theory should have validity.
metals.(22*23’
parameters reach limiting values of these solut.es ranging
from
a
upon the pot,ency
of bhe solute obstacle. There is no reason to believe b.c.c.
metals
cause
“rapid”
with
that’ interstitials
(or most, tetragonal hardening,)15) motion
a thermally
distortions
in
which
are such large barriers
t,hat, they should not also conactivatable
barrier.t
However
they are larger barriers than the spherically
symmetrical
of the parameters
H and 1’ as a function
such as substitutional
in close-packed
few to tens of an at.O,. depending
because
rcctivated dejormntio~~ n/echrinism
obstacles
solutes
The above mentioned
stitute
thermally
values will
This type of behavior is often observed for hardening
to dislocat,ion 4.1. Dislocntiorl-ijzterstitlnl
the lower
at which these limiting
quite
ones menConed low interstitial
above.
one must
concentrations
deal
in order
to observe the t’ypical alloy hardening behavior depict,ed in Fig. 9. Figures 3 and 4 demonstrate that this is the case for oxygen in niobium Low.‘6)
Lawley
et aZ.(25) for
ct
d..(g)
interstitials
in
all b.c.c.
metals.
least below (and perhaps ening behavior
is observed.
Smialek
molybdenum, In these systems,
interstitial
much below)
atomic ppm are necessary
and
iron.
tungsten and tantalum respectively. and probably
as do Stein and
Koo@)
levels
at
a few hundred
before normal
alloy hard-
For higher concentration
t In fact, advocates of intrinsic hardening theories have generally argued that interstitials are not potent enough hardeners t’o account for much of t,he low temperature strength.‘2*)
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~~_ -
FIG. 9.
----
PURE METAL
-
SOLID SOLUTIONS
----
LIMITING
VALUE
Schematicplots of T* and (AT/A In 3) as a function of temperature
and v and H as a function of effective stress for ideal solid solution hardened systems.
alloys, those used in the vast majority of investigations on b.c.c. “metals”, the parameters depicted in Fig. 9 tend toward their limiting values (the dotted lines) and give the appearance of a compositionindependent change of these parameters as a function of temperature or effective stress. 4.2. Dislocation-interstitial
interaction
is the
dominant low temperature hardening mechanism in b.c.c. metals
The present results suggest that if niobium has an intrinsic (Peierls) strength at O”K, it must be below the strength of the “pure” niobium used in the present experiments (~30 kg mm-2) and possibly as low as the strength of the unsoftened material suggested by the extrapolations shown in Fig. 4 and displayed as the dotted line (bottom) in Fig. 5 (~15 kg mm-2). Both of these apparent Peierls stresses are significantly below values quoted for zone refined and lesser purity niobium from other studies, e.g. >50 kg mme2. Obviously these latter values do not represent true Peierls stresses. The question we consider now is whether or not the 0°K strength of the niobium studied here represents the Peierls stress of niobium. We argue that it does not. Our base material contained approximately 20 at. ppm of residual nitrogen and perhaps as much carbon. From changes of lattice parameters,f2Q Snoek peak relaxation strengths(27*W) and room temperature hardening rates(29) with solute addition, both of these solutes
should be significantly more potent low temperature hardeners than oxygen. In Fig. 5, beyond the softening range of compositions, the hardening rate of oxygen is of the order of 10 p at 77°K. Thus, higher rates of hardening and the above stated compositions for nitrogen and carbon can easily account for the major part of the remnant hardening of niobium. Hence, the base niobium containing ~40 at. ppm of residual interstitials is “dirty” relative to the potency of the interstitial types remaining. Stein and LOW(~)have observed that reducing the residual carbon level in iron from >200 at. ppm to less than 0.03 at. ppm lowered the 0.2 per cent offset *yield stress by ~50 per cent from ~25 kg mm-2 to -13 kg mm-2 at 78°K. For such large apparent hardening rates ( 2 10 ,u) at very low concentrations, it is understandable that Stein and Low’s results have not easily been duplicated and that many investigations on other b.c.c. metals containing interstitials with comparable hardening rates have failed to disclose the decreased temperature dependence of T* to be observed at lower concentrations than the tens (more often hundreds or thousands) of at. ppm which the materials contained. 4.3. Alloy softening by interstitials solutes ix b.c.r. metals is caused by solute association (scavenging) Three theories have been advanced to account for the occurrence of solution softening in b.c.c. alloys at low temperatures. These are (1) a decrease in the
RAW
Ah‘D
STRESGTH
GIBALA:
OF
~IOBIU~f-OSYGES
Peierls potential as a result of alloying,c12*30) (2) an increase in the mobile dislocation density brought about by the presence of solute atoms,(ll*lg) and (3) interstitial gettering (in substitutional alloys).@‘.ss) A reduction in the Peierls stress has been visualized to occur in a number of ways. One proposalo2) is that the Peierls stress is reduced as a result of solute atoms producing a local disorder in the periodic lattice of the solvent metal. This tends to smear out the Peierls energy barrier resulting in dislocation mobility at lower stress levels than in an unperturbed crystal. The Peierls stress T@ can be expressed as(33)
Pb 7pNuzexp
-
2V12
( 1 c
(3)
where a can vary between $ and 1, b is the Burgers vector, y is the shear modulus, c fz b for screw dislocations and A is the “width” of the dislocation. According to this expression, TV will be reduced if p is reduced by alloying.~30) Jones et u&(~*)have determined the elastic constants of high purity niobium and a niobium-O.3 at.% oxygen alloy. The decrease they observed in the shear modulus due to alloying is too small to account for the amount of alloy softening observed in the Nb-0 system. The concept of dissociated dislocations has also been invoked to explain alloy softening.oa) Solute atoms can be visualized to affect dissociated dislocations in two distinct and opposite ways. If the temperature dependence of the strength of b.c.c. metals is due to dissociated dislocations which are sessile, then for sessile t.o glissile transformation to occur the dislocat,ions must constrict along their length. The presence of solute atoms is thought to aid such a constriction. making possible dislocation mobility at lower stress levels than in the absence of constricting sites represented by the solutes. The constriction, in effect. represents an increase of the stacking fault energy or a reduction of the width of the dissociated dislocation. Alternatively, solute addition could reduce the stacking fault energy, thus increase the width of the dislocations and according to equation (3) lower the Peierls stress. According to rate theory, the shear strain rate + that results upon the application of a stress to a crystal is given by (4)
Conradc2) and others’ll*lg) have suggested that for the same Peierls st,ress, a change in the pre-exponential
SOLID
SOLX-TIOSS
631
factor PO can change the shear stress r*. Since it(l = pAb, a change in POcan occur through a change in p. Nakada and Keh’il’ propose that alloy softening in the Fe-N system is a result of an increase in the mobile dislocation density due to alloying. ChristIss) and Christ and Smithog) have developed a model for solution softening in interstitial alloys by proposing that an increase in the mobile dislocation density occurs through solute atom induced cross slip of dislocations. This mechanism is presumed to occur when the repulsion between fixed interstitial dipoles in certain orientations and screw dislocations leads to movement of the dislocations onto cross-slip planes. The two theories, a reduction in t,he Peierls stress and an increase in the mobile dislocation density can not account for all the observed experimental facts associated with the phenomenon of solution softening. A reduction in the Peierls Stress assumes that the low temperature deformation of the unalloyed metal is governed by the intrinsic strength of the lattice. It was demonstrated in Section 4.2 that most of the remnant strength of the purified niobium can be accounted for by dislocation-in~rstitial interaction. An increase in the mobile dislocation density by alloying should not affect the activation enthalpy but should only change the pre-exponential factor PO. This is not found to be the case for Nb-0 solid solutions (Figs. 7, 8). An increase in the mobile dislocation density by cross slip of dislocations due to repulsion between interstitials and dislocations assumes that the minimum distance of approach between the dislocation and the interstitial for maximum repulsion is of the order of a few Burgers vectors. However at these distances the dislocation-interstitial interaction energy is at a substantial fraction of its maximum value ; hence interstitial pinning is probably more favorable than interstitial repulsion. Perhaps the strongest argument against intrinsic alloy softening is that the softening is not a reversible phenomcnon.‘14) Whereas addition of oxygen to purified niobium results in alloy softening at low tern. peratures (Fig. 4), removal of interstitials involving the same composition range does not result in alloy softening (Fig. 2). If alloy softening is due to a modiovation of an intrinsic phenomenon such as a reduction of the Peierls stress or a change in the mobile disloeation density. then it should occur when solutes are added or removed from the solvent. In the light of these facts, the mechanism that best explain the occurrence of solution softening is a scavenging type of process.(14) Alloy softening in substitutional systems has
632
ACTA
METALLURGICA,
sometimes been attributed to get&ring or scavenging of interstitials by substitutionals.‘31*32) This interpretation can readily be applied to substitutional alloys of metals in which the alloying element is a strong carbide, nitride, or oxide former. This theory has been rejected by most investigators as an explanation for alloy softening in other substitutional systems and in interstitial alloys for which interstitial compounds are unlikely to form. However, scavenging need not mean only the formation of a chemical compound. Association between substitutional and interstitial solutes and between interstitials of the same or different types in solid solution can readily occur and cause softening. Internal friction studies,(27.36~37) electron diffraction and microscopy,(as*39) field ion microscopy,‘*0) superconducting properties’41*42) and thermodynamic studies(46) attest to the occurrence of solute association (clustering) in the b.c.c. metals. Interstitial solute association in solution can lead to a decrease in the low temperature strength by a modification of the properties of the randomly distributed solute in the following ways: (1) Solute association reduces the total number of thermally activatable obstacles to dislocation motion over that which would be available if the interstitials were randomly distributed. (2) The effective hardening rate ar*/aC,, where Ci is the obstacle concentration of an associated pair (or larger cluster) of interstitials can be smaller than the sum of the hardening rates of the unassociated interstit,ials. This is reasonablesince clustering occurs with a reduction in the total free energy of the solid so1ution,‘36) of which strain energy should be a dominant contribution. (3) The temperat,ure dependent contribution to strengthening by associated interstitials can be different (and not necessarily athermal) from the temperature dependent contribution of unassociated interstitials. The change in the strain rate sensitivity peak temperature with alloying in Fig. 6 is a strong indication of the modification of the interstitial contribution to the temperature dependent part of the shear stress by clustering. An ideal alloy hardened system would display an increase in the strain rate sensitivity peak at about the same temperature with increasing solute as in Fig. 9. In the niobium-oxygen system the peak heights for the unalloyed niobium and the niobiumoxygen alloys remain relatively constant, while the peak shifts first to higher temperatures wit,11alloy softening (reflecting the different hardening rate of the associated solutes) and then to lower temperatures with alloy hardening. Corresponding changes in the
VOL.
18,
1910
activation enthalpy and the activation volume reflect the changes in the strain rate sensitivity and results in a deviation of the activation parameters from the ideal behavior depicted in Fig. 9. There remains the question regarding the specific nature of the clusters envisaged to occur and cause the alloy softening. Clearly, oxygen atoms must be clustering with each other or with other interstitials such as nitrogen, carbon or hydrogen. The residual substitutional levels of the purified niobium are too small to account for the observed magnitude and the rate of alloy softening by substitutional-oxygen association. Oxygen clustering alone probably can not account for the observed alloy softening either. The oxygen Snoek peak heights and the residual resistivity increase monotonically with increasing oxygen concentration. These parameters give no indication that a large amount of the oxygen is present in clusters or pre-precipitates at the lower concentration levels. This would be required to explain the immediacy, the magnitude and the large rate of alloy softening with initial additions of oxygen relative to the subsequent alloy hardening at a given temperature (e.g. 113’K in Fig. 5). It would seem that, oxygen, present as the dominant impurity, also associates with other residual interstitials which are more potent hardeners, i.e. nitrogen or carbon, and reduces their effectiveness as thermally activatable obstacles to dislocation motion. Of these, oxygen-nitrogen association in niobium has been observed by internal friction techniques,‘37*43) is thought to be characterized by a larger binding enthalpy than oxygenoxygen association,(43) and represents the most probable cause of the alloy softening. Hydrogen can also associate with oxygen, but is not likely to be a potent enough hardener to cause appreciable softening when clustered. 4.4. Alloy softening by substitutional also a scavenging process
solutes is
This conclusion was drawn in an earlier papero4) on the basis that the extent of alloy softening caused by substitutional solutes was determined largely by the purity, relative to residual interstitials, of the base materials used to prepare the alloys. Here we have noted in addition that both the activation enthalpies and the activation volumes for these systems are functions of solute concentration(ia*sO.sii in the same manner observed for an interstitial alloy, niobiumoxygen. and as expected for a scavenging-type of process. These results are not well explained by intrinsic softening mechanisms, whether by a reduction
RAVI
ASD
STRESGTH
GIBALA:
OF
of the Peierls stress or an increase in the mobile location density. Finally we re-emphasize stitial association association)
dis-
that substitutional-inter-
(as well as interstitial-interstitial
in solid
phenomenon
SIOBIUM-OSYGES
solution
is a well documented
from internal friction,‘44) field and thermodynamic studies.(46’
microscopy(J5) solute association
to cause alloy softening
necessary that the t,wo species undergoing cause appreciably
different hardening
ion For
it is only association
rates and tem-
perature dependences of the flow stress of the base metal. This is virtually always true in b.c.c. metals containing
gradual
hardening
rapid hardening interstitials. that substitutional
The exceptions
substitutional-interstitial
l-4
occur when the
int,eraction
small or the alloy is relatively
b.c.c.
and
softening is almost always observed
in b.c.c. alloy systems.
4.5. Gonclusio)~s
substitutionals
Thus, it is not surprising
is
extremely
free of interstitials.
nre true in general
for
all
metals
It might be argued that the present results obt,ained for
niobium-oxygen.
metalf3”)
among
involving
the b.c.c.
relativel;v low potency in
an
elastically
transition
metals
soft and a
obstacle among the interstitials
b.c.c.
metals,?
are
not
hardened
systems.
We
believe
typical
of
interstitial
otherwise.
Indeed,
t,he results suggest that if a relatively
weak hardener
in a soft metal
r*, then more
potent
can effect a sizeable
hardeners
vanadium.
tantalum,
and tungsten
observed
iron,
harder metals such as
chromium,
An indication
of the correctness of this
is the very large amount in tungsten.‘“‘)
of alloy softening which
and chrom-
The large amounts
infer that the residual interstitials
hardeners.
rhenium
of alloy softening
molpbdenum’48)
ium(4s) when allored with rhenium. are pot,ent
molybdenum
should effect larger r*‘s at much lower
concent’rations. position
in elastically
softens
For to
tungsten
in particular,
concentrations
than 20 at.Ob. large dislocation-interstitial
of
more
interaction
energies (-1 eV) are expected and appreciable association between rhenium and oxygen is well docu-
REFERENCES 1. J. E. DORS, Dislocation Dynamics, p. 27. McGraw-Hill (1968). 2. H. COSR~D, The Relaiion Between the Structure and Xechanical Propertias of MetaEs, p, 475. H.M.S.O. (1963). 3. J. E. DORS and S. RAJSAK, Trans. Am. Inst. Min. Engrv 230, 1052 (1964). 4. J. W. CHRISTIAS and B. C. XASTERS, Proc. R. Sot. A281, 223 (1964). 5. R. L. FLEISCHER, Acta Uet. 15, 1513 (1967). 6. D. F. STEIS and J. R. Low, JR., Acta Xet. 14,1183 (1966). 7. H. CONRAD, J. Metala 16. 582 (1964). 8. R. C. Koo, Acta Met. 11, 1083 (1963). 9. A. LAUZET. J. TAS DER S~PE and R. MADDIS, J. In&. Metals 91, 23 (1962-3). and S. J. 10. W. C. LESLIE, R. J. SOBER, S. G. Ba~coclc GREES. Trans dm. Sot. Netals. 62. 690 (1969). T. XA&ADA and A. S. KEH. Acta ;cet. 16; 903 ‘( 1968). :;: P. L. RAFFO and T. E. MITCHELL, Trans. Am. Inst. Mirr. Engrs 242, 905 (1968). 13. R. J. ARSESA~LT, Acta Net. 17. 1291 (1969). 14. Ii. V. RA~I and R. GIBALA, Scripta Met. 3, 54i (1969). 15. K. V. RAVI, Ph.D. Thesis, Case Western Reserve Univer-
sity, September 1969.
16. T. E. MITCHELL, R. A. FOXALL and P. B. HIRSCH, Phil. Maq. 8, 1895 (1963). 17. 11. $. DUESBER~ and P. B. HIRSCH, Dislocation Dynamics, p. 57. McGraw-Hill (1968). See also R. A. FOSALL, Ph.D. Thesis, Univ. of Cambridge (1967). 18. H. CONRAD and R. HATES, Trans. dm. Sot. hfetala 56, 128 (1963). 19. B. W. CHRIST and G. T-. SXITH, >1bm. Sciettt. Revue M&all. 65, 208 (1968). 20. A. A. HENDRICGSON, B. C. PETERS and R. A. STRAHL, U.S. Atomic Energy Commission Technical Report No. COO-916-13, March 1968. B. C. PETERS and A. A. HENDRICKSON, J. Xetals 21, 116A (1969). 21. G. C. DAS and R. J. ARSENATJLT, Scripta 111et. 2, 495 (1968). 22. K. R. EVASS and R. F. FLAKAGAS, Phil. .liag. 18, 977 (1968). 23. D. T. PETERSOX and R. L. SKAQGS, Trails. Am. Inst. Min. Engrs 242, 922 (1968). 24. J. W. CHRISTIAN, Scripta Met. 2, 569 (1968). 25. R. L. SJIIALEI(. G. L. WEBB and T. E. MITCHELL, Scripta xet. 4, 33 (1970). 26. A. TAYLOR and S. J. DOYLE, J. less-common Metals 13,
313, 399 (1967). 27. R. W.
POWERS and M. 1.. DOYLE, J. appl. Phys.
30, 514
(1959).
28. P. M. ROBINSOH and R. RA\~LIXQS, Iron Steel 31, 3, 66, 97 (1958). 29. Z. C. SZKOPIAIC.J. less-common Metals. 19, 93 (1969). R. J. ARSESAULT, Acta filet. 15, 501 (1967). 7:: H. H. KRAXZLEIX, RI. S. BURTON and G. V. %ITH, Trans.
Am. Inst. Min. Engrs 233,69
(1965).
Trans. Am. Inst. Min. Engrs 253, 1500 (1965). J. WEERTMAN and J. R. WEERTMAS, Elementary Dislocation Theory, p. 159. Macmillan (1964). K. A. JOKES, 8. C. Moss and Ii. M. ROSE, Acta Met. 17, 365 (1969). H. W. CHRIST, Acta Met. 17. 131i (1969). R. GIB~LA and C. A. WERT, Arta Net. 14, 1095 1105
32. S. S. STOLOFF, R. G. DAVIES and R. C. Ku, 33. 34. 35. 36.
Rep. 19, 505 (1964). 38. J. VAN LAXD~X-T and S. A~IELIXCKX, Sppl.
ACKNOWLEDGEMENT
work
633
SOLI-TIOSS
(1966). 37. D. J. VAX O~IJEX and A. 6. PAX DER GOOT, Philips
mented.‘J5)
This
SOLID
was supported
At,omic Energy _4T (ll-l)-1676.
Commission.
by
the
under
United Contract
States pie.
t Kate for example from the tabulation of Snoek relaxation stren@hs . . by Powers and .. Doyle’*” that the strength .of Nb-0
relaxat,lon 1s the smallest among the systems mvolvmg the __ . . __ group 1 -A metals and won and the mterstitials 0, N and C.
Rss.
Phys. Lett. 4,
15 (1964). 39. J. VAS LAKDUTT, Phyls. Status Solidi 6, 957 (1964). 40. S. XAI
236,924 (1966). 44. P. M. BUNN, I~. G. C’UMMINOS and H. W. LEAVESWORTH, *JR., J. appl. Phy.7. 33, 3009 (1962).
634
ACTA
METALLURGICA,
45. D. T. NOVICKand E. S. MACHLI~-, Trans. Am. Sot. Metals 61, 777 (1968). 46. H. I. AARONSON,H. A. DOMIANand G. M. POVSD, Trans. Am. Inst. Min. &?r8 256, 753, 768 (1966).
VOL.
18,
1950
47. P. L. RAFFO, J. less-common Metals 17, 133 (1969). 48. D. L. DAVIDSONand F. R. BROTZEB,J. hlelds 21, 116A (1969). 49. A. GILBERT,J. Metak 21, 89A (1969).
STRENGTH
OF NIOBIUM-OXYGEN K. V.
RAVIt
and
R.
SOLID
SOLUTIONS*
GIBALAt
Single crystals of high purity niobium and niobium-oxygen alloys were deformed in tension at temperatures between 4.2”K and 55O’K. At room temperature and above, the yield stress increased monotonically with oxygen content over the entire range of compositions investigated (to -4300 at. ppm). At lower temperatures alloy softening was observed, i.e. the yield stress and its temperature dependence decreased with initial additions of oxygen to approximately 500 at. ppm. With larger solute contents. the yield stress increased monotonically, similar to results at higher temperatures. Correspondmgly, the activation enthalpies and volumes at a constant effective stress first decreased then increased as These results coupled with others in the literature demonstrate a function of oxygen concentration. that dislocation-interstitiel solute interaction is the dominant deformation mechanism of b.c.c. metals et low temperatures. Furthermore, both interstitial and substitutional alloy softening are thr result of solute msocietion (scavenging) processes in solid solution. RESISTANCE MECANIQUE DES SOLUTIONS SOLIDES SIOBIVM-OSYGESE Des monocristaux de niobium de haute pm&C et d’slliages niobium-oxygtine ont &e d&form& en traction entre 4,2OK et 550’K. A temperature ambiant,e et au-dessus, la limite elestique augmente rCguli&ement avec la concentration en oxyg&ne pour tout le domaine de compositions Btudi& (jusqu’it 4300 ppm at.). Aux temp&atures infbrieures, les auteurs ont observe une d&consolidation de l’alliage, c.a.d. que la limite Qlastique et sa variation avec la tempbrature diminuent avec des additions initiales d’oxygkne allant jusqu’b 500 ppm et. environ. Pour des concentrations de solute plus &levees, la limite Blsstique augmente r&uli&ement, comme pour les temp&atures BlevBes. En m6me temps, les enthalpies et les volumes d’activation pour une contrainte effective constante diminuent d’abord puis augmentent en for&ion de la concentration d’oxyde. Ces r&ultats. joints & d’autres donnes par la litterature montrent que l’intersction dislocation-atome de solute est le m&canisme de deformation predominant dens les m&aur C.C.C.aux basses temp&atures. En outre, la d&consolidation it la fois des alliages interstitiels et de substitution est le resultat de processus d’association (raasemblement) du solute dans la solution solide. DIE FESTIGKEIT VOX ?;IOB-SAUERSTOFF-LEGIERUSGES Hochreine Siob-Einkristalle und Siob-Sauerstoff-Legierungen wurden bei Temperaturen zwischen 4,2OK und 550°K im Zugversuch verformt. Bei Zimmertemperat,ur und dariiber nimmt, die Streckgrenze im gesamten untersuchten Zusammensetzungsbereich (bis etwa 4300 ppm) monoton mit dem Sauerstoffgehalt zu. Bei tieferen Temperaturen wurde Legierungsentfestigung beobechtet, d.h. die Streckgrenze und ihre Temperaturabhangigkeit nehmen mit wachsendem Sauerstoffgehelt bis etwa 500 ppm ab. Bei griifieren Sauerstoffgehalten nahm die Streckgrenze iihnlich wie bei den Hochtempereturversuchen monoton zu. Entsprechend nahmen die Aktivierungsenthalpien und -volumina bei konst,anter effektiver Schubspannung als Funktion der Sauerstoffkonzentration zuerst ab und dann zu. Unsere Resultate zusammen mit. anderen veriiffentlichten Ergebnissen zeigen, daR die Wechselwirkung zwischen Versetzungen und Zwischengitter-Ve runreinigungen der entscbeidrnde Verformungsmecha.. nismus kubisch-raumzent,rierter Metalle bei tiefen Tempereturen ist.
experimental
1. INTRODUCTION
The large increase b.c.c.
metals
with
in the yield
decreasing
of the melt,ing temperature of interest rate ture
While
dislocation
at
mechanisms
the yield
possible
have
strength,“)
thought.
in recent
of viewpoints On one
maintained
t,here has been
hand,
many
a majority a
barrier
to
dislocation
mechanism
work intended t’o demoneach
of the Peierls-type (or flow)
of
the
models
of the thermal stress.
viewpoints.
have generally component
the activation
volume (i) to interstitial cont.ent in as evidence that an intrinsic
prevails.
long range, athermal
Solute
atoms
obstacles
are regarded
as
which do not partici-
pate in a direct way in the rate controlling tion
motion.
of
enthalpy,
of experiments
have
of the meta1,(2-4) viz. the Peierls-
potential
of
is an
investigators strength
and theoretical correctness
and t,he activation
been
into two distinct schools of
that the low temperature
property
Nabarro
years
the
cited the insensitivity
a great deal
several
strate
Proponents
~0.2
to account for the origin of the low tempera-
polarization
intrinsic
temperature
has generated
and controversy.
controlling
proposed
and flow &ress of
deforma-
mechanism.
deformation
Advocates of impurity controlled mechanisms have argued that most such
experiments
have been performed on relatively impure
Others assert that the high strength is caused by the
materials and that, in a fc>winstances where high purity
interaction
between
materials
have
interstitial
impurities
mentioned
parameters
There
exists
moving
dislocations
and residual
in solution.(s*6)
in the
literature
an abundance
of
* Received August 22, 1969. t Division of Metallurgy and Materials Science School of Engineering, Case Western Reserve University, Cleveland, Ohio 44106, U.S.A. ACTA
METALLURGICA,
VOL.
18, JUXE
1970
623
been
examined,‘6***s)
the
are concentration
above
dependent.
The presence of residual interstitials. of course, is a necessary precursor for strengthening according to these
theories
which
ranged, thermally motion
consider
activatable
interstitials obstacles
as short
to dislocation
in the b.c.c. metals, as for solutes in all other
ACTA
624
Both
classes of materials.
sides have offered
ment between experimental of simplified
and/or
ItfETALLURGICA,
agree-
results and the predictions
flexible
quantitative
theories
as
18. 19iO
VOL.
machined
crystals were chemically
surface damage introduced
These specimens were further purified by outgassing in a valveless and well trapped vacuum system which
support of their positions. The intentional alloying
of b.c.c. metals with either
employed
a diffusion
substitut,ional or interstitial solutes has unfortunately not yet greatly helped resolve this controversy.
of -lO4-1O-g
Yart of the problem
outgassing
arises in that the solutes have
judged
of
~500°K
impurity
hardening
would
consider
impure.
have been found to increase as
well as decrease the low temperature
strength of b.c.c.
metals
such
depending
temperature, strength
upon
variables
and strain rate.‘10-15)
with
initial
addition
The decrease
of
termed alloy or solution softening
as purity,
solute
has
pump.05’
torr
temperature
usually been added to base materials which advocates Also, solute additions
polished to remove
during machining.
of ~2600°K.
the
from the differences for zone-melted
Controlled
base niobium
(2.5 kg/mm2) and outgassed
amounts of oxygen were introduced
the purified crystals by forming
in
different
thicknesses
on the specimen
surfaces
been
diffusing
the oxygen
into the crystals.
The films of
and constitutes
an
various
thickness
were formed
This conclusion
low tempera-
in
b.c.c.
was based on preliminary dependence
in interstitial
(Nb-0)
results which
heated
for 2 hr at ~1300°K.
affected
alloys.
study of the effects of small and controlled of interstitial of high
oxygen
purity
of grain size.
The crystals
center of the stereographic permit comparison
a more
on the mechanical
niobium.
were used in the investigation
reports
Single
crystals
to eliminate the effects were oriented
quenched after.
under
a vacuum
near the
triangle for easy glide to
with the results of other investiga-
experiments
mechanical on
specimens
diffusion
the alloys were homogenized
beam
z3ne
outgassing.
because
it can be purified
melting
and
Oxygen
was selected
solute because
high
by electron
vacuum
oxygen
state
in niobium.
ease with which it can be added in controlled property
solid
as the interstitial
of its high solubility
and the availability
for
t,he
amounts.
of several physical and mechanical
calibrations
of these alloys as a function
of
concentration.
The
compositions
identically room
of
prepared
the niobium-oxygen
polycrystalline
temperature/liquid
the crystals. comparing
Some
pure niobium
external
dia. rods were used to grow oriented
cm
single crystals
in an elect’ron beam zone melter under a vacuum
checks
of
h1O-6
ratios of
were made analysis
Table 1 lists the solut,c
All tensile tests were conducted
on a floor model
Instron machine at a nominal strain rate of 8.3 x lo-* set-r except for strain rate sensitivity which were obtained set-’
and
8.3 x lo-* set-‘.
t,ween 42°K controlled
by cycling
and 550°K temperature
Test
during
the
deformation
8.3 x 1O-5
temperatures
by employing
liquid
Temperatures
baths.
thermocouple All shear stress-
were carried out on a Univac
area being
and
orientation
incorporated
the
computations. proportional
in length with a gage length of 2 cm and a dia. of 2 mm
absent
and as the lower yield
were spark-machined
upper and lower yield points were present.
The
changes into
as the
machine.
be-
were obtained
by an iron-constantan
shear strain calculations 110X computer,
determinations
between
tori-. The crystals were seeded to obtain specimens oriented in the middle of the stereographic triangle for single slip. Tensile specimens 5.5 cm on a ,Servomet
by and
of various specimens.
placed at the center of the specimen.
PROCEDURE
in the form of ~0.3
alloys
wires and by the
helium resistivity
t’hese results with chemical
were measured
2. EXPERIMENTAL
Commercially
for
to check that
were determined from the heights of the Snoek peaks in
concentrations
these studies
annealed
and were not contamin-
weight gain after anodization.
as the base metal
therc-
ated during the diffusion anneal.
tions on such crystals of different residual interstitial was chosen
and were
tests and internal friction
levels.
Niobium
torr
All specimens
into water and tested immediately
Routine
by
acid at
in outgassed
of ~10~*
different time intervals were performed
(Nb-W)
properties
paper
substitutional
complete additions
The present
of
phosphoric
The oxygen was then diffused into
tubes
the extent of alloy and
of 10%
quartz
of the yield strength
high purity niobium and influenced softening
in a solution
metals.
showed that the level of residual interstitials the temperature
anodizing
and
on the specimens
the crystals while they were encapsulated
mechanism
into
anodic oxide films of
different voltages.
strengthening
be
stress at
crystals.
that impurity ture
may
in the athermal
important aspect of the present investigation. In a previous investigation(l*) the authors concluded hardening is the dominant
pressures
for 8 hr at a
The effect of the vacuum
in purifying
(0.5 kg/mm2)
Typically,
were maintained
The critical resolved shear stress (CRSS) was takrn limit
when
yield
points
were
stress when distinct Because
RXVI TABLE
1.
ASD
GIBALA:
STRESGTH
OF
SIOBIUM-OXYGES
Concentrations of niobium 8nd niobium-oxygen”,“’ Oxygen@’ iat. %,
Treatment
M8terial
0.0025 ( 0.0028’~’ 0.017 0.023 0.0432
Siobium
Outgassed
Niobium-oxygen Niobium-oxygen Niobium-oxvnen
Anodized Anodized Anodized
Siobium-oxygen j?iiobium-oxygen Niobium-oxygen
Anodized 50 V Anodized 50 XVtwice” Anodized 50 T- four times’”
Niobium-oxygen
Anodized
5 J* 10 *(’ 20 V
50 V six times”’
0.072 0.141 0.2lW 0.282 10 .-393W’ 0.432
(a) Nitrogen, carbon and hydrogen contents were determined by internal friction (S) a&l chemical analysis (C, H) to be 0.002 & 0.001 at.%, 0.002 + 0.001 at. %, and <0.009 at.%, res ectively. 1!?1Substitutional impurity contents determined by Ledoux and Company to be <450 wt. ppm. Jlajor impurities included tantalum ((200 wt. ppm,), tungsten ( < 100 wt. ppm), h8fnium (
24
SOLID
SOLVTIOSS
of the well defined onset of macroscopic flow for virtually all of the specimens, none of the trends estab~shed in this investigation were itiuenced b\ the particular choice of the strength parameter. 3. EXPERIMENTAL 3.1.
RESULTS
Effects of puri$cation
The stress-strain curves of the purified niobium at various t,emperat~es are presented in Pig. 1. Three stage hardening is observed at temperatures between 250°K and 550°K. Above and below room t,emperature the yield stress changes with an attendant change in the work hardening characteristics. With increasing temperature above 298’K the lengt,h of Stage I decreases, the rate of work hardening in Stage IT increases and the stress at the onset. of Stage 311 increases. At temperatures below -250% the stressstrain curves are parabolic. with yielding followed b> relatively rapid non-linear work hardening. No yield points or twinning were observed down to 42°K;. In Fig. 2 the temperature dependence of the resolved shear stress of the high purity niobium is compared
-I
16
a
I
0. I
i
0.2
(
6
250°K
4-
2
. 298 lK ~~~~
0
FIG.
0
415-K 550’ K 0.4
625
0.8 1.2 SHEAR STRAIN 1. Resolved shrsr stress-shear strain curves at several temper8tures for high purity niobium single crystals.
ACTA
626
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40
4
CHRlSTl4N AND MASTERS (4)
Q
YlTCnCLL
e
sv, al. (161
DUESBERY f&D H1RSCW I171 PRESENT STUDY
VOL.
18,
1970
level at higher concentrations, a result typically observed in alloy hardening systems. In Fig. 2 there is an indication of such a saturation effect from the similarities in the temperature dependences of zone refined niobium from different laboratories.(**16) It appears that many investigators have mistakenly interpreted interstitial hardening as strictly athermal when in fact they have studied hardening in impure materials, i.e. above a few hundred at. ppm interstitials, for which the changes in temperature dependence had leveled off. 3.2. EflecEsof oxygen additions
100
I
200 TEMPERATURE
--
300 ’ K
FIG. 2. Temperature dependence of the resolvod shear stress for niobium single crystals of different base pu&,ies.“.l@,l7)
with that of niobium of different base purities, as compiled from several investigations. The data of Christian and Masters(*) and Mitchell et oZ.‘16) show the temperature dependence of the yield stress of zone refined single crystals of niobium which we estimate to have ~300-400 at. ppm of residual int,erstitials. Duesbery and Hirsch(17j have purified niobium by outgassing in vacuums of the order of 1O-1o torr. For these crystals the residual interstitial level is probably in the vicinity of 10 at. ppm. The purification clearly reduces the temperat’ure dependence of the shear stress of niobium over the entire temperature range. The differences in the shear stress increase with decreasing temperatures. One indication of the potency of the interstitials to act as hardeners is the observation that the shear stress at, 42°K (29 kg mm-2) of the niobium employed in this work is lower than that of the zone refined crystals(q*l”) at 77’K (~33 kg mms2). Stein and Low’@ have reported a similar dependence of the low temperature strength on interstitial concentration in high purit,y iron. From the compilations of Conrad and Hayes it is known that the temperature dependence of t.he flow stress does not change appreciably with interstitial
Typical stress-strain curves of purified niobium and niobium-oxygen alloys deformed at three of several temperatures used in this investigation are given in Fig. 3. More complete data are presented elsewhere.(16) Yield points were observed only in a few instances in the high concentration niobium-oxygen alloys at low temperatures. At room temperature and above, a monotonic increase in the yield stress with increasing oxygen content occurred and was a~eompanied by an increase in the lengt,h of Stage I, a decrease in the rate of work hardening in Stage II and an increase in the stress at the onset of Stage III. At low temperatures (<298”K) the shear stress of the dilute niobium-oxygen alloys is lower t$han that of the purified niobium. Coupled with this alloy softening the work hardening behavior undergoes a
FIG. 3. Resolved shear stress-shear strain curves for niobium-oxygen single crystals at 77”K, 185°K and 298°K.
RAT-I
\ \ \ \ - \
STREXGTH
GIBALA:
OF
NIOBIUM-OXYGES
\
\
*
aoo2s
0
0.011
‘1
0.141
IO432
\
X
\
EXTRAPOLATED WEAR STRESSES ,FROY FIG.51
\
e >
\
a
\
\ 1:\A i ’
\\ I\\\ 0
\\
v
10 -
\
\
‘A
~
x’, ,p\y \\
\
\x
:\\_
0.
0
I
100
p\ A -&_.-
\ \‘\ -A x\
‘-i_.
zoo
.I.
--_-_*--rC
I
300
SOLID
SOLCTIOA-S
627
of solute concentration at various t,emperatures, Fig. 5. At room temperature the shear stress increases with the square root or cube root of the oxygen conAt t,emperatures below 298°K alloy centration. softening is manifested as an initial decrease in the shear stress with oxygen addition followed by nonlinear hardening at, the higher oxygen levels. The maximum amount of alloy softening is observed to occur at different oxygen concentrations at different temperatures, the minimum generally moving toward lower concentrations with decreasing t.emperature. The rate of hardening following the alloy softening at 77°K reflects the large temperature dependence of the strength of niobium-oxygen alloys at, t~m~r~tures below ~100°K. This was also observed in Fig. 2 for niobium of different. impurity levels controlled in amount but not in type of interstitial.
AT%0
20 N
ASD
400
FIG. 4. Temperature dependence of the resolved shear stress of niobium and niobium-oxygen single crystals. The dotted curve represents extrapolated zero-oxygen stresses taken from Fig. 5.
The stress-strain curves change change with allo$ng. from it parabolic form exhibited by the unalloyed niobium to curves resembling a three stage form with yielding followed by low rates of work hardening in the dilute alloys of niobium. At high oxygen concentrations the shear stress increased with increasing oxygen and the stress-strain curves returned to a parabolic form and occasionally exhibited yield drops. At, 77°K only the first. addition of oxygen (0.017 at,.:,; cent) resuhed in B decrease in the shear stress. At, higher oxygen concentrations twinning was observed along with an increase in the yield stress. The temperature dependence of the shear stress of niobium and three of the niobium-oxygen alloys is shown in Fig. 4. At 298°K and above all the alloys display a higher shear stress than the pure niobium. At temperatures below 298°K the majority of the alloys display a lower shear stress than the unalloyed niobium at, some temperatures. For example, the niobium-O.017 at.YO oxygen alloy has a lower “yield stress than the pure niobium over the entire temperature range below ~250%, whereas the 0.141 at.94 oxygen alloy is stronger than the base niobium above -230°K and below ~1OO’K. This latter behavior is similar t,o behavior observed for Fe-N alloys(1r**9) which display alloy softening. The phenomenon of alloy softening is most strikingly displayed when the shear stress is plotted as a function
3.3. l’hermal activation analysis The activation enthalpy and activation volume for low temperature deformation behavior are often determined by the methods out.lined by Conrad(r) and others. These parameters are given by
and
Al.
X OXYSEN
Fro. 5. The concentration depcudence of the resofved shear stress of niobium-oxygen single crystals in the temperature range 77-3WK
ACTA
628
METALLURGICA,
VOL.
18,
1970
AT
4 0 + 0 *
0 0
I
I
100
200 TEMPERATURE
X
0
0.0025 0.017 0.023 0.141 0.282
300
4
0
‘K
FIG. 6. The temperature dependence of the strain rate sensitivity of niobium and niobium-oxygen single crystals.
respectively. Here T* = TV- r,,, T,,is the shear stress, T# is the long range athermal stress, r* is the thermally activatable portion of the total stress, 9 is the shear strain rate and &, = pvAb, where p is the mobile dislocation density, v is the attempt, frequency, A is the area swept out by each successful t’hermal fluctuation and b is the Burgers vector. k and T have their usual meaning, and the subscript p refers to T,, having t,he temperature dependence of the shear modulus p. The partial derivatives in equations (1) and (3) are l.5r
Al
Y 0
4 0.0025 -0 0.017 +
I.0 -
0.072
-& 0.432
(
determined from the strain rate sensitivity of the shear stress at constant temperature and the temperature dependence of the shear stress at constant strain rate. Representative examples of the temperature dependences of the strain rate sensitivit’\r of niobium and niobium-oxygen alloys are given in Fig. 6. Figures 7 and 8 show the activation enthalpy and the activation volume as a function of the effective stress T*. These parameters are functions of oxygen concentrations at low temperatures and at, all effective stress levels. The composition dependence of t,he activation parameters parallels the composition dependence of strength, i.e. the addition of small amounts of oxygen results in alloy softening and a reduction of thtb activation parameters (an increase in the peak temperature in the case of the strain rate sensitivity). whereas at the higher oxygen concentrations an increase in the shear stress is accompanied by a corresponding increase in the activation parameters and a decrease in the strain rate sensitivity peak temperature. These results are quite similar to those obt,aincd in comparable studies of alloy-softened substitutional allovs.(13*20*21) Since the skain rate sensitivitp and t,hcb activation parameters are a function of solute concentration, the rate determining step for the low temperature deformation of niobium is a function of the oxygen concentration of the metal. 4. DISUSSION
FIG.
7. Activation enthalpy vs. effective stress for niobium and niobium-oxygen single crystals.
In studies of the low temperature strength of the b.c.c. metals experimental results are often analyzed in terms of the predictions of various models of lattice
RAY1
ATD GIBALA:
STRESGTH
OF
SIOBIUM-OSTGES
SOLID
+ -0. 0 9*
6%
SOLI-TIOSS
AT YO 00025 0.017 0.0432 0.141 0.282
FIG. 8. Activation volume vs. effective stress for niobium and niobium-oxygen single crystals.
or impurity
hardening.
as evidence
for or against
mechanisms. theories such
These analyses are then used
Apart from objections
and the logic
analyses
investigation distorts
one of these two general
seem
inappropriate
the temperature To
enumerate
present arrived
and established
approach, the
present
of alloy softening
and compositional expected
depend-
for either type of
our discussion
and discuss separately
conclusions
for
because the occurrence
ences from that normally mode1.t
to the individual
of this analytical
succinctly,
we
each of the major
at in this study from our results
information
from the literature.
hardening.
The more potent the hardener,
the concentration be reached. “gradual”
solute interactio?l in
h.c.c. metals c.ol,stitutes n low temperawe
and interstitial
for concentrat,ions
The variations hardened
material
In general. Ar*lA ature.
is given
increasing
schematically
in Fig. 9.
purit,y results in a lower
In 1; and H and a higher v at, constant In
practical
terms.
these
not decrease or increase indefinitely concentration; values
T*, AT/A In 3,
of T(or T*) of an ideal solution
and
give
the
rrppenrance
temper-
paramet#ers as a function
they u-ill reach or approach of
T*, will of
limiting
athermal-type
t Nevertheless, some of these analyses were made as part
of this investigation.‘15’ It was found that the experimental results were in only modest agreement with lattice hardening Dorn-Rajnak)‘3’ and impurity hardening (e.g. (e.g. Fleischer)‘s’ models and then only at higher concentrations (22000 at. ppm), where neither type of theory should have validity.
metals.(22*23’
parameters reach limiting values of these solut.es ranging
from
a
upon the pot,ency
of bhe solute obstacle. There is no reason to believe b.c.c.
metals
cause
“rapid”
with
that’ interstitials
(or most, tetragonal hardening,)15) motion
a thermally
distortions
in
which
are such large barriers
t,hat, they should not also conactivatable
barrier.t
However
they are larger barriers than the spherically
symmetrical
of the parameters
H and 1’ as a function
such as substitutional
in close-packed
few to tens of an at.O,. depending
because
rcctivated dejormntio~~ n/echrinism
obstacles
solutes
The above mentioned
stitute
thermally
values will
This type of behavior is often observed for hardening
to dislocat,ion 4.1. Dislocntiorl-ijzterstitlnl
the lower
at which these limiting
quite
ones menConed low interstitial
above.
one must
concentrations
deal
in order
to observe the t’ypical alloy hardening behavior depict,ed in Fig. 9. Figures 3 and 4 demonstrate that this is the case for oxygen in niobium Low.‘6)
Lawley
et aZ.(25) for
ct
d..(g)
interstitials
in
all b.c.c.
metals.
least below (and perhaps ening behavior
is observed.
Smialek
molybdenum, In these systems,
interstitial
much below)
atomic ppm are necessary
and
iron.
tungsten and tantalum respectively. and probably
as do Stein and
Koo@)
levels
at
a few hundred
before normal
alloy hard-
For higher concentration
t In fact, advocates of intrinsic hardening theories have generally argued that interstitials are not potent enough hardeners t’o account for much of t,he low temperature strength.‘2*)
ACTA
630
METALLURGICA,
VOL.
18,
1970
~~_ -
FIG. 9.
----
PURE METAL
-
SOLID SOLUTIONS
----
LIMITING
VALUE
Schematicplots of T* and (AT/A In 3) as a function of temperature
and v and H as a function of effective stress for ideal solid solution hardened systems.
alloys, those used in the vast majority of investigations on b.c.c. “metals”, the parameters depicted in Fig. 9 tend toward their limiting values (the dotted lines) and give the appearance of a compositionindependent change of these parameters as a function of temperature or effective stress. 4.2. Dislocation-interstitial
interaction
is the
dominant low temperature hardening mechanism in b.c.c. metals
The present results suggest that if niobium has an intrinsic (Peierls) strength at O”K, it must be below the strength of the “pure” niobium used in the present experiments (~30 kg mm-2) and possibly as low as the strength of the unsoftened material suggested by the extrapolations shown in Fig. 4 and displayed as the dotted line (bottom) in Fig. 5 (~15 kg mm-2). Both of these apparent Peierls stresses are significantly below values quoted for zone refined and lesser purity niobium from other studies, e.g. >50 kg mme2. Obviously these latter values do not represent true Peierls stresses. The question we consider now is whether or not the 0°K strength of the niobium studied here represents the Peierls stress of niobium. We argue that it does not. Our base material contained approximately 20 at. ppm of residual nitrogen and perhaps as much carbon. From changes of lattice parameters,f2Q Snoek peak relaxation strengths(27*W) and room temperature hardening rates(29) with solute addition, both of these solutes
should be significantly more potent low temperature hardeners than oxygen. In Fig. 5, beyond the softening range of compositions, the hardening rate of oxygen is of the order of 10 p at 77°K. Thus, higher rates of hardening and the above stated compositions for nitrogen and carbon can easily account for the major part of the remnant hardening of niobium. Hence, the base niobium containing ~40 at. ppm of residual interstitials is “dirty” relative to the potency of the interstitial types remaining. Stein and LOW(~)have observed that reducing the residual carbon level in iron from >200 at. ppm to less than 0.03 at. ppm lowered the 0.2 per cent offset *yield stress by ~50 per cent from ~25 kg mm-2 to -13 kg mm-2 at 78°K. For such large apparent hardening rates ( 2 10 ,u) at very low concentrations, it is understandable that Stein and Low’s results have not easily been duplicated and that many investigations on other b.c.c. metals containing interstitials with comparable hardening rates have failed to disclose the decreased temperature dependence of T* to be observed at lower concentrations than the tens (more often hundreds or thousands) of at. ppm which the materials contained. 4.3. Alloy softening by interstitials solutes ix b.c.r. metals is caused by solute association (scavenging) Three theories have been advanced to account for the occurrence of solution softening in b.c.c. alloys at low temperatures. These are (1) a decrease in the
RAW
Ah‘D
STRESGTH
GIBALA:
OF
~IOBIU~f-OSYGES
Peierls potential as a result of alloying,c12*30) (2) an increase in the mobile dislocation density brought about by the presence of solute atoms,(ll*lg) and (3) interstitial gettering (in substitutional alloys).@‘.ss) A reduction in the Peierls stress has been visualized to occur in a number of ways. One proposalo2) is that the Peierls stress is reduced as a result of solute atoms producing a local disorder in the periodic lattice of the solvent metal. This tends to smear out the Peierls energy barrier resulting in dislocation mobility at lower stress levels than in an unperturbed crystal. The Peierls stress T@ can be expressed as(33)
Pb 7pNuzexp
-
2V12
( 1 c
(3)
where a can vary between $ and 1, b is the Burgers vector, y is the shear modulus, c fz b for screw dislocations and A is the “width” of the dislocation. According to this expression, TV will be reduced if p is reduced by alloying.~30) Jones et u&(~*)have determined the elastic constants of high purity niobium and a niobium-O.3 at.% oxygen alloy. The decrease they observed in the shear modulus due to alloying is too small to account for the amount of alloy softening observed in the Nb-0 system. The concept of dissociated dislocations has also been invoked to explain alloy softening.oa) Solute atoms can be visualized to affect dissociated dislocations in two distinct and opposite ways. If the temperature dependence of the strength of b.c.c. metals is due to dissociated dislocations which are sessile, then for sessile t.o glissile transformation to occur the dislocat,ions must constrict along their length. The presence of solute atoms is thought to aid such a constriction. making possible dislocation mobility at lower stress levels than in the absence of constricting sites represented by the solutes. The constriction, in effect. represents an increase of the stacking fault energy or a reduction of the width of the dissociated dislocation. Alternatively, solute addition could reduce the stacking fault energy, thus increase the width of the dislocations and according to equation (3) lower the Peierls stress. According to rate theory, the shear strain rate + that results upon the application of a stress to a crystal is given by (4)
Conradc2) and others’ll*lg) have suggested that for the same Peierls st,ress, a change in the pre-exponential
SOLID
SOLX-TIOSS
631
factor PO can change the shear stress r*. Since it(l = pAb, a change in POcan occur through a change in p. Nakada and Keh’il’ propose that alloy softening in the Fe-N system is a result of an increase in the mobile dislocation density due to alloying. ChristIss) and Christ and Smithog) have developed a model for solution softening in interstitial alloys by proposing that an increase in the mobile dislocation density occurs through solute atom induced cross slip of dislocations. This mechanism is presumed to occur when the repulsion between fixed interstitial dipoles in certain orientations and screw dislocations leads to movement of the dislocations onto cross-slip planes. The two theories, a reduction in t,he Peierls stress and an increase in the mobile dislocation density can not account for all the observed experimental facts associated with the phenomenon of solution softening. A reduction in the Peierls Stress assumes that the low temperature deformation of the unalloyed metal is governed by the intrinsic strength of the lattice. It was demonstrated in Section 4.2 that most of the remnant strength of the purified niobium can be accounted for by dislocation-in~rstitial interaction. An increase in the mobile dislocation density by alloying should not affect the activation enthalpy but should only change the pre-exponential factor PO. This is not found to be the case for Nb-0 solid solutions (Figs. 7, 8). An increase in the mobile dislocation density by cross slip of dislocations due to repulsion between interstitials and dislocations assumes that the minimum distance of approach between the dislocation and the interstitial for maximum repulsion is of the order of a few Burgers vectors. However at these distances the dislocation-interstitial interaction energy is at a substantial fraction of its maximum value ; hence interstitial pinning is probably more favorable than interstitial repulsion. Perhaps the strongest argument against intrinsic alloy softening is that the softening is not a reversible phenomcnon.‘14) Whereas addition of oxygen to purified niobium results in alloy softening at low tern. peratures (Fig. 4), removal of interstitials involving the same composition range does not result in alloy softening (Fig. 2). If alloy softening is due to a modiovation of an intrinsic phenomenon such as a reduction of the Peierls stress or a change in the mobile disloeation density. then it should occur when solutes are added or removed from the solvent. In the light of these facts, the mechanism that best explain the occurrence of solution softening is a scavenging type of process.(14) Alloy softening in substitutional systems has
632
ACTA
METALLURGICA,
sometimes been attributed to get&ring or scavenging of interstitials by substitutionals.‘31*32) This interpretation can readily be applied to substitutional alloys of metals in which the alloying element is a strong carbide, nitride, or oxide former. This theory has been rejected by most investigators as an explanation for alloy softening in other substitutional systems and in interstitial alloys for which interstitial compounds are unlikely to form. However, scavenging need not mean only the formation of a chemical compound. Association between substitutional and interstitial solutes and between interstitials of the same or different types in solid solution can readily occur and cause softening. Internal friction studies,(27.36~37) electron diffraction and microscopy,(as*39) field ion microscopy,‘*0) superconducting properties’41*42) and thermodynamic studies(46) attest to the occurrence of solute association (clustering) in the b.c.c. metals. Interstitial solute association in solution can lead to a decrease in the low temperature strength by a modification of the properties of the randomly distributed solute in the following ways: (1) Solute association reduces the total number of thermally activatable obstacles to dislocation motion over that which would be available if the interstitials were randomly distributed. (2) The effective hardening rate ar*/aC,, where Ci is the obstacle concentration of an associated pair (or larger cluster) of interstitials can be smaller than the sum of the hardening rates of the unassociated interstit,ials. This is reasonablesince clustering occurs with a reduction in the total free energy of the solid so1ution,‘36) of which strain energy should be a dominant contribution. (3) The temperat,ure dependent contribution to strengthening by associated interstitials can be different (and not necessarily athermal) from the temperature dependent contribution of unassociated interstitials. The change in the strain rate sensitivity peak temperature with alloying in Fig. 6 is a strong indication of the modification of the interstitial contribution to the temperature dependent part of the shear stress by clustering. An ideal alloy hardened system would display an increase in the strain rate sensitivity peak at about the same temperature with increasing solute as in Fig. 9. In the niobium-oxygen system the peak heights for the unalloyed niobium and the niobiumoxygen alloys remain relatively constant, while the peak shifts first to higher temperatures wit,11alloy softening (reflecting the different hardening rate of the associated solutes) and then to lower temperatures with alloy hardening. Corresponding changes in the
VOL.
18,
1910
activation enthalpy and the activation volume reflect the changes in the strain rate sensitivity and results in a deviation of the activation parameters from the ideal behavior depicted in Fig. 9. There remains the question regarding the specific nature of the clusters envisaged to occur and cause the alloy softening. Clearly, oxygen atoms must be clustering with each other or with other interstitials such as nitrogen, carbon or hydrogen. The residual substitutional levels of the purified niobium are too small to account for the observed magnitude and the rate of alloy softening by substitutional-oxygen association. Oxygen clustering alone probably can not account for the observed alloy softening either. The oxygen Snoek peak heights and the residual resistivity increase monotonically with increasing oxygen concentration. These parameters give no indication that a large amount of the oxygen is present in clusters or pre-precipitates at the lower concentration levels. This would be required to explain the immediacy, the magnitude and the large rate of alloy softening with initial additions of oxygen relative to the subsequent alloy hardening at a given temperature (e.g. 113’K in Fig. 5). It would seem that, oxygen, present as the dominant impurity, also associates with other residual interstitials which are more potent hardeners, i.e. nitrogen or carbon, and reduces their effectiveness as thermally activatable obstacles to dislocation motion. Of these, oxygen-nitrogen association in niobium has been observed by internal friction techniques,‘37*43) is thought to be characterized by a larger binding enthalpy than oxygenoxygen association,(43) and represents the most probable cause of the alloy softening. Hydrogen can also associate with oxygen, but is not likely to be a potent enough hardener to cause appreciable softening when clustered. 4.4. Alloy softening by substitutional also a scavenging process
solutes is
This conclusion was drawn in an earlier papero4) on the basis that the extent of alloy softening caused by substitutional solutes was determined largely by the purity, relative to residual interstitials, of the base materials used to prepare the alloys. Here we have noted in addition that both the activation enthalpies and the activation volumes for these systems are functions of solute concentration(ia*sO.sii in the same manner observed for an interstitial alloy, niobiumoxygen. and as expected for a scavenging-type of process. These results are not well explained by intrinsic softening mechanisms, whether by a reduction
RAVI
ASD
STRESGTH
GIBALA:
OF
of the Peierls stress or an increase in the mobile location density. Finally we re-emphasize stitial association association)
dis-
that substitutional-inter-
(as well as interstitial-interstitial
in solid
phenomenon
SIOBIUM-OSYGES
solution
is a well documented
from internal friction,‘44) field and thermodynamic studies.(46’
microscopy(J5) solute association
to cause alloy softening
necessary that the t,wo species undergoing cause appreciably
different hardening
ion For
it is only association
rates and tem-
perature dependences of the flow stress of the base metal. This is virtually always true in b.c.c. metals containing
gradual
hardening
rapid hardening interstitials. that substitutional
The exceptions
substitutional-interstitial
l-4
occur when the
int,eraction
small or the alloy is relatively
b.c.c.
and
softening is almost always observed
in b.c.c. alloy systems.
4.5. Gonclusio)~s
substitutionals
Thus, it is not surprising
is
extremely
free of interstitials.
nre true in general
for
all
metals
It might be argued that the present results obt,ained for
niobium-oxygen.
metalf3”)
among
involving
the b.c.c.
relativel;v low potency in
an
elastically
transition
metals
soft and a
obstacle among the interstitials
b.c.c.
metals,?
are
not
hardened
systems.
We
believe
typical
of
interstitial
otherwise.
Indeed,
t,he results suggest that if a relatively
weak hardener
in a soft metal
r*, then more
potent
can effect a sizeable
hardeners
vanadium.
tantalum,
and tungsten
observed
iron,
harder metals such as
chromium,
An indication
of the correctness of this
is the very large amount in tungsten.‘“‘)
of alloy softening which
and chrom-
The large amounts
infer that the residual interstitials
hardeners.
rhenium
of alloy softening
molpbdenum’48)
ium(4s) when allored with rhenium. are pot,ent
molybdenum
should effect larger r*‘s at much lower
concent’rations. position
in elastically
softens
For to
tungsten
in particular,
concentrations
than 20 at.Ob. large dislocation-interstitial
of
more
interaction
energies (-1 eV) are expected and appreciable association between rhenium and oxygen is well docu-
REFERENCES 1. J. E. DORS, Dislocation Dynamics, p. 27. McGraw-Hill (1968). 2. H. COSR~D, The Relaiion Between the Structure and Xechanical Propertias of MetaEs, p, 475. H.M.S.O. (1963). 3. J. E. DORS and S. RAJSAK, Trans. Am. Inst. Min. Engrv 230, 1052 (1964). 4. J. W. CHRISTIAS and B. C. XASTERS, Proc. R. Sot. A281, 223 (1964). 5. R. L. FLEISCHER, Acta Uet. 15, 1513 (1967). 6. D. F. STEIS and J. R. Low, JR., Acta Xet. 14,1183 (1966). 7. H. CONRAD, J. Metala 16. 582 (1964). 8. R. C. Koo, Acta Met. 11, 1083 (1963). 9. A. LAUZET. J. TAS DER S~PE and R. MADDIS, J. In&. Metals 91, 23 (1962-3). and S. J. 10. W. C. LESLIE, R. J. SOBER, S. G. Ba~coclc GREES. Trans dm. Sot. Netals. 62. 690 (1969). T. XA&ADA and A. S. KEH. Acta ;cet. 16; 903 ‘( 1968). :;: P. L. RAFFO and T. E. MITCHELL, Trans. Am. Inst. Mirr. Engrs 242, 905 (1968). 13. R. J. ARSESA~LT, Acta Net. 17. 1291 (1969). 14. Ii. V. RA~I and R. GIBALA, Scripta Met. 3, 54i (1969). 15. K. V. RAVI, Ph.D. Thesis, Case Western Reserve Univer-
sity, September 1969.
16. T. E. MITCHELL, R. A. FOXALL and P. B. HIRSCH, Phil. Maq. 8, 1895 (1963). 17. 11. $. DUESBER~ and P. B. HIRSCH, Dislocation Dynamics, p. 57. McGraw-Hill (1968). See also R. A. FOSALL, Ph.D. Thesis, Univ. of Cambridge (1967). 18. H. CONRAD and R. HATES, Trans. dm. Sot. hfetala 56, 128 (1963). 19. B. W. CHRIST and G. T-. SXITH, >1bm. Sciettt. Revue M&all. 65, 208 (1968). 20. A. A. HENDRICGSON, B. C. PETERS and R. A. STRAHL, U.S. Atomic Energy Commission Technical Report No. COO-916-13, March 1968. B. C. PETERS and A. A. HENDRICKSON, J. Xetals 21, 116A (1969). 21. G. C. DAS and R. J. ARSENATJLT, Scripta 111et. 2, 495 (1968). 22. K. R. EVASS and R. F. FLAKAGAS, Phil. .liag. 18, 977 (1968). 23. D. T. PETERSOX and R. L. SKAQGS, Trails. Am. Inst. Min. Engrs 242, 922 (1968). 24. J. W. CHRISTIAN, Scripta Met. 2, 569 (1968). 25. R. L. SJIIALEI(. G. L. WEBB and T. E. MITCHELL, Scripta xet. 4, 33 (1970). 26. A. TAYLOR and S. J. DOYLE, J. less-common Metals 13,
313, 399 (1967). 27. R. W.
POWERS and M. 1.. DOYLE, J. appl. Phys.
30, 514
(1959).
28. P. M. ROBINSOH and R. RA\~LIXQS, Iron Steel 31, 3, 66, 97 (1958). 29. Z. C. SZKOPIAIC.J. less-common Metals. 19, 93 (1969). R. J. ARSESAULT, Acta filet. 15, 501 (1967). 7:: H. H. KRAXZLEIX, RI. S. BURTON and G. V. %ITH, Trans.
Am. Inst. Min. Engrs 233,69
(1965).
Trans. Am. Inst. Min. Engrs 253, 1500 (1965). J. WEERTMAN and J. R. WEERTMAS, Elementary Dislocation Theory, p. 159. Macmillan (1964). K. A. JOKES, 8. C. Moss and Ii. M. ROSE, Acta Met. 17, 365 (1969). H. W. CHRIST, Acta Met. 17. 131i (1969). R. GIB~LA and C. A. WERT, Arta Net. 14, 1095 1105
32. S. S. STOLOFF, R. G. DAVIES and R. C. Ku, 33. 34. 35. 36.
Rep. 19, 505 (1964). 38. J. VAN LAXD~X-T and S. A~IELIXCKX, Sppl.
ACKNOWLEDGEMENT
work
633
SOLI-TIOSS
(1966). 37. D. J. VAX O~IJEX and A. 6. PAX DER GOOT, Philips
mented.‘J5)
This
SOLID
was supported
At,omic Energy _4T (ll-l)-1676.
Commission.
by
the
under
United Contract
States pie.
t Kate for example from the tabulation of Snoek relaxation stren@hs . . by Powers and .. Doyle’*” that the strength .of Nb-0
relaxat,lon 1s the smallest among the systems mvolvmg the __ . . __ group 1 -A metals and won and the mterstitials 0, N and C.
Rss.
Phys. Lett. 4,
15 (1964). 39. J. VAS LAKDUTT, Phyls. Status Solidi 6, 957 (1964). 40. S. XAI
236,924 (1966). 44. P. M. BUNN, I~. G. C’UMMINOS and H. W. LEAVESWORTH, *JR., J. appl. Phy.7. 33, 3009 (1962).
634
ACTA
METALLURGICA,
45. D. T. NOVICKand E. S. MACHLI~-, Trans. Am. Sot. Metals 61, 777 (1968). 46. H. I. AARONSON,H. A. DOMIANand G. M. POVSD, Trans. Am. Inst. Min. &?r8 256, 753, 768 (1966).
VOL.
18,
1950
47. P. L. RAFFO, J. less-common Metals 17, 133 (1969). 48. D. L. DAVIDSONand F. R. BROTZEB,J. hlelds 21, 116A (1969). 49. A. GILBERT,J. Metak 21, 89A (1969).