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Oxidation Characteristics

of Ti=14Al-21Nb Alloy

Sankara N. Sankaran,

Ronald IC Clark,

JaIaiah Unnam,

and Karl E. Wiedemann

i Ti-14ll-21Nb (NA_A-lP-3012) INGOT_X!DATI_NLLOY C_4AOACTCFKISTICS (NASA) 24 p OF

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NASA

Technical

Paper

3012

1990

National Aeronautic,sand

Space Administration

Office of Management

Scientific and Technical

Information Divi_sion

Oxidation Characteristics

of Ti-14A1-21Nb Alloy

Sankara N. Sankaran

Analytical Services & Materials, Inc.

Hampton, Virginia

Ronald K. Clark

Langley Research Center

Hampton, Virginia

Jalaiah Unnam

and Karl E. Wiedemann

Analytical Services & Materials, Inc.

Hampton, Virginia

Summary

The static oxidation kinetics of Ti-14A1-21Nb

(percent weight) ingot alloy were studied in air over

the temperature range from 649°C to 1093°C in a

thermogravimetrie apparatus. The oxidation products were characterized by X-ray diffraction, electron

microprobe analysis, energy dispersive X-ray analysis, and Auger electron spectroscopy. Cross sections

of the oxidized samples were also examined using

light and scanning electron microscopy. The oxidation rate was substantially lower than the conventional alloys of titanium, but the kinetics displayed

a complex behavior involving two or more oxidation

rates depending on the temperature and duration of

exposure. The primary oxide formed was TiO2, but

this oxide was doped with Nb. Small amounts of

A1203 and TiN were also present in the scale. The

diffusion of oxygen into the alloy was observed, and

the diffusivity seemed to be dependent on the microstructure of the metal. A model has been presented to explain the oxidation behavior of the alloy

in terms of the reduction in the oxygen diffusivity in

the oxide caused by the modification of the defect

structure of TiO2 by Nb ions.

Introduction

Titanium aluminides are candidate materials for

use in hypersonic structures because of their high

specific strength and superior high-temperature properties compared with conventional titanium alloys

and superalloys. These exceptional properties are the

result of the ordered microstrueture. However, like

most intermetallies, these alloys suffer from a lack of

adequate room-temperature ductility. Variations in

composition through the addition of other elements

such as niobium, vanadium, and molybdenum are being attempted to overcome this problem. The alloy

Ti-14A1-21Nb (percent weight) is one such modification of the intermetallie Ti3A1 in which Nb has been

added to improve the room-temperature ductility.

Hypersonic structures are expected to be exposed

at high temperatures to a variety of environments ineluding oxygen, and thus oxidation resistance of the

candidate alloys is of importance. Although some

mechanistic studies on the role of aluminum and niobium on the oxidation of titanium have been reported

(refs. 1-5), only a few limited studies (refs. 6 and 7)

on the oxidation behavior of titanium aluminides

have appeared in the open literature. The purpose of

the present study was, therefore, to characterize the

oxidation behavior of Ti-14A1-21Nb ingot alloy.

The oxidation of metals such as titanium and zirconium is characterized by both oxide formation and

concurrent dissolution of gas into the metal (ref. 8).

Models applied to partition the total oxygen uptake

during oxidation exposure into these two components

have been described by several authors (refs. 9 12).

Even though these models are substantially true in

explaining the oxidation of the alloys of titanium as

well, they do not provide an adequate basis to account for the role of the alloying elements in controlling the oxidation rate. Accordingly, in the present

study, a theoretical analysis has been developed to

integrate Wagner's theory (ref. 13) of oxidation into

a model for titanium oxidation developed by Unnam

et al. (ref. 12) so that the defect structure of the oxide lattice can be taken into account. This approach

enables one to determine the effect of alloying elements as dopants in controlling the oxidation kinetics of alloys. The model has been applied to predict

the oxidation behavior of Ti-14A1-21Nb alloy, and the

results have been compared with experimental data.

Oxidation kinetics were measured in laboratory

air using thermogravimetry and covered a temperature range from 649°C to 1093°C. The oxidized specimens were examined using X-ray diffraction (XRD),

light and scanning electron microscopy (SEM), microprobe X-ray analysis, energy dispersive X-ray

analysis (EDS), and Auger electron spectroscopy

(AES) to characterize the oxidation products and

microstructural changes in the alloy.

Theory

Figure 1 shows a schematic diagram of the oxidation process arid emphasizes that, at equilibrium,

the chemical potential of oxygen is continuous across

a two-phase interface even though the concentration

gradients can show a discontinuity. The formation of

an oxide scale of finite thickness on the metal implies

that the oxide growth rate is faster than the dissolution process. Therefore, the total oxidation rate is

essentially governed by the transport characteristics

of the oxide compound (refs. 13 and 14), so that

d m) _kr (_= MO

where

d

( )total

Z

kr

m

Mo

Z2

(1)

flux of oxygen uptake, g/cm2/sec

instantaneous oxide thickness, cm

rational rate constant defined by

Wagner (ref. 14), eq/cm/sec (where

eq in the unit denotes equivalents)

equivalent weight of oxygen, g/eq

atomic weight of oxygen, g

valency of oxygen ions in the oxide
Understeady-stateconditions,kr is a constant and

is a fundamental property of the oxide compound.

On the other hand, the thickness of the oxide formed

on metals mmh as titanium and their alloys, which

have a high solubility for oxygen, is governed by

the partitioning of the oxygen toward the dissolution process. Thus, a complete description of the total oxygen uptake during the oxidation of such metals/alloys requires a knowledge of two parameters:

the rational rate constant kr and the instantaneous

oxide thickness z.

Combining the defect structure properties of the

oxide compound with its thermodynamic and electrochemical properties, and making certain simplifying

assumptions, allowed Wagner (ref. 14) to derive the

rational rate constant kr as

can be writtcn where

#_ standard chemical potential of

oxygen

a O activity of oxygen

PO_ partial pressure of oxygen in gas

phase with which the solid is in

equilibrium

Accordingly, kr can be redefined as

/_O2 _2 ( Wl )

kr - C_ P_'

2 _Z_ D1 +D2 dlnPo2

where

(4)

where

- RT _ _D1 + D2 dp 0

p,!

02

(2)

p if

02

partial pressure of oxygen at

oxide/gas interface

dissociation pressure of oxygen

corresponding to oxide/metal (or

outer-oxide/inner-oxide) equilibrium

Ceq

Zl

Z2

!

#0

tl

#o

#o

D1

02

R

T

equilibrium concentration of nmtal

ions or oxygen ions, eq/cm 3

valence of metal ions

valence of oxygen ions

chemical potential of oxygen at

oxide/gas interface

chemical potential of oxygen at

oxide/metal interface (or outeroxide/inner-oxide interface in the

case of multiple oxides)

chemical potential of oxygen at any

arbitrary position in the oxide

self-diffusion coefficient of metal

ions in the oxide, cm2/sec

self-diffusion coefficient of oxygen

ions in oxide, cm2/sec

universal gas constant, cal/mol-K

absolute temperature, K

If one expresses the chemical potentials in terms

of activities and notes that activities can be expressed

in terms of the partial pressure of oxygen for a solid

in equilibrium with a diatomic gas, the equation

PO = g_) + RT ln aO

pl/2,_

(ao --_ 02 J (3)

The significance of this definition for the rate constant kr is that it can be predicted from independently measured values of the self-diffusion coefficients of the metal ions and oxygen ions in the oxide

compound and also from the thermodynamic properties of the oxide. Unfortunately, since migration of

ions in ionic compounds occurs through defect sites

in the lattice, and the defect concentrations are dependent on the nonmetal activity (ref. 15), the diffusion coefficients are a function of the oxygen pressure

and such data are not readily available. However, for

many oxides, tracer-diffusion-coefficient data after

equilibration to a specific oxygen pressure (usually

pure oxygen at 1 atm) are available and the defect

structure data are also known; thus, from a knowledge of the defect equilibria and the dominant defect

in the oxide, the oxygen pressurc dependence of the

diffusion coefficients of the ions can be established.

TiO2 is known to be an n-type semiconductor

(refs. 15 and 16), but disagreements exist on whether

interstitial cations or oxygen vacancies are the predominant defects in the oxide. Based on studies of

nonstoichiometry of TiO2 as a function of temperature and oxygen partial pressure, Kofstad (ref. 16)

concluded that the defect structure of TiO2 is composed of both types of defects. He further determined

that the oxygen vacancies are doubly ionized and

the interstitial ions are composed of two kinds: Ti 3+

and Ti 4+. However, marker experiments performed

during the oxidation of titanium between 800°C and

1200°C (refs. 6 and 8) reveal that oxygen ions are the

mobile species in TiO2, implying that doubly ionized

2
oxygenion vacanciesarethepredominantdefectin

the oxideandthat the diffusivityof oxygenionsis

largerthanthat of titanium ions(i.e., D2 >> D1).

Prom a mechanistic point of view, D2 can be related

to the concentration of doubly ionized oxygen ion as

where

D V diffusivity of oxygen ion vacancies

in TiO2

[Vo] concentration of oxygen ion

vacancies

The designation of the symbol V0 for doubly ionized oxygen vacancies follows the notation developed

by Kr5ger and Vink (ref. 17). In this notation, the

symbol denotes the defect species, the subscript denotes the site, and the superscript designates the effective charge of the defect relative to the lattice. For

dilute concentration of vacancies, it can be assumed

that no defect interactions occur and that D V is independent of the oxygen pressure. Thus, equation (4)

can be further reduced to

kr = DV J Po_ Y 0 d In PO2

(6)

Thus, any oxygen pressure dependence of the oxidation kinetics of oxygen-deficient oxides is brought

about by the oxygen-pressure-dependent concentration of oxygen ion vacancies in the lattice.

The formation of doubly ionized oxygen vacancies

in the TiO2 lattice can be described through the

following equilibria (ref. 15):

lo} Oo _v_+ _ 2

v_ _ y6 + e' (7)

v6-- vo + e'

and by combining the above equations and using the

laws of mass action:

, }

O0_V 0+_02+2e

K--_ p)/202 " [Vo] "n2

(8)

where

O O oxygen ion occupying a normal site

in the oxide

V_ neutral oxygen ion vacancy in which

both electrons are associated with

the vacancy

V0 singly ionized oxygen ion vacancy

in which only one of the electrons is

associated with the vacancy

V6 doubly ionized oxygen ion vacancy

from which both electrons are freed

e' electron

n concentration of electrons

K equilibrium constant

The electrical neutrality requirements of the crystal

dictate that the positive and negative charges be balanced, i.e., that2 [170] = n. Thus, the concentration

of the vacancies can be written as

K1/3

[Vo]- 41/3 PO:/6 (9)

Since K is a constant at any given temperature, this

relationship establishes the oxygen pressure dependeuce of the concentration of oxygen ion vacancies in

a Pure TiO2 crystal.

In a doped crystal of TiO2, the oxygen pressure

dependence of the vacancies can be altered by the

dopant ion, depending on its valency and concentration in the oxide. Thus, if Nb205, a higher-valent

cation, dissolves in TiO2, the process results in the

annihilation of oxygen vacancies through the reaction

Nb205 + V0 _ 2NbTi + 500 (10)

Here, Nb_ i is the Nb 5+ ion occupying a normal Ti 4+

site, whic_ has a unit positive charge relative to the

lattice. The electroneutrality condition for this case

becomes

2 [170] + [Nb,ipi] =n (11)

W on[,%]>> [Nbd thocon

centration of vacancies is obtained from equation (8)

after appropriate substitution. Thus,

K p-l/2

P°]-tTJ N --i (12)

Onthoother if[.0]>>INsTil,

the native defect concentration is greater than the

concentration of the dopant, then 2 [170] _ n, from

3
equation(11),andthevacancyconcentrationisgiven

byequation(9). Thus,in thiscasethedopantdoes

notcontrolthedefectstructureoftheoxide.

In ananalogousmanner,it canbeshownthat the

additionof a lower-valentcationto theTiO2lattice

will resultin the creationof oxygenion vacancies.

Forexample,if A1203is dissolvedintoTiO2,it will

resultin thegenerationof oxygenvacanciesthrough

thereaction

A1203_ 2Al_ri + 300 + V0 (13)

Here, Al{ri is the A13+ ion occupying a normal Ti 4+

site that has a unit negative charge relative to the

lattice. The electroneutrality requirement for this

case can be written as

(14)

Again, two possibilities arise: if the dopant concentration is too small, i.e., if n >> [Al{ri ], then

_2 IV0] and the vacancy concentration is the n

same as that of a pure TiO2 crystal as given by equation (9). However, if [A1wi q >> n, then

[V0] = [Al_ri ] (15)

Equations (9), (12), and (15) provide a basis to

determine the equilibrium concentration of oxygen

ion vacancies in undoped or doped TiO2. In principlc, if one knows the concentration of oxygen ion

vacancies, the rate constant for oxidation can be predicted from equation (6). However, a major difficulty

in evaluating the rate constant from equation (6) is

that D V (the diffusion coefficient for the vacancies in

TiO2) is not a measurable parameter. However, since

independent measurements of tracer diffusion coefficient of oxygen in pure TiO2 equilibrated to 1 atm

oxygen are available in the literature, D V can be extracted from such data using equations (5) and (9)

as

D V -- 41/30; (16)

fK1/3

where

D2

/

tracer diffusion coefficient of oxygen

in TiO2

correlation factor for diffusion

Making a simplifying assumption that D V is not

significantly affected by the small concentration of

dopant ions in the TiO2 lattice allows the rational

4

rate constant to be determined for each of the following three cases by appropriate substitutions into

equation (6) and integration:

(1) For pure TiO2,

3Ceq [(p/, _-1/6

kr - f D_ [_ O2J _ (rO2)-1/6 j.,] (17a)

(2) For TiO2 doped with Nb 5+,

kr = 41/3CeqD_ K2/3 [(p_)-l/2

f [NbTi] 2

(17b)

-(pb2)-l/2]

(3) For TiO2 doped with A13+,

-1/6CeqD2 [AIITi]* [ P']02

kr 4 ] K1/3 In [ O2J

(17c)

Since the dissociation pressure of titanium oxides is

extremely low, i.e., P/o2 >> P/'O2' the oxide growth

rate becomes independent of the oxygen pressure in

the environment.

Having established a relationship to compute the

rational rate constant kr in terms of the fundamental

properties of the oxide compound, equations (17) in

combination with equation (1) can be used to predict

the oxidation rate of titanium and its alloys. However, solution of equation (1) still requires a knowledge of the instantaneous thickness of the oxide.

With suitable conversion of units, the oxide thickness can be expressed in terms of the weight gain

components:

Z---- (__) [(A)tota]- (A)a]loy ] (18)

where

(m/A)total total weight gain associated

with oxidation process, g/cm 2

(m/A)alloy component of weight gain due

to dissolution process, g/cm 2

V equivalent volume of oxide

compound, cm3/eq

Substitution of equation (18) into equation (1) yields

the differential equation governing the oxidation rate
ofmetals/alloysthat exhibitconcurrentoxideformation andoxygendissolution:

d m2kr

_ (A)t°tat= V[(_)total- (_)alloy ]

(19)

The solution of equation (19) to determine the

time dependence of the oxidation rate in terms of

the known parameters requires some simplifications

to be made. First, assuming that diffusive processes

are rate controlling, the oxidation rate follows a

parabolic behavior:

A total = ]

d_d (Am) total 2t]/2 kp /

(20)

where

kp weight-gain parabolic-rate constant,

g/cm2/sec 1/2

t duration of oxidation exposure, see

Second, as adequately described by Unnam et al. in

reference 12, an independent solution is available for

determining the oxygen dissolution in titanium and

its alloys in terms of the fundamental properties of

the metal/alloy; consequently,

m

(_)alloy ----kdtl/2 (21)

with

where

kd= 2fo(Csl-Co) (_-) 1/2

kd

Csl

Co

Io

D8

rate constant pertaining to oxygen dissolution in metal/alloy,

g/cm2/secl/2

concentration of oxygen in the alloy

at alloy/oxide interface, atomic

percent

oxygen concentration deep in the

alloy, atomic percent

conversion factor to express concentrations in grams per cubic

centimeter

oxygen diffusivity in the alloy,

cm2/sec

duration of exposure, sec

Substitution of equations (20) and (21) into (19)

leads to the solution

kd-l-(k2-t-_)l/2tl/2

total 2

(22)

Thus,

kp= 2

The significance of the development of equation (22) is that the overall weight-gain parabolicrate constant kp is directly related to fundamental constants that describe the transport properties in the metal/alloy and in the oxide compound.

Two limiting conditions can be observed from equation (23): in the first case, if k d >> kr, then

kp = kd (24a)

In this case the oxygen-uptake-rate constant is

solely governed by the dissolution process in the

metal/alloy. In the second case, if kr >> kd, then

(24b)

In this case the oxygen-uptake-rate constant is essentially controlled by the transport through the oxide

compound.

How well the kinetic model described up to this

time adequately explains the oxidation behavior of

Ti-14Al-21Nb alloy can be corroborated by comparing the experimental rate constant kp determined

from the TGA data with that computed from equations (23) and (24).

Experimental Details

Specimens and Materials

The chemical analysis of the hot-rolled Ti-14A1-

21Nb ingot alloy is shown in table I. Thermogravimetric analysis (TGA) samples (approximately 1 cm

by 1.5 cm by 0.2 cm with a nominal surface area of

4 cm 2) were cut from the ingot alloy using a slowspeed diamond saw. A 0.16-cm-diameter hole was

drilled at one end of the sample to facilitate hanging from the microbalance during the oxidation exposure. Prior to the test, the samples were ground to

a uniform surface finish by grinding with silicon carbide papers ranging from 240 grit (60 #m was the average grinding particle size) through 600 grit (15 ttm
was the average grinding particle size). After grinding, the dimensions of the samples were measured to

the nearest 0.001 cm: The samples were then thoroughly cleaned with soap solution, acetone, and ethyl

alcohol, and then they were dried in air. The initial

weights of the cleaned samples were recorded to the

nearest 0.05 mg.

Oxidation Studies

Oxidation tests were conducted in laboratory air

over the temperature range from 649°C to 1093°C in

two types of TGA apparatus. For the temperature

range from 649°C to 871°C, the tests were performed

in an apparatus in which the weight of the sample

was monitored with a Cahn R-100 microbalance. Although this apparatus had the sensitivity required for

even small weight changes, it was not designed for use

at temperatures above 900°C. Therefore, for the temperature range from 871°C to 1093°C, the tests were

performed in a Netzsch simultaneous thermal analysis (STA) unit, which is a combination TGA/DTA

(differential thermal analysis) apparatus. Correlation of the data obtained from the two apparatus

units was confirmed through the overlapping data

from both apparatus units at 871°C.

In conducting these tests, the following sequence

of steps was followed: first, the sample was introduced into the TGA setup; second, the microbalance

zeroed; and third, the furnace temperature raised

from room temperature to the test temperature over

a period of about 30 minutes. The change in the

weight of the sample was monitored from the start

of the furnace heat-up. The temperature of the sample was monitored through a separate thermocouple

located about 0.1 cm below the sample.

At each temperature, duplicate or triplicate tests

were performed to ensure reproducibility in the data.

Most samples were oxidized for a period of about

25 hr at temperature, but in some cases the oxidation

runs were continued for a period of about 120 hr to

study the effects of long-term exposure. At the end

of each test, the sample was cooled rapidly by cutting

off the electrical power to the furnace. At the end of

each test, the final weight of the sample was recorded

to the nearest 0.05 mg.

Characterization Studies

After oxidatlon, selected samples were analyzed

by XRD, electron microprobe analysis, EDS, and

AES. Thc cross sections of a few samples were

examined using convention£i light microscopy and

SEM. Cross-sectional metallographic samples were

prepared by first wrapping them in aluminum foil,

hot mounting them in a special edge-retaining resin,

6

and utilizing the usual polishing techniques. A few of

the samples were etched with Kroll's reagent (2 percent HNOa + 1 percent HF + 97 percent H20)

to study the effect of oxidation on the alloy microstructure. Oxide thicknesses were determined from

photomicrographs of the cross sections.

Results and Discussion

Weight Gain Kinetics

Figure 2 shows a typical output of the TGA

data plotted in parabolic coordinates, (with weight

gain normalized to unit area of the sample as a

function of the square root of time). The apparent

weight loss and nonlinearity in the data early in the

exposure period result from transient temperature

and buoyancy effects. These effects vanish within

1 hr from the start of a test.

Because of these transient effects, the TGA data

for the first hour were discarded and the time scale

was adjusted by replacing the heat-up time with an

equivalent time at oxidation temperature, computed

using an estimated activation energy for the oxidation process. Further, the total weight gain determined from the initial and final weights of the sample was used to adjust the TGA data to thc true

weight gain. Figure 3(a) shows the final TGA data

for the samples oxidized from 649°C to 871°C, and

figure 3(b) shows the data for the temperature range

from 871°C to 1093°C.

From figure 3 it can be surmised that the oxidation of this alloy follows parabolic kinetics, but it is

characterized by two or more slopes separated from

each other by distinct transitions. It appears that

these transitions are marked by specific weight gain

values. Also, the presence or absence of these transitions is dependent on the temperature and duration

of exposure at temperature. For instance, the first

transition at 0.25 mg/cm 2 is distinctly noted for the

samples exposed at 760°C and 816°C, whereas both

the first and second transitions (at 0.5 mg/cm 2) are

pronounced only for the sample exposed at 871°C.

None of these transitions is observable for the samples oxidized at 649°C and 704°C; it is evident from

figure 3(a) that the oxidation duration at these temperatures was not long enough to approach the specific weight gain needed for these transitions. For the

samples oxidized at higher temperatures (982°C and

1093°C), the weight gain rates were so high (fig. 3(b))

that these transitions could not be identified from the

TGA data. It is safe to assume that these transitions

must have occurred at the very early stage of oxidation at these temperatures, with the result that the

kinetics of these samples were dominated by the third

slope.
The measuredrateconstants,i.e.,the slopesof

the weightgaindatafrom figure3, as wellasthe

otherrelevantdetailsfor eachof thetestsarelisted

in tableII andplottedin figure4. Theresultshighlight the complexityin the oxidationbehaviorof

this alloy and suggesthe existenceof severaloperatingmechanismsdependingon the temperature

anddurationof exposure.Forexample,thestageI

rateconstantshowanactivationenergyof approximately85kcal/moloverthetemperaturerangefrom

649°Cto 760°C,but they approacha valueof approximately54kcal/moloverthetemperaturerange

from760°Cto 871°C.Ontheotherhand,thestageII

and stageIII kineticsshowactivationenergiesof

54kcal/moland64kcal/mol,respectively.

Characteristicsof Oxide Layer

Samplesexposedto temperaturesup to 760°C

retainedthe specularfinish of the metal surface,

but the colorof the metalchangedgraduallyfrom

goldento blackasthetemperatureof exposurewas

increased.Ontheotherhand,thecolorof thescales

changedfrom blackto dark grayas the exposure

temperaturewasincreasedfrom816°Cto 982°C.The

samplesoxidizedat 1093°Cspalled,perhapsduring

thecoolingcycleafteroxidationsincetheTGAdata

did not revealanyabnormalitiesusuallyassociated

with spalling.Theoutersurfaceof thespalledoxide

appearedgrayin color,but the innersurfaceof the

spalledoxideand the residualoxideon the metal

appearedblackin color.

XRD analysisof the scalesfrom a fewsamples

representingthe threekinetic stagesof figures3

and 4 is summarizedin table III. TiO2, TiN, and

A1203weretheprincipalphasesobservedin theoxide. The relativeproportionsof thesecompounds

in the scalewerecalculatedfrom the measuredXraypeakheightsafterapplyingtheappropriatelinearreflectivitycoefficients.Thelatterparameterwas

calculatedfromthestructurefactorsfor eachmaterial andreflection.In performingthesecalculations,

it wasassumedthat theoxideswerethin relativeto

theworkingdepth;andbecausethewidthof theentranceslit of thediffractometerwasproportionalto

sin8, it wasnot necessaryto correctfor the oxide

thickness.It wasfurther assumedthat the oxide

crystalswererandomlyoriented. Thecalculations

yieldedcompositionsintermsofvolumepercent,and

theresultingnumbersWereroundedoffto thenearest5percent,theconfidencelimit of theestimation.

It is significantto notethat althoughthepredominantoxidationproductin allthecaseswasTiO2,the

samplesfrom stageI showedverylittle of reaction

products,stageII showedsomeTiN, andstageIII

showedA1203.It is alsointerestingto notethat the

sampleoxidizedat 871°C(whichdisplayedall three

stagesof oxidation)showedbothTiN andA1203to

bepresentin theproduct.Theresultsalsoseemto

indicatethat theproportionof TiN in thescaledecreasesandthat ofA1203increaseswith anincrease

in thetemperatureof oxidation.

Oxidethicknessmeasurementsafter oxidation

were obtainedfrom cross-sectional metallography.

However,therangeof temperaturesfor whichthese

datacouldbeobtainedwaslimitedbecausethesamplesthat oxidizedat lowertemperatures(suchas

649°C,704°C,and 760°C(shortterm)) showedno

discerniblereactionproductsin the cross-sectional

metallographs.Also,sincethe samplesexposedat

1093°Cspalled,theoxidethicknessdatacouldnot

beobtainedforthistemperatureofexposureaswell.

ThelimiteddatafromtheremainingexposureconditionsarepresentedintableIV.Usingtheoxidethicknessmeasurementsandthe total weightgainvalues

obtainedfromweightmeasurementsbeforeandafter

exposureallowedtheisolationofthecontributionsto

oxideformationandmetaldissolution.Theseresults

arealsopresentedin tableIV.

Both SEMandEDSwereusedto furthercharacterizethemorphologyandcompositionof theoxides. Again,thesestudieshad to be restrictedto

samplesoxidizedat 871°Candabovesincethesamplesoxidizedatlowertemperatures(649°Cto816°C)

showedverylittle reactionproducts,andthe cross

sectionof thesesamplescouldnotberesolvedin the

SEM.Figure5presentsthesecondaryelectronimage

of thecrosssection,andshowstheEDSmapsof Ti,

A1,Nb, andO for a sampleoxidizedat 982°Cfor

24hr. Theoxideis characterizedby layeredporosity andfingersof particlescloseto the outeredge.

Consistent with the XRD data, these studies also indicate that the scale is predominantly composed of

TiO2. The A1 and Nb maps indicate that although

the oxide is free of A1 in the bulk, it has incorporated

a substantial amount of Nb; the concentration of this

element is seen to decrease from the inner edge to the

outer edge. The finger regions show an enhancement

of the A1 signal intensity and a corresponding loss

of the Ti intensity; a slight enhancement of the O

intensity at these regions suggests these particles to

be A1203. Since the XRD data did not detect any

niobium oxides in the scale, Nb must be present in

the TiO2 as a dopant.

A similar conclusion can be drawn from the sample exposed at 1093°C for 24 hr and which had

spalled. Figure 6 shows the appearance and the

X-ray maps of A1 and Ti of the outer side of the

spalled oxide; an Nb map was not taken because

of very low intensity. It is particularly significant

to note that A1 is distributed discretely along the

7
oxidesurface,andalsoto notethat in thoseregions

whereA1is prcscnt,Ti is absent.Figure7 shows

the morphologyandX-ray mapsof Ti and Nb of

theundersideofthespalledoxidc.TheEDSdid not

detectsignificantA1ontheundersideof thespalled

oxide. The predominantfeaturesof the underside

area smallgrainsizeof theoxideanda fairly uniformdistributionof Ti andNb. TableV showsthe

resultsof themicroprobeanalysisof thespalledoxide. Principally,the analysisindicatestheoxideto

bepredominantlyTiO2;Nbis seenontheinnersurfaceandA1on the outer surface. Also,the data

showareasonablematchbetweentheanalysisofthe

innersurfaceof theoxideandtheresidualoxideon

themetal,indicatingthat spallingwasnotalongany

chemicalinterface.Calculations show the composition at the inner surface of the spalled oxide to be

90 percent TiO2 and 10 percent Nb2Os, which is the

maximum solubility for Nb205 in TiO2 (refs. 7, 18,

and 19).

The appearance of layered porosity in the sample

cxposed to 982°C (fig. 5) and the occurrence of

spallation of the oxide on samples exposed to 1093°C

indicate that the layered porosity may be the major

cause for spallation.

In order to characterize the oxidation products

formed on the alloy surface after low-temperature

exposures (649°C to 760°C (short term)), a few of

those samples were examined by AES. A typical

spectrum is shown in figure 8. A notable feature

of the spectrum is the existence of a chemical shift

of the low-energy aluminum lines, indicating the

presence of A1203. The presence of titanium lines

in the spectrum indicates the simultaneous existence

of TiO2 with A1203. However, reliable quantitative

data on the distribution of these phases could not be

obtained from the AES analysis.

It can be concluded from the characterization

studies that TiO2 is the main component of the

scale and that a continuous protective layer of A1203

does not form on these alloys at all temperatures

and exposure durations reported in this study. It

has been suggested by Zelenkov and Osokin (ref. 1)

that this alloy attempts to form a protective layer of

A1203 at early stages of oxidation, but the depletion

of A1 in the alloy leads to the eventual breakdown of

this layer and the subsequent formation of TiO2. In

support of this argument, Perkins et al. (ref. 5) have

shown that the diffusivity of A1 in the alloy is not

rapid enough to compensate for the depletion of the

clement through oxide formation.

Characteristics of the Metal

Figure 9 shows the microstructure of the alloy in

the as-received condition (no oxidation) and after ex8

posure to a few selected oxidation conditions. The asreceived alloy appears to consist of two phases of almost equal proportion. The composition of these two

phases as determined by EDS analysis is presented

in table VI; for comparison, the nominM analysis of

the alloy is also presented. There is a partitioning

of the elements, particularly Nb, between the two

phases. Prom the Ti3A1-Nb phase diagram quoted

by Strychor et al. (ref. 20), hexagonal a2 and bodycentered cubic _ are the expected phases in the alloy

at all temperatures covered in this study. On the

other hand, XRD data of the base alloy indicated

the predominant structure to be hexagonal although

some peak splitting was observed in all the peaks except in the (001) type. This result suggests that both

phases are the same type but differ in composition.

A1 in Ti produces contraction of the lattice principally along the "a" axis, whereas Nb has very little

effect in the Ti lattice; therefore, segregation of Nb

would not be readily apparent from the diffraction

data.

It is evident from figure 9 that the heat treatment associated with the oxidation process has significantly changed the microstructure of the alloy. In

the alloy bulk, grain refinement and grain growth

can be noted. In the vicinity of the oxide layer, as

clearly seen for the sample oxidized at 982°C, there

is a formation of a single-phase case and a skirted region very close to the metal/oxide interface. An EDS

analysis of the skirted region showed an enhancement

in the A1 and Nb intensities and a corresponding decrease in the Ti intensity, with the composition of

this layer approaching that of TiA1. On the other

hand, the presence of only a single phase case indicates the stabilizing effect of oxygen on one of the

two phases reported earlier.

Figure 10 summarizes a portion of the the XRD

data corresponding to a 2t? range of approximately

37 ° to 40°; this range was selected to observe the effect of oxidation exposure on the (002) reflection of

the alloy before and after oxidation exposure. Compositional broadening of the (002) line can be observed for the sample oxidized at 649°C (fig. 10(b)).

For the samples oxidized at 760°C, only line shifts

could be noted (fig. 10(c)). The existence of a line

shift rather than line broadening for this exposure

condition suggests that the metal is saturated with

oxygen within the X-ray penetration depth of the alloy (approximately 9 #m). For the sample exposed

to a high temperature (982°C), the oxide thickness

was large enough that only weak intensities of the

base-alloy reflections could be obtained (fig. 10(d)).

Also, the (002) peak was overlapped by a series of

peaks corresponding to a TiA1 phase. From the compositionally broadened peak for the sample exposed
to 649°C(fig. 10(b))andusingthe analyticaltechniquedevelopedbyWiedemannandUnnam(ref.21),

a very precise composition-depth profile was constructed as shown in figure 11. From this analysis,

the maximum solubility of oxygen for this exposure

was determined as 12.5 atomic percent. These results indicate that oxygen dissolution into the alloy

is a significant process during oxidation exposures.

Oxygen in titanium alloys has been shown to lower

ductility, and this problem is likely to exist with the

Ti-14A1-21Nb alloy also.

Suggested Mechanisms

The existence of at least three slopes in the oxidation kinetics of Ti-14A1-21Nb alloy with corresponding changes in activation energies (figs. 3 and 4) reflects the complexity of the oxidation behavior of the

alloy. Before attempting to identify the causes for

this behavior, a comparison of the measured rate constant with the theoretical predictions can be made

to gain additional information. Figure 12 compares

the predicted and measured rate constants kp for

Ti-14A1-21Nb alloy and commercial purity titanium

(Ti-A55); the experimental data have been reproduced from figure 4 and the predicted values have

been calculated according to equation (23). Since

TiO2 was the predominant oxide in the scale formed

on Ti-14A1-21Nb and the oxide had a relatively low

A1 content but a substantial Nb content, kr was calculated according to equation (17b) for the alloy and

according to equation (17a) for commercial purity

titanium. The parameters required for the computation of kr and k d and the sources for the data

(refs. 12, 15, 18, 22-26) are listed in table VII. In

spite of the multiple sources for the data for the theoretical estimate (table VII), the agreement between

the predicted and measured values is reasonable, implying the validity of using the defect-structure-based

model for predicting the oxidation of Ti-14A1-21Nb

alloy.

The deviation of kp from the predicted values associated with variations in the activation energies

(figs. 4 and 12) indicates the dominance of different

operating mechanisms as a function of temperature.

Indeed, this behavior was expected because examination of equation (23) shows that the rate constant

kp is in fact comprised of two temperature-dependent

terms, kr and kd, and the dominance of one term or

the other on the oxidation kinetics at a given temperature can give rise to the observed effect. However, the existence of at least three slopes during the

isothermal oxidation of the alloy (fig. 3) suggests that

the duration of exposure also causes changes in the

operating mechanism. The existence of such transitions has been noted by several past investigators

(refs. 8, 9, 12, and 27). Kofstad et al. (ref. 8) and

Wallwork and Jenkins (ref. 9) have suggested that

the first-stage kinetics during the oxidation of titanium is predominantly controlled by the process of

oxygen dissolution into the metal; the transition at

the end of this stage is attributed to the onset of the

oxide-formation process. The variations in the kinetics during the second and subsequent stages of oxidation have been attributed to cracking and healing

processes in the oxide. An understanding of whether

the observed transitions in the Ti-14A1-21Nb alloy

are associated with changes in the metal-dissolution

kinetics or oxide-formation kinetics will help in characterizing the oxidation behavior of this alloy.

In order to determine if the dissolution process

had the major contribution to the observed kinetics, the oxygen diffusivity in the alloy was backcalculated from the oxidation data through equation (22) using the data from tables III and IV and

taking Cst = 12.5 atomic percent (figs. 10 and 11).

Using the XRD data of the alloy to compute the density of the alloy-O solid solution allowed the conversion factor in equation (22) to be computed as fo =--

0.0172 g/cm2/atomic percent-cm. Figure 13 shows

the Arrhenius plot of the data; the 649°C data point

was determined from the composition-depth profile

in figure 11. Also shown in figure 13 is the plot of

the oxygen diffusivity data from Wiedemann et al.

(ref. 25). It is clear that the diffusivity values computed from oxidation kinetics data are up to an order

of magnitude lower than that measured by Wiedemann et al. (ref. 25). On closer scrutiny of the data,

it appears that the diffusivity data in figure 13 may

be composed of two slopes as indicated by the dashed

line.

The similarity in the trends between the firststage kinetics (fig. 4) and the oxygen diffusivity in

the alloy (the dashed curve in fig. 13), and the minimal oxide formation as evidenced by metallographic

data, suggests that this stage of the oxidation process

may be dominated by the oxygen dissolution kinetics. Taking cognizance of the variations in the alloy

microstructure with exposure conditions (fig. 9), it

seems probable that the variation in the activation

energy during the first stage is due to the morphological effects. Thus, the samples exposed from 649°C to

approximately 760°C may have alloy oxygen transport through a two-phase field, each having different

diffusivities and solubilities. On the other hand, the

dominance of one of these phases at higher temperatures (approximatly 760°C to 982°C) may be the

cause for the drop in the activation energy. Eventually, the alloy surface forms a single phase case and

this event possibly corresponds to a critical oxygen

= 9
contentrthc first transitionmaybea markof this

event.

Thekineticsof thesecondandsubsequent sages

of oxidationof this alloyarepossiblycontrolledby

anoxideformationratecharacterizedby therationalrateconstantkr. Although XRD analysis shows

small amounts of A1203 and TiN to be present in the

scale, TiO2 is the predominant oxide that is formed

on the alloy surface for all the exposure conditions

reported in this study. Hence, the kinetics during

these stages are dominated by the transport properties of the TiO2 layer. However, the TiO2 compound

is doped with Nb ions with the result that the oxygen

transport through the oxide is limited by the reduction in concentration of oxygen ion vacancies caused

by doping. At higher thicknesses of the oxide, layered

porosity begins to occur because of alternating cracking and healing processes the second transition may

have been caused by this cracking.

Further oxidation studies on this alloy after suitable heat treatment to stabilize the microstructure

are necessary to discern the microstructural effects.

Conclusions

The static Oxidation kinetics of Ti-14A1-21Nb ingot alloy were studied in air over a temperature range

from 649°C to 1093°C in a thermogravimetric apparatus. From this study the following conclusions are

presented:

1. Ti-14A1-21Nb alloy has an oxidation resistance

that is superior to that of commercial purity

titanium or to other titanium alloys. However,

the oxidation kinetics are complex; two distinct

transitions separating three stages of differing

activation energies are marked by specific weight

gain values.

2. The oxide phase is predominantly composed of

TiO2, but small amounts of A120 3 and TiN are

also present. The TiO2 has a substantial amount

of Nb ions dissolved in it. Calculations show

that Nb ions in Ti02 reduce the concentration

of oxygen ion vacancies, thereby reducing the

diffusivity of oxygen through it. A reasonable

agreement between the measured and calculated

rate constants for oxidation supports this view.

3. High oxidation temperatures and long-term exposures at these temperatures result in layered

porosity in the oxide---the onset of porosity is displayed by a change in the oxidation kinetics. It

appears th_at healing of the outer layers results in

continuously decreasing rate constants for samples exposed to very high temperatures (>__982°C).

Also, the layered porosity may be responsible for

spalling of the oxide.

4. The alloy has a substantial solubility for oxygen

(_12.5 atomic percent). In the absence of a prior

high-temperature heat treatment after processing, oxidation exposures lead to changes in the

microstructure of the alloy, Diffusivity of oxygen

into the metal seems to be dependent on the alloy

microstructure.

5. The alloy region close to the interface is enriched

in A1 and Nb; the role of this layer on oxidation

kinctics needs further investigation.

NASA Langley Research Center

Hampton, VA 23665-5225

May 21, 1990

References

1. Zelenkov, I. A.; and Osokin, E. N.: Oxidation Resistance

of the Compound Ti 3Al and Its Alloys at Temperatures of

700 and 800 C. Poroshkovaia Metallurgiia, vol. 14, no. 10,

Oct. 1975, pp. 72 75.

2. Zelenkov, I. A.; and Martynchuk, E. N.: Heat Resistance

of Alloys of the Compound TiAl With Niobium at 800

and 1000 C. Metallofizika, no. 42, 1972, pp. 6366.

3. Chaze, A. M.; and Coddet, C.: Influence of Alloying

Elements on the Dissolution of Oxygen in the Metallic

Phase During the Oxidation of Titanium Alloys. J. Mat.

Sci., vol. 22, Apr. 1987, pp. 1206 1214.

4. Streiff, R.; and Poize, S.: Oxidation of Aluminide Coatings on Unalloyed Titanium. High Temperature Corrosion, Proceedings of Int. Conf., Nat. Assoc. Corrosion Engineers (R. A. Rapp, ed.), 1983, pp. 591 597.

5. Perkins, R. A.; Chiang, K. T.; and Meier, G. H.: Formation of Alumina on Ti-A1 Alloys. Scr. MetaU., vol. 21,

no. 11, Nov. 1987, pp. 1505 1510.

6. Hauffe, Karl: Oxidation of Metals. Plenum Press, 1965,

pp. 2i7 225.

7. Chen, Y. S.; and Rosa, C. J.: High-Temperature Oxidation of Ti-4.32 wt.% Nb Alloy. Oxid. Met., vol. 14, no. 2,

Apr. 1980, pp. 147 165.

8. Kofstad, P.; Anderson, P. B.; and Krudtaa, O. J.: Oxidation of Titanium in the Temperature Range 800-

1200°C. J. Less-Common Met., vol. 3, no. 2, Apr. 1961,

pp. 89 97.

9. Wallwork, G. R.; and Jenkins, A. E.: Oxidation of Titanium, Zirconium, and Hafnium. J. Electrochem. Soc.,

vol. 106, no. 1, Jan. 1959, pp. 1(_14.

Wallwork, G. R.; Smeltzer, W. W.; and Rosa, C. J.:

The Parabolic Oxidation Kinetics of Alpha-Zirconium

at 850°C. Acta Metall., vol. 12, no. 4, Apr. 1964,

pp. 409 415.

Pemsler, J. P.: Studies on the Oxygen Gradients in

Corroding Zirconium Alloys. Jour. Nuclear Materials,

vol. 7, no. 1, 1962, pp. 16-25.

Unnam, J.; Shenoy, R. N.; and Clark, R. K.: Oxidation

of Commercial Purity Titanium. Oxid. Met., vol. 26,

no. 3/4, Oct. 1986, pp. 231 252.

10.

11.

12.

10
13. Wagner, Carl: Theory of the Tarnishing Process. Z.

Physik. Chem., vol. B, no. 21, 1933, pp. 25 41.

14. Wagner, Carl: Diffusion and High Temperature Oxidation

of Metals. Atom Movements, American Soc. for Metals,

c.1951, pp. 153 173.

15. Kofstad, Per: Nonstoichiometry, Diffusion, and Electrical

Conductivity in Binary Metal Oxides. John Wiley & Sons,

Inc., c.1972, pp. 15 46.

16. Kofstad, P.: High- Temperature Oxidation of Metals. John

Wiley & Sons, Inc., 1966, p. 169.

17. KrSger, F. A.; and Vink, H. J.: Relations Between

the Concentrations of Imperfections in Crystalline Solids.

Solid State Physics--Advances in Research and Applications, Volume 3, Frederick Seitz and David Turnbull, eds.,

Academic Press, Inc., "1956, pp. 307 435.

18. Goldschmidt, H. J.: An X-Ray Investigation of Systems

Between Niobium Pentoxide and Certain Additional Oxides. MetalIurgia, vol. 62, no. 373, Nov. 1960, pp. 211 218.

19. Roth, R. S.; and Coughanour, L. W.: Phase Equilibrium

Relations in the Systems Titania-Niobia and ZirconiaNiobia. J. Res. National Bureau Stand., vol. 55, no. 4,

Oct. 1955, pp. 209 213.

20. Strychor, R.; Williams, J. C.; and Sofia, W. A.: Phase

Transformations and Modulated Microstructures in

Ti-AI-Nb Alloys. Metall. Trans., vol. 19A, no. 2,

Feb. 1988, pp. 225-234.

21. Wiedemann, K. E.i and Unnam, J.: A Method of

Rapidly Obtaining Concentration-Depth Profiles From

X-Ray Diffraction. J. Appl. Phys., vol. 58, no. 3, Aug. 1,

1985, pp. 1095-1101.

22. Shewmon, Paul G.: Diffusion in Solids, McGraw-Hill

Book Co., Inc., 1963, p. 102.

23. Liu, Z.; and Welsch, G.: Communication: Literature Survey on Diffusivities of Oxygen, Aluminum, and Vanadium

in Alpha Titanium, Beta Titanium, and in Rutile. MetaU.

Trans, vol. 19A, no. 4, pp. 1121-1125.

24. Kubaschewski, O.; and Alcock, C. B.: Metallurgical Thermochemistry, Volume 2_, Fifth ed. Pergamon Press Ltd.

(England), c.1979, p. 383.

25. Wiedemann, K. E.; Clark, R. K.; and Sankaran S. N.:

Emittance, Catalysis, and Dynamic Oxidation of Ti-14Al21Nb. NASA TP-2955, 1989.

26. Wiedemann, K. E.; and Unnam, J.: An X-Ray Diffraction

Study of Titanium Oxidation. NASA CR-176263, 1984.

27. Shenoy, R. N.; Unnam, J.; and Clark, R. K.: Oxidation

and Embrittlement of Ti-6Al-2Sn-4Zr-2Mo Alloy. Oxid.

Met., vol. 26, no. 1/2, Aug. 1986, pp. 105-124.

11
Table I. Chemical Analysis of Ti-14A1-21Nb Alloy

Element

A1

Nb

Mo

Fe

V

C

O

N

H

Zr

Ti

Composition,

Nominal

14

21

Balance

percent weight

Actual

14.1

21.30

<.01

.09

<.01

.02

.13

.015

.0065

<.01

Balance

Table II. Thermogravimetric Analysis Results

Sample

TA-5

TA-7

TA-9

TA-1

TA-2

TA-3

TA-4

TA-6

TA-8

TA-18

TA-21

TA-10

TA-11

TA-12

TA-16

TA-17

TA-14

TA-15

TA-19

Exposure

conditions

Temperature,°C

649

649

649

704

704

704

760

760

760

760

760

816

816

816

871

871

982

1093

1093

Time, hr

102

112

112

24

24

25

25

24

24

116

117

24

24

24

26

24

24

24

2O

Total weight

gain, mg/cm 2

0.087

.072

.062

.084

.092

.080

.250

.210

.210

.370

.390

.410

.380

.380

.740

.730

2.820

Spalled

Spalled

Weight gain rate, a

#g/cm2-vrh--r, for stage--

2.53

3.24

3.62

11.63

11.58

9.64

45.6

37.5

41.2

33.0

35.0

89.6

84.5

84.8

135.5

126.6

II III

40.1

29.0

30.8

30.7

61.1

62.6

59.9

103.5 138.6

114.0 142.9

483.4

1367.2

1442.2

aThe columns refer to the corresponding stages in figure 3.

12
Table III. Oxide Composition From X-Ray Diffraction Data

Estimated composition,

Exposure conditions

Temperature, °C

760

76O

816

871

982

Time, hr

25

117

24

26

24

Stage

I

II

II

III

III

TiO2

70

75

75

80

75

percent volume

A1203 I

10

25

TiN

30

25

25

10

Table IV. Oxide Thickness and Partitioning of Oxygen

Sample

TA-21

TA-10

TA-16

TA-14

Exposure conditions

Temperature, °C

760

816

871

982

Time, hr

117

24

26

24

Oxide

thickness, a

#m

V_Teight gain,

mg/cm 2

Oxide

1.65

1.45

2.95

13.65

0.28

.25

.50

2.32

Metal

0.11

.16

.24

.49

aAverage of five measurements.

Table V. EDS Analysis of Spalled Oxide

Composition, percent weight a

Element Outer surface Inner surface Residual oxide

Ti 50 (26) 40 (24) 40 (24)

Nb 0 (0) 20 (6) 22 (7)

0 44 (69) 40 (70) 37 (68)

A1 6 (5) 0 (0) 1 (1)

aAtomic percent is given in parentheses.

13
TableVI. DistributionofElementsAmongPhasesin Ti-14A1-21NbAlloy

Composition,percentweighta

Element Nominal Dark phase Light phase

Ti 65 (66) 64 (65) 69 (66)

Nb 21 (11) 23 (12) 16 (8)

Al 14 (25) 13 (23) 15 (26)

aAtomic percent is given in parentheses.

Table VII. Data for Theoretical Estimation of Rate Constant

Parameter Value Reference

Ccq 0.21 eq/cm 3

f 0.78 22

D_ 2.0 x 10 -3 exp(-60000/RT) 23

1.7 × 10 -2 exp(-66000/RT) 23

Po_ 1.1 x 1010 exp(-146000/RT) 24

(for undoped TiO2)

Cst

3.7 x 109 exp(-156 000/RT)

(for TiO2 doped with Nb)

24

K 4.0 x 102 exp(-i05 000/RT) 15

0.2 ,8

26

08

25 atomic percent (pure titanium)

12.5 atomic percent (Ti-14A1-21Nb)

2.7 exp(-51 400/RT)

(pure titanium)

0.0234 exp(-45 090/RT)

(Ti-14A1-21Nb)

0.0178 (pure titanium)

0.0172 (Ti-14A1-21Nb)

f0

25

26

25

12

Remarks

__P02calculated for Ti305/TiO2

equilibrium since Ti305 is

next lowest stable phase

p f!02 calculated for Ti3Os/TiO2

equilibrium since Ti305 is

unstable in presence of Nb

Assumes maximum solubility

Calculated from X-ray data

14
E

O_

8

¢)

O

::Z

1⁄2

gas

Oxide Metal

Csl _ CO

I

1

i

i

I i ..

i i

Distance

Figure 1. Schematic representation of oxidation process where oxygen distribution is presented in terms of

concentration and chemical potential.

.4

.3--

Temperature

¢.-

I

I

0 I

, Weight gain

_2/ i I I I I

0 1 2 3 4 5

Time,

100o

-- 800

-- 600

-- 400

-- 200

0

6

Figure 2. Typical output of TGA data for Ti-14A1-21Nb alloy.

?

0_

E

15
Figure 3.

E

E

ff

E_

E0J

1.0 m

(a) 649°C to 871°C.

.8 -- Temperature Stage

.6 -- / _71°C

II[

.... -'_ o- " II

.4 -- ,f -,, 816 C

,2 ..................... I

(b) 871°C to 1093°C.

10

8

6

4

2

760°C

649°C

1 I

-- f Temperature

J'1093oc

_ / Stage

__ / II1

/ / 982°C

2 4 6 8 10

Time,

I J

12 14

TGA results of Ti-14AI-21Nb alloy presented in parabolic coordinates. The two plots differ in scale.

i

OJ

E

r-

¢-

0

f_

_)

n.-

10 -1 __

10-2 ___

10 -3

10-4 ___

10 -5

10 -6

.6

O Ti-A55 (ref. 12)

[] Ti-6AI-2Sn-4Zr-2Mo (ref. 27)

Ti-4.32Nb (ref. 7)

• Ti-14AI-21Nb; stage !

• Ti-14AI-21Nb;stage II

• Ti-14AI-21Nb; stage llI

I I 1 I I

.7 .8 .9 1.0 1.1

Temperature, K -1

.I

1.2 x 10 -3

Figure 4. Arrhenius plot of oxidation rate constants for Ti-14AI-21Nb and other titanium alloys.

16
ORIGINAE PAGE

3LACK AND WHITE PHOTOGRAPH

o..

E

E

t-

°_

E

<

0

o.

£

E

-t

o

25

.o

Z

E

t--

)<

0

r_

,,.o

<I

0

r_

%

b.D

0

g

*d

m

e_

0

ORIGINAL PAGE IS

OF POOR QUALITY ].'7
ORIGENAE PAGE

BLACK AND WHITE PHOTOGRAPN

=

i _ 22

Titanium map

Micrograph

i I

20pm

Aluminum map

Figure 6. Secondary electron image and elemental maps of outer surface of spalled oxide of Ti-14Al-21Nb

sample exposed at 1093°C for 24 hours.

Micrograph

Titanium map ,_ .

10 p.m

Niobium map

Figure 7. Secondary electron image and elemental maps of underside of spalled oxide of Ti-14A1-21Nb sample

exposed at 1093°C for 24 hours.

18
O

AlinAI203

Ar

AI

r--

Relative

intensity ....

Ti

Electronic energy, eV

Figure 8. Typical AES of outer surface of Ti-14AI-21Nb sample during stage I kinetics.

ORIGINAL PAGE

BLACK AND WHITE PHOTOGRAPH

No oxidation , _ 816°C

501_m

704°C 982°C

Figure 9. Cross-sectional micrographs of Ti-14AI-21Nb alloy after various oxidation exposures.

ORIGINAL PAGE IS

OF POOR QUALITY

19
(a) Unexposed alloy.

(b) 649°C for

(c) 760°C for 116 hr.

• ,, .-

h

(d) 982°C for 24 hr. / _ f_ ,,...

I l l J '*"

36 37 38 39

J

4O

28 range, deg

Figure 10. XRD data in 28 range from approximately 37 ° to 40 ° for Ti-14Al-21Nb alloy after various oxidation

exposures. Compositional broadening in part (b) and peak shift in part (c) are due to oxygen uptake. The

new peaks seen in part (d) are from TiA1 phase,

E

Figure 11.

112 hr,

15 -

¢--

._o

E

o 10 .

C

_- 5 O

O

t"-

x

O

0 2 4 6 8 10

Depth, g m

Oxygen depth profile computed from XRD data of Ti-14A1-21Nb alloy after exposure at 649°C for

2O
J

(M

E

10"2

10"3

Temperature, °C

1200 1000 800 600

O')

- 10-4

C

(D

c

O

-5

_ 10 -

rr

10-6

.6

I I I I

__Ti-A55

Ti-14Al-21 Nb "_o ?

!

I l I I 1

.7 .8 .9 1.0 1.1

Temperature, K -1

.2x 10 -3

Figure 12. Comparison of measured and predicted oxidation rate constants for Ti-14AI-21Nb alloy and Ti-A55.

Symbols are experimental points from figure 4 and solid lines are from theoretical estimation.

(.}

(D

09

E

10 -9 =

10-10 _

10 11 _

m

10 -12 _

m

m

10-13 ___

m

10-14

.6

1200

Temperature, °C

1000 800 600

I I I

,_ Wiedemann et al. (ref 25)

Present study i_

XN

I I I I I

.7 .8 .9 1.0 1.1 1.2x 10 -3

Temperature, °C

Figure 13. Comparison of oxygen diffusivity in Ti-14A1-21Nb alloy measured from the present study and the

XRD data. Dashed line represents an alternate fit of the data. Diffusion coefficient (D) of the present study

was 0.013 em2/sec exp (4810° c/_m°l-K).

21
Report DocumentationPage Natlonai Ae_or,aul_c s and

Space Adm_nqstraT_on

1. Report No. / 2. Government Accession No. 3. Recipient's Catalog No.

NASA TP-3012 1

4. Title and Subtitle

Oxidation Characteristics of Ti-14A1-21Nb Alloy

7. Author(s)

Sankara N. Sankaran, Ronald K. Clark, Jalaiah Unnam ,

and Karl E. Wiedemann

9. Performing Organization Name and Address

NASA Langley Research Center

Hampton, VA 23665-5225

12. Sponsoring Agency Name and Address 13.

National Aeronautics and Space Administration

Washington, DC 20546-0001 14.

5. Report Date

July 1990

6. Performing Organization Code

8. Performing Organization Report No.

L-16658

10. Work Unit No.

505-43-71-01

11. Contract or Grant No.

Type of Report and Period Covered

Technical Paper

Sponsoring Agency Code

15. Supplementary Notes

Sankara N. Sankaran, Jalaiah Unnam, and Karl E. Wiedemann: Analytical Services & Materials,

Inc., Hampton, Virginia.

Ronald K. Clark: Langley Research Center, Hampton, Virginia.

16. Abstract

The static oxidation kinetics of Ti-14A1-21Nb (percent weight) ingot alloy were studied in air over

the temperature range from 649°C to 1093°C in a thermogravimetric apparatus. The oxidation

products were characterized by X-ray diffraction, electron microprobe analysis, energy dispersive

X-ray analysis, and Auger electron spectroscopy. Cross sections of the oxidized samples were

also examined using light and scanning electron microscopy. The oxidation rate was substantially

lower than the conventional alloys of titanium, but the kinetics displayed a complex behavior

involving two or more oxidation rates depending on the temperature and duration of exposure.

The primary oxide formed was TiO2, but this oxide was doped with Nb. Small amounts of A1203

and TiN were also present in the scale. The diffusion of oxygen into the alloy was observed, and

the diffusivity seemed to be dependent on the microstructure of the metal. A model has been

presented to explain the oxidation behavior of the alloy in terms of the reduction in the oxygen

diffusivity in the oxide caused by the modification of the defect structure of TiO2 by Nb ions.

17. Key Words (Suggested by Authors(s))

Oxidation

Oxidation kinetics

Titanium-aluminide alloys

19. Security Classif. (of this report)

Unclassified

NASA FORM 1626 OCT 86

18. Distribution Statement

Unclassified--Unlimited

Subject Category 26

20. UnclassifiedSecurity Classif. (of this page) 21. No. 22 of Pages [ 22. Price A03

NAS A-Langley, 1990

For sale by the National Technical Information Service, Springfield, Virginia 22161-2171