Room temperature fracture processes of a near-a titanium alloyfollowing elevated temperature exposure

A. L. Pilchak • W. J. Porter • R. John

Received: 28 February 2012 / Accepted: 13 June 2012 / Published online: 6 July 2012

� Springer Science+Business Media, LLC (outside the USA) 2012

Abstract Near-a titanium alloys are used at higher

temperatures than any other class of titanium alloys. As a

consequence of thermal exposure, these components may

develop locally elevated oxygen concentrations at the

exposed surface which can negatively impact ductility and

resistance to fatigue crack initiation. In this work, mono-

tonic and fatigue fracture mechanisms of Ti–6Al–2Sn–

4Zr–2Mo–0.1Si samples exposed to laboratory air at

650 �C for 420 h were identified by means of a combina-

tion of quantitative tilt fractography, metallographic sec-

tioning, and electron backscatter diffraction. These

mechanisms were compared and contrasted with those

operative during similar tests performed on material is the

as-received condition with uniform oxygen content. While

faceted fracture was not observed during quasi-static

loading of virgin material, locally elevated concentrations

of oxygen near the surfaces of exposed samples were

shown to change the fracture mode from ductile, microvoid

coalescence to brittle facet formation and grain boundary

separation at stresses below the macroscopic yield point.

Similar features and an increased propensity for facet for-

mation were observed during cyclic loading of exposed

samples. The effects of this time-dependent degradation on

monotonic and cyclic properties were discussed in the

context of the effect of oxygen on crack initiation and

propagation mechanisms.

Introduction

Near-a titanium alloys, like Ti–6Al–2Sn–4Zr–2Mo

(Ti-6242) and Ti-5.68Al-4.05Sn-3.65Zr-0.68Nb-0.52Mo-

0.33Si (IMI-834), are known for maintaining strength at

elevated temperatures. Because of its high solid solubility in

a-titanium, inward diffusion of interstitial oxygen occurs

simultaneously with oxide growth during elevated temper-

ature excursions. The oxide scale is typically made up of

TiO2 although Ti3AlN has been found in IMI-834 upon

exposure to temperatures exceeding 750 �C [1]. Although

protective at room temperature, the TiO2 scale formed by

high temperature exposure grows to be thicker and may spall

during repeated thermal excursions, which limits the maxi-

mum service life.

In solid solution, oxygen can strengthen the a-phase under

quasi-static deformation rates [2]; however, it can be detri-

mental to fatigue and fracture properties through its effect on

slip behavior. It has been demonstrated that oxygen strongly

affects slip character causing a shift from wavy to planar

glide with increasing oxygen concentration [2, 3]. Physi-

cally, planar slip results in a decrease in tensile ductility and

toughness during monotonic loading [4, 5]. With regard to

fatigue, Mahoney and Paton [6] have shown that oxygen

does not strongly affect fatigue crack growth behavior over a

wide range of DK when its concentration is relatively small

(0.06–0.18 wt %). In contrast, Bache et al. [7] have shown

that high concentrations of O (0.51 wt %) in mill-annealed

Ti–6Al–4V resulted in a 5- to 10-fold increase in the long

crack growth rate in air at DK \ 12 MPa m1/2 for R = 0.01

and a smaller, but measureable, increase at larger DK. At a

load ratio of 0.5, material with moderate (0.165 wt %O) and

high (0.4 wt %O) oxygen content both exhibited increased

crack growth rates compared to low oxygen material

(0.09 wt %O) over the entire range of DK investigated

A. L. Pilchak (&) � W. J. Porter � R. John

Air Force Research Laboratory, Materials and Manufacturing

Directorate, Wright Patterson Air Force Base, OH 45433, USA

e-mail: adam.pilchak@wpafb.af.mil

W. J. Porter

University of Dayton Research Institute, Dayton, OH 45469,

USA

123

J Mater Sci (2012) 47:7235–7253

DOI 10.1007/s10853-012-6673-y

suggesting that roughness-induced crack closure effects may

overshadow O effects in the lower-R regime.

In practice, facet formation is commonly observed near

crack initiation sites during cyclic loading [8–10], but not

during quasi-static loading, even at very low temperatures

in the presence of notches which should exacerbate the

tendency for brittle faceting [11]. Ward-Close and Beevers

[12] have shown that fatigue crack growth can be an order

of magnitude faster for faceted growth versus striation

growth for a given value of DK. Accordingly, it has been

proposed [13] that the mechanism by which oxygen exerts

its influence on fatigue crack growth is through increased

facet formation resulting in more rapid crack propagation

rates. Indeed, Bache et al. [7] have observed an increased

number of facets on the fracture surface of high-oxygen

samples. Sarrazin-Baudoux et al. [13] have suggested that

such observations are due to the increased propensity for

planar slip and subsequent fracture along the slip bands, but

this was not demonstrated explicitly. Said another way, it

seems that higher oxygen concentrations permit faceted

fracture over a wider range of grain orientations and per-

haps over a wider range of DK.

In the present paper, observations regarding the effects

of oxygen on crack propagation by facet formation during

monotonic and cyclic loading are presented. It is shown

that high levels of oxygen promote extensive facet for-

mation under both tensile and cyclic loading. The mecha-

nism of crack advance is shown to be a combination of

transgranular and intergranular growth. It is noteworthy

that this work is distinct from previous investigations in

that we are investigating the effects of simulated service-

induced oxygen ingress resulting in a gradient in oxygen

concentration as opposed to measuring properties of

materials produced with varying bulk oxygen contents.

Materials and experimental procedures

Materials

Ti-6242 ? 0.1Si (Ti-6242S) sheet, approximately 2.29-mm

thick, was acquired from Timet (Henderson, NV). The

actual composition of the alloy was determined to be

Ti–5.93Al–2.01Sn–4.05Zr–1.88Mo–0.12Si in weight per-

cent, while the bulk oxygen, carbon, iron, and nitrogen

contents were 0.107, 0.023, 0.05, and 0.001, respectively.

Metallic element composition (Al, Si, Zr, Mo, and Sn) was

determined by atomic absorption and inductively coupled

plasma mass spectrometry, while oxygen and nitrogen

content were determined by inert gas fusion; carbon con-

tent was measured by infrared absorption following com-

bustion. Rectangular blanks 24 mm 9 160 mm were

sectioned from the plate, thoroughly cleaned with acetone,

dried, and subjected to elevated temperature exposure at

650 �C for 420 h in laboratory air.

Mechanical testing

Flat dogbone specimens for tensile and fatigue tests were

prepared from the as-received sheet and the thermally

exposed blanks by wire EDM. Tensile properties were

assessed by means of constant displacement rate tests

resulting in an approximate strain rate of 10-4 s-1. Rep-

licate tests were performed on material in the as-received

condition and in the exposed condition. In addition, two

tests were interrupted at a total strain of 0.01 to investigate

the early stages of deformation. Finally, two tests were

performed on thermally exposed material that had the

brittle oxide layer, but not the oxygen-enriched layer,

removed by mechanical polishing in the longitudinal

direction. A total of ten fatigue tests, five in the as-received

condition and five on exposed samples, were performed at

a maximum stress of 800 MPa with a load ratio of 0.05 at a

frequency of 20 Hz. This stress level was chosen to pro-

duce failures in the low cycle fatigue range. The as-frac-

tured samples were cleaned with acetone before examining

in the SEM, but otherwise were given no additional surface

preparation.

Electron microscopy

Metallographic specimens were extracted from the exposed

blanks by wire electrical discharge machining. Sectioning

of tested samples with a slow speed saw was performed as

needed to understand the deformation mechanisms asso-

ciated with loading. The specimens were metallographi-

cally prepared by grinding with progressively finer SiC

paper through 600 grit followed by polishing with 9, 6, 3,

and 1-lm diamond paste with glycol extender. Final pol-

ishing was performed by polishing overnight in a vibratory

polisher containing 0.05-lm non-crystallizing colloidal

silica. Metallographic and fractographic inspections were

performed in a field emission source scanning electron

microscope (SEM) operated at 20 kV with a probe current

of approximately 2 nA. Ultrahigh resolution (UHR) images

of the fracture surface were taken using a scintillator

detector mounted inside the pole piece of the SEM with a

bias of 20 V. The microstructure near the surface and in the

bulk of the samples was examined in the SEM. Electron

backscatter diffraction (EBSD) was used to assess the

macrotexture and degree of microtexture present in the

material. In this context, the term microtexture implies

large regions of near-constant crystallographic orientations

that appear fully recrystallized via standard optical and

SEM observation. For this purpose, the SEM was operated

7236 J Mater Sci (2012) 47:7235–7253

123

at 20 kV with a probe current of *50 nA. An area

equivalent to the cross section of the tensile samples

was scanned at a 5-lm step size by a combination of beam

control scans and automated stage movements [14] to

determine macrotexture. The individual tiles of EBSD data

were combined to produce a seamless image by means of

the open source software AnyStitch [15]. Additional scans

were performed at higher resolution (0.15 to 1-lm step

size) to obtain detailed information about the micro-

structure.

The spatial and crystallographic orientations of various

features on the fracture surface were determined by a

combined EBSD/quantitative tilt fractography technique.

This method has been covered extensively in the literature

[8, 9, 16–18], but is briefly outlined here for thoroughness.

First, the spatial orientation of a given feature, a facet for

instance, on the fracture surface is determined by acquiring

images of it at the same magnification at two different tilt

angles. Next, an origin and three common features on the

facet are identified. Using the equations developed by

Themelis et al. [16], vectors connecting these features

(which lie in the plane of the facet) can be determined.

Finally, the cross product of two vectors in the facet plane

yields the facet normal vector. The crystallographic ori-

entation of the fractured grain is obtained directly by EBSD

without any additional preparation of the facet surface.

This is generally possible although extensive effort is often

needed to ensure that other parts of the fracture surface do

not shadow the EBSD detector. This may include using

longer working distances or tilt angles other than 70�.

These two pieces of information are combined to identify

the crystallographic plane of fracture by plotting an inverse

pole figure (IPF) of the EBSD-deduced orientation centered

on the facet normal vector.

The depth of oxygen ingress during the thermal excur-

sion was determined indirectly with optical measurements

made on oxalic-etched samples and directly via wavelength

dispersive spectroscopy [19, 20]. To this end, the intensity

of the O Ka peak as a function of depth from the exposed

surface was measured using an electron microprobe ana-

lyzer (EPMA) operated at 15 kV with a probe current of

approximately 100 nA. Care was taken to prepare the

sample with minimal edge rounding which was verified by

scanning white light interferometry. The probe was focused

to a spot and 100 line scans were collected orthogonal to

the exposed surface at a step size of 0.3 lm. The depth-

dependent O Ka intensity of these 100 lines was averaged

to produce the final concentration profile. The energy res-

olution of the microprobe (*5 eV) was sufficient to

resolve the O Ka peak from the Ti La peak, but due to

uncertainties with quantifying light elements via EPMA, no

quantitative measurements were made and instead the

values were simply normalized to the bulk O content.

Results

The primary results of this research included correlations

between microstructure, crystallography, and fracture

topography as elucidated through the use of quantitative

tilt fractography, EBSD, and high resolution SEM. The

microstructure and texture of the as-received materials and

the changes that occurred during exposure are described

briefly followed by more rigorous analysis and discussion of

the cracking behavior.

Microstructure and texture

The microstructure of the as-received sheet was typical for

an a ? b-processed material consisting of a high volume

fraction of globular a grains, approximately 7–10 lm in

diameter with smaller regions of fine secondary a with very

little retained b (Fig. 1). The volume fraction of the glob-

ular phase was approximately 90 pct., and the remainder

was accounted for by the a ? b transformation product and

the retained b phase. The as-received material exhibited

basal and transverse texture components (Fig. 2) of similar

intensity, approximately 2–3 times random in the 0001 pole

figure (*5 times in the orientation distribution function

(ODF)). A stronger intensity (pole figures: *5.7 times

random; ODF: *14 times random) that aligned [0001]

with the rolling direction (RD) was also observed. The

material was also moderately microtextured, consisting of

narrow bands of grains with similar orientations that

extended over several millimeters in the as-rolled sheet

(Fig. 3). The similar orientation of the grains within the

bands gave rise to a large population of low angle

boundaries as indicated by the misorientation distribution

in Fig. 4. The microtextured regions were plate shaped,

approximately 25-lm thick, and were parallel to the rolling

plane.

Oxygen ingress and surface oxidation during exposure

The overall depth of oxygen ingress during exposure is a

function of both time and temperature. Measurements from

the oxalic-etched samples of the present exposure, 650 �C

for 420 h, yielded a depth of 35 ± 5 lm. This measure-

ment is based on sixty measurements made at varying

positions along the surface of a *25.4-mm long sample.

The higher sensitivity of the EPMA revealed the gradient

in oxygen concentration and also showed that the oxygen

penetrated further into the material than was evident by

mean measurement by the oxalic etching method (Fig. 5a).

Moreover, the data showed good correlation with microh-

ardness measurements made on the same lot of exposed

material [21]. The concentration of major alloying ele-

ments varied somewhat due to the partitioning of the a- and

J Mater Sci (2012) 47:7235–7253 7237

123

b-stabilizing elements to their respective phases, but no

large, systematic losses from the surface were noted

(Fig. 5b). The oxide scale formed on the surface had a

complex morphology consisting of lenticular and nodular

features both with multiple length scales (Fig. 6). The

former constituent, which covered the majority of the

exposed surface, coarsened slightly with exposure time.

Similar morphology has been observed in another near atitanium alloy, IMI-834 [1]. X-ray diffraction analysis of

the surface scale of the IMI-834 indicated that it was

comprised entirely of Ti and TiO2 when the exposure

temperature was less than 750 �C. Similarly, the only oxide

detected by X-ray diffraction analysis of the surface scale

in the present investigation was TiO2 (results omitted to

save space). In cross section, the scale was approximately

1-lm thick and was rigidly adherent to the substrate.

Consistent with previous observations [1], no spallation

Fig. 1 Backscattered electron (BSE) image showing the microstruc-

ture of the as-received Ti-6242S

Fig. 2 Equal area projections showing the texture of the as-received

Ti-6242S; the contours are in levels of 1.0 times the probability of a

random distribution (maximum intensity: 5.7). The rolling and

transverse directions are parallel to X and Y, respectively, while the

sheet normal is orthogonal to the plane of the projection

Fig. 3 Crystal orientation map depicting the spatial distribution of

grain orientations. Here, grains are colored with respect to the sample

rolling direction (into the plane of the image) which also corresponds

to the loading direction of the tensile and fatigue samples. The sample

reference frame is equivalent to that in Fig. 2 (Color figure online)

Fig. 4 Scalar misorientation angle distribution for the as-received

Ti-6242S material calculated using neighboring grain method (as

opposed to neighboring pixel method) with a minimum a misorien-

tation angle of 2�

7238 J Mater Sci (2012) 47:7235–7253

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was detected during the isothermal exposure or during

subsequent mechanical testing.

Mechanical test results

The bulk elastic modulus and yield strength were essen-

tially unaffected by the presence of the oxygen-enriched

layer. However, a slight decrease in ultimate tensile

strength and significant decrease in tensile elongation to

failure were observed following exposure for 420 h. The

tensile tests have been summarized in Table 1. The strain

gauges on the samples in the as-received condition failed

at a strain of approximately 0.15, while the samples

continued to elongate to approximately 0.215 based on

Fig. 5 a Background corrected and normalized oxygen Ka peak

intensity measured by EPMA and b major alloying element concen-

tration profiles as measured by wavelength dispersive spectroscopy.

Note that all data were obtained using background subtraction

methods. The range of oxygen penetration depths measured by the

oxalic etching technique as well as a microhardness profile are also

shown in (a) (Color figure online)

Fig. 6 Secondary electron images showing the morphology of the oxide developed on the sample surface during exposure at 650 �C for times of

a 118 h and b 420 h

J Mater Sci (2012) 47:7235–7253 7239

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the crosshead displacement at failure. Conversely, the

exposed samples failed at significantly smaller strains

under the same nominal loading conditions. The exposed

samples without the oxide layer failed at just a slightly

higher level of strain compared to the exposed (with

oxide), which was still much less than the as-received

samples.

Fatigue test results from the as-received samples are

compared with those from the exposed samples in

Table 2. These data reveal that the exposure caused a

significant decrease in lifetime over an order of magni-

tude in some cases. There was also less scatter associated

with the exposed samples compared to the non-exposed.

In fact, three of the exposed samples failed within 216

cycles of each other with a total range of only 2,551

cycles. In contrast, much wider scatter, with a range of

approximately 60,000 cycles, was noted in the as-received

material.

Fractography

The fracture surfaces of the exposed samples deformed

under monotonic tension and continuous cycling are

compared and contrasted with samples tested in the

as-received condition.

Monotonic loading

The fracture surfaces of samples tested in the as-received

condition were typical of ductile metals exhibiting dimples

and tear ridges over the entire fracture surface. As evident

by comparing Fig. 1 with Fig. 7a, the size of the ductile

dimples and spacing of the tear ridges was of the order of

the primary a grain size. The fracture surfaces of the

exposed samples, on the other hand, had facets (Fig. 8)

along each surface of the sample that was in contact with

air during the thermal cycle and the dimensions of the

individual facets were also consistent with the primary

grain size. This region extended approximately 40–50 lm

in from the surface of the sample which correlated well

with depth of oxygen penetration. The facet surfaces

(Fig. 8) had features reminiscent of river markings on

cleavage facets [22] and these markings changed directions

abruptly at grain boundaries suggesting that the early

stages of crack growth were strongly crystallographic in

nature. At higher magnifications and resolutions afforded

by UHR mode (Fig. 9), the facet surface topography

became more evident. At the smallest crack lengths,

nearest the exposure surface, the facets appeared micro-

scopically smooth, but their roughness increased with

increasing crack length. This surface roughness was man-

ifested as increases in the height and density of tear ridges

on the facets. This increased roughness, consistent with

increased local strain to failure, correlated well with the

decaying oxygen concentration with increasing depth into

the sample. The cracking mode transitioned to ductile

microvoid coalescence outside of the oxygen-enriched

region at which point it was indistinguishable from the

fracture surfaces of the samples tested in the as-received

condition (Fig. 7b).

Beneath the fracture surface, but still within the gauge

section of the specimen, there was a periodic array of

cracks oriented nominally orthogonal to the loading

direction (Fig. 10a) [21]. Although not shown, a similar set

of periodic cracks were observed in the tensile samples that

Table 1 Tensile test results

Specimen

ID

Modulus,

GPa

Yield

strength,

MPa

UTS,

MPa

Failure

strain

(mm/

mm)

Comments

T1 121 952 1010 *0.215b As-received

T2 120 955 1016 *0.215b As-received

T3a 121 950 [950 [0.010 650 �C/

420-h

T4 120 950 982 0.052 650 �C/

420-h

T5a 121 958 [958 [0.010 650 �C/

420-h

T6 123 947 990 0.069 650 �C/

420-h

T7 123 952 996 0.072 650 �C/

420-h

(oxide

removed)

T8 124 947 985 0.066 650 �C/

420-h

(oxide

removed)

a These tests were intentionally interrupted just after the yield pointb Extensometer limit exceeded

Table 2 Results of cyclic fatigue testing on as-received (C) and

exposed (CE) specimens

Specimen ID Nf Specimen ID Nf

C1 37,610 CE1 4,548

C2 47,207 CE2 5,739

C3 57,908 CE3 5,647

C4 66,303 CE4 7,099

C5 96,452 CE5 5,523

Mean: 61,096 Mean: 5,711

Maximum stress = 800 MPa, R = 0.05, Nf = number of cycles to

failure

7240 J Mater Sci (2012) 47:7235–7253

123

were tested after the oxide had been removed. SEM

examination of the open cracks showed that these fracture

surfaces also exhibited extensive facet formation

(Fig. 10b). Transverse sections of the cracks (made on

longitudinal sections of the sample gauge) revealed that

those cracks which did not propagate to failure showed that

they were approximately 50-lm deep, similar to the

depth of the faceted regions from the surface of the sam-

ple (Fig. 11a). Crystal orientation mapping via EBSD

(Fig. 11c) revealed that the crack followed both inter-

granular and transgranular paths. The orientations of the

grains on both sides of the crack face are reported in

Fig. 11d which shows that the crack was growing through

grains belonging to both of the major texture components

(Fig. 2). Thus, it can be inferred that neither microtexture

nor macrotexture influenced the crack path. Several sub-

micron-sized voids associated with phase and grain

boundaries (Fig. 11b) were observed near the arrested

cracks. The density of voids was higher near the surface of

the sample, where the O content was higher. In general, as

determined by EBSD, void formation occurred at triple

points and also grain boundaries with highly misoriented

c-axes (usually between 60� and 90�, regardless of the

orientation of the c-axes with respect to the loading axis),

but was occasionally observed at lower angle boundaries.

Because commercial Ti alloys generally contain no hard

second phases or non-metallic inclusions, voids are

nucleated at slip band intersections with interfaces,

including grain and phase boundaries, or because of plastic

incompatibility at these locations [23, 24]. In this case,

void formation within the oxygen-enriched region was

Fig. 7 Secondary electron images showing ductile microvoid coalescence in the a as-received and b exposed samples; a is representative of the

entire fracture surface while b is representative of the regions of material which had the bulk oxygen content

Fig. 8 Secondary electron image of faceted fracture in the oxygen-

rich, near-surface region of the sample subjected to monotonic

tension. The numbers identify facets that were investigated with the

quantitative tilt fractography/EBSD technique

J Mater Sci (2012) 47:7235–7253 7241

123

presumably because high levels of oxygen inhibited dis-

location motion which would normally have maintained

plastic compatibility.

Cyclic loading

The fracture surfaces of the as-received (unexposed)

cyclically loaded samples (Fig. 12) were typical of

relatively high maximum stress fatigue tests. Using the

facet surface roughness as an indication of crack length

[25], it was determined that crack initiation occurred from

within a single grain at the surface of the sample and

propagated inward and along the sample surface. A few

facets were formed at small crack lengths where the cyclic

stress intensity factor range was too small to permit suffi-

cient crack tip opening displacement for striation growth.

Fig. 9 UHR secondary electron images of the facet surfaces

7242 J Mater Sci (2012) 47:7235–7253

123

This is likely because the cracks are propagating near the

basal plane which implies that striation formation would

require operation of hc ? ai slip systems which have

considerably higher critical resolved shear stress than the

hai type slip systems. To date, no threshold value for the

transition from faceted to striation growth has been defined

as this value would depend on both grain orientation and

crack length. In general, these facets were widely separated

and did not form a contiguous path. While appearing

brittle, i.e., flat and lacking dimple formation, at low

magnification, the facets exhibited evidence of localized

plastic deformation on their surfaces, an example of which

is shown in Fig. 13a, c. Owing to the apparent lack of

grains with their basal planes suitably oriented for faceted

growth [12], however, other crack advance mechanisms

were operative at small crack lengths including striation

growth, ductile tearing, and fluting [11, 26]. Flutes are

equivalent to classic ductile dimples, except that the void

grows asymmetrically because of the plastic anisotropy of

the a phase [27]. It is important to note that although the

flute walls appear smoother, suggesting that they are

formed by a more brittle process, there is actually a sig-

nificant amount of strain expended during their formation

due to slip on intersecting prism hai slip systems [11, 27,

28]. Shallow fatigue striations were observed in isolated

grains at relatively short crack lengths (\250 lm) which

became more prominent and deeper with increasing crack

length. One extreme example, where striations were evident

\75 lm from the crack initiation site, is shown at high

magnification in Fig. 14. Some secondary cracking

between striations was also evident. Grain 6, on the other

hand, formed a flute when it fractured (Fig. 13b, d). It is

worth nothing that the features that look like striations on

the flute walls (Fig. 13) are actually due to the intersection

of slip bands with the fracture surface, also known as

serpentine glide, as discussed by Chesnutt and Williams

[29], and thus do not indicate progressive crack growth like

typical fatigue striations. As a result, there is no knowledge

regarding the number of cycles required to form a flute on a

fracture surface during cyclic crack growth. Aside from

these few isolated features, much of the fracture surface

was non-descript exhibiting microtear ridges and dimples,

the spacing of which were much less than a grain diameter.

The fast fracture region of the sample exhibited large-sized

dimples, tear ridges, and other features characteristic of

ductile fracture where the dimple spacing was of the order

of a grain diameter.

Considerable differences were noted within *150 lm

of the initiation sites on the fracture surfaces of the exposed

and cyclically loaded samples compared to samples in the

as-received condition. For instance, the extent of faceted

growth was far greater following exposure with intercon-

nected facets extending from the surface to depths of

approximately 40 lm. In the as-received condition, crack

initiation occurred within a single a grain, whereas a line

origin was observed along the surfaces of the exposed

samples. In this context, a line origin implies that crack

initiation occurred from multiple locations that collectively

Fig. 10 Secondary electron images showing a orthogonal, periodic cracking of the gauge section and b facets inside an open crack

J Mater Sci (2012) 47:7235–7253 7243

123

form the periodic cracks shown in Fig. 10a. A typical

region within the line origin is shown in Fig. 15. Around

the initiation sites, the fracture surface was essentially

indistinguishable from that of the tensile sample with both

exhibiting contiguous faceted features. Cracking of the

gauge section on planes orthogonal to the stress axis was

also observed although the extent was far less than during

monotonic loading, i.e., these cracks were not observed to

extend over the entire gauge section. This is most likely

due to the lower stress level although a rate-sensitivity

effect cannot be ruled out. Consistent with the previous

cases, facet surface roughness increased as the crack

length increased (Fig. 16). Beyond the faceted region, the

dimensions of which corresponded well with those of the

oxygen-affected region, the fracture surface was indistin-

guishable from that of the as-received sample exhibiting

striation growth and ductile tearing.

Fracture plane determination

The spatial and crystallographic orientations of the facets

on various fracture surfaces were investigated by the

Fig. 11 a Typical BSE image of a longitudinal section of the exposed

and monotonically loaded sample showing the depth of the periodic

surface cracking (the loading direction is horizontal), b higher

magnification BSE image showing the presence of subsurface voids at

grain and phase boundaries, c crystal orientation map illustrating the

occurrence of both inter- and transgranular fracture, and d distribution

of grain orientations along the crack path (Color figure online)

7244 J Mater Sci (2012) 47:7235–7253

123

quantitative tilt fractography/EBSD technique. The crys-

tallographic orientations of the fractured grains are pre-

sented on inverse pole figures with respect to the loading

direction and the facet normal direction in Figs. 17, 18, 19.

The spatial orientations of the facets are summarized in

Fig. 20. Recall that a point on an IPF identifies the {hkil}

plane normal to a particular direction in the sample refer-

ence frame. This technique reveals the actual crystallo-

graphic fracture plane to an accuracy of *3� [30] through

the use of a geometric correction to account for the spatial

orientation of the facet surface without knowledge of its

inclination a priori.

Monotonic loading

Under quasi-static loading conditions, no facets were

observed on the as-received sample, but extensive facet

formation was observed in the exposed sample. The ori-

entations of 12 such fractured grains are shown in Fig. 17.

These data reveal a rather wide distribution of grain ori-

entations in which the c-axes were inclined between 15�and 70� from the tensile direction. Despite the wide range

of grain orientations sampled none of the facets formed on

the basal plane (Fig. 17). In fact, none of the facet planes

were parallel with low-index crystallographic planes or the

expected a titanium slip systems (basal, prism, first or

second order pyramidal), rather were coincident with a

variety of high order {hkil} planes. No obvious correlation

between grain orientation and fracture plane existed. For

example, grains 3 and 5 have significantly different ori-

entation with their [0001] axes inclined *17� and 60� to

the loading direction, respectively, yet they fractured on

essentially the same crystallographic plane (when crystal

symmetries are considered). In addition, grains 2, 9, and

10a had similar orientations, yet fractured on significantly

different planes.

Cyclic loading

The high stress level used in the fatigue test precluded

significant facet formation in the as-received samples; so,

very few grains were available to interrogate with the

quantitative tilt fractography/EBSD technique. As a result,

the orientations of other non-faceted features were also

investigated. The results, shown in Fig. 18, indicate that

grains 1, 2, and 4 are all fractured on the basal plane and

consistent with typical faceted growth during fatigue [8, 9].

The c-axes of these grains were inclined between 29�, 27�,

and 7� from the loading direction, respectively. The c-axes

of grains 3 and 6 were nearly perpendicular to the loading

direction, and therefore these grains were oriented for

prismatic hai slip and thus would be expected to fracture by

striation growth [9]. Shallow striations were indeed

observed on the fracture surface corresponding to grain 3,

but fluted fracture was observed in grain 6. This may have

been caused by differences in the microstructure sur-

rounding the latter grain. Van Stone et al. [27, 28] have

argued in favor of a constraint requirement for flute for-

mation. Another contributing factor is likely the orientation

of the grain relative to the local crack front. For exam-

ple, fluted fracture has previously been observed in

highly textured Ti–6Al–4V where grains were nominally

{10�10}[0001] oriented, where {hkil} is perpendicular to

the loading axis and [uvtw] is orthogonal to the local crack

front [31], whereas classic striation growth was observed

in {10�10}\11�20[- and {11�20}\10�10[-oriented samples.

Grain 5, which fractured near the {2�1�11} plane, also

exhibited striation formation. In general, it is not possible

to obtain EBSD patterns from striated regions because of

the large amounts of plastic deformation associated with

ductile crack extension at moderate to high DK [9]. In this

case, however, the crack was small when these striations

were formed such that subsurface deformation was limited

and therefore EBSD patterns were obtained directly from

the as-fractured surface. The occurrence of isolated faceted

and striation growth at similar crack lengths is a clear

indication of the orientation dependence of fracture mode

and thus in the local crack growth rate.

In contrast to samples in the as-received condition, there

were many facets on the fracture surface of the exposed

and cyclically loaded samples. A total of 14 facets were

Fig. 12 Secondary electron image of the crack initiation site in the

as-received and cyclically loaded sample. The numbers identify

grains that were investigated with the quantitative tilt fractography/

EBSD technique

J Mater Sci (2012) 47:7235–7253 7245

123

analyzed in detail and the results are presented in Fig. 19.

Although a wide range of grain orientations were probed

with c-axes inclined between *14� and 85� from the

loading direction; none of the fracture planes coincided

with slip planes and only one (grain 1) was found to be

nearly coincident with a low-index {2�1�10} plane. In

general, the grains which were aligned for easy prismatic

hai slip (grains 1, 3, 9 and 13) tended to fracture on a plane

nearly perpendicular to the c-axis, whereas the remaining

grains (which were more suitably aligned for basal slip if

dislocation glide indeed occurred) fractured on planes

inclined between 15� and 50� from (0001).

Fig. 13 UHR secondary electron images of a facet and a flute at low (a) and (b) and high (c) and (d) magnifications, respectively. The direction

of local crack propagation in denoted by the arrows

7246 J Mater Sci (2012) 47:7235–7253

123

Discussion

The salient results from this research included observations

of the microstructure and texture in the as-received sheet

and their relation to features observed on the fracture sur-

faces. Based on these observations, the mechanisms of

crack initiation and propagation in the presence of elevated

concentrations of oxygen are discussed below.

Microstructure and texture

The microstructure of the as-received and exposed samples

at the constituent length scale was essentially the same

because the exposure temperature was too low to change

the microstructure, at least at the constituent length scale.

The texture of the rolled sheet consisted of the typical basal

and transverse components that would be expected from

subsolvus rolling or plane strain compression [32]; how-

ever, the presence of the texture component with [0001]

parallel to the rolling direction is interesting as this ori-

entation is not predicated as a stable orientation within the

crystal plasticity framework. This is important because the

texture of the rolled sheet and plate is often generalized as

only containing basal or transverse texture components [33,

34]; thus, the texture of plate and sheet products is not

always measured as a matter of course during production.

A similar texture component, in which the c-axis aligns

with the primary material extension direction during

processing, has also been observed following subsolvus

heat treatment of extruded a ? b titanium [35, 36] and also

within secondary a (or a0) transformation texture in hot-

rolled Ti–6Al–4V [32, 35]. The transformation texture can

be ruled out as causing the [0001] || RD component in the

present investigation due to the low volume fraction of

secondary a and the lack of the intensities at *45� to both

the rolling and transverse directions, which would be

expected based on transformation from the typical

(110)[001] BCC rolling texture which forms during both

sub- and supersolvus rolling [32, 33, 37]. Since material

texture affects the formability of titanium products and also

impacts the mechanical properties at the component level,

further investigation into the formation of textures with

preferred [0001] directions that are not predicted by plas-

ticity theory is warranted.

Fig. 14 UHR secondary electron image of the fracture surface

showing shallow striations in grain 5. The arrow denotes the direction

of local crack propagation

Fig. 15 Secondary electron image of the crack initiation site in the

exposed and cyclically loaded sample. The numbers identify grains

that were investigated with the quantitative tilt fractography/EBSD

technique

J Mater Sci (2012) 47:7235–7253 7247

123

Fractography, crystallography, and the mechanisms

of facet formation

Perhaps, the most notable experimental result from the

present investigation was the occurrence of extensive facet

formation in the exposed specimen under quasi-static

deformation rates. In general, facets are normally associ-

ated with fatigue crack initiation. It is generally accepted

that this type of initiation during continuous cycling of

material with standard oxygen content occurs by the for-

mation of intense slip bands followed by crack initiation

and subsequent cycle-by-cycle propagation back along the

slip band [9, 38–40]. This process continues at the grain

level over the entire crack front until there are no grains

suitably oriented for facet formation, or the cyclic stress

intensity range, DK, at a particular location is sufficiently

large to permit enough crack tip opening displacement to

blunt the crack, which is subsequently resharpened upon

closing, to form a fatigue striation. At this stage, crack

propagation then occurs either by striation growth or by a

combination of cyclic ductile tearing and/or microvoid

formation at phase boundaries permitting stable ductile

crack advance. The cyclically loaded as-received sample

exhibited precisely this type of behavior. There was minor

Fig. 16 UHR-SEM images of facets a 1 and b 5 from the exposed and cyclically loaded sample

Fig. 17 Results from the facet

crystallography analysis for

exposed and monotonically

loaded samples. Note that 15�grid spacing was used and

inversion symmetry was notenforced (Color figure online)

7248 J Mater Sci (2012) 47:7235–7253

123

facet formation due to the lack of grains with suitably

oriented basal planes (Fig. 2) [10, 12, 41], and the transi-

tion to striation growth occurred at a relatively small crack

length due to the low load ratio (R = 0.05) combined with

the relatively high peak stress and therefore large DK. In

contrast, the tensile and fatigue fracture surfaces of the

exposed samples contained many facets. Since the sample

orientation was the same, the increase in the extent of

faceted fracture could not be attributed to the differences in

texture or microstructure morphology and therefore could

only be due to the oxygen enrichment. The physical

mechanisms for the degradation of properties are discussed

next. First, however, it is important to point out that the

orientation of the test specimens relative to the texture of

the as-received sheet was such that there were grains in a

variety of hard and soft orientations, both elastically and

plastically. This suggests that the macrotexture did not

affect the orientations of the grains that fractured during

testing. More specifically, the mechanism of crack

Fig. 18 Results from the facet

crystallography analysis for

cyclic loading in the as-received

sample. Note that 15� grid

spacing was used and inversion

symmetry was not enforced

(Color figure online)

Fig. 19 Results from the facet

crystallography analysis for

exposed and cyclically loaded

samples. Note that 15� grid

spacing was used and inversion

symmetry was not enforced

(Color figure online)

Fig. 20 Summary of facet normal angles for the various tests

J Mater Sci (2012) 47:7235–7253 7249

123

initiation and crack growth was not altered by the abun-

dance, or lack, of available slip systems. Grains that would

tend to deform predominantly by basal hai, prism hai and

pyramidal hc ? ai slip were all present.

Crack initiation

Crack initiation occurred at multiple locations along the

surfaces of the exposed tensile samples resulting in exten-

sive faceted growth. In contrast to crack initiation during

cyclic or dwell loading [8, 25], for instance, the initiation

process could not be linked to a particular grain or region

within the faceted zone and was instead determined to be a

line origin. Periodic cracks orthogonal to the tensile axis

formed on the gauge section during both monotonic (both

with and without the surface oxide present) and cyclic

loading suggested that crack initiation was not sensitive to

the local microstructure, instead was dominated by the

macroscopic applied load. The transverse metallographic

sections made through the arrested cracks revealed two

initiation mechanisms. The first was cracking of the oxide

followed by subsequent penetration of the crack into the

material. An example of this mechanism is shown in

Fig. 16a. The cracked oxide is visible on the left hand side

of the image while the smooth facet is on the right. In the

second, there was evidence of grain boundary microvoid

nucleation at locations within the oxygen-enriched region

of the sample, but at depths greater than the thickness of the

actual TiO2 oxide layer (Fig. 11). Both micromechanisms

of fracture were operative to cause the periodic cracking

observed on the sample scale in the as-exposed condition. It

is also noteworthy that similar cracking behavior was

observed in samples where the surface oxide was removed,

indicating that the oxygen-enriched region is also suffi-

ciently brittle to sustain long range, rapid cracking. Similar

voids were not observed in the bulk regions of the material

unaffected by oxygen, meaning that their formation must be

related to the higher interstitial content in this region.

Although void formation is typically associated with

enhanced ductility, this occurred at stresses well below

macroscopic yield [4, 42], thereby severely degrading the

overall tensile elongation to failure of the sample because

the cracks propagated at low applied stresses.

With regard to crystallographic orientation, none of the

facets analyzed that bordered the surface of either the

exposed tensile sample (Fig. 8: #’s 7 and 9) or the exposed

fatigue sample (Fig. 11: #’s 1–3, 9, 12) exhibited a low

index crystallographic fracture plane. Thus, crack initiation

was neither the result of classical cleavage [22] nor con-

ventional faceted growth [40]. The former occurs via

atomic separation along low-index crystallographic planes,

while the latter occurs along slip bands (which are them-

selves typically parallel to low-index planes).

Crack propagation

Regardless of the mechanism of crack initiation, the

method of propagation was similar within the oxygen-

enriched regions of exposed tensile and fatigue samples as

evidenced by their similar appearance and similar crystal-

lographic and spatial characteristics. In near-a and a ? balloys with standard oxygen content, facets do not form

under quasi-static loading, and facets formed during con-

stant amplitude cyclic fatigue are typically parallel to the

basal plane [7, 9, 18, 43]. Faceted growth gradually tran-

sitions to striation growth as the crack length, and thus

DK increases. A larger DK correlates with a larger crack tip

plastic zone size which, in turn, implies increased defor-

mation within each subsequent grain as the crack grows

until eventually there is sufficient crack tip opening dis-

placement to form fatigue striations.

The propagation mechanism was altered substantially,

however, by the elevated temperature exposure and inward

diffusion of oxygen. First, the extent of faceted growth was

governed by the extent of oxygen enrichment as opposed to

DK. In addition, there were no indications of progressive or

incremental crack growth through the faceted region.

Transverse cross sections through the periodic cracks in the

gauge section (Fig. 11) revealed that the crack tip was

blunted and arrested after growing out of the oxygen-

enriched region. This can be rationalized on the basis of the

effect of oxygen on slip character. A crack tip is usually

blunted as it grows through ductile material. In the pres-

ence of oxygen, however, slip is highly localized and

constrained into relatively few slip bands ahead of the

crack tip which inhibits blunting. As a result, the crack tip

remains sharp and therefore has a strong driving force for

further extension through the enriched material. With

increasing depth from the surface, the facet surfaces appear

increasingly rough which is consistent with more localized

plasticity during fracture due to lower oxygen content.

The metallographic cross sections also showed that the

crack path was both transgranular and intergranular. This

observation can explain the occurrence of fracture on high

order {hkil} planes following exposure (Figs. 17, 19).

Consider the following scenario in the context of the

schematic microstructure presented in Fig. 21. Upon

reaching a critical stress level in the elastic regime, the

oxide cracks and voids are formed at the boundaries and

triple points between highly misoriented grains within the

oxygen-enriched region. Based on its macroscopic

appearance and orientation, cracking in the oxygen-rich

region is clearly mode-I dominated (Figs. 10, 11), but

details of the local microstructure and arrangement of voids

will govern the grain-level crack path through the micro-

structure. In the hypothetical case (Fig. 21), the crack in

the oxide encounters grain A and seeks to relieve the stress

7250 J Mater Sci (2012) 47:7235–7253

123

field associated with the void at triple point ABC resulting

in transgranular failure under nominally mode-I loading.

Next, the crack propagates along grain boundary BC (the

path of least resistance) to triple point BCD. The next

growth increment is transgranular through grain D because

growing along either boundary to the closer voids would

result in too far of a deviation from mode I growth. This

tradeoff between intergranular/transgranular growth con-

tinues until the crack tip has grown out of the oxygen-

enriched region at which point more conventional growth

mechanisms are observed: striation growth for cyclic

loading and void growth and coalescence for tension.

While the non-crystallographic fracture has been

explained, it is necessary to describe why this type of non-

conventional growth is preferred over slip band cracking

and/or cleavage. Grain and phase boundary diffusivity have

estimated to be several orders of magnitude faster than bulk

diffusivity in the present alloy at 650 �C [20]. This implies

that the grain and phase boundaries have higher oxygen

concentration compared to the a grain interiors. Conse-

quently, the grain boundaries were preferentially embrittled

which explains the propensity for void nucleation at this

location. Void growth, on the other hand, is suppressed

because oxygen strengthens the grain interiors making the

grain boundaries the weakest link. As a consequence, the

voids cannot grow and the crack tip follows the path of

least resistance to link the voids whether it is transgranular

or intergranular. Evidence of microplasticity on the facet

surfaces (Fig. 9, for example) indicates that this is not an

entirely brittle process and it is likely that there is highly

localized dislocation activity at the crack tip facilitating its

growth. This is being investigated in more detail by means

of site-specific foil extraction and transmission electron

microscopy analysis.

Implications of oxygen ingress on properties

Based on the results presented above, it can be concluded

that the inward diffusion of oxygen alters the grain-level

cracking mechanisms operative in Ti-6242S during quasi-

static and cyclic loading. The change in fracture mode has

potentially significant implications on the lifing of com-

ponents that are exposed to high temperature during service

because of the occurrence of inter- and transgranular

cracking under nominally elastic loads. Fractography

revealed the extent of plastic deformation associated with

grain-level facet formation following exposure which was

much lesser than that during cyclic loading in the

as-received material. A normal-stress controlled fracture

criterion therefore governs the material’s properties in the

oxygen-rich region of the sample. In fact, Liu and Welsch

[44] have shown that Ti–6Al–2V (close to the a phase

composition in Ti–6Al–4V) with a bulk oxygen level of

0.65 wt %O fractures at stresses as low as *300 MPa

which is well below the macroscopic yield point in material

with only 0.07 wt %O. In the present case, the high-

oxygen levels are concentrated near the surface, and thus

the brittle, normal-stress controlled fracture event is limited

to this region which forms a sharp precrack at nominally

elastic strains for the bulk of the material. This may lead to

anomalous growth rates under cyclic loading due to small

crack effects [45–49] or via room temperature creep crack

growth [50] under nominally elastic loads. At crack lengths

exceeding the dimensions of the oxygen-enriched region,

the fracture surface of the exposed sample was indiscern-

ible from the as-received samples, which consisted of

fatigue striations and some evidence of microvoid forma-

tion at grain and phase boundaries. Thus, the debit in

tensile ductility and fatigue life following exposure can be

attributed to a decreased resistance to crack initiation and

accelerated crack growth rates in the oxygen-rich region.

As evident in Table 2, the scatter in fatigue lifetime for

the exposed samples was much less than the material in the

as-received condition. This was attributed to the fact that

the embrittled oxygen region cracked under nominally

elastic loading to the depth of oxygen penetration. This

essentially negated the number of cycles to crack initiation

by forming a ‘‘precrack’’ at the beginning of the test. Since

the depth of the ‘‘precrack’’ was controlled by the oxygen

penetration depth, which was consistent from sample to

sample, the total fatigue lifetime in the exposed samples

was dominated by the crack growth regime. In contrast, the

lifetimes of the samples in the as-received condition

included the stochastic nature and uncertainty associated

with time to crack initiation resulting in a wider range of

fatigue lifetimes. This interpretation is consistent with the

results of Jha et al. [50] who showed that uncertainty in

crack initiation lifetime was the major source of total

Fig. 21 Schematic representation for the occurrence of trans- and

intergranular failure. The gray circles represent the intergranular

voids depicted in Fig. 11

J Mater Sci (2012) 47:7235–7253 7251

123

lifetime variability in fatigued samples exhibiting bimodal

lifetime behavior.

Conclusions

The micromechanisms of crack growth during quasi-static

and cyclic loading of Ti–6Al–2Sn–4Zr–2Mo following

long-term elevated temperature exposure to laboratory air

have been investigated. Tensile elongation to failure and

fatigue life were severely compromised by the ingress of

oxygen. The following conclusions were reached:

– There was a change in the fracture mechanism from

classic ductile dimple failure to brittle facet formation in

the regions of the tensile sample, that were enriched with

oxygen during the thermal exposure. The change was not

due solely to the formation of a brittle TiO2 oxide.

– Void formation was observed at high angle grain

boundaries and interphase boundaries in the oxygen-

rich region following tensile loading. Void growth

(ductility) was restricted by the high-oxygen content,

and thus fracture occurred along grain boundaries and

non-crystallographic planes spatially oriented perpen-

dicular to the loading direction.

– Faceted, striation, and fluted fractures, which all

occurred by slip processes on low-index, rational

crystallographic planes, were observed near the crack

initiation sites on the fracture surface of the as-received

fatigue samples. In contrast, extensive transgranular

facet formation on high-index crystallographic planes

and brittle, intergranular fracture were observed near

the crack initiation sites on the exposed samples. At

longer crack lengths away from the oxygen-enriched

region, the fracture surfaces were indistinguishable

from samples in the as-received condition.

– Although the appearance of the fracture surface is

consistent with brittle fracture, cracking did not occur

along low-index crystallographic planes and thus was

not classical cleavage. Nevertheless, cracking in

the oxygen-rich region was normal-stress controlled.

The correlation between oxygen concentration and the

critical stress level for cracking warrants further

investigation.

Acknowledgements This work was performed as part of the

in-house research activities of the Air Force Research Laboratory,

Materials and Manufacturing Directorate, AFRL/RXLM, Wright

Patterson Air Force Base, OH. The financial support of the Air Force

Office of Scientific Research through Task No. 09RX24COR,

Dr. David Stargel, Program Manager, is gratefully acknowledged. Two

of the authors were partially supported under onsite Air Force contracts

FA8650-07-D-5800 (ALP), Dr. Ali Sayir, Program Manager, and

FA8650-09- D-5223 (WJP) during the time this work was completed.

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