NBSIR 76-836

LOW TEMPERATURE FRACTURE BEHAVIOR OF

A Ti-6AI-4V ALLOY AND ITS ELECTRON BEAM WELDS

R L Tobler WmmM81?

Cryogenics Division

Institute for Basic Standards

National Bureau of Standards

Boulder, Colorado 80302

April 1976

Prepared for

Kirtland Air Force Base

Alb uquerque, New Mexico 87117

NBSIR 76-836

LOW TEMPERATURE FRACTURE BEHAVIOR OF

A Ti-6AI-4V ALLOY AND ITS ELECTRON BEAM WELDS

R.L. Tobler

Cryogenics Division

Institute for Basic Standards

National Bureau of Standards

Boulder, Colorado 80302

April 1976

Prepared for

Kirtland Air Force Base

Alb uquerque, New Mexico 87117

U.S. DEPARTMENT OF COMMERCE, Elliot L. Richardson, Secretary

James A. Baker, III, Under Secretary

Dr. Betsy Ancker-Johnson , Assistant Secretary for Science and Technology

NATIONAL BUREAU OF STANDARDS, Ernest Ambler, Acting Director

CONTENTS

Page

1. INTRODUCTION 1

2. EXPERIMENTAL 1

2.1 Material 1 2.2 Electron Beam Melds 2 2.3 Specimens 2 2.4 Orientation 3 2.5 Microstructure 3 2.6 Procedure 3 2.6.1 Fracture 4 2.6.2 Fatigue 5 2.6.3 Static Load Cracking 5

3. RESULTS 6

3.1 Fracture Toughness 6 3.2 Fatigue Crack Growth Rates 8 3.3 Static Load Cracking (SLC) 8

4. SUMMARY 10

5. DISCUSSION 11

6. RECOMMENDATIONS 13

7. ACKNOWLEDGMENTS 14

8. REFERENCES 15

iii
ABSTRACT

The effects of electron beam (EB) welding on the fracture

behavior of a recrystallization annealed, extra-low-interstitial

Ti-6A1-4V alloy have been investigated at temperatures in the

ambient-to-cryogenic range. Plane strain fracture toughness

(KT ) and subcritical crack growth parameters were measured

Ic using compact specimens 10 to 25.4 mm thick. These parameters

can be used to predict the safe operating lifetimes of cryogenic

pressure vessels and other welded Ti-6A1-4V structures.

Although EB welding transforms the base metal microstructure extensively, its effects on the material's fatigue

crack propagation resistance at intermediate stress intensity

factors are negligible. The growth rates, da/dN, of fatigue

cracks sited in the fusion and heat-affected zones of

weldments were temperature insensitive and nearly equivalent

to rates for the base metal. However, welding introduces a zone of low fracture toughness at the heat-affected-zone/fusionzone boundary. The K value for this boundary zone at liquid

3/2 nitrogen temperature (76 K) was 45 MN/m , 16% lower than the

base metal. The base metal fracture toughness increases

between 4 and 295 K, with an abrupt transition to higher K

values occurring at temperatures between 76 and 125 K. Static

load cracking, temperature effects, and specimen orientation

effects on the fracture behavior of this titanium alloy are central topics of discussion.

Key words: Electron beam welding; fatigue; fracture toughness; low temperature tests mechanical properties; titanium alloy.

i V
1. INTRODUCTION

As discussed by Johnson [1], Ti-6A1-4V is employed in critical cryogenic applications

related to aerospace. Favored for its high strength-to-weight ratio, this alloy has become

the principal construction material for the cryogenic pressure vessel components of electrical,

propulsion, and reaction control systems. Such pressure vessels are designed using

low safety factors (y 1.5), and may experience service temperatures as low as 20 K (-423°F);

they are usually spherical, and are constructed by welding hemispherical forgings at the

equator.

Although welding is essential in fabricating these and other large engineering structures,

weld zones represent a potential source of structural weakness. Microcracking or porosity

may occur during weld metal solidification, and heat conducted to the base metal may induce

microstructural transformations with resultant mechanical property degradation. In the case

of Ti-6A1-4V, low ductility and toughness in gas metal-arc or gas tungsten-arc welds can be

attributed to oxygen and nitrogen contamination [2] . Thicker sections are necessary to guarantee the structural reliability of a weld if its mechanical properties undermatch those

of the base metal.

Disregarding its high cost, electron beam welding offers significant technical

advantages compared to arc welding. The kinetic energy of high velocity electrons is con- verted to intense heat upon collision with the structural component. The highly localized

heating produces narrow, moderately tapered, deeply penetrating fusion zones. Consequently,

factors such as distortion, residual stress, total heat input, and heat-affected zone sizes are minimized. Also, electron beam welding is performed in vacuum, reducing the possibility

of impurity pickup [3].

Tensile and fatigue data for electron beam welds in 25.4 to 50.8 mm thicknesses of

Ti-6A1-4V are known from tests at room temperature, but most results describe unflawed

specimen behavior [4,5]. Fracture mechanics parameters are needed to predict the low temper- ature service performance of large structures which cannot be guaranteed to be flaw-free.

Such data for an extra-low- interstitial Ti-6A1-4V alloy in the recrystallization annealed

condition are presented here. Fracture toughness (K^) , fatigue crack growth rate (da/dN)

and static load cracking parameters are reported for the unwelded and the electron beam

welded alloy, at selected temperatures between 295 K (70°F) and 4 K (-453°F)

2 . EXPERIMENTAL

2.1 Material

The extra- low- interstitial (ELI) Ti-6A1-4V alloy tested in this study was donated by

Beech Aircraft Corporation, Boulder, Colorado. The material was received in the form of a

1 m diameter hemisphere with a 51 mm wall thickness. The hemisphere had been recrystallization

annealed at 1200 K (1700°F) for 4h, furnace cooled to 1033 K (1400°F) in 3h, cooled to 756

(900°F) in 3/4 h, and air cooled to room temperature.
A commercial chemical analysis of this alloy is listed in Table 1. Mechanical properties

of interest are listed in Table 2. The tensile properties at room and liquid nitrogen

temperatures were measured using established techniques [6,7]. The elastic properties,

measured by an ultrasonic method [8], were equivalent to the published values for a similar

heat.

2.2 Electron Beam Welds

Several pairs of arc-shaped bars were machined from stock and electron beam welded by

Chem-Tronics , Inc. , Santee, California. Most of these bars were 90 cm in arc length and

3.8 x 5.1 cm in cross-sectional area. Two bars having 18 x 5 cm cross-sectional areas were

also welded, but the larger integral heat sink offered by this geometry did not affect the

fracture properties.

The double-rabbeted, self-backing, self-aligning, square-groove joint used for these

welds is shown in Figure la. The workpieces were machined, chemically cleaned, and mounted

between aluminum chill plates in a rotary indexing fixture. Circumferential welding was

-2 -4 accomplished in vacuum at 4 x 10 Pa (3 x 10 torr) , using tack, penetration, and cosmetic

passes. The workpieces were rotated under a vertical electron beam, at a gun-to-work distance

of 32.4 cm. Other machine settings, determined from trial runs on scrap pieces, are listed in

Table 3. The completed welds were annealed in air at 811 K (1000°F) for 50h, simulating a stress-relief heat treatment that is applied to welded pressure vessels.

A weld cross-section that was etched in an 851^0 tlOHFrSHNO^ solution is shown in

Figure lb. Large prior beta grains give the 30 mm long, 5 mm wide fusion zone a mottled

appearance. The included angle of the fusion zone at central thickness is approximately

5 degrees. Heat-affected zones extend approximately 2 mm above and below the fusion zone.

The fracture specimens were machined such that the flawed seam at the weld root (which

is typical for this type of joint preparation) was removed by notching. A compact specimen

of the ML orientation (described below) is shown in Figure lc. Based on commercial ultrasonic

inspections and metallographic evidence, these laboratory specimens are representative of

production quality electron beam welds.

2.3 Specimens

The compact specimen was chosen because it offered a high K measurement capability for

its size [9], and a geometry compatible with that of the welded stock. This standard specimen

is described in the ASTM method of test for plane strain fracture toughness of metallic

materials, ASTM E-399-74 [10].

Base metal specimens were machined from the unwelded forging; weld specimens, from

either of two welded-bar samples. Specimen thickness, B , ranged from 10 to 25.4 mm, and

widths, W, ranged from 28 to 50.8 mm, as listed in Table 4. Without exception, the planar

specimen dimensions were proportional to W, as described in E-399-74, and knife edges for

clip gage retention were machined integral to the notch. The knife edges were usually located

at the specimen edge. An exception is the series of base metal specimens referred to in

Figure 6 below, where loadline deflection is reported.

2
1 . - I ris- tatitr.

Reference axes for specimen orientation are defined in Figure 2. The letter R

represents the radial direction of the original forging, while M and L represent circumferential (meridional and latitudinal) directions. Base metal specimens were machined in as many as six orientations, designated RL, LR, RM, MR, LM, or ML. The first letter of each

designation identifies the direction normal to the fracture plane; the second letter

identifies the direction of crack propagation. Weld specimens were tested only in MR and ML orientations where the cracks propagate radially or equatorially

2.5 Micrtstrutture

The recrystallization annealed base metal microstructure of primary alpha and intergranular beta (or transformed beta phase) is shown in Figure 3. Although nearly equiaxed, the

zi:r: = :r;:;.r; viewed alsr.g the R axis is distinguishable from :'r.:se of the M and L axes. The latter microstructures are equivalent, having an average primary alpha grain diameter of

approximately 0.013 msn.

A -isrcstructural traverse cf the weld is shown in Figure 4. At distances of 4 mm or greater from the fusion zone centerline, the base metal microstructure was unaltered by

veldir.g. The outer heat-affected zone (3 to 3.5 mm from the fusion line) etches darkly, while

the center of the heat-affected zone shows primary alpha in a highly permutated matrix. Near

the heat-affected-zone/fusion-zone boundary, the structure consists of aged acicular alpha.

The fusion zone is entirely aged martensitic alpha.

Three critical weld locations were selected for fracture testing: 1) FZ, the fusion

zone centerline, 2) HAZ, the center of the heat-affected zone, and 3) HAZB, the macros copically visible heat-affected-zone/fusion-zone boundary. Notches parallel to the

fusion line were sited in the desired zones, after etching each specimen blank. The fatigue

cracks subsequently propagated parallel to the notch plane, except in the case of HAZB

specimens. The cracks in HAZB specimens propagated along the heat-affected-zone/fusion-zone

boundary, even though this was at a 2 to 3 degree angle from the notch plane. In any case, no difficulty was encountered in maintaining the fatigue cracks within their intended zones.

2 6 ?r: cedure

The 100 kN capacity servo—hydraulic testing machine and cryostat that were used to conduct these tests are shown in Figure 5. As previously described [11,12], a vacuum

insulated dewar containing a cryogen encapsilates the specimen and clip gage at low temperatures. Test environments included unconditioned laboratory air at 295 K (70°F), ethanol and

dry ice at 195 K (-108°F), liquid nitrogen at 76 K (-323°F) , and liquid helium at 4 K (-453°F)

Temperatures between 195 (-108 °F) and 108 K (-200°F) were also obtained, by admitting cold

nitrogen vapor to the cryostat. A chromel-constantan thermocouple was attached to the

specimen during these tests, and a servo-mechanical temperature controller maintained the

specified test temperature to within + 3 K. Clip gage sensitivity changed by no more than

1.5% between 295 and 4 K, and the E-399-74 linearity requirements were satisfied.

3
2.6.1 Fracture

The fracture toughness specimens were precracked at their K test temperatures and

3/2

fractured at a stress intensity factor increase rate of approximately 1 MN/m per second.

The critical stress intensity factors were calculated according to the solution for compact

specimens

K

Q

= P Q

B~V1/2 [f(a/W)] (1)

where

f(a/W) = 29.6 (a/W) 1/2

- 185.5 (a/W) 3/2 + 655.7 (a/W) 5/2

- 1017.0 (a/W) 7/2

9/2 + 638.9 (a/W) ' (2)

In these equations, is the load defined by the secant offset procedure, and a is the

average of three crack length measurements at 25, 50, and 75% of specimen thickness (B) [10].

The specimen width, W, is defined in Figure lc.

A quantitative measure of fracture resistance under linear-elastic conditions, K^, is

defined operationally by the E-399-74 method. The measurements constitute data only

if plasticity prior to fracture is negligible, and all other E-399-74 requirements are

satisfied. In the tests reported here, the crack front uniformity was satisfactory. However,

in tests at room temperature, the specimen sizes failed to satisfy the requirement that:

B > 2.5 (Kp/O

y

) 2 (3)

where is the uniaxial yield strength of the material. Nonlinear behavior at room

temperature proved unavoidable, because specimens of a thickness meeting Eq. (3) could not be

manufactured from the available stock.

Nevertheless, using an elastic-plastic fracture criterion based on the J-integral,

meaningful values were derived for the base metal. For nonlinear-elastic materials, J is the rate of change of potential energy with respect to crack area. The critical value,

Jj- c

, is the value of J required to initiate crack extension. Since the J-integral applies

to nonlinear load-deflection behavior, the size criterion of Eq. (3) need not be satisfied

for a valid result.

Experimentally, J was evaluated as an energy proportional to the area, A, under the

load-loadline deflection curves of 20.3 mm thick compact specimens. The extrapolation

procedure outlined by Landes and Begley [13] was used. According to their methods, six

nearly identical specimens (a/W = 0.58 + 1%) were loaded to cause decreasing amounts of stable

crack extension. These six specimens were then unloaded, heat-tinted, and fractured into

halves for crack extension measurements. The J values at points of unloading were calculated

using the equation [14]:

4
(4)

where b is the specimen ligament (W-a) , and A is a coefficient greater than 2 and dependent

on a/W [14]. The six values of J obtained were plotted versus crack extension and

extrapolated to determine J T . Finally, KT was obtained from the expression:

Ic Ic

K

lc = fl ' J lc

(5)

1-v

where E is Young's modulus and v is Poisson's ratio. Thus, a room temperature value was estimated, but only for base metal specimens of the LR orientation.

2.6.2 Fatigue

Considerable amounts of fatigue crack growth data were obtained during the precracking

of fracture toughness specimens where the maximum fatigue stress intensity factor did not

3/2 exceed 34 MN/m . Additional specimens were tested at 295 and 76 K, solely for fatigue crack

growth rates at higher stress intensity factors.

The fatigue tests were conducted under controlled load, using a sinusoidal load cycle at frequencies, F, from 10 to 30 Hz. The ratio, R, of minimum/maximum load was 0.1. A digital

indicator was used to continuously monitor the peak loads, which were accurate to + 2% of the

specified values.

Fatigue crack growth was monitored by elastic compliance or crack opening deflection

measurements, an approach based on the fact that deflection per unit load, 6/P, increases

with increasing crack length. Correlations between crack length and 6/P were obtained

experimentally, using the average crack length measurements of fractured specimens. Changes

in clip gage sensitivity and Young's modulus with temperature were accounted for by plotting

the base metal data versus EB 6/P, or by shifting the weld specimen room temperature

calibration curves [12]. A correlation curve applicable to 20.3 mm thick compact specimens

of the base metal (with 6 measured at the loadline) is shown in Figure 6.

The crack growth rate measurement procedure involved recording 6-versus-P periodically.

Crack lengths, inferred from the empirical correlations, were plotted versus cycles, N. Computer programs were used to fit the a-versus-N curves with a third-order polynomial, to differentiate the curves for da/dN data, and to calculate the corresponding stress intensity

factor range, AK, from the peak fatigue stress intensity factors:

AK = K - K . (6) max min

2.6.3 Static Load Cracking

A series of specimens were sustained in inert environments at constant load to determine

whether time-dependent crack growth culminating in fracture could occur. These base metal

(B = 25.4 mm) and weld specimens (B = 10 mm or 17 mm) were precracked as for K testing.

Each specimen was loaded within 15 seconds to a specified stress intensity factor. Many

5
specimens were instrumented with a clip gage so that cracking could be monitored by

deflection measurements.

For the base metal specimens, meaningful static load cracking rates were measured by

unloading after moderate crack extensions of no more than 2 mm. The test procedure involved:

1) sustaining the specimens at constant load for hold times, At, ranging from 15 to 10

seconds, 2) unloading and heat-tinting the specimens to oxidize their crack surfaces, and

3) fracturing the specimens into halves to measure the crack extension. Tinting was accomplished by heating with an acetlyene torch. After specimen fracture, increments of static

load crack extension were identifiable as straw-blue oxidized areas located between the

fatigue crack and the untinted ligament.

An average crack extension, Aa, was obtained from three measurements with a traveling

microscope at 25, 50, and 75% of specimen thickness. The cracking rate, Aa/At, for each test was then calculated and plotted versus the nominal value of K at the initial crack length

and static load, P : s

K = P B"

1W~

1/2

f £ (7)

s W

•

For compact specimens, K increases continuously with crack extension under constant load,

and the Aa/At-versus-K relationships are only approximate. However, the calculations were

performed for Aa values from 1 to 2 mm, and differences between the initial and final K values were not greater than 5%.

3. RESULTS

3.1 Fracture Toughness

Fracture test records for two base metal specimens are shown in Figure 7. These loaddeflection curves are typical of weld specimen behavior also. Linearity prior to fast tracture

occurred at 76 K, where all KT measurements were valid. Nonlinearity and stable crack

Ic ' extension occurred at room temperature, and all tests failed the requirement of Eq. (3). The

maximum loads at 295 K exceeded 1.1 Fn , which also disqualifies the K„ values as K results.

V Q Ic Tables 5-9 list the results at each temperature, along with observations relating to validity

requirements. In several cases a/W exceeds 0.55. Although this deviates from the E-399-74

specifications, the results were not in disagreement with valid data, as indicated by tests at 76 K. This near equivalence of results is indicated in the tables by the statement

Q Ic J- integral data for the base metal at 295 K are shown in Figure 8. Apparently, J varies linearly at crack extensions greater than 0.12 mm. As discussed by Landes and

Begley [13], the J measurement point might be selected by several different methods:

1) The curve could be linearly extrapolated to Aa = 0, identifying J on the

ordinate axis as approximately 80 kJ/m2

.

2) A line, J/2a , may be constructed, its intersection with the J-Aa curve yielding ^ y 2 J T = 88 kJ/m

Ic
3) A 1% crack extension may be defined as critical. In this case, the J-Aa curve ^ / 2 indicates that J T = 100 kJ/m . Ic Summarizing, J for the LR orientation of the base metal at room temperature lies in the

§ 2 range 80 to 100 kJ/m (457 to 571 in-lb./in ), depending on the method applied to select the

measurement point from Figure 8. According to Eq. (5), the corresponding K values for this

3/2 alloy at room temperature are 99 to 111 MN/m

Fracture toughness results for the base metal are plotted in Figure 9, which includes the room temperature data derived by J-integral tests. The ML and LR orientations exhibit a similar temperature dependence: between 295 and 125 K, remains at relatively high shelf

levels; but, as the temperature is lowered between 125 and 76 K, a 33 to 45% decrease of K^

c occurs. On further temperature reductions to 4 K, the data for the LR orientation indicate

only a slight decline of K^. Thus, the recrystallization annealed and furnace cooled

Ti-6A1-4V alloy exhibits a fracture toughness transition at low temperatures which resembles

the classical fracture transitions of other metals and alloys having body-centered cubic or hexagonal crystal structures. The abrupt decrease of in the interval 76 K < T < 125 K suggests that a change in the micromode of fracture occurs, but there was no attempt to verify this f ractographically

Some anisotropy in the fracture toughness data is evident at temperatures above the

transition region. The shelf K values for the LR orientation range from 100 to 110

3/2 < MN/m , while shelf values for the ML orientation average approximately 15% lower. These

orientation effects become immeasurable at temperatures within the transition range or below.

The data for six orientations at 76 K are listed in Table 7 and bracketed in Figure 9. At

this temperature any effects due to orientation are indistinguishable from the scatter of

replicate tests, although it may be significant to note that the RL orientation yielded two

of the lowest K values. Neglecting any minor effects, the nineteen results at 7*6 K can be

3/2 statistically described: the mean value of K at this temperature is 60.8 MN/m , the K

3/2 values range between 54.6 and 65.9 MN/m , and the standard deviation is 3.0.

Other results in Table 8 indicate that stress-relief heat treatments such as applied to welded structures (533 or 811 K, 50h, air cool) do not lower the base metal fracture

toughness. However, prestraining at room temperature can reduce the toughness of Ti-6A1-4V

by as much as 50%, while virtually eliminating the temperature dependence of over the

ambient-to-cryogenic range. This point is emphasized in Figure 10 and clarified in discussion.

Electron beam welding also lowers the fracture toughness of this T1-6A1-4V alloy. The

weld data are listed in Table 9, and the average and K^

c values at 295 and 76 K are compared in Figure 11. The fusion and heat-affected zones at 76 K are nearly equivalent, with

3/2 averaging 4 to 6% lower than the 60.8 MN/m value for the base metal. However, a zone

of minimum toughness is located at the heat-affected-zone/fusion-zone boundary; there, K

3/2 for three specimens at 76 K ranged from 47.5 to 53.1 MN/m ', averaging 16% lower than the

base metal.

7
3.2 Fatigue Crack Growth Rates

Fatigue crack growth rates for the base metal at 295, 195, 76 and 4 K are shown in

Figures 12 and 13. These data were obtained from specimens of the LR and ML orientations, as

noted. Additional data for RL, LM, MR and RM orientations were obtained at 76 K during

precracking of the K test specimens of Table 8. These additional data, obtained at AK values

3/2 between 22 and 32 MN/m , fall within the scatterbands of Figures 12 and 13. A comparison

of all results shows no measurable specimen orientation or temperature effects over ;he

range of AK investigated.

Fatigue crack growth in electron beam welds was investigated at two temperatures. As indicated in Figures 14 and 15, cracks were propagated in the fusion, heat-affected, and

boundary zones at 295 and 76 K. Again, it appears that any temperature effects are insignificant compared to the scatter observed in replicate tests.

The scatterbands of Figures 12-15 are superimposed for comparison in Figure 16. The

unshaded band represents the base metal results of Figures 12 and 13 combined. The bands for

weld data are nearly equivalent in width, and nearly coincide with thr base metal band.

Thus, despite the microstructural transformations that are introduced, electron beam welding

has little effect on the fatigue crack propagation resistance of this alloy. The fatigue

crack growth rates of T1-6A1-4V are virtually structure insensitive for the AK ranges

investigated.

The linear tread of data on logarithmic coordinates is in accord with the equation

proposed by Paris and Erdogan [15], namely:

§ = C(AK)

n

. (8)

Here C and n are empirical constants which can be determined graphically. A line

approximation conforming to this equation and favoring the upper bound of crack growth rates

is shown in Figure 16. The value of C, the intercept on the ordinate axis at AK = 1, is

_9

2.7 x 10 mm/cycle; the slope, n, is 3.7. Of course, other linear approximations could

be chosen for specific applications of these data.

3.3 Static Load Cracking (SLC)

Several aspects of static load cracking in the base metal are illustrated by the test record and fracture surface of i igure 17. The deflection at constant load reflects timedependent crack extension which is most pronounced at the center of the specimen thickness

(plane strain conditions). Crack extension at the specimen edges (plane stress conditions) is

retarded. Judging from the regular increase of deflection, cracking in the base metal began

without delay; i.e., no incubation periods were observed in tests at 295, 195, or 130 K. In two tests at 295 K, cracking led to complete fracture when the initial load was maintained

indefinitely. The phenomenon occurs at stress intensity factors well below the "critical"

levels associated with P^ or the maximum loads of E-399-74 fracture test records.

The base metal cracking rates are highly sensitive to the applied stress intensity

factor. The rates can be described as power law functions of K, having the form:
(9)

Here, B and m, empirical constants that depend on temperature, were graphically determined

from Figure 18, using SI units. At room temperature,

|| = 1.9 x 10"56

(K) 27

(10)

and, at 195 K,

|| = 4.3 x lCf

72

(K) 36

(11)

As indicated in Figure 18, the rates at 195 K exhibit less scatter and are faster than

those at room temperature. To demonstrate that these effects are not due to stress corrosion,

the chemistry of the environment was alternated. The cracking rates for two specimens

submerged in ethanol at 295 K were not higher than those in room air. Similarly, the data at 195 K are in excellent agreement, regardless of whether the environment was nitrogen vapor or

a mixture of alcohol and dry ice. Therefore, the environments used here are believed to be

chemically inert.

Although the data at 130 K are few, they clearly indicate a decrease in static load

cracking rates between 195 K and 130 K. Thus, the rates do not vary monotonically with temperature, but peak at some temperature between 295 K and 130 K. One possible explanation for

this trend is a change in the micromode of cracking, as suggested in the discussion.

Three room temperature specimens tested at constant stress intensity factors ranging from

3/2

52 to 60 MN/m showed no evidence of cracking after hold times of 116h. Accordingly, 60

3/2 MN/m can be taken as the apparent room temperature value of K ^, the threshold stress intensity factor below which static load cracking was not observed in 116h tests. Since all

of the base metal SLC specimens were of the ML orientation, the ratio of K^/K at room temperature is estimated to be 0.7.

Note that the static load cracking rates are less sensitive to temperature variations

at lower K. If the linear trends of Figure 18 are extrapolated, they appear to converge at % 3/2 K = 63 MN/m . This prompts a speculation that K^,^ may not vary greatly with temperature,

but additional experiments would be necessary to establish K values with confidence.

Static load tests were also performed on specimens submerged in liquid nitrogen at 76 K. Three base metal specimens were loaded for 10 to 16h at K/KT ratios from 90 to 96%. As

Ic verified by heat-tinting, crack growth did not occur under these conditions. The absence of

cracking at 76 K may be viewed as a continuation of the declining trend observed below

195 K. Such results are consistent with the hypothesis that static load cracking is a thermally activated process which can be suppressed at sufficiently low temperatures.

9
Fusion zone and heat-affected-zone boundary specimens from electron beam welds were

subjected to static load tests at 76 K and 295 K. As noted earlier, the stock dimensions

dictated the use of specimens smaller than the 25.4 mm thickness favored for the base metal,

and a 10 mm weld specimen thickness was selected.

At 76 K, duplicate specimens from the weld zones were sustained for up to 16h at K/K^

ratios between 90 and 98%. The results showed no evidence of static load cracking at this

temperature.

The weld specimens did exhibit cracking at room temperature, but the behavior differed

from that of the base metal. Table 10 summarizes the results. For example, an unknown

amount of stable crack extension occurred during loading to the initial K value. A crack extension of 0.07 mm was measured after heat-tinting one fusion zone specimen which had

3/2 been subjected to a K value of 64 MN/m and promptly unloaded. Thus, not all of the crack

extension in static tests could be attributed to static load cracking. Also, crack arrest occurred in tests of the fusion zone, while fast fracture occurred at the HAZ/FZ boundary.

For these reasons, meaningful crack growth rates could not be calculated for weld specimens.

The heat-affected zone boundary seems critical with respect to static load cracking

resistance. One room temperature HAZB specimen fractured at an initial stress intensity factor

3/2

of 60.4 MN/m , a level which had proved innocuous to base metal and fusion zone specimens.

Presumably, the greater susceptibility to static load cracking at the heat-affected zone boundary relates to the fact that the and values are lowest there as well. Unfortunately,

the scatter in results for HAZB specimens at room temperature is considerable. Good estimates

of require hold times greater than lOOh.

Finally, there is evidence that an incubation period precedes cracking in the HAZ/FZ

3 boundary specimens. A delay of 10 s was noted in one test during which there was no sign of

any crack growth; subsequently, an increase of clip gage deflection signaled the onset of

cracking and fracture followed within an additional 20 s. Another specimen showed no timedependent crack extension at all, but fractured spontaneously after 20 s at a stress intensity

3/2 factor of 69.9 MN/m . These observations suggest that different mechanisms may govern static

load cracking in base and weld metal specimens.

4. SUMMARY

The fracture behavior of a recrystallization annealed and furnace cooled extra-lowinterstitial Ti-6A1-4V alloy and its electron beam welds is summarized as follows:

1) Fracture toughness : At temperatures between 4 and 295 K, K for this alloy

3/2

ranges from 54 to 105 MN/m . An abrupt transition in K values occurs between

76 to 125 K. There is a slight degree of fracture anisotropy, but only at temperatures above the transition region. Electron beam welding leads to a 16%

decrease in K^.^ at 76 K, the heat-affected-zone/fusion-zone boundary being the

region of minimum fracture toughness in an electron beam weld.

10
2) Fatigue crack growth rates : Within the scatter of replicate tests, and over the IK ranges

investigated, the fatigue crack propagation resistance of this alloy is insensitive to specimen orientation, temperature, and microstructural transf crmations due to welding. A single equation favoring the upper bound of the fatigue crack growth rates (mm/ cycle)

3/2 as a function of AK (MN/m ) is:

% = 2.7 x 10"9 (AK)

3 " 7 (12) dN

3) Static lead era-king : Precracked base and weld metal specimens sustained at stress intensity factors approaching did not exhibit static load cracking in liquid

nitrogen environments at 76 K. However, static load cracking did occur in inert

environments at temperatures between 130 and 295 K. Cracking in the base metal begins

3/2 _

without delay a- K values greater than 6Z M' - ~. .he HAT. rl 5:e:it=:.= exhibit an incubation period prior to cracking, suggesting that a different mechanism is operative.

The degradation in fracture resistance due to the possibility of fracture under static

loading can amount to 30%.

5. DISCUSSIOM

Campbell's review [16] of low temperature fracture behavior reveals few valid data for

titanium alleys, The existing literature dees net associate transitional fracture hehavier

with Ti-6A1-4V alloys, but asserts that decreases uniformly at low temperatures.

Certainly, the normal grades having relatively high carbon, oxygen, nitrogen, and hydrogen

3/2 contents exhibit low K_

c values (35 to 43 MN/m ) and mild temperature dependences [12,17]. But why were transitions for extra-low-interstitial TI-6A1-4V alloys not previously

reported? Among other factors, the interval 76 K < T < 125, was never previously investigated,

and invalid results predating current testing standards sometimes obscured the true trend of K

Ic [18-20].

Invalid fracture data must be cautiously interpreted, for it has been demonstrated

that Kq results are specimen size dependent, and of limited use for comparative or design

purposes. Some evidence for titanium alloys, including Ti-6A1-4V, indicates that Kq increases

with specimen dimensions [21-23] . The present results for the base metal at room temperature,

when compared directly with Annis' data [24], show that K increases from 67 to 83 to 93 3/2 ^ MH/m as compact specimen thickness (W/B = 2) increases from 12.5 to 20.3 to 25.4 mm. In view of this size dependence, the finding that room temperature values realistically

rank the fracture toughness of welds, as shown in Figure 11, may be fortuitous. Only K^

c data legitimately establish a ranking.

Although specimen, orientation effects were noted in the K__ data at temperatures above

the transition region, this alloy is nearly isotropic compared to most rolled or forged alloys.

Irr.si taring tne ecuiaxed nature of one retry srallized rase metal nier : structure as revealed in Figure 3, one appreciates that several orientations tested here were virtually equivalent.

11
Nevertheless, we can only assume that slight microstructural dissimilarities account for the

anisotropy noted in Figure 9. It is not inconsistent to report no measurable orientation

effects at low temperatures, in view of the implied change in fracture mode. Analogously,

fracture anisotropy in steels is most pronounced at temperatures above the transition

region [25].

The virtual independence of fatigue crack growth rates with respect to temperature is

supported by other work on Ti-6A1-4V alloys. The mill-annealed normal grades tested by

Fowlkes and Tobler [12] and Wei and Ritter [26] show no measurable differences in rates at temperatures ranging from 4 to 563 K. Pittinato's data [27] for an ELI grade tested in

gaseous helium also show little effect of temperature in the range 145 to 295 K. Evidently,

there is no correlation between K and da/dN; the former property varies sensitively with

temperature, while the latter does not.

Microstructural effects on fatigue crack growth in titanium alloys are poorly understood.

Yoder, Cooley, and Crooker [28] reported that a beta-annealed commercial purity T1-6A1-4V base

metal microstructure offers superior resistance, with rates that are up to 10 times slower

than the same material in the mill-annealed condition. Figure 16 demonstrates that the EB weld fusion zone and base metal fatigue crack growth rates are virtually identical, despite

the fact that these two zones contain aged martensitic alpha and equiaxed alpha microstructures

respectively. Similarly, Irving and Beevers [28] tested five microstructures of a T1-6A1-4V

alloy and concluded that the fatigue crack growth rates were structure-insensitive at AK

3/2 values greater than 12 MN/m . The present results are in harmony with Johnson and Paris' statement [30] that structure exerts only a secondary influence on da/dN. In any event,

microstructure and temperature effects may become quite significant at AK values below

3/2

12 MN/m , or at relatively high AK values approaching K [29,31]. Therefore, the statements

in this report regarding fatigue crack growth behavior apply strictly to the intermediate

range of AK investigated.

Static load cracking in titanium alloys at temperatures below the secondary creep

regime (i.e., at test temperature/melting-point ratios, T/T , of less than 0.4) is not completely understood [32-35]. A successful theory must rationalize several observations:

that the resistance to fracture under static loading is reduced by electron beam welding,

that different mechanisms are involved in base and weld metal specimens, and that

cracking rates at low temperatures (195 K) may be more severe than at room temperature.

At the 50 ppm level of the present alloy, interstitial hydrogen may facilitate cracking

by diffusing to and concentrating at regions of high stress [35-38]. Local fracturing is assisted by reductions of fracture surface energy and cohesive strength. It is also known

that primary creep occurs in Ti-6A1-4V alloys, at stress levels as low as 70% of the

conventional yield strength [39]. Thus, primary creep in the crack-tip plastic zone of a compact specimen at room temperature is expected. If we accept the theories that interstitial

hydrogen and creep effects are responsible for static load cracking in Ti-6A1-4V, it follows

from the present results at 130 K that these processes r must be operative at T/T ratios as m

low as 0.07.

12
If static load cracking is governed by thermally activated processes involving diffusion

or creep, one might expect crack growth rates to decrease with decreasing temperature.

Clearly, the base metal data at temperatures between 295 K and 195 K (Figure 18) do not follow

the expected trend.

Rationalizations for the increase in static crack growth rates between 295 and 195 K postulate a change in the micromode of fracture. Static load cracking in T1-6A1-4V at room temperature involves mixed modes of dimpled rupture and cleavage [35]. Possibly, an increased constraint to plasticity at 195 K favors cleavage and higher cracking rates. Varying tendencies to crack branching may also play a role. Such effects, superimposed on those of a thermally activated process, would explain an otherwise unpredicted temperature

dependence.

6. RECOMMENDATIONS

Gowan and Stein described the fracture of a Ti-6A1-4V pressure vessel which was destructively tested at room temperature [40] . There was a sizeable volumetric expansion

prior to fracture, and postmortem tests confirmed that the burst-tank material had been

strain-hardened [41]. Fracture tests of material taken from the deformed and fractured tank were shown in Table 6 and Figure 9. Clearly, the burst-tank material properties are not representative of the Ti-6A1-4V alloy in its annealed condition, and the data for prestrained

material are useless in fracture mechanics evaluations of existing structures.

Electron beam welding reduces the and values of the material in a region near the

macroscopically visible HAZ/FZ boundary. The minimum cycle-life of a welded structure of constant geometry and thickness would be calculated by assuming that a flaw exists at this

boundary zone. However, if thick weld lands are employed to reduce the stresses in welds,

the base metal properties may be critical. Thus, structural failures may not necessarily

initiate in the welds [40].

At temperatures of 76 K or lower, is a conservative design criterion, since static

load cracking is improbable. At room temperature, neither K^

c nor are conservative design

criteria since static load crack growth in flawed structures occurs at lower stress intensity

factors. To account for this possibility, fracture mechanics evaluations at room temperature

should employ K^^, rather than K^, as the critical stress intensity never to be exceeded in structures during service.

The reliability of a K^^ measurement increases at very long hold times [34]. In the

absence of rigorously determined values for the base metal, eqs. (10) and (11) can be used

to make conservative predictions of static load cracking resistance. Assuming these equations

hold at lower K values, the time required to produce measurable crack growth at a given

3/2 temperature and stress intensity factor level can be calculated. For K < 55 MN/m , the time

required for significant crack growth should exceed the service lifetimes of most Ti-6A1-4V

structures.

13
Since the attempt to develop static load crack growth rate expressions for weld specimens

failed, an alternative test procedure aimed at direct determination of is recommended.

The procedure described by Novak and Rolfe [42], whereby numerous specimens are bolt-loaded for

long periods, could be adopted for this purpose. Testing should concentrate on the

heat-affected zone boundary region where values are evidently lowest.

The possibility of optimizing the fracture resistance of this alloy should also be

considered. It proved feasible to enhance the utility of steels by shifting or eliminating

their transition temperatures by grain size, heat treatment, or compositional modifications.

Similar effects for T1-6A1-4V might be investigated with the intention of obtaining higher

K^^ values at temperatures below 125 K, while improving the resistance to static load cracking

at higher temperatures. Data for related titanium alloys suggest that an increased cooling

rate from the annealing temperature may influence these properties substantially [43,44]

7 . ACKNOWLEDGMENTS

This final report describes work sponsored by Kirtland Air Force Base, under Kirtland

Order AFWL 75-177. The assistance of H. Vail (Chem-Tronics , Inc.) in providing the electron

beam welds and related technical information is gratefully acknowledged. The assistance of

E. Ballinger and S. Naranjo in fracture testing and data reduction is also deeply appreciated.

14
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lb
40. J. G. Gowan and C. Stein, Metallographic analysis of fracture surfaces of the 20.7 Cu.

Ft. pressure vessel for the ALL cryogenic storage assembly, presented at the Titanium

Materials Properties Seminar, Kirtland AFB, Albuquerque, New Mexico (February, 1975).

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Materials Properties Seminar, Kirtland AFB, Albuquerque, New Mexico (February, 1975).

42. S. R. Novak and S. T. Rolfe, Modified WOL specimen for KT „„ environmental testing, J. Mater. JMLSA 4, 701-728 (1969)

43. R. H. Van Stone, J. R. Low, Jr., and J. L. Shannon, Jr., The effect of microstructure on

the fracture toughness of titanium alloys, NASA Tech. Rept. No. 2-Ti, Department of

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on fracture behavior of mill-annealed Ti-6A1-4V, presented at the Ninth National

Symposium on Fracture Mechanics, University of Pittsburgh (August, 1975).

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CM

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01o.

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cm in co

r»l H

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26
cowO

u

OJ

to I i s:

to10 0) O -C

S- 4-> u

4—

0) o —- a;Q. C oo O

i— -i-> +-> <_i A3 rt3 s- a. « E

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27
Figure 2. Notation for reference axes and specimen orientations.

28
29
30
Figure 5. Low temperature fatigue and fracture testing apparatus.

31
'H1H9NH1 >DV*D 3GVH3AV
<

o

60

40

DEFLECTION, 8, in

0.2 0.4 0.6

Ti-6AI-4V (ELI)

B = 25.4 mm

W/B= 2

T=295 K

a/W=.500

T = 76 K

a/W=.515

ML Orientation

15

10

-5 -1

0.5 1.0 1.5

DEFLECTION,B, mm

2.0

Figure 7. ASTM E 399-74 fracture test records for the base metal at

295 and 76 K.

33
u !/qi- u ! 'ivaoaiNi-r

oooooooo vO^CMOOO^O^CN

3

ui/ r >| 'ivaoaiNi-r

34
TEMPERATURE, °F

-453 -323 -100

100

CO

o

o

<

at

80

60

No. Tests ( )

T

32 7 0 —i r

Ti-6AI-4V (ELI)

LR Orientation,

B = 20.3 mm

ML Orientation,

B=25.4 mm

- 80

100

CO

CO

o

0*

60 ^

<

a:

100 200

TEMPERATURE, K

300

Figure 9. The temperature dependence of Kj c for the base metal, indicating

transitional behavior at 76 K < T < 125 K.

35
TEMPERATURE, K

Figure 10. The effect of a room temperature prestrain in biaxial tension

on the fracture toughness of the base metal

36
BM HAZ HAZB FZ

Figure 11. The average fracture toughness as a function of crack location

in the "M-6A1-4V base metal and its electron beam welds.

37
STRESS INTENSITY FACTOR, AK, Icsiin"2

5 10 50

V

E

E

•» Z

IQS

<

o

o

u

<

cm

u

10

5

III! 1 1—I I I I

Ti-6AI-4V (ELI)

Base Metal

4 K

76 K

200 K

295 K

o

A

o o

„

Oo

Of

OO"

#

i

O

.o

-10

J_L J I

I I I I I

1©'

w

10

<

5

o

o

<

a*

u

<

5 10 50 100

STRESS INTENSITY FACTOR, AK, MN/m 3/2

Figure 12. Fatigue crack growth rates for base metal specimens of the

LR orientation [11].

38
AK, ksi.ir/2

10 50

-2

10

id

3

io

4

Ti_6Al-4V(ELI)

Base Metal

O 76K

A 195K

• 295K

10

-3

10

-4

10

-5

10

100

v2

w

>«

uC

0

<

o

o

<

LUD

o

STRESS INTENSITY RANGE, AK, MN/m

13. Fatigue crack growth rates for base metal specimens, primarily

of the ML orientation.

39
AK, k si* i

5 10 50

| ! I I I

|

1

1 I |

I I I II

5 10 50 100

STRESS INTENSITY R ANGE, AK, MN/m^

Figure 14. Fatigue crack growth rates for electron beam weld fusion

zone specimens of the MR orientation.

40
Figure 15. Fatigue crack growth rates for heat-affected zone specimens

of the MR orientation with cracks sited near the center (HAZ)

and boundary ( HAZ B )

41
2 10 100

STRESS INTENSITY RANGE, AK, MN/mTM

Figure 16. A comparison of fatigue crack growth rate scatterbands for

base metal and weld zones.

42
Q

<

o

DEFLECTION, 5, in

0.2 0.4 0.6

60

T

1 r

Hold Time,At=37 s T=195 K, a/W=0.55

- 10

C9

o

<

5 2

0.5 1.0 1.5

DEFLECTION, 8, mm

2.0

Figure 17. Static load fracture behavior of the base metal, showing the

compact specimen fracture surface.

43
MSE

i

o

<

5

O

o

<

u

10

STRESS INTENSITY FACTOR, K, ksi.in 1 '2

20 50 90

10

-2

10

-3

10r4

10

,-5

10

10

1 III

Ti-6 AI-4V (ELI)

i i i

T=295 K

air

alcohol X

T=195 K

o N2 • alcohol, dry ice t

T=130K

a N2 Id

Llf

no growth

fill

- 10

- 10

:5

- 10

10

ltf

20 50 100

STRESS INTENSITY FACTOR, K, MN/m3/2

MSc

<

<

5

O

O

<

u

Figure 18. Static load cracking rates for the base metal at room and

cryogenic temperatures.

44
NBS-114A (REV. 7-73)

U.S. DEPT. OF COMM. BIBLIOGRAPHIC DATA

SHEET

1. PUBLICATION OR REPORT NO. NBSIR 76-8362. Gov't Accession

No.

3. Recipient's Accession No.

4. TITLE AND SUBTITLE

Low Temperature Fracture Behavior of a Ti-6A1-4V Alloy and its

Electron Beam Welds

5. Publication Date Anril 1976

6. Performing Organization (-ode 275.03

7. AUTHOR(S)

R.I. Tnhler

8. Performing Organ. Report No.

9. PERFORMING ORGANIZATION NAME AND ADDRESS

NATIONAL BUREAU OF STANDARDS

DEPARTMENT OF COMMERCE

WASHINGTON, D.C. 20234

10. Project/Task/Work Unit No. 2750431

11. Contract/Grant No.

AFWL/75-P0-177

12. Sponsoring Organization Name and Complete Address (Street, City, State, ZIP)

Kirtland AFB

Albuquerque, NM 87117

13. Type of Report & Period Covered

FY 1975

14. Sponsoring Agency Code

15. SUPPLEMENTARY NOTES

16. ABSTRACT (A 200-word or less (actual summary of most significant information. If document includes a significant

bibliography or literature survey, mention it here.) The effects of electron beam (EB) welding on the fracture behavior of a recrystall izatioi

annealed, extra-low-interstitial Ti-6A1-4V alloy have been investigated at temperatures

in the ambient-to-cryogenic range. Plane strain fracture toughness (Ki c

) and subcritical

crack growth parameters were measured using compact specimens 10 to 25.4 mm-thick. These

parameters can be used to predict the safe operating lifetimes of cryogenic pressure vessels and other welded Ti-6A1-4V structures.

Although EB welding transforms the base metal microstructure extensively, its effects

j on the material's fatigue crack propagation resistance at intermediate stress intensity factors are negligible. The growth rates, da/dN, of fatigue cracks sited in the

fusion and heat-affected zones of weldments were temperature insensitive and nearly

equivalent to rates for the base metal. However, welding introduces a zone of low

fracture toughness at the heat-affected-zone/fusion-zone boundary. The Kj c value for

this boundary zone at liquid nitrogen temperature (76 K) was 45 MN/nr 's 16% lower than

the base metal. The base metal fracture toughness increases between 4 and 295 K, with

an abrupt transition to higher Kt

c values occurring at temperatures between 76 and 125

K. Static load cracking, temperature effects, and specimen orientation effects on the

fracture behavior of this titanium alloy are central topics of discussion.

17. KEY WORDS (six to twelve entries; alphabetical order; capitalize only the first letter of the first key word unless a proper name; separated by semicolons)

Electron beam welding; fatigue; fracture toughness; low temperature tests; mechanical

properties; titanium alloy.

18. AVAILABILITY ? Unlimited

T

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