NASA Contractor Report 185219

The Effect of Hydrogen and Microstructure on

the Deformation and Fracture Behavior of a

Single Crystal Nickel-Base Superalloy

William S. Walston

Carnegie-Mellon University

Pittsburgh, Pennsylvania

April 1990

Prepared for

Lewis Research Center

Under Grant NAG3-463

National Aeronautics and

Space Administration

(NASA-Cr_- 1G521 9)

MTC_,J%TRUE Tt.IRIE ON

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NTCKtI.-_,A¢_ SUP_ALLOY final Report

Thesis (Carne_i e-Mel Ion Univ-) ?54

TH_ EFFcCT _F HYDROGLN AND

THE DFFORMATION AND

_F A SINGL_ CRYSIAL

Ph.D.

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ABSTRACT

A study has been conducted on the effects of internal hydrogen and microstructure

on the deformation and fracture of a single crystal nickel-base superaUoy. In particular,

room temperature plane strain fracture toughness and tensile tests were performed on

hydrogen-free and hydrogen charged samples of PWA 1480. The role of microstructure

was incorporated by varying the levels of porosity and eutectic T/q( through hot isostatic

pressing and heat treatment.

The room temperature behavior of PWA 1480 was unusual because precipitate

shearing was not the primary deformation mechanism at all strains. At strains over 1%,

dislocations were trapped in the y matrix, and an attempt was made to relate this behavior to

compositional differences between PWA 1480 and other superalloys. Another unique

feature of the tensile behavior was cleavage of the eutectic _,/'¢, which is believed to initiate

the failure process. Fracture occurred on {111 } planes and is likely a result of shear

localization along these planes. Elimination of the eutecfic T/'_ greatly improved the tensile

ductility, but porosity had no effect on tensile properties.

Large quantities of hydrogen (1.74 at.%) were gas-phase charged into the material,

but surprisingly this was not a function of the amount of porosity or eutectic T/T' present.

Desorption experiments suggest that the vast majority of hydrogen is at reversible lattice

trapping sites. This large, uniform concentration of hydrogen dramatically reduced the

tensile strain to failure, but only slightly affected the reduction in area. Further evidence of

strain localization due to hydrogen was observed in the deformation structure. Available

hydrogen embrittlement models were examined in light of these results and it was found

that the hydrogen enhanced localized plasticity model can explain much of the tensile

behavior.

KIC fracture toughness tests were conducted, but it was necessary to also perform

Jlc tests to provide valid data. Unexpectedly, the values from these tests did not agree;

further, the crack growth behavior and fractography differed greatly. The eutectic y/y'

behaved differently in these two types of tests, and the possible effect of the unloading

present in JIC tests on this constituent was examined.
ACKNOWLEDGEMENTS

Theresearchdescribed herein was made possible by the financial support of NASA

Lewis Research Center. The support and assistance of project monitor Dr. R.L.

Dreschfield is gratefully acknowledged. I would like to thank my advisors, Drs. I.M.

Bernstein, A.W. Thompson and J.C. Williams. _They made the changing of the guards as

easy as possible, and I am indebted to each of them for the guidance and encouragement

they have provided along the way.

Special thanks arc due to Dr. N.R. Moody of Sandia National Laboratories, who

graciously performed the hydrogen charging of all samples. His cooperation and

assistance, as well as many helpful discussions on hydrogen effects am greatly appreciated.

Thanks are also extended to Dr. J.P. Lucas of Sandia for performing the short rod fracture

toughness tests. Through the course of this work, there were many who provided insight

and assistance, including Dr. M. Dollar of Illinois Institute of Technology, Dr. J. Chen6 of

Universit6 Paris Sud, Dr. L.G. Fritzemeier of Rocketdyne and Dr. W.M. Garrison, Jr. of

Carnegie Mellon University.

Finally, I would like to thank my family for their encouragement throughout this

endeavor. My parents were wise enough to instill in me the desire to learn and the will to

accomplish my goals. Above all, I owe my deepest appreciation to my wife, Courtney,

who not only understood the rigors and crazy hours of graduate school, but also was

willing to use her skills in proofreading this document.

ii
TABLE OFCONTENTS

1 INTRODUCTION ........................................................................... 1

2 BACKGROUND..................... ....................................................... 5

2.1 Single Crystal Nickel-Base Superalloys...................................... 5

2.1.1 GeneralBackground.................................................. 5

2.1.1.1 Solidificationprocess.................................... 6

2.1.1.2 Heat treatment and hot isostatic pressing .............. 10

2.1.1.3 Composition effects ...................................... 14

2.1.1.4 Anisotropy ................................................. 15

2.1.2 Deformation Behavior ................................................ 17

2.1.2.1 General characteristics ................................... 18

2.1.2.2 Flow stress modeling .................................... 20

2.1.2.3 Tension/Compression asymmetry ...................... 27

2.1.3 Fracture Behavior ..................................................... 28

2.1.3.1 Crystallographic fracture ................................. 28

2.1.3.2 Relation to microstructure ............................... 31

2.1.4 Effect of Hydrogen on Superalloys ................................. 32

2.2 PWA 1480 ....................................................................... 36

2.2.1 Relevant Studies ....................................................... 37

2.2.2 Effect of Hydrogen ................................................... 39

2.3 Hydrogen Embrittlement Processes ........................................... 40

2.3.1 Trap Theory ............................................................ 41

2.3.1.1 Determination of the trap binding energy .............. 43

2.3.2 Hydrogen Enhanced Localized Plasticity ........................... 47

2.4 Fracture Toughness of Polycrystal and Single Crystal Superalloys ...... 49

2.4.1 Micromechanistic Fracture Initiation Models ...................... 51

2.5 Background Summary .......................................................... 53

3 EXPERIMENTAL PROCEDURES ...................................................... 55

3.1 Material Preparation ............................................................. 55

3.1.1 As-Received Material ................................................. 55

3.1.2 Laue X-ray Analysis .................................................. 55

3.1.3 Heat Treatment ........................................................ 56

3.1.4 Hot Isostatic Pressing ................................................ 59

3.1.5 Hydrogcn Charging, Analysis and Desorption .................... 60

o,°

111
4

3.2 Microstructural Characterization ............................................... 61

3.2.1 Metallography ......................................................... 61

3.2.2 Quantitative Metallography ........................................... 62

3.3 Mechanical Testing .............................................................. 65

3.3.1 Tensile Testing.. ...................................................... 65

3.3.2 Compression Testing ................................................. 66

3.3.3 Kxc Testing ............................................................ 67

3.3.4 JIc Testing ............................................................. 68

3.4 Deformation and Fractography Characterization ............................ 70

3.4.1 TEM Foil Preparation and TEM Procedure ........................ 70

3.4.2 Slip Band Analysis .................................................... 72

3.4.3 Sectioning Experiments .............................................. 73

3.4.4 Pre-Strain and Pre-Charge Tensile Tests ........................... 74

3.4.5 Quantitative Fractography ............................................ 74

3.4.6 Crystallographic Fracture Orientations ............................. 75

3.4.7 Fractographic Analysis Techniques ................................. 77

RESULTS .................................................................................... 80

4.1 Microstructural Characterization ............................................... 80

4.1.1 Alternate Heat Treatment ............................................. 84

4.1.2 Hot Isostatic Pressing ................................................ 85

4.2 Hydrogen Trapping Characteristics ........................................... 89

4.3 Mechanical Properties .......................................................... 93

4.3.1 Tensile Properties ..................................................... 93

4.3.1.1 Effect of hydrogen ........................................ 93

4.3.1.2 Effect of eutectic T/q( ..................................... 99

4.3.1.3 Effect of porosity ......................................... 101

4.3.2 Compression Properties .............................................. 103

4.3.3 Fracture Toughness ................................................... 103

4.3.3.1 Klc results ................................................. 103

4.3.3.2 JIc results .................................................. 104

4.4 Deformation Behavior in Tension and Compression ....................... 106

4.4.1 Development of Deformation Structure ............................. 106

4.4.1.1 Effect of hydrogen ........................................ 113

4.4.1.2 Effect of eutectic T_ ..................................... 118

4.4.2 Slip Character .......................................................... 119

4.4.3 Double Notch Tensiles and Sectioned Tensiles .................... 125

iv
4.4.4 Pre-Strain and Pre-Charge Tensile Tests ........................... 127

4.5 Tensile Fracture Behavior ...................................................... 128

4.5.1 General Fracture Characteristics ..................................... 128

4.5.2 Effect of Hydrogen ................................................... 135

4.5.3 Effect of Alternate Heat Treatment and HIP'ing ................... 137

4.6 Fracture Toughness ............................................................. 137

4.6.1 Fracture of Kit Samples ............................................. 137

4.6.2 Fracture of JIc Samples .............................................. 142

5 DISCUSSION ............................................................................... 149

5.1 Hydrogen Solubility and Trapping ............................................ 150

5.2 Tension and Compression Properties ......................................... 155

5.3 Deformation and Fracture Behavior without Hydrogen .................... 162

5.4 Deformation and Fracture Behavior with Hydrogen ........................ 170

5.5 Analysis of Strengthening ...................................................... 175

5.6 Fracture Toughness ............................. ................................ 177

6 CONCLUSIONS ............................................................................ 184

7 SUGGESTIONS FOR FUTURE RESEARCH ......................................... 189

8 REFERENCES .............................................................................. 192

APPENDICES

APPENDIX A

APPENDIX B

APPENDIX C

APPENDIX D

APPENDIX E

APPENDIX F

Metallography Data and Tensile Properties ........................... 202

Deformation and Fracture Behavior of < 130> Samples ....... 207

Short Rod KIC Testing ................................................. 213

Hydrogen-induced Crack Growth ................................. 222

Tritium Autoradiography ............................................... 234

Quantitative Results of Reppich CRSS Model ....................... 244

CHAPTER 1

INTRODUCTION

Nickcl-basc superalloys must possess a favorable balance of high temperature

properties, such as creep strength, thermal fatigue resistance and oxidation resistancc, for

use in gas turbine and rocket engine applications. 1 It is possible to attain this wide army of

mechanical properties by adjusting the complex compositions of these alloys, which in turn

affect the characteristics of the microstructure produced. Nickel-base superaUoys consist of

a nickel-rich fcc y phase and an ordered fcc (Ll2-type) _/' phase based on Ni3AI. The _/'

phase precipitates from the 7 phase by an aging reaction, with modem superalloys usually

containing a high volume fraction (-60-70%) of cuboidal 'le precipitates. These 7'

precipitates are typically coherent with the 7 matrix having mismatches of less than 0.5%.

The _/' precipitate strength can more than double with temperature up to 700°C, 2 which

gives nickel-base superalloys their excellent high temperature strength.

The compositions of nickel-base superalloys generally are very complex, and many

of the properties are controlled by the partitioning of these elements to the "_or _' phase. 3

A1, Ti, W, and Ta preferentially segregate to the _/' phase, and thus are added to increase

the 7' volume fraction and provide solid solution strengthening. 4 Solid solution

strengthening of the y phase is accomplished by additions of Cr and Mo, as well as W and

A1. 3 Cr and A1 also are very important in providing oxidation and corrosion resistance. 1

In order to obtain a finite difference between the 3/solvus temperature and the incipient

melting temperatures, elements such as Co are added. 5 There are many other elements that

are typically added to nickel-base superalloys, and the reader is referred to several reviews

of compositional effects for further details. 1.3,4

It is important in achieving maximum engine performance to operate at the highest

possible temperature. Thus, increasing the working temperature of superalloys has been a
major driving force behindsuperalloy design and development. Initially, compositional

and microstructural variations provided higher temperature benefits, but increases during

the last 20 years have primarily stemmed from processing advances. Directional

solidification and single crystal casting techniques have pushed the operating temperature of

superalloys past 1100°C (about 85-90% of the melting temperature). 6 Pratt and Whitney

pioneered the development of directional solidification, which culminated in 1980 with the

production of single crystalline PWA 1480. 7 PWA 1480 and other single crystal

superalloys have demonstrated superior performance over conventionally cast superalloys

primarily because of their higher melting temperatures and _/' solvus temperatures. These

advantages stem from the elimination of grain boundary strengthening elements necessary

in polycrystals. 7

Based on the superior performance of single crystal superalloys under ambient

conditions, their potential use in hydrogen-fueled engines is under consideration, s Many

potential single crystal alloys have been screened using notch tensile tests and other

mechanical tests performed in hydrogen atmospheres designed to simulate service

conditions, s These tests provided valuable initial data, but did not completely evaluate the

alloy's resistance to hydrogen embritflement. To this end, NASA began a program at

Carnegie Mellon to fundamentally characterize the effect of hydrogen on single crystal

superalloys. PWA 1480 was one of the alloys chosen based on promising results of tests

conducted in hydrogen atmospheres.

This study concentrated primarily on room temperature properties since hydrogen

embrittlement is generally most severe near this temperature. As-received PWA 1480

contained substantial amounts of porosity and eutectic _'/T', and the effects of these were

systematically studied because of their possible role as hydrogen trapping sites. Hydrogen

was introduced into the material by gas-phase charging resulting in a uniform internal

concentration, followed by testing in air. Gas-phase charging was used, as opposed to

cathodic charging, 9"1° because it is easier to assess hydrogen's effects when a uniform
concentrationis present,incontrasto athinsurfacelayer. Therehavebeenonly twoother

investigationsof the effect of a uniform concentrationof hydrogen on singlecrystal

superalloys,9-11and both of thesewere performedon CMSX-2. Therefore, it was

importantto f'trstdeterminetheroleof hydrogentrappingin this typeof alloy bymeansof

tritium autoradiographyandhydrogendesorptionexperiments. To studythe effect of

internalhydrogenandmicrostructure on the deformation and fracture behavior, series of

tensile tests were performed, followed by TEM studies and the use of a variety of

fractographic examination techniques.

The fracture toughness of these types of alloys has never been studied before,

although it could be an important design criterion should a crack develop in service. Thus,

the crack growth behavior and fracture toughness were studied on compact tension

specimens as a function of hydrogen and microstructure. In a parallel study covered in

Appendix C, Dr. J.P. Lucas of Sandia National Labs performed short rod fracture

toughness tests using fractured halves of full size toughness samples that had been

previously tested at Carnegie Mellon. This study was done to expand the knowledge and

characterization of this novel testing technique to a new class of alloys.

Another important situation that could be encountered in service is crack

propagation at values below Kic due to the presence of a hydrogen atmosphere. Again,

this area has not been studied on single crystal superalloys, and hydrogen-induced crack

growth experiments in a hydrogen gas atmosphere were performed at Sandia under the

direction of Dr. N.R. Moody. The background, description and results of these tests are

given in Appendix D.

In summary, the main goals of this work were:

1) To characterize the extent of hydrogen trapping and its effect on mechanical

properties.
2) To characterize the effect of hydrogen and microstructure on the tensile

deformation and fracture behavior. To relate the observed behavior to available

hydrogen embrittlement models.

3) To evaluate the crack growth behavior and fracture toughness as a function of

hydrogen and microstructure. To relate the fracture toughness to

microstructural crack initiation sites through micromechanistic models.

4) To evaluate the resistance to hydrogen-induced crack growth and to compare the

results to tests on internally charged specimens.

5) To compare all observations to similar work on polycrystal and single crystal

superalloys in an attempt to make generalizations concerning the effect of

hydrogen and microstructure on this class of alloys.
Chapter 2

BACKGROUND

2.1 Single Crystal Nickel-Base Superalloys

2.1.1 General Background

The technology that led to single crystal superalloys is primarily responsible for the

increase in the use temperature of superaUoys over the past twenty years. 12 Single crystal

processing research was pioneered by Pratt and Whitney in the 1970s, culminating in the

production of PWA 1480 in 1980. 7 Major advantages of single crystals are the increases

achieved in melting and solution heat treatment temperature, as well as in the 7' solvus

temperature. These increases stem from the elimination of the grain boundary

strengthening elements; boron, carbon, hafnium, and zirconium, which are necessary

additions in polycrystals to eliminate grain boundary failure. Elimination of these elements

also prevents borides and carbides from acting as fatigue initiation sites, which results in

improved single crystal fatigue properties. Single crystals also possess superior high

temperature creep strength and ductility compared to polycrystals and directionally

solidified superalloys. 13 However, single crystal superalloys are not perfect, and many of

their problems are a direct result of the crystal growth process. Among these are localized

recrystallization, segregation, incipient melting, anisotropy and porosity. Through

extensive research, these problems are now becoming understood and can be controlled, as

the next few sections discuss.

First, the single crystal solidification process will be covered, outlining the

formation of all phases and constituents. Local elemental segregation occurs during

solidification, and the discussed heat treatment attempts to reduce this segrcgation. The

5
heattreatmentalsomaximizestheq(precipitatesize and volume fraction, which control the

mechanical properties. Also as a result of the solidification process, porosity is formed,

and its elimination by Hot Isostatic Pressing (HIP'ing) will be discussed. Finally, the

effect of the complex composition and the inherent anisotropy on several physical and

mechanical properties wiU be covered.

2.1.1.1 Solidification process

The single crystal solidification process plays a key role in affecting subsequent

properties. Therefore, it is important to understand the origin of such deleterious

phenomena as interdendritic porosity, eutectic y/y' and segregation, to better assess their

roles in the performance of single crystal nickel-base superalloys.

A schematic of the generic single crystal solidification process is shown in Figure

2.1. The molten superalloy is poured into a ceramic mold, and dendritic solidification

begins at the the water cooled chill plate in the form of several grains of various

orientations. The mold temperature 7 is usually around 1500-1600°C, and the thermal

gradient is approximately 70°C/cm. The <001> orientation has been shown to have the

highest growth rate, and after a certain distance, the <001> grains overtake all others and

proceed into the helical grain selector. The most favorably oriented of these grains emerges

and f'tlls the mold. The resulting single crystal is usually within 10 ° of the <001>

orientation with dendrites aligned along the three orthogonal <001> directions.

Within the microstructure, the y phase has formed as well as any other constituents,

such as eutectic _,/q(, or mu and sigma phases. In PWA 1480, the only other constituent

besides _, and q( is the eutectic y-q(. Upon solidification, the y phase forms first, but as the

temperature drops, and there is still molten metal in the interdendritic region, the eutectic

T/q( is nucleated. It has been noted that there are two types of eutectic T/q( present in

superalloys: 15 The eutectic T/q( formed at higher temperatures, which is finer and lacelike
HEAT

I_DIATION

Figure 2.1 Schematic of the single crystal solidification process.14

in appearance, and that formed at lower temperatures, which appears more uniform and

consists of mostly _/'. Yaoxiao, et al., 15 also divide the eutectic T/_/' into three distinct

regions: nucleation, core and cap. The appearance of each kind of eutectic T/_f in PWA

1480 will be discussed later, as will the significance of each regime.

There also is porosity present in the interdendritic region along with the eutectic

T/_'. Porosity is found in cast superalloys due to the molar volume decrease that occurs

upon solidification. Campbell 16 was the fLrst to treat this problem in depth and proposed

an equation to describe pore nucleation:

Pg + Ps = P* (2-1)

where Pg is the gas pressure in the liquid, Ps is the shrinkage pressure and P is the

pressure to nucleate a pore. Pg will be very small in superalloys due to the low solubility

of gases in nickel. 17 This analysis has been extended 18 to include variables that may either

7
be controlled during the single crystal solidification process or changed through

compositional variations. One variable that relates to this thesis is the thermal gradient, and

it has been observed that the amount of porosity decreases as the gradient is increased, is,19

It also was found that increasing the amounts of AI, Ti and Co increase porosity, while

higher amounts of Cr decrease porosity. 16

Campbell 16 stated that for Fe and A1 alloys, P* is very high and cannot be attained

by homogeneous nucleation. Therefore, nuclei will most likely be needed to facilitate

nucleation, and as an example, he suggested inclusions containing poorly-wetting solid and

liquid phases. This suggestion of heterogeneous nucleation of pores is of particular interest

for PWA 1480 due to the relationship found to exist between pores and eutectic _//_/', as

discussed later.

Thermal stability of the pores was studied by Anton and Giamei 17 in a single crystal

nickel-base superalloy. They found that even a small amount of time at the solution

temperature would dramatically increase the volume fraction and slightly increase the size

of the as-cast porosity. They explained these increases by the Kirkendall mechanism of

pore growth, which relies on compositional gradients and subsequent associated vacancy

diffusion. The increase in growth rate and volume fraction would be expected to level off

after a certain time because of the competition between KirkendaU-induced migration and

loss of vacancies to the surface. In fact, it was found that the increase in size and volume

fraction was only significant at times up to 1.2 hour, and at times longer than 4 hours the

pores began to shrink and disappear. 17 This is an important result for this study because

time spent at the solution temperature was used as a variable.

Besides porosity, the solidification process also leads to macroscopic and

microscopic segregation of elements. On a macroscopic scale, there are composition

gradients associated with the dendritic macrostructure, which are determined by the

partition coefficients of each alloy. In a study on CMSX-2, 2° a similar single crystal

superalloy, it was found that Ta and W segregated heavily to the dendrite core and arms,
while Co andCr showeda slight increasein this regionover theaveragecomposition.

Therewasvery little changein thecompositionprofiles of A1or Ti. Elementsalsowill

partition betweenthe_,or7' phase. Table 2.1 showsthe resultsof anatommicroprobe

analysisonPWA 1480.21A1,Ta,Ti andW primarily segregateto the7'phase,while Cr

andCosegregateto the_,phase.

Table2.1 Microprobeanalysisof thecomposition(at%)of the7and7' phasesin

PWA 1480.

A1 Ti Cr Ni

ymatrix 2.79 0.17 34.03 49.26

7' precipitate 15.4 2.85 2.17 70.10

Co Ta W

11.01 0.56 2.15

3.10 4.59 1.68

An unusual result was observed in the y' phase of CMSX-2. 21 A local

concentration gradient was observed in the 7' near the y/7' interface. This most likely

occurred due to the growth during the second aging heat treatment of the 7' particles that

were nucleated in the first aging treatment. The two stage aging heat treatment will be

explained in the next section. This local concentration gradient is quite narrow in

comparison to the total 7' size, but it may prove important in terms of dislocation behavior

in as much as a changing composition can affect the APB energy and therefore the stress to

move dislocations in this region. These results for CMSX-2 should be quite similar to

those that would be found in PWA 1480 due to the similar compositions, microstructures

mad two-stage aging heat treatment.
2.1.1.2 Heattreatmentandhotisostaficpressing

Theheattreatmentof singlecrystalnickel-basesuperalloysinvolvesasolutionheat

treatmentstagefollowed by anagingtreatmentandis similar to otherprecipitationheat

treatmentsin alloys suchasA1-Cu. The solution heat treatment in PWA 1480 is designed

to homogenize the segregation that occurred during solidification and to dissolve any

unwanted constituents, such as primary _/' or eutectic _//q(.22 The solution heat treatment

temperature is in the equilibrium ), phase field; above the _/' solvus and below the incipient

melting temperature 0MT). Incipient melting is the local melting of any low melting

temperature alloy constituents. This is a serious problem in many single crystals because

the heat treatment window (difference between the y solvus and the IM'D is normally very

small. Careful control and ramping of the temperature is thus essential to avoid incipient

melting.

Figure 2.2 shows a normal solution heat treatment cycle with the resulting incipient

melting temperature (IMT). The IMT increases in this figure due to the diffusion and

homogenization of any low melting temperature alloy constituents such as the eutectic _'/'l'.

This heat treatment cycle characterizes the type used in this work. It can be seen from this

figure that any slight temperature fluctuations could lead to incipient melting. It is estimated

that in the as-cast condition, PWA 1480 has an incipient melting point of 1290°C-12920C

when this type of heat treatment cycle is followed. _ This is only slightly above the

solution heat treatment temperature of 1288°C used in this study, so incipient melting was a

real concern.

10
Temp.

T'lme

Figure 2.2 Typical solution heat treatment cycle for PWA 1480 and the corresponding

incipient melting temperature 0MT).

The solution heat treatment stage is always followed by rapid cooling to avoid

unwanted precipitation of the _/' phase. A two-stage aging heat treatment follows with the

nucleation of the "1(phase controlled in the f'trst aging heat treatment stage. There also is

some growth of the "y' precipitates during the f'trst aging treatment, while in the second

aging heat treatment the growth process continues until the desired volume fraction and size

of y has precipitated. The second aging treatment was originally designed for diffusion

bonding of coatings to the alloy for service needs; however, this step is now commonly

used on uncoated samples to obtain better control of the q( precipitation process.

Pratt and Whitney has received a recent U.S. patent on a modified solution heat

treatment, which claims a more homogeneous material flee of incipient melting. 24 In fact,

it is further claimed that material which has undergone incipient melting can be healed by

this treatment with no loss in mechanical properties. _ The novel approach calls for the

IMT to actually be exceeded. Figure 2.3 shows the proposed solution heat treatment cycle.

The key is that sufficient time must be spent at the f'mal temperature of 1302°C to allow the

damage of incipient melting to heal by substantial diffusion. As it is held at this

temperature, the IMT is continuously increasing so that the material is soon below the IMT.

The purposeful exceeding of the IMT and the subsequent healing guarantee that the

11
Temp.

IMT

T'ui_

v

Figure 2.3 Modified heat treatment for PWA 1480 showing how IMT is exceeded.

resultant microstructure will be more compositionally homogeneous. This technique has

recently been successfully demonstrated on PWA 1480 by Rocketdyne. 25 Unfortunately,

this modified solution heat treatment could not be used in this work, because its existence

was not known until the majority of samples had already been heat treated. It is reported

here simply for completeness, as well as to contribute to the understanding of the

phenomenon of incipient melting.

The removal of porosity through Hot Isostatic Pressing (HIP'ing) will now be

discussed. It is well known that porosity resulting from solidification shrinkage in

superalloy castings can serve as crack initiation sites, reducing the useful life of the alloy

under conditions of fatigue or creep. HIP'ing has been used for many years on polycrystal

superaUoys to eliminate porosity and improve mechanical properties. 2628 Studies also

have been carried out on single crystals that show improved properties after

HIP'ing. 19,28,29 While there are associated drawbacks to HIP'ing, including altered

microstructures and unwanted carbides, 28'3° modifications to the HIP'ing cycle and postHIP heat treatments can usually overcome these. With single crystals, special care is

needed to avoid recrystallization, particularly as the absence of grain boundary

strengthening elements would result in severe degradation of mechanical performance.

12
It has been observed that when conventional HIP'ing cycles were employed with

single crystals, local recrystallization usually occurred, 31'32 believed due to the deformation

associated with closing the pores. Conventional HIP'ing cycles have been modified to

avoid such recrystallization. 32'33 A conventional cycle would involve increasing the

pressure and temperature up to desired levels and then holding for a period of time. A less

severe pressure and temperature versus time schedule was needed for single crystals, and

Eridon and Dalai 32 have developed a cycle at Howmet Corp. that involves a two-step

increase in pressure. During heating to a value near the normal solution temperature,

pressure is f'_rst applied below the level needed to close the pores. The pressure is provided

by a gaseous atmosphere, usually argon. Once the desired temperature is reached, it is held

for a period to assure a uniform temperature throughout the sample. Finally, the pressure

is increased at a moderate rate up to the pressure necessary to close the pores. The samples

are left under these conditions for a period of time to allow diffusional processes to occur.

The authors have outlined the specific conditions needed for PWA 1480 as foUows: 32

1) Heat to 1288°C under pressure of 34.5 MPa.

2) Hold for 1-2 hours.

3) Increase pressure at a rate of 0,4-0.55 MPa/min. to 103.5 MPa.

4) Hold for 4 hours.

Rocketdyne 33 has developed a similar cycle, which has been used on the samples in

this work. Unfortunately the details of that cycle currently are restricted by a U.S. Patent

Secrecy Order, although it is likely similar to the one cited above. It is claimed for PWA

1480 that this HIP'ing cycle completely removes the casting porosity, 33 while avoiding

recrystallization. Furthermore, the microstructure is said to be better homogenized due to

the extended time at the solution temperature during HIP'ing. The same study showed no

effect of HIP'ing on the tensile properties but an increased suess-rupture life and improved

fatigue properties. 33 Crack initiation now occurred at surface defects, eutectic )'/7' or

isolated small carbides instead of subsurface pores.

13
2.1.1.3 Composition effects

The occurrence of compositional segregation is important due to the effect that

composition has on many physical and mechanical properties. In terms of this thesis, it is

most relevant to know the effect of certain elements on the solid solution strengthening of _,

and _/', the stacking fault energy, and the antiphase boundary energy. These quantities are

known for some alloys of simpler compositions, and as a rough estimate it may be inferred

from the compositional differences of PWA 1480 what these values may be, or at least the

trend that should be followed.

PWA 1480 is an advanced superalloy with some significant compositional

differences from most other superalloys. Of course, grain boundary strengthening

elements such as C, B, Zr and Hf are not present. A key characteristic is the large amount

of Ta added to give high creep strength and oxidation resistance. 7 Cr and AI levels also are

relatively high to ensure oxidation and corrosion resistance. The W level is lower than in

many alloys, as is the Co level. It also should be mentioned that Mo and V are not present,

particularly as comparisons will be made in the thesis to CMSX-2, which contains a slight

amount of Mo, a lower Ta level and a higher W level with all other elemental levels being

similar.

In these single crystal high volume fraction _/' superalloys, strengthening below

60% of the melting temperature is due to overcoming the Anti-phase Boundary (APB)

energy, the presence of a high volume fraction of particles and solid solution strengthening.

The largest contribution is the APB energy, 2 and thus knowledge of compositional effects

on this quantity is important. It also is important to determine the effect of composition on

the Stacking Fault Energy (SFE), since dislocations splitting in the _,matrix would have a

profound effect on the deformation behavior. A low SFE would make it harder to constrict

the partials and shear the "/precipitates. The information concerning compositional effects

on stacking fault energy (SFE) and antiphase boundary energy (APBE) are not very

14
conclusive. It is believedthatNi hasahighSFE,34buttheresultanteffecton thisvalueof

manyelementsactingin combinationis notknown. It hasbeenreportedthatTi strongly

increasesthe SFE in Ni3A1 alloys, while Cr and Co slightly increase it. 1 There is no

evidence on the effect of other alloy elements, and no available data on the SFE of PWA

1480 or similar alloys.

There has only been one detailed study of the effect of various elements on the

APBE of Ni3Al-based alloys. 35 Brown and Ham concluded that Cr increases and Ti

decreases the APBE. This is not especially relevant for PWA 1480 considering Cr strongly

segregates to the 7phase, and Ti is only present in the alloy in a small amount. Nguyen 36

reports, without reference, that Ta should increase the APBE, but Decker 3 states that Ti and

Ta should not significantly change the APB energy. Therefore, it is clear that the trend in

the SFE or APBE cannot be estimated a priori for PWA 1480.

In examining the effect of composition on solid solution strengthening, it is clear

that one of the main differences between PWA 1480 and other superalloys is the increased

level of Ta. It is well known that Ta is a strong solid solution hardener, and it has been

shown to strongly segregate to _.21 Other compositional differences are not as great, and

it is not thought they would play a significant role in altering solid solution strengthening in

comparison to similar alloys. Many studies have shown that small amounts of Ta

dramatically increase the yield stress of Ni3Al-based alloys. 4'37"39 Thus, it is probable that

PWA 1480 has a higher y yield stress than most other superalloys.

2.1.1.4 Anisotropy

Single crystal superalloys are very anisotropic, and therefore it is necessary to know

the value of physical constants as a function of orientation. Elastic constant data will be

important in subsequently analyzing the fracture toughness data and the dislocation

behavior in the 7 and y phases. To begin with, the Young's modulus, E, for cubic

15
materialscanbecalculated from: _

1/E = $11 - 2(Slt - $12 - 1/2 S4n)(1121_+ 12422+ 13112)2 (2-2)

where Sij are the elastic compliances and lij are the direction cosines. In many cases, the

elastic data are reported in the form of the stiffness, Cij, which can be convened into elastic

compliance. The shear modulus can be calculated from: al

1/G = 4/3(S]1 - $12) + 1/3 S_ on (111)[1 I0] (2-3)

Thus, the necessary elastic constant data can be obtained from elastic compliance

values. Many authors concerned with the elastic properties of superalloys have merely

estimated the superalloy constants from the results of Ni and Ni3A1.12,42 The purpose of

this section is to take advantage of the measured data from superalloys and determine the

best estimate for PWA 1480. While there have been relatively few actual measurements on

T/_/' alloys, 43n5 considerable data exists for both Ni and Ni3AI. 4°.46-n8 Perhaps the best

study on Ni3AI was done by Kayser and Stassis, 4s and their data are shown in Table 2.2

along with data on Ni. n° The data on Ni are much older and may be less reliable.

Wallow, et al., 43 performed an excellent study on a few different T/_ alloys

examining the effect of temperature, composition and volume fraction of ]/. They analyzed

Nimonic 105 with _/' volume fractions of 10% and 50%. From this data they extracted the

elastic constants of the T phase by assuming that Cij was linear with the volume fraction of

T'- The 7' elastic constants were calculated from an alloy made of the 7' composition.

These results are shown in Table 2.2 as 105_[and 105T'.

The data of Wallow, et al., compare favorably with other T/T' alloys of different

compositions, nn.n5 From this information, it may be concluded that composition only

affects the elastic constants slightly. The closeness of the T and _ values also suggests that

modulus hardening is not important and any modulus effects on dislocation behavior

should be negligible.

16
Table2.2 ElasticCompliancesof VariousAlloys (10"6MPa-I)

$11 $12 $44

Ni 7.99 -3.12 8.,44

Ni3M 9.52 -3.79 7.99

105T 7.86 -3.02 7.28

105_/' 8.03 -3.08 7.80

PWA 1480A 9.06 -3.59 8.13

PWA 1480B 7.98 -3.06 7.64

Nimonic 105: (at.%) Ni, 10AI. 18Co, 16Cr, 1.STi, 3Mo, plus minor elements

The elastic compliances for PWA 1480 have been calculated from a rule of mixtures

method using 70% _/'. The results also are shown in Table 2.2 as PWA 1480A and PWA

1480B, where A was calculated from Ni and Ni3A1, and B was calculated from 105y and

105"y'. The elastic compliances for A and B differ somewhat, but the difference in the

Young's modulus turns out to be fairly small. The two extremes of the Young's modulus,

<001> and <111> have been calculated:

A

El00 = 110 GPa

B

Eloo = 125 GPa

A

E111 = 300 GPa

B

E111 = 316GPa

All further calculations in this study, will be made using an average value from method A

and B above. Thus, in the [001] orientation, Young's modulus is 118 GPa and the shear

modulus is 54.5 GPa.

2.1.2 Defomlation Behavior

The deformation of single crystal nickel-base superalloys is mainly controlled by

the nature of the T/Y' microstructure. Most importantly, there is a coherent interface

between the _, matrix and the "y' precipitate with a typical misfit of less than 0.5%. This

allows for matrix dislocations to penetrate the interface and shear the ordered 7' precipitates

17
provided this is energetically favorable. Furthermore, since there is a large volume fraction

(>60%) of y' precipitates, dislocation bypass is not possible, except at high temperatures. 49

Therefore, at the temperatures of interest in this work, shearing of the q( precipitates by

dislocations should be the primary mode of deformation.

2.1.2.1 General characteristics

Since the _/' precipitate is an L12-type phase, it is necessary for pairs of dislocations

to travel through the phase so that the second dislocation may restore the order that the first

dislocation altered. Before the pair can enter the 'y' precipitate from the matrix, the

antiphase boundary (APB) energy of the _ precipitate must be provided by the applied

stress. "5° There also are other factors to consider that are examined in the next section. The

normal shearing process that occurs in superalloys involves a pair of a/2<110> dislocations

lying on a { 111 } plane separated by a distance determined by the APB energy. This type

of behavior has been observed in many single crystal superalloys at the low temperatures

(<400°C) used in this work. ll.Sz-s3 The only difference between deformation in single

crystals versus polycrystals would, of course, be due to grain boundaries. In a single slip

orientation, there would be no obstacles to dislocation motion in a single crystal.

However, in this study, [001] samples primarily were used, which undergo multiple slip in

tension. From simple Schmid factor calculations, it is apparent that there are multiple

systems that are equally stressed, which can quickly lead to dislocation-dislocation

interactions and "polycrystal" type behavior.

Although shearing of the _' precipitate by pairs of a/2<110> dislocations occurs in

most superalloys at low temperatures, it is not the only deformation mechanism operating.

Single dislocations, as well as loosely associated pairs, are free to move through the

narrow ), phase. 49 However, large amounts of deformation cannot be obtained in this

manner, and shearing of the )/precipitate invariably occurs as the deformation proceeds, lj

There are exccptions to this, and it has been observed in some superalloys that dislocations

18
arenot ableto easilypenetratethe7' precipitate,if at all.x1.54Caremust beexercisedin

acceptingthis conclusion,sinceCopleyandKearhaveshownthatevenwhendislocations

wereprimarily observedin they phase,massiveshearingof the7' hadstill occurred.2 This

wasrationalizedbyconsideringtherateof dislocationmovementin eachphase,andit was

concludedthatthedislocationsspentthemajority of timein the7phase.Therefore,it could

beexpectedto observedislocationsprimarily in the_'phaseevenin alloyswhereprecipitate

shearing was the main deformation mechanism. In other high volume fraction _,'

superalloys,shearingwasevidencedby long, straightscrewdislocationsextendingacross

the_/and7' Phases.11Thus,it cannotbeassumedthat sheafinghasnot occurredsimply

becausedislocationsarenotobservedin the7,phase.

It also is possible to find a variety of stacking faults in superalloysthat may

participate in the deformationprocess. Inside the '1e precipitate it is possible to have

superlattice intrinsic stacking faults (SISF's), superlattiee extrinsic stacking faults (SESF's)

and other complex faults (CF's). 55 It has been shown that the energy of SISF's and

SESF's should be lower than the APB energy in _//_{ superaUoys, 55 and the relative

appearance of stacking faults and APB's depends upon their relative energies in the system

of interest. 49.56 The most common faults observed in L12-type alloys and _/Pt¢ superalloys

are the SISF and APB faults.

briefly will be covered now.

dislocation as follows: 57

APB faults have previously been mentioned, and SISF's

An SISF can form by dissociation of an a<101> type

a[IlO] _ a/31211] + a/3[I2I] (2-4)

This is only one of a number of possible reactions, and many others have been reported. 58

60 In a 'y/y' alloy, SISF's can form by a reaction at the _,/y' interface 61 or by two

intersecting APB-type pairs in the 7, precipitate. 62 After formation, the SISF can play a

role in the deformation process. 57'63 It should be re-emphasized that while SISF's have

bccn reported in supera/loys, deformation under tensile-type strain rates is invariably

19
dominatedby APB-typepairs. 49.62 This is not to say that SISF's cannot play a role in and

affect the overall deformation behavior. The micromechanisms of the formation of each

type of fault will not be covered in further detail, and the reader is referred to reviews by

Pope and Ezz, 49 and Kear, et al.55

2.1.2.2 Flow stress modeling

Utilizing models developed to predict the critical resolved shear stress (CRSS) can

assist in interpretation of observed deformation behavior. Models of the yielding process

and (CRSS) in low volume fraction _/' alloys are well-developed 35.64-69 and are based

upon Fleischer-Friedel 7° statistics, which deal with dislocations interacting with a dilute

array of point obstacles. Extended models to treat high volume fraction _/' must not only

account for the higher concentration of particles, but also the fact that the dislocation pair

can be in the same particle simultaneously. Figure 2.4 illustrates the difference in

deformation behavior between low volume fraction and high volume fraction _/' alloys. It

is not enough to simply alter the low fraction theories accounting for a higher volume

fraction as Ardell, et al., did. 71 Ardell, et al.,Tl modified the equation by Brown and

Ham 35 for a low volume fraction _/' alloys. The dislocation configurations are different for

high volume fraction T' alloys and this was not considered in this study. Ardell, et al.,

obtained good agreement between experimental results and their CRSS equation, but the

highest volume fraction used was only 36%. 71 This equation has not been applied to

higher volume fractions, 69 and it is believed that different models should be used, which

incorporate the differences in the dislocation configurations that are present in high volume

fraction T' alloys.

Different approaches must be considered and the two that are commonly used are

that of Copley and Kear, 2 and Reppich. 73 Copley and Kear 2 were the first to develop a

dynamic model for yielding in alloys containing coherent, stress-free, ordered panicles.

20
Figure 2.4

,o 0 0 'JO _° u

--6" 0 o

o © 0 n © _.

(b

(a) (b)

Dislocation structure in (a) low volume fraction and (b) high volume fraction

1/ alloys. 72

Their model is based upon a dislocation dynamics approach where dislocations were

assumed to move according to an exponential velocity-stress function. 74 This velocitystress function can be evaluated if the alloy structure, APB energy of the ordered phase,

and the drag stresses of the matrix and precipitate phases are known. They tested this

approach in the polycrystalline alloy MAR-M 200, which contains about 60 volume percent

7' dispersed in a _' matrix with about 5 volume percent carbides present. The 11'particles

were spherical and had a diameter of approximately 0.25 I.tm, while the carbides were

found to have no effect on the yielding behavior.

The model develops a force balance for a pair of dislocations attempting to enter a "/'

precipitate. The forces that are incorporated are:

1) force Xb; X = resolved shear stress.

2) force -T;, 7 = APB energy.

3) repulsive force, c/Ax, due to trailing dislocation; c= l.tb2/27z.

4) force T(x) due to line tension; x = displacement of dislocation segment.

In these expressions, I.t is the shear modulus and b is the Burgers vector. From these

forces, velocity components for the leading and trailing dislocations are calculated in the

21
matrix and precipitate. All of the above forces are incorporated in the velocity-stress

function for the precipitate, while the APB energy term is omitted in the matrix equations.

The average velocities of the dislocation pairs were calculated for various positions relative

to the matrix-particle interface, and it was found that penetration of the ftrst dislocation into

the particle was the rate-limiting step during yielding. They obtained the following

expression for the critical resolved shear stress (CRSS), xc, of the _/#/' material:

), T K'(Xm + %)

+ (2-5)

't;c- 2b br 2

(In v,/v) D m + Dp

K = (2-6, 2-7)

K 2('tb + T/r) - ),

! .._.

where r is the radius of curvature of the particles, x.m and xp are the CRSS of the matrix and

precipitate, v is a velocity near the shear wave velocity, v is the average dislocation

velocity, and Dm and Dp are the drag stresses of each phase. Copley and Kear also derived

a yield stress equation based now on static arguments that is similar in form to equation 2-5

with K' equal to unity.

The application of equation 2-5 requires the calculation or estimation of a number of

parameters. First, Copley and Kear measured the drag stresses and CRSS of each phase

by testing alloys of similar composition to the _' and _/' phases, as shown in Figure 2.5.

Then the CRSS of each phase, '_m and xp, was calculated as a function of temperature. The

drag stresses for each phase were calculated using the equation:

O

D = in(v/v)x c (2-8)

The APB energy calculation used by Copley and Kear yielded a value of 164 ergs/cm 2 on

{ 111 } planes. For the line tension, they assumed the approximation T = l.tb2/2, however,

this has been shown to be a poor approximation for these types of alloys. 69,75 The rest of

the par,'uneters are experimentally obtainable.

22
.ag

o3

0C

(O

40

3O

2O

Figure 2.5

0

0 4OO 80O 1200 1600

T_F)

CRSS of y and 7' phases as a function of temperature. 2

Obtaining values for the dislocation dynamics is difficult, and most investigators

use the version of the model based upon static arguments with K' = 1.76,77 This has been

successfully done by Dollar and Bernstein 77 to predict the temperature dependence of the

flow stress in CMSX-2 and PWA 1480. In PWA 1480, a small modification had to be

made to account for the unusual behavior in which most of the dislocations were trapped in

the 3' matrix and did not shear the precipitates. This led to a work hardening term being

added to the Copley and Kear equation for the temperatures where trapping occurred: 77

3' T

xc = 2b br + 0"41('l:m + '_p) + °tl'tbpl/2 (2-9)

where ot is a constant and I3 is the dislocation density in the 3' matrix. The term 0.41 was

added to account for Schmid's Law. 78

23
Theothermodelgenerallyappliedto highvolumefractiony' alloys is by Reppich.

Reppich 73 used the Schwarz and Labusch 79 correction for high volume fraction "1(alloys

and applied it to the low volume fraction model by Huther and Reppich. 8°,81 The latter is

designed to predict the increase in the CRSS due to a low volume fraction of coherent,

stress-free, ordered particles.

Huther and Reppich 8°'81 begin with the premise that for low volume fractions of

small particles, the models based upon the work of Fleischer and Friedel, 7° which propose

that the CRSS will increase with particle size are valid. However, at some critical particle

size the dislocations will change from being loosely coupled to being tightly coupled,

resulting in a hyperbolic decrease in the cutting stress. An equation for this change in

cutting stress, A'r,o, was derived:

Twfl/2/

A,Co = 0.86..._____.[ 1.28wd__7T _ 1) 1/2 (2-10)

where f is the volume fraction, d is the mean particle diameter and w is a term accounting

for the elastic repulsion between paired dislocations outside the particles. By equating this

equation with the one for small particles, the peak stress and a critical diameter for the peak

stress, de, can be calculated:

A'l:p_ = 0.56(_)f 1/2 (2-11, 2-12)

Finally, depending upon the shape of the particle, Orowan looping may become favorable

at larger particle sizes. Figure 2.6 sums up the Huther-Reppich model for spherical

particles.

Schwarz and Labusch 79 looked at simulated motions of a dislocation through a

random array of weak obstacles of finite size, and coined a parameter, rio, which is

proportional to the ratio of the average angle that the dislocation makes by zigzagging

24
CRSS

f_

\

D_ D* porticle

diameter

Figure 2.6 Huther-Reppich model for spherical particles. 8° Plot of CRSS vs.

particle size. Curve 1 is equation 2-10 in text.

through the obstacles to the maximum angle that the dislocation can sustain without cutting.

rlo is called the normalized obstacle depth and for spherical particles can be expressed as:

(2fT'] 1/2

no = \ "-F-'_./ (2-13)

where F is the dislocation obstacle interaction force. When rio <<1, Fleischer-Friedel type

models are valid. However, Reppich 73 substituted realistic values for nickel-base T/7'

alloys and concluded that 1"1ois almost never << 1. Finally, Schwarz and Labusch derived

an expression that relates the increase in CRSS due to a high volume fraction of

particles, A%, to the increase found due to small volume fractions, A't_: 79

A't o = 0.95(I + C'qo)AXFF (2-14)

where A,r_ is the CRSS calculated from the Fleischer-Friedel model and C is a calculable

constant. For large particles and high volume fractions, F = 2T and rio = (f/2r0.

Now, we are in a position to apply the Huther-Reppich hardening model in

combination with the Schwarz-Labusch correction for high volume fraction _' alloys. The

modified values of CRSS vs. particle size for spherical particles are shown in Figure 2.7.

25
A To

¢i-"

15

O5_

PAIRWISE

'_ OROWAN

CUTTING

! nm

10(3 200 300 400 500 _

_ _._r f _ ,50 o/.

f

4OO

3OO

2OO

IO0

0

2 3 4 S 6 7 8 g _0 IJ CZ

--din

roo

60o

50o

Figure 2.7 Reppich model accounting for large Vf of spherical particles. Curve 1: cutting

stress for weakly-coupled dislocation pairs. Curve 2: cutting stress for

strongly-coupled dislocation pairs. Curve 3: Orowan stress. Curve 4:

penetration stress. 73 Thin lines are results of low volume fraction model and

thick lines are correction.

The thin lines represent the conventional model, and the thicker lines employ the SchwarzLabusch correction.

A minimum value of C was used in these calculations, and obtaining C seems to be

a difficult task. Further, Huther and Reppich admit that their value of g and w could

contain large errors. 81 A serious concern is that in practice w becomes a term determined

by fitting the experimental data. In other words, "it contains all the uncertainties of the

theory."8° Further, Ardel169 points out that there is a large spread in the values obtained by

different investigators for the constants C and w. 75'82"85 Ardel169 concludes that the use of

the Labusch and Schwarz model may hold some promise, but currently there are too many

unknown variables to reliably predict the CRSS.

26
2.1.2.3 Tension/Compressionasymmetry

It has only recently been discovered that Ni3Al-based alloys exhibit a

tension/compressionflow stressasymmetry.86-9° It wasgenerally found that for the

<001> orientationthe flow stressin tensionexceedsthatin compression.For mostother

orientations,the compressiveflow stresswasgreaterthanthe tensileflow stress. D.P.

Popeand co-workers86"89haveled researchin this areastudyingNi3Al-basedalloys, as

well asPWA 1480. Themostrecentpaperby HerediaandPope88providesanoverview

of thepreviousmodelsandexperimentalobservations.

Paidar,et al.,86werethe first to describea model to explain this asymmetry. It

actually is a generalizedversion of the Takeuchiand Kuramoto (TK) model, 6° which

explained the anomalous increase in flow stress with temperature. The TK model worked

well in explaining this effect, but incorrectly predicted the tension/compression asymmetry.

Paidar et al., looked at the effect of the applied stress tensor on the dislocation core width

and not just on cross slip, as was the case for the TK model. After deriving an activation

energy for cross slip and relating this to the flow s_'ess, they were able to show how the

direction of applied stress could affect the flow stress. This supports the premise that

dislocation core constriction or expansion plays a larger role than the stress driven cross

slip in determining the tension/compression flow stress asymmetry in these alloys.

Shah and Duhl 9° observed the same flow stress asymmetry in PWA 1480 but

considered the splitting of dislocation partials as the explanation. They find that in the

<001> orientation an applied tensile stress causes the partials to constrict and a compressive

stress expands them. Since it is necessary to bring the partials together before cross slip

occurs, it is argued that the samples pulled in tension will be stronger than those under

compression. This analysis also was successfully used to predict the orientation at which

no asymmetry occurred in the yield strength.

The Paidar, et al., 86 model has been experimentally proven not only in their own

work using previous data on Ni3(AI,Nb) single crystals, 89 but also in separate work on

27
Ni3(AI,Ta) singlecrystals87and on PWA 1480. 88 Each of these studies report the same

orientation dependence of the flow stress asymmetry, as well as similar temperature effects.

Therefore, it can be concluded that this model utilizing the effect of the applied stress tensor

on the dislocation core is widely applicable. However, the similar, yet simpler model of

Shah and Duhl also predicts the correct trends, and a final decision on appropriateness of

each is not yet possible.

2.1.3 Fracture Behavior

2.1.3.1 Crystallographic fracture

The fracture behavior of single crystal nickel-base superalloys is very different from

their polycrystal counterparts. In single crystal superalloys, it is very common for the

fracture to occur along specific crystallographic planes and to have a cleavage-like

appearance. This type of fracture has been observed in other fcc single crystals 91-_ and is

thought to be a general phenomenon of fcc single crystals. All of these studies report

failure along {111} planes, however {001} failure has been reported in some

instances. 9'95'96 This behavior has been primarily studied under fatigue conditions, 9'1"99

although it also is observed in monotonic loading. 9,t°°,l°l

In most cases the fracture surface has the appearance of cleavage at low

magnifications, but with evidence of moderate amounts of fractographic plasticity at higher

magnifications. 9,95,1°° Physical measurements such as reduction-in-area and crack-tip

strain also indicate that this is a plastic process. For failure along {111 } planes, it is

commonly observed that more than one {111 } variant is involved. 95'96'97'100 Figure 2.8

illustrates this type of fracture in an [001] compact tension specimen with a transverse

orientation of [010]. The determination of which planes are active in the fracture is not a

difficult process for single crystals and is treated elsewhere 1°2, as well as in section 3.4.6.

28
Chart and co-workers 97']m'1°3 have studied this phenomenon of crystallographic

failure from a viewpoint of the stresses involved. This subject is now discussed, as well as

in section 2.4 because of its importance in linear elastic fracture mechanics analysis.

Figure 2.8

,1'1'

°°*L7

L

W

[

Schematic of { 111 } crystallographic fracture in an [001] CT specimen. 95

In the crystallographic failure of fcc single crystals, the crack path occurs along planes

which can have large shear stresses; however the fracture surface has a cleavage-like

appearance. Koss and Chan 1°1 found that large normal stresses build up ahead of the

crack tip when fracture is occurring along a coplanar slip band. These normal stresses are

about three times what one would find ahead of a crack tip when uniform multiple slip is

involved. It is believed that the large normal stresses do not allow the crack to deviate, and

this results in the fracture path remaining exactly along the coplanar slip band.

This same analysis also shows that for fcc single crystals, fracture along a planar

slip band is stable and secondary slip is very difficult. 101 They obtained an expression for

S $

r_,/rp, where rp is the plastic zone size for a coplanar slip band and rp is the plastic zone size

in the plane of the crack due to secondary slip. This ratio is an indication of whether

uniform multiple slip, or non-uniform coplanar slip should occur at the tip of a moving

crack. When r_,/rp <<1, then coplanar slip is expected and vice versa. Koss and Chan's

29
results showed that for fcc single crystals this ratio would be 0.07, so coplanar slip would

be stable.

Koss and Chan 1°1 also suggest that the fracture along coplanar slip bands is due

solely to the "elastic-plastic" state of Stress involved and that the dislocation configuration

ahead of the crack tip is of secondary importance. Further, this type of fracture is expected

to be associated with a low fracture toughness, but this degradation of toughness can be

altered by means of tortuous crack paths, crack branching or renueleation of a shear band at

a boundary.

The formation of ridges on the fracture surface also was examined by Chan, 103

who showed that they are the result of fracture along more than one { 111 } plane. In

previous work, 1°1 Chan showed that coplanar slip did not relax the normal stresses ahead

of the crack tip, and thus concluded that cross slip alone cannot reduce the normal stresses

ahead of the crack tip. 103 The result is that large normal stresses can occur on the coplanar

slip plane, as well as the cross slip plane. The combination of large normal stresses and

localized shear on both planes can thus explain the observation of simultaneous {111}

cross-slip planes leading to ridges on the fracture surface. While Chan did not extend this

to include two pairs of cross-slip systems, it would seem logieal to do so. This could then

explain the presence of all four { 111 } variants on a fracture surface. 95 An important

consequence of this analysis is that the crack growth direction is predicted to be <110>,

which is common to both cross-slip planes. Support for this is provided in a fatigue study

on MAR-M 200,1°3 where fatigue striations were observed perpendicular to the <110>

direction proving that <110> was the direction of crack growth.

The presence of facets also was analyzed by Telesman and Ghosn 95 from the point

of view of the resolved shear stress on these planes. This approach also was used by

Chan, Hack and Leverant 97 and was successful for predicting which planes would be

present on the fracture surface, but only when a single plane traversed the entire fracture.

When there were fracture ridges with two planes active, the model was not applicable.

30
The TelesmanandGhosnmodel95calculatesthe resolvedshearstresson all the

possibleslip systemsand thenpredicts which systemsshouldoperateand when. The

limiting valueof thisresolvedshearstressisdefinedas:1°4

Krss = lim Xrss_/2r_r (2-15) toO

where r is the distance to the crack tip and Xrss is defined as the projection of the stress

tensor, t_, on a plane whose outward normal is n in the direction of slip B. Cracking is

assumed to occur along the slip systems with the highest Krss value. While it is difficult to

assess whether the authors are implying crack growth across the facets or parallel to them,

it will be assumed that growth is perpendicular to the facets based on Chart's evidence 1°3,

as well as evidence in this work. Telesman and Ghosn suggest that crack growth begins

on one system, and after some unspecified distance, the crack then will turn and propagate

on another system of high Krss. 95 The criteria for choosing which { 111 } system, or

exactly why and when the crack changes direction, is not given. They also consider that

the same behavior occurs for the other pair of { 111 } systems and, hence, development of

the fracture ridges directly follows. Unfortunately, this model is not intuitively obvious,

and there are other systems with similar Krss values that also appear to be valid possibilities

for crack growth. Further, it also is difficult to imagine how the crack advances using this

model. It seems better at this point to use Chan's simplified cross slip argument 1°3 and

assume that the two pairs of slip systems merely intersect and that there is not a change in

crack direction.

2.1.3.2 Relation to microstructure

The relationship between the T/'/' microstructure and the fracture surface has been

studied by several researchers. 9'11.95.105 When fracture occurs on {001 } planes, it is quite

easy to see the _//q/outline on the fracture surface. 9'95 Further, it is clear that the matrix is

the more ductile phase as evidenced by tear ridges through the "¢phase, while the _ has a

31
morebrittleappearance.Howeverit is notclearif thefinal fractureisoccurringthroughthe

"/matrix or at the _//_/' interface, 9 although some evidence suggests fracture through the 3,

phase. 11 Telesman and Ghosn 95 believe that this type of fracture suggests that the only

deformation occurring is in the 2, phase since {001} cube slip is not observed at room

temperature. 52 This leads one to believe that the fracture across the y precipitate is a rapid,

brittle fracture. Sherry supports this conclusion by stating that quasi-cleavage of the 7' is

occurring with alternate shear of the y by octahedral slip) °5

When fracture occurs on { 111 } planes, as is more typical, it is difficult to ascertain

the role of the )'/y microstructure. However, the same mechanisms may apply, with the

exception of failure of the 7' on {111 } planes. The fracture of the 7' is still brittle-looking,

but is now along the slip plane. This brittle-type fracture can be rationalized using the

above argument of Koss and Chan, 1°1 identifying the larger than usual normal stresses

ahead of the crack tip in this type of fracture.

2.1.4 Effect of Hydrogen on Superalloys

The effect of hydrogen on superalloys has been studied for over twenty years,

primarily driven by the environmental conditions present in the Space Shuttle Main Engine

(SSME). A large amount of data currently exists on hydrogen embrittlement, mainly from

screening tests on a wide variety of superalloys; much of this information has been

compiled by Fritzemeier and Chandler. 8 Tests on superalloys most often were performed

under a 34.5 MPa hydrogen atmosphere to simulate the SSME conditions with comparative

tests performed in a 34.5 MPa helium atmosphere. But there were some tests done on

intemaUy charged specimens tested in air or a hydrogen atmosphere. 8t°

No attempt will be made to analyze results outside the scope of this thesis, and the

reader is referred to reference 8 and its references for further details. The review 8 covered

ovcr 35 supcralloys, and compiled much of thc work done in the area of tensile properties,

32
creep rupture,fatigue and fracture mechanics. Therewas no obvious trend between

wrought, conventionally cast and single crystalline material as far as hydrogen

embrittlementwasconcemed.It wasfound,not unexpectedly,thattheworstembrittlement

occurredat room temperatureasopposedto cryogenic or elevated temperatures. It also

was concluded that trends seen in other alloys tested in hydrogen, such as strain rate,

frequency and pressure, were equally observed in superalloys. In summary, there were no

superalloys found to be totally resistant to hydrogen-induced degradation of properties

when subjected to severe enough conditions. Specific results in the literature that are

applicable to this work will now be discussed.

The effect of single crystal orientation on hydrogen embrittlement has been studied,

and in PWA 1480 it was noted that the <001> orientation appeared to be the most severely

embrittled orientation in terms of notch tensile strength. 1°6 Despite this result, the

advantages of the <O31> orientation in other areas, particularly creep behavior, continue to

make it the most attractive orientation for rocket engine applications. Generally the notch

tensile strength of [001] single crystals tested in a hydrogen atmosphere at room

temperature was reduced by at least 40%, and the reduction of area (R.A.) in smooth

specimens was reduced by more than 20% in all specimens. 8 But the R.A. may not be a

valid indicator of hydrogen embrittlement based on results in this thesis and those of

Fritzemeier. 25 Fritzemeier found that most single crystal superalloys tested in hydrogen

showed a much larger corresponding loss of strain to failure compared to the decrease in

R.A. It was generally found that the more severe the embrittlement for a given alloy, the

larger this discrepancy became. It is believed that these differences reflect strain

localization, a manifestation of hydrogen-enhanced localized plasticity, as will be discussed

later in the work.

Various observations will now be mentioned including an effect of microstructure

was observed in the CMSX series of alloys, where larger, more uniform cuboidal

precipitates improved the alloys' resistance to hydrogen embrittlement. 8 It is possible that

33
the homogeneousdeformation associatedwith this microstructure, comparedto the

localizedslip observedin theCMSX alloyswith smaller,non-uniformprecipitates,led to

theimprovedproperties.Crackswereobservedonthetensilegagesurfacein MAR-M 246

testedin hydrogen,which initially startedon {001} planesand then propagatedon to

{111} planes.25 SurfacecrackingalsohasbeenseenoninternallychargedCMSX-2tested

in air, however, no analysisof the crack orientation was performed.9 In separate

experiments,theeffect of testing pre.-charged samples in air and hydrogen was studied on

MAR-M 246, and it was observed that the combination of internal and external hydrogen

produced slightly worse properties than for either ease tested separately, s

Only a few studies have carefully analyzed the structure/property relations in the

hydrogen embritflement of superalloys, 9'1°'2"5'77 and these will be discussed now. In

CMSX-2, the role of microstructure in the hydrogen embrittlement of cathodically charged

and gas-phase charged tensile specimens has been carefully studied. 9,1°,1°7 It was

concluded that gas-phase charging produced more severe embrittlement, with the difference

directly related to the concentration profiles resulting from the different charging

conditions. Under both charging conditions, the role of hydrogen trapping at

heterogeneities on properties was studied using desorption experiments and tritium

autoradiography. 9,1°,1°7 It was found that the pores accumulate hydrogen in cathodic

charging, but probably not in the gas-phase charged specimens, where the charging

temperature of 350°C, is thought to be too high for effective trapping. 1° Hydrogen also

was observed to accumulate in the eutectic 2'/7', either due to direct trapping or because of a

higher solubility in the eutectic 2'/7' compared to the rest of the 2'/7' microstructure. 1° A

microprobe analysis using a deuterium probe, attempted to determine the trapping and

solubility characteristics of the 2' and 7' phases, 1°8 and while no consistent conclusions

could be drawn from this work, evidence suggests there is a higher concentration of

hydrogen at the 2'/y' interface. This work on CMSX-21° also revealed that there was no

effect of tensile strain rate on embrittlement over a range of 5 x 10 -a to 5 x 10 .2 sec -l. This

34
is contraryto resultsin manyotheralloys1°9wheretheeffectof hydrogenwasfoundto be

lesssevereat high strainratesbecauseof a reducedinteractionbetweendislocationsand

hydrogen. In CMSX-2,it wassuggestedthat this argumentdoesnot applybecauseof the

high, uniformconcentrationof hydrogenin thesample,thusminimizing dynamiceffects.

FurtherstudiesonCMSX-2ll performedatCarnegieMellon showedsimilarresults

to theabovestudies.Thework to bediscussednow focusedonthedislocationstructureas

a function of hydrogenand strain in an attempt to correlate theseobservationswith

companiontensile propertiesand fracture behavior. The presenceof hydrogenled to

dislocationsbeingtrappedin they phaseat all strainsinsteadof the extensive shearing of

precipitates observed in uncharged specimens. This dislocation trapping led to strain

localization and failure in the T matrix, as revealed directly by fractographie evidence. II

Strain exhaustion is the most likely explanation as the work hardening also increased in

hydrogen charged samples, prompted by the high concentration of dislocations in the

narrow 3' phase. Hydrogen also led to a change from fracture along { 111 } planes to an

{001 } fracture where the 3' and T' phases could easily be seen on the fracture surface.

Apparently, hydrogen is causing the majority of deformation to occur by {111 } slip in the 3'

matrix leading to failure in this phaseJ t For those dislocations present in the '1( phase at

room temperature, hydrogen-induced cross slip onto the {001} planes was observed. 11

This was rationalized using the Takeuchi and Kuramoto model 6° in terms of hydrogen

retarding the "dynamical break away" process described in this model. All of the above

observations are important to this work because they involve the effects of internal

hydrogen on single crystal superalloys. Thus, the results of this work will be heavily

compared to the above studies in hopes of further defining the effect of hydrogen on this

type of alloy.

35
2.2 PWA 1480

Exceptfor theCMSX-2 work, moststudieshaveconcentratedon screeningalloys

for hydrogen application. A notable exception has been PWA 1480 and since its

introductionin 1980,it hasbeenoneof themoststudiedsinglecrystalsuperalloys.In an

effort to evaluateits potentialfor usein jet androcketengines,therehavebeenmanytests

performedboth in air and hydrogenenvironments. Howeverevenhere manyof these

studies have concentratedon data accumulationand have ignored structure/property

relations,anddeformationandfracturebehavior.

The major testing efforts have been in the areas of low and high cycle

fatigue, 19,76,110-115 fatigue crack growth,95.113-115 creep and stress

rupture7.113,114,116,117 and elevatedtemperaturetension testing. 23,76.77,90,113,116,118,119

There have been many tests performed in hydrogen gas at 34.5 MPa to simulate the space

shuttle main engine operating conditions, 1°6A13'115'118'12° but there have only been two

studies concerned with the effect of internal hydrogen on mechanical properties. 25A12

Porosity was found to play a large role in many of these tests, but there have been few

published results on hot isostatically pressed material. 19,z3'112 Thus, it appears that the

,areas of hydrogen and microstructural effects are in need of further research.

Some of the points to come out of these studies that are useful to this work are

discussed below and include the highly anisotropic nature of the mechanical properties due

to the anisotropy of the elastic modulus. 9°'114'118 It also was observed that fractures were

very crystallographic at low temperatures and non-crystallographic at higher

temperatures 11,_ probably due to the decrease in predominance of octahedral slip. In the

low cycle fatigue studies, it was found that fatigue crack initiation sites were invariably the

casting porosity. 23'114'115 Like other nickel-base alloys, 8,121 hydrogen similarly degraded

the low cycle fatigue, notch tensile strength and tensile ductility of PWA

14_qO.25'10_''113'118,120 ftowever, the adverse effects of hydrogen were often less severe

36
in comparison to other single crystal and directionally solidified superalloys. The

following sections will focus on properties and observations of direct relevance to the

research reported in this work.

2.2.1 Relevant Studies

The first study to carefully characterize the microstmcture and deformation behavior

of PWA 1480 was by Shah and Duhl, 9° who examined the effects of orientation,

temperature and _' size on yield strength. They tested three major orientations, [001],

[011] and [111], in tension and compression and noted a tension/compression asymmetry.

They attributed this to the role of thermally activated cube cross-slip, as discussed in

section 2.1.2.3. The [001] orientation was found to be the strongest in tension and

weakest in compression at all temperatures, with a slight drop in the yield strength at

400°C. Schmid's law was not obeyed, 78 which was explained by differences in the

propensity for cross slip in each orientation. At low temperatures, the smallest 3" size

microstructure was the strongest, but at intermediate temperatures the differences became

small. In the discussion of these results, Shah and Duhl assumed that dislocations sheared

the T' precipitates at all temperatures, however there was no TEM evidence provided to

support this contention.

In a later study, Milligan and Antolovich 76 also looked at the effect of temperature

on yield strength of PWA 1480 containing -60% volume fraction ofT' precipitates with an

average size of 0.5 l.tm. The alloy also had a volume fraction of porosity of 0.15-0.20%,

and 4.5-5.0% of eutectic 3'/3('. The 3'/T' mismatch had been measured from extracted 3"

precipitates in a related study and was reported to be 0.28%. 122 Milligan and Antolovich

performed interrupted tensile tests to fixed plastic strains and analyzed the deformation

substructure of [001] oriented samples. They observed that at low temperatures, (<760°C),

shearing of the T' was the dominant deformation mechanism. At higher temperatures, T'

37
bypass occurred and eventually became dominant. It should be mentioned that

temperatures between 20°C and 705"C were not tested, so the previously observed drop in

strength around 400°C 9° was not observed.

Milligan and Antolovich 76 attempted to model the measured critical resolved shear

stress (CRSS), but with limited success. They found that although the Copley and Kear

model 2 apparently described PWA 1480 at low temperatures, one of the assumptions in the

theory was violated, namely that Imnetration of the leading dislocation into the 7' precipitate

was the rate-controlling step. In contrast to this assumption, it was observed that 8-10

dislocations could reside in the 7' before the first dislocation completely traversed it.

Therefore, for PWA 1480, they concluded that the rate controlling step is instead the drag

stress of the precipitate, and further that there is not a currently available CRSS model to

predict the deformation behavior.

In a later paper, 123 Milligan and Antolovich studied the yield drop between room

temperature and 400°C and linked this behavior with their observation of a large density of

Superlattice Intrinsic Stacking Faults (SISF's) in the 7' precipitates at room temperature.

They proposed that these faults act as barriers to glide and raise the CRSS at room

temperature. The density of SISF's is reduced at higher temperatures, and thus the drop in

CRSS up to 400°C. Above 400°C, no SISF's were observed and the rise in CRSS is

credited to the thermally activated cube cross-slip mechanism. 1_

In another study of PWA 1480, Dollar and Bernstein 77 attempted to describe in

more detail the effect of temperauire on the development of the deformation structure. They

examined tensile properties and deformation structures at temperatures of 20°C, 400°C, and

730°C, and at various plastic strain levels. At room temperature, superdislocations were

observed trapped in the 3' matrix in the early stages of plastic deformation. With increasing

plastic strain, the dislocation density increased in both the precipitate and matrix but still

was subStztr_tially higher in the matrix. They argued that since the yield stress of the y'

ph,_sc is considerably less than the yield stress of the superalloy, 2 superdislocations can

38
move into the _' below the overall yield stress. However, when they enter the _, phase they

cannot move back into the _ because of the additional APB energy required. It should be

mentioned that there was not a large density of SISF's reported in the 7' precipitates at

room temperature, in contrast to the study of MiUigan and Antolovich. 12-3

At 400°C and 730°C, the dislocations were no longer trapped in the matrix at low

strains. The superdislocations observed at 400°C were now long, straight a/2<110> screw

dislocations, but at 730°C the screw dislocations were no longer straight. These dislocation

configurations were attributed to the cross slip of superdislocations, as originally proposed

by Kear and Wilsdorf. lzs Dollar and Bernstein also noted a decrease in strength up to

400°C. In order to explain the yield drop at 400°C and to model the CRSS in this

temperature regime, they used the deformation structure observations to modify Copley and

Kear's model. This analysis is detailed in section 2.1.2.2.

In yet another study on PWA 1480, Fritzemeier determined the effects of high

thermal gradient casting, hot isostatic pressing (HIP'ing) and heat treatment on the

mechanical properties. 19.z3 It was found that using a higher thermal gradient during the

solidification process yielded a material that was easier to HIP with fewer and smaller

pores, as well as a more compositionally homogeneous casting. While it was possible to

eliminate all pores by HIP'ing, this had no effect on subsequent tensile properties. The

primary aim was to improve the fatigue properties, which was accomplished by IMP'ing.

The tensile ductilities are higher than many other reported values, and this could be due to

microstructural variations or more likely, orientation effects.

2.2.2. Effect of Hydrogen

The majority of tests designed to evaluate the susceptibility of PWA 1480 to

hydrogen embrittlement have been performed in a hydrogen atmosphere of 34.5 MPa.

()fter_ a comparison of the notch tensile strength in high pressure helium versus hydrogen

39
was done, and PWA 1480 generally performed better than other alloys in its

class.25.106.118,120 Fritzemeierhasmadesomeobservationson tensilesamplestestedin

hydrogengas,25which areof particularinterestto thiswork. In general,theslip lineson

thetensilegagesurfaceof the samplestestedin heliumweremorehomogeneousthanon

thosesamplestestedin hydrogen. Furthermore,thereductionin areawasnot affectedas

drasticallyby testingin hydrogenaswasthestrainto failure. Bothof theseobservations

leadto the conclusion that hydrogen-inducedlocalizedplasticity126,127is occurringin

PWA 1480. Similar observations were made on other single crystal nickel-base

superalloys,25asreportedin section2.1.4.

2.3 HydrogenEmbrittlementProcesses

There have been many proposed processes or mechanisms for hydrogen

embrittlement of metals, but none can yet satisfactorily explain all of the experimental

results reported in the literature. While a given process may explain the mechanical

behavior in the same system, the same process may not operate in the same or another

system under different conditions of hydrogen fugacity, temperature, stress state, etc. The

most widely accepted hydrogen embrittlement processes will be discussed, along with the

concept of hydrogen trapping. These can be broken down into four categories:

1. Hydride precipitation

2. Decohesion

3. Trap theory

4. Hydrogen enhanced local plasticity

The first of these relates to systems that form hydrides or martensitic phases. 128

This phase is normally brittle, ,and failure occurs along or through this phase. While pure

nickel has been rcported to form hydrides, it is not possible at the hydrogen concentrations

40
or fugacitiesusedin this work.129Further,thereis noevidenceof hydrideformation in

nickel-basedalloysunderanyconditions. Therefore,this mechanismisof little interestin

relationto thisthesis.

The decohesionmechanismproposes that the presenceof hydrogen solutes

decreasestheatomicbond strengthof the metal causinga brittle fracturewhenthe local

stressexceedstheatomic bond strength. 13°,131 This brittle failure process can occur in

conjunction with general plasticity in the sample, and therefore is in competition with other

failure mechanisms. This theory has come under recent criticism because of the lack of

consistent experimental evidence and a sound theoretical basis. 127 However, it remains a

popular mechanism and may be the only explanation for hydrogen embrittlement in some

systems. We believe the two most relevant mechanistic possibilities in single crystal

superalloys are the last two, and these will be considered in more detail.

2.3.1 Trap Theory

The trap theory states that a crack will initiate at a hydrogen trapping site when the

quantity of trapped hydrogen at the site exceeds a "critical concentration. ''132,133 Trapping

of hydrogen at microstructural sites has been observed in superalloys, 9A° so this hydrogen

embrittlement process will be covered in detail. The theory does not usually directly

consider any particular transport mode for the hydrogen to reach the trap, but recognizes

that diffusion along short circuit paths or by dislocation transport would shorten the time to

reach the critical concentration. There is no evidence of "critical concentrations" nucleating

cracks, and this theory will be approached from the viewpoint that it can explain hydrogen

trapping. Thus, it is not a bonafide hydrogen embrittlement mechanism but rather a

process by which embrittlement may occur.

The number and strength of traps, and whether the traps are reversible or not, play

a key role in the effect of hydrogen on the material. For this reason, it is important to

41
identify thoseinternal defectsthat could act as hydrogentrapping sites,as well asto

analyzetheirrelativetrappingstrengths.Thereis experimentalevidencethatsupportsthe

ideathatcracksnucleateat trappingsites,1°7'132'134howeverit is notknown if thisis due

to a "critical concentration"beingreached. Direct experimentalobservationof trapping

sitescan beobtained by the useof tritium autoradiography. Furtherdetails of tritium

autoradiographyandits applicationto anumberof othersystemsaregivenin Appendix E.

A thorough review of observed trapping sites in many different alloys and their relation to

hydrogen embrittlement has been given by Lacombe, Aucouturier and Chene. 1°7

Trapping sites can include grain boundaries, inclusions, dislocations, voids and

interstitial sites in the lattice. A classification approach has been used, 132 where all

trapping sites can be broken down into physical traps, attractive traps or a combination of

both. Physical traps such as voids or inclusions are modifications to the crystal lattice and

generally are considered moderate to strong trapping sites; attractive traps are regions where

hydrogen atoms are acted upon by an attractive force, such as the stress field of a

dislocation. These traps are invariably more reversible than physical traps, because the

probability of the hydrogen atom diffusing from this region or being left behind by a

moving dislocation is greater. The trap strength and reversibility of most traps often can be

determined by desorption experiments. 9'135'136

Trapping can also be evaluated in terms of internal versus external hydrogen

embrittlement and reversible versus irreversible traps. 132 Pressouyre states that some of

the differences found in hydrogen embrittlement under different testing conditions can be

explained by the types of traps involved. 132 Irreversible traps always will act as sinks, and

reversible traps can be a source or sink depending upon the testing conditions and their

interaction with other traps. It has been shown that in samples with both types of traps, it

is easier to reach a critical concentration of hydrogen at a potential crack nuclei when the

specimen is internally charged rather than if the same specimen is tested in hydrogen gas.

This is belicved to occur because all traps act as sinks in the external case, while reversible

42
traps may actually contribute hydrogen to the potential crack nuclei by dislocation transport

in internally charged specimens. Pressouyre points out many cases where the above

argument is borne out in experinmnts. 132

While the trap theory of hydrogen embrittlement explains hydrogen trapping in the

microstmcture, it is rather unsatisfactory in explaining how crack propagation and failure

occurs. Therefore, the trap theory is not complete in itself, and one of the other

mechanisms, such as decohesion, must be invoked to explain how the crack is nucleated.

2.3. I. I Determination of the trap binding energy

As mentioned, the trap strength and reversibility of the traps are important in

characterizing the effect of hydrogen on material behavior. There are several methods to

determine the binding energy of hydrogen to microstructural features, and the desorption

method will be covered here. This technique involves heating a pre-charged sample and

measuring the rate or amount of hydrogen diffusing out of the free surface. The most

accurate desorption profiles are obtained by using deuterium- or tritium-charged samples

and continuously monitoring the release rate. In all cases, it is important to have reliable

diffusion data. Table 2.3 gives an overview of available hydrogen diffusion data in nickelbase alloys. 13714° It can be seen that the data scatter is extremely large, especially in the

7/]/alloys. Estimates for CMSX-29A° suggest that the data of Robertson 137 for pure Ni

may best fit this type of alloy. Robertson's data is a compilation of literature values, and

the resulting diffusivity value is smaller than almost all of the others in Table 2.3. It is

apparent that quality diffusion studies of hydrogen in high volume fraction ,/' alloys are

needed.

43
Table 2.3 Diffusion data of hydrogen in nickel-base alloys.

Alloy

IN718

IN 903

IN 903

Waspalloy

Ren6 41

MAR-M 246

Ni

Ni

Do (cm2/s)

1.07 x 10 -2

2.46 x 10 .2

4.00 x 10 -2

15.5 x 10 -2

3.90 x 10 d

2.0 x 10 .2

4.76 x 10 -3

6.44 x 10 -3

Q (cal/mole)

- 11,900

-12,590

-7,715

-8,278

-8,668

-6,954

-11,411

-9,610

Reference

140

140

138

138

138

138

139

137

Hirth 141 has compiled the large body of existing data of trap energies in iron, but

there has been substantially less work published on fcc alloys. 1°7'136 Hirth reports a

binding energy of 0.61 eV for hydrogen trapping on internal free surfaces, which is

slightly less than the estimated value of hydrogen trapped at pores in CMSX-2. 9 In nickel,

the trap binding energy of thoria particles has been estimated to be 0.30 eV, 135 and the

binding energy to dislocations in nickel is around 0.10-0.15 eV. 136 It is clear from these

few examples that there is a wide range of binding energies to crystal defects and

microstructural features.

A general drawback in using desorption data to calculate binding energies is that it

is very difficult to determine the exact relationship between a given binding energy and the

specific microstructural features. All microstructural components, as well as dislocations in

the material, can contribute to the observed binding energy, and it requires many tests to

separate out the effects of each microstructural site. However, desorption tests are

44
relatively simpleto perform,andthetrappingdataobtainedareusefulin understandingthe

effectof hydrogenonthematerial'sbehavior.

To bemostuseful,thepresenceandstrengthof gappingsitesshouldbedetermined

by desorptionexperiments,but asyet thereis no standardmethodof convertingsuchdata

into trap binding energies.The number of methodscurrently used 9'135'136 should be

applicable to the data obtained in this work, namely, the amount of hydrogen desorbed as a

function of time and temperature..

Lee and Lee 135 have proposed an expression for obtaining the trap binding energy

by performing desorption experiments as a function of temperature:

In -1 =ln_o-o+ R_!

(2-16)

where NT and NL are the concentration of trap sites and normal lattice interstitial sites, CT

and Co are the trapped and equilibrium concentration of hydrogen, EB is the trap binding

energy, Eo is the heat of solution, R is the gas constant, and TH is the desorption

temperature. Estimates are made for values such as NT, NL and Eo, and the desorption

curves yield CT and Co. Finally, by plotting Equation 2-16 the binding energy is obtained.

This method requires very accurate desorption curves over a wide range of temperatures.

Further, estimating NT is a difficult task and can severely influence the binding energy.

Thus, although this approach has some success in predicting the binding energy of

hydrogen to thoria particles in nickel, 135 it is not known if it can easily be applied to other

desorption data, such as in this work.

Thomas 136 has used measurements of the desorption of deuterium in nickel to

obtain values of the binding energies of hydrogen to dislocations and to Fe or Ti impurities.

The desorption was performed as a function of time and temperature resulting in detailed

45
plots that,asabove,were thenusedto fit thefollowing expressionfor aone-dimensional

case:

dc(x,t) drDdC(x't)] dCT

dt £xL d--i- (2-17)

where c(x,t) is the lattice concentration of hydrogen and dcT/dt is the trap rate minus the

detrap rate.

dCT Dc(x,t)[ [_ "--(ED + ET) ]

dt - 7L2 NT--fcT] -- (CTU)ex kT (2-18)

In this expression, CT is the trapped concentration of hydrogen, D is the diffusivity, _. is the

jump distance, NT is the density of traps, u is the attempt frequency (1013 sec'l), ET is the

trap binding energy, ED is the lattice activation energy, k is a constant and T is the

temperature. Again, estimates had to be made for some parameters, while others were

calculated by data fitting. Instead of using the entire equation, as proposed by Thomas, it

appears possible that the detrap rate portion could simply be used to analyze the desorption

data:

dCT . . F--(ED+ET)q

= -tCT.;expL _ "/

(2-19)

A problem in this approach is that the attempt frequency, "o, is so large it dominates the

equation, and unrealistically small binding energies are calculated. However, the more

general method used by Thomas is physically appealing, and application of this method to

the desorption data in this work was attempted.

A much simpler method to determine binding energies from desorption data was

used by Baker, et al., 9 to obtain the binding energy of hydrogen to pores in CMSX-2 by

46
usinganequationproposedby Hirth141andsolvingfor thebindingenthalpy,Hb:

= (kpl/X)ex (2-20)

where c is the atomic fraction of trapped hydrogen, k is a solubility constant, p is pressure,

Hs is the heat of solution, R is the gas constant and T is temperature. The solubility data

was taken from a similar alloy, 9.14° and the pressure was the equilibrium partial pressure

of hydrogen in air at 1 atm. The amount of flapped hydrogen can be estimated from the

desorption data and a binding energy found. This method is perhaps the least accurate, but

its ease of application make it the most promising for use in this study.

2.3.2 Hydrogen Enhanced Localized Plasticity

The final explanation for hydrogen embrittlement to be covered is based on the

premise that hydrogen embrittlement is a macroscopically britde fracture process, which

results from a hydrogen-enhanced localized plasticity (H.E.L.P) at the crack tip. 126 Again,

this is a process by which hydrogen embrittlement can occur and not a complete

mechanism. Birnbaum and co-workers have studied the H.E.L.P. mechanism using an

environmental cell TEM with in situ deformation of thin foils. 127.142-147 Bulk specimens

also have been used to substantiate the occurrence and importance of H.E.L.P. 14815°

Fritzemeier 25 has preliminary results that suggest H.E.L.P. is occurring in single crystal

superalloys, and thus this hydrogen embrittlement process will be covered in detail.

Macroscopic embrittlement as measured by decrease in strain to failure or reduction

in ,area can be caused by localization of high plastic strains to the region of the crack tip. It

is believed that the strains are localized because hydrogen reduces the stress to generate and

move dislocations, and thus the dislocation activity remains in a local area, leading to high

47
local strains. Hydrogencaninteractwith thedislocationsby eitherenteringthroughthe

cracktip, diffusing to it or beingpresentfrom charging.Thiseffectalsocanbemanifested

in adecreasein theflow stresssincedislocationmotionandgenerationiseasier.Theeffect

of hydrogenpressureon dislocationvelocity is schematicallyshownin Figure2.9,127an

analysisapparentlyvalid for both screwor edgedislocations,as well as for loops and

Frank-Readsources. It was developedby using the conceptsof dislocation shielding

wherehydrogenatomsshield the elasticinteractionof one dislocationwith another.127

This meansthat two-edgedislocationsthatnormallyrepel,could insteadattractandsmall,

atomic-scaleslitsopenupalongtheslip plane.127

Dislocation

velocity

0

_..j dislocations

HydrogenPressure

v

Figure 2.9. Schematic of the effect of hydrogen pressure on dislocation velocity.

The experimental confirmation of H.E.L.P. in a number of alloy systems has

mainly come from observations of in situ deformation of thin foils. Evidence was obtained

that suggests hydrogen aids ira the generation and mobility of dislocations in Fe, 142 Ni, 143

IN 903,144 SS 310,127 (zTi 14__ high purity AI, 146 7050 A1 and 7075 A1147, and

48
microscopicfracture alongslip planesalso hasbeenobserved.127There also has been

supportive evidence based on flow stress and fractographic observations made in bulk

samples of various alloys, t4815° The macroscopic fracture generally reflects the external

stress systems and other variables, and the correlation with the observed microscopic slits

on the slip plane remains to be done. From results on single crystal superalloys, 9,1°.25 it

appears that the trap theory and H.E.L.P. may be useful in explaining some of the results

of this work.

2.4 Fracture Toughness of Polycrystal and Single Crystal SuperaUoys

While the fracture toughness of superalloys is generally considered to be of less

importance than fatigue behavior in their design and application, an existing crack in service

can amplify the importance of crack initiation toughness under initial static loading

conditions. Fracture toughness also can be an important quantity used to determine the

initial stress levels in hydrogen-induced crack growth tests, as was done in this work.

Fatigue crack growth of superalloys is currently a well-developed field, but there has been

little work in the area of fracture toughness, 151156 particularly the effect of

hydrogen. 152'153A55 Table 2.4 gives the available fracture toughness values for some

nickel-base alloys and iron-based superalloys. To the best of the author's knowledge,

there are no published data on high volume fraction y superalloys, let alone single crystal

superalloys, and for that reason alone obtaining the fracture toughness of PWA 1480 as a

function of microstructure and hydrogen was a primary goal of this work. Determining the

fracture toughness of a single crystal presented certain problems that are addressed below.

In as much as earlier fatigue studies showed that porosity can be a favored crack initiation

site,19,J 14,157 it was believed that some micromechanistic fracture models could be used to

describe the fracture process in PWA 1480. Therefore, these also are reviewed.

49
Table2.4. FractureToughnessValuesof Nickel-BaseAlloys andIron-BaseSuperalloys.

IN 600

IN 800

IN 903

JBK 75

A 286

A 286

IN718

IN718

IN 625

t

tt

KIC (MVaqm)

90

139

140-165

98

78-122

92-122

57

Jc (kPa-m)

382

854

Kill (MPa4m)

94 tt

84-111 tt

63 tt

Reference

151

151

152

153

154

155

156

155

155

internally charged to 5000 appm for IN 903 and 8000 appm for JBK 75.

tested in hydrogen gas at 34.5 MPa pressure.

The crystallographic failure mode of single crystals was described in section 2.1.3,

and perhaps the most pertinent question in this work is whether the Mode I Linear Elastic

Fracture Mechanics (LEFM) approach is valid. The observed cracks in single crystals are

always inclined to the stress axis and thus can involve all three modes of crack opening ( I,

II, III). Chan and Cruse 42 used a two dimensional boundary integral equation (BIE)

method and analyzed the multi-mode problem using a compact tension specimen design.

To apply this method, it is necessary to measure the angles that the macroscopic crack

makes with the width and thickness of the specimen, as shown in Figure 2.10. It was

found that as long as the angle between the crack and the perpendicular of the stress axis,

_, was less than 30 °, then using a Mode I solution introduced negligible errors. This also

assumed that a projected crack length, a', is used where a' = a cos]3 + 1.

50
©

Figure 2.10. Schematic showing dimensions needed for measurement of a'.

This is a very important and simplifying Finding for this work particularly because

the use of Mode I solutions for this type of single crystal previously had been questioned in

the literature. 1°° Mode I solutions are commonly used in the analysis of single crystal

superalloy crack growth data, 95.114,115 and will be used in this study as well.

Another important point to examine when dealing with nickel-base single crystals is

the effect of anisotropy. Chan and Cruse 42 examined this effect and concluded that the

anisotropy typical in nickel-base single crystals was not sufficient to significantly alter the

expressions for KI. Therefore, it appears that the effect of anisotropy can be conveniently

ignored in using conventional KI expressions designed for isotropic specimens. Thus, in

subsequent calculations of stress intensities, the value of Young's modulus in the loading

direction is simply used.

2.4.1 Micromechanistic Fracture Initiation Models

Micromechanistic fracture initiation models will be used to correlate the fracture

toughness with the microstructure, in hopes of determining a microstructural crack

initiation site. In particul,'u-, models 158164 that relate the crack tip opening displacement

(CTOD) to microstructural parameters will be discussed. The following expressions were

mainly developed for specimens that failed by microvoid coalesence, however, their

51
adaptation to other ductile fracture processes is possible. In order to calculate SIC . the

expression relating the various fracture toughness parameters is used: 158

KI2¢ 5i¢ o"c

JIc= E - d n

(2-21)

where E is Young's modulus in plane strain, Oo is the flow stress, and dn is a term related

to Oo/E and the work hardening exponent n.

Rice and Johnson 159 used continuum equations for void growth 165 to obtain a

relation between _c and microstructural parameters. They considered the growth of a preexisting void of radius Ro, which is at a distance Xo from the crack tip and obtained an

expression for 5Ic:

_Slc= Xo F(XJRo) (2-22)

where F is a weakly varying function of Xo/Ro. Xo and Ro can be related by the

expression for the average three dimensional nearest neighbor spacing between initiation

sites: 166

Xo = 0.89Ro(f) -u3

where f is the volume fraction.

ttahn, et al., 16° showed that:

(2-23)

Using this equation and the work of Rice and Johnson;

KIC = (f)-l/6(E_yRo) -1/2 (2-24)

It has been found that this equation does not work well for a wide variety of materials 167

most likely because failure is assumed to occur by impingement of voids, which is not the

case in most materials. However, this model was a significant step towards modeling

fracture initiation toughness by incorporating the microstructure.

52
Morerecenteffortshavefocusedontheso-calledcharacteristicdistancemodel.161-

164The reasoningbehindthis modelis that theequivalentplasticstrain,gr,,mustexceed

thecriticalfracturestrain,El, over some characteristic distance, 1o; expressed as:

_IC = 13}1_ (2-25)

However, physically interpreting 13} and 1o can be difficult. Values of 13} have been

obtained in a variety of ways, but there is no conclusive evidence that any one method is

correct. 167 It also is unclear from existing experimental results exactly what the

relationship of lo is with regards to the microstructure.

Regardless of these shortcomings, this concept has been extended to include the

growth of the initiation sites during deformation. A few models have employed this

concept t58A64A68 and seek to substitute the fracture strain with a measurable quantity

obtained from the fracture surface. Garrison t64 has proposed an expression that relates _c

to the radius of the initiation site before fracture, Ri, and after fracture, R,,:

_IC -- Xo(Rv/Ri) (2-26)

This expression has been shown to hold for some materials, but not others. 169 Application

of this model was attempted here for PWA 1480, assuming that the pores are the fracture

initiation sites. Their growth, if any, during the fracture process was assessed and used in

the above expression.

2.5 Background Summary

It has been seen that while superalloy single crystals are microstructurally simple,

they present many complexities when studying their behavior. The individual single crystal

superalloy composition plays a large role in controlling the properties, and therefore much

more research is needed in understanding this area. Some of these compositional effects

arc known, but much of the knowledge is only empirical. Based upon the literature in this

53
area,it appearsthatPWA 1480shouldhavecomparablepropertiesto similar alloys,with

thepossibleexceptionof havinga higher'y' yield strength due to the presence of increased

amounts of Ta. There also appeared to be a possible relation between the pores and the

eutectic "//y' based upon the solidification literature, and this was explored. Incipient

melting is a very real problem, as is the success of HIP'ing such materials, and these

concerns were explicitly monitored.

It was found that it is possible the deformation of PWA 1480 differs from other

superalloys in that dislocation shearing is not the primary deformation mechanism at room

temperature; this was studied with the aid of flow stress models to assist in the

interpretation of the tensile properties and deformation. Also, the manner in which

hydrogen affects the deformation behavior is not known, and thus was a major goal of this

work. Some studies have suggested that hydrogen can induce localized plasticity in these

alloys, and this was examined in depth. The role of hydrogen trapping or of H.E.L.P on

subsequent flow and fracture was also studied. The desorption method was used to obtain

binding energies for hydrogen traps in an attempt to understand better the effect of

hydrogen trapping sites on subsequent mechanical properties.

Superalloy single crystals tend to fracture along { 111 } planes, and this has been

explained in terms of the stress state present ahead of the crack tip. It is possible that the

presence of internal hydrogen could alter this fracture mode in some way. Finally, it was

discussed that the fracture toughness of single crystal superalloys has not been studied, and

this appears to be a very fruitful area of research. Even polycrystal alloys, which are

reported to have a relatively high fracture toughness, have had very little reported on

fracture mechanisms. It was found that Mode I solutions can be used on these single

crystals, as has been successfully done in fatigue crack growth studies of these type of

alloys. Finally, it is hoped that the use of micromechanistic fracture initiation models will

:lid in interpretation of the data and possible help identify any microstructural crack

it_iti_ttion _itcs.

54
Chapter 3

EXPERIMENTALPROCEDURES

3.1 Material Preparation

3.1.1 As-Received Material

Twenty single crystal slabs of PWA 1480 produced by TRW Metals Division in

December 1983 were obtained from NASA Lewis Research Center in December 1986.

They were approximately 6.4 cm wide, 15.2 cm long, and 1.6 em thick with the long

direction being the growth axis. The samples were received in the solution heat-treated

condition described below with orientation data provided for a few bars. It is believed that

the ends of the bars had probably been sandblasted at TRW both to check the orientation

and look for any small angle grain boundaries. The bulk composition was provided by

TRW and in weight percent is as follows:

Major Elements

AI Co Cr Ta "ri w Ni

4.83 5.35 10.43 11.86 1.29 4.07 Bal.

Fe S

0.08 0.0015

Trace Elements

Si Mn P C Zr Hf B

<0.02 <0.02 <0.015 42ppm <75ppm <75ppm < 15ppm

3.1.2 Laue X-ray Analysis

The growth axi_ orientation of most of the slabs was determined by the Laue backreflection technique using a Siemens Kristalloflex 2 x-ray source. Best results were

55
obtainedwhentheendsof theslabswereelectropolishedor mechanicallypolisheddownto

0.05lam A1203 slurry. A two-axis goniometer capable of being mounted to a wire saw

was used when certain non-conventional orientations were desired. All Laue pictures were

taken with a 3 cm distance between the sample and the Polaroid Type 57 film. Typically,

35 kV and 20 mA for 5 minutes gave a good pattern. Table 3.1 shows the growth axis

orientation of all slabs and the transverse orientation of most slabs.

3.1.3 Heat Treatment

The normal heat treatment for PWA 1480, used by the majority of

researchers, 23,76.9°,114,116 was followed. This consists of a solution heat treatment

followed by a two-stage aging heat treatment. The first aging stage allows '1( precipitate

nucleation, and the second stage allows some growth and maximizes the '_ volume

fraction. All of the material was given the full heat treatment, even though TRW had

previously given the material a partial heat treatment. This was done to achieve a uniform,

desired _' size and volume fraction. The heat treatment used was:

4 hours at 1288°C, air cool to room temperature

4 hours at 1080°C, air cool to room temperature

32 hours at 875°C, air cool to room temperature

At the beginning of the study, the samples were heat treated after machining and were

encapsulated in quartz tubes back filled with 1/3 atm. argon. It was decided to discontinue

this practice, however, due to dimensional changes, varying cooling rates and problems

with the quartz tubing. The vast majority of all subsequent samples were heat treated as

as-received slabs prior to machining. Some minimal oxidation occurred that was machined

56
Table3.1 Solidificationaxisandtransverseaxisorientationof allPWA 1480single

crystals.

single crystal slab # growth axis Wan_vcrse axis

16 2 ° off [001] 5 ° off <130> towards <110>

28 1° off [001] 2 ° off <120> towards <110>

32 1° off [001] <230>

41 0 ° off [001 ] <230>

60 1 ° off [001] 5 ° off <010>

78 4 ° off [001]

80 3 ° off [001]

107 4 ° off [001]

114 5 ° off [001] 4 ° off <230> towards <110>

119 3 ° off [001 ] 2 ° off <230> towards <110>

131 4 ° off [001] 4 ° off <010>

134 7 ° off [001] 4 ° off <010>

166 4 ° off [001 ]

167 4 ° off [001]

169 2 ° off [001] <110>

177 4 ° off [001]

179 2.5 ° off [001]

180 4 ° off [001] <140>

183 2 ° off [001]

192 3 ° off [001] <130>

57
off. Further,thecooling rateat thecenterof thesampleswasmonitoredandfound to be

faster than the l l5°/min suggestedby Pratt and Whitney24 to guaranteeminimal

precipitationof the"[ phaseduringcooling.

An isolatedproblemencounteredduringthesolutionheattreatmentwasincipient

melting. The temperaturedifference betweenthe 7' solvus and the incipient melting

temperature(IMT) is only a fewdegreeswhenthisheattreatmentprocedureis used.24 To

avoid overshooting above 1288°C, a ramp of 10°C/min to 1150°C followed by a 3°C/rain

ramp to 1288°C was used; however isolated instances of incipient melting of the eutectic

T/7' still occurred. Figure 3.1 shows an example of such incipient melting.

Figure 3.1

lOOgm

Incipient melting that infrequently occurred during heat treatment.

In PWA 1480, the eutectic T/T' will melt if the IMT is exceeded, but it can also

dissolve below the IMT if sufficient time is allowed for diffusion to occur. An alternate

heat treatment was developed in order to minimize or eliminate the eutectic T/7' phase

because of its undesirable properties described in a later section. The only difference

between this alternate heat treatment and the normal heat treatment is that more time is spent

58
at thesolutiontemperatureto encouragetheeutectic7/7'to completelydissolve. While the

time neededvariedwith thematerial'sprior history,it wasfoundthat20hourswasneeded

fortheas-receivedmaterial. A normal two-stage aging heat treatment was then followed.

3.1.4 Hot Isostatic Pressing

Selected samples were Hot Isostatically Pressed (HIP'ed) to study the effect of

grown-in porosity on mechanical properties. During heating to a level near the normal

solution temperature, pressure is applied to a level that is below that needed to close the

pores. The pressure is provided by a gaseous atmosphere, usually argon. Once the

desired temperature is reached, it is held for awhile to assure a uniform temperature

throughout the sample. Finally, the pressure is increased at a moderate rate up to the

pressure necessary to close the pores. The key is to slowly ramp the temperature and

pressure to avoid any recrystallization. The samples are left under these conditions for a

period of time to allow diffusional processes to occur. The f'u'st attempt at HIP'ing was

performed at Sandia National Laboratories in Livermore, Calif. The HIP'ing cycle was a

single ramp to 1285°C and 103 MPa pressure. This level was held for one hour, followed

by a quick helium cool. This produced a very slight reduction in porosity but no effect on

mechanical properties. It was subsequently learned that Rocketdyne had developed a

HIP'ing cycle specifically for PWA 1480, and this was attempted 33 at the Howmet facility

in Whitehall, Mich. While the Rocketdyne cycle is under a U.S. patent secrecy order, it is

known to be very close to a patented cycle developed at Howrnet. 32

spccific conditions for PWA 1480 are as follows:

For the latter case, the

1) Heat to 1288°C under a pressure of 34.5 MPa.

2) Hold for 1-2 hours.

3) Increase pressure at a rate of 0.42 to 0.55 MPa/min. to 103 MPa.

4) I told for 4 hours.

59
It was found that this cycleonly removedabout half of theporosity,and it was

believedthat alongertimeat theHIP temperatureandpressurewouldcompletelyremove

theporosity. A secondattemptwasmadeat Howmetnow holding thesamplesfor 8 hours

ratherthan4 hoursat 1288°Cand 103MPa. This increasewas successfulin virtually

eliminatingall porosity. Sincethecoolingratein theHIPautoclaveis insufficientto avoid

_/precipitation,thesamplesmustbere-solutionizedandheattreatedafterHIP'ing.

3.1.5 HydrogenCharging,AnalysisandDesorption

All hydrogen-containing samples used in this study were gas-phase charged at

Sandia National Laboratories in Livermore, Calif., under the direction of Dr. Neville

Moody. The samples were charged for 15 days at 350°C and at a pressure of 103.4 MPa

followed by a rapid cool while still under pressure. These severe conditions were

necessary due to the low diffusivity of hydrogen in PWA 1480. Gas-phase charging was

preferable to cathodic charging that has been used on similar alloys 9,1° because of the

uniform concentration of hydrogen obtained by gas-phase charging. After charging, the

samples were packed in dry ice and shipped by overnight express to Carnegie Mellon

where all samples were stored at 0°C until tested.

Hydrogen contents were analyzed by Luvak Inc. using vacuum hot extraction at

900°C. Samples were shipped to Luvak overnight in dry ice and stored at 0°C until

analyzed. At least two samples in the same condition were analyzed, and the results for the

two samples were always within 5%.

Hydrogen desorption experiments were performed by outgassing small cubes with

an edge length of 3.8 mm in air at various temperatures and times. These were gas-charged

along with the mechanical testing samples in the Sandia charging vessel. The cube size is

believed sufficiently large to account for any variability in hydrogen trapping site density.

Further, the use of two to three cubes for each analysis also provided more reliable

60
concentrationdata. A baselinewasfirstestablishedaftereachchargingrunby inunediately

analyzingthehydrogencontentof afewsamplesafterreceipt.Thelossof hydrogenat0°C

androom temperaturealso wasmeasuredasa function of time to accountfor extended

periodsatthesetemperatures.

3.2 MicrostructuralCharacterization

3.2.1 Metallography

Most metallographicspecimenswerecold mountedusingQuickmountself-setting

resin. When edge retention was important, specimenswere mounted in blue diaUyl

phthalate, which was hot-compactedin a Buehler Powermet Press. Mounting was

followed by sequentialhandgrinding through240,320,400 and 600 grit silicon carbide

papers. The sampleswere thenpolishedusing 1i.tm,0.3I.tmand0.05I.tmA1203 slurry.

To avoid streaking and dimensional distortion of the pores, it was necessary to rotate the

specimen opposite the wheel rotation while polishing. Further, to check any effect of the

A1203 slurry on pore dimensions, polishing also was done with a diamond polishing

compound. No quantitative differences resulted, so the easier method with the A1203

slurry was exclusively used.

Different etches were used depending upon which microstructural feature was to be

viewed. The best etch to use when viewing the dendritic macrostructure or eutectic 3,/,/'

was found to be 33% acetic acid, 33% nitric acid, 33% distilled water and a few drops of

hydroflouric acid applied for 10 seconds. Glyceregia (45% hydrochloric acid, 40%

glycerol and 15% nitric acid) applied for 60 seconds was the best overall general etch and

worked especially well for the "//'/' microstructure.

Samples were viewed optically on a Reichert Ultrastar metallograph using normal

bright-field illumination. Occasionally, polarized light using Nomarski interference was

61
usedto betterview thedendriticmacrostructureor eutecticT/q/,andall opticalphotographs

weretaken using Polaroid type 52 film.

3.2.2 Quantitative Metallography

Quantitative metallography was performed following the methods reviewed by

Underwood. los Direct measurements of between 50 and 100 features were taken for each

sample to determine the dimensions of the pores, eutectic _,/'_ and _/' precipitates. The

pores were determined to be fairly spherical, with an aspect ratio of 1.10. Generally, the

eutectic _'/T' were oblong, so the longest and shortest dimension were measured. The

average aspect ratio was 2.18, so the eutectic )'/T' were classified as oblate spheroids

following Underwood's nomenclature. To calculate the average size of the eutectic ),/'_', all

measurements were averaged. The edge length of the T' precipitates were measured, and

pictures were also taken to determine any T' size differences between interdendritic and

dendrite core areas, as well as regions bordering the eutectic ),/y.

A grid with spacings of 1 cm was used to measure the volume fraction and nearest

neighbor spacings of the pores and eutectic '),/_. A 200x magnification was used for all

pictures so that the grid size corresponded to the size of the object being measured, as

suggested by Underwood. A minimum of 50 measurable objects were used for

calculations, generally requiring between 10-12 photos for each sample. A point count

approach was used to calculate the volume fraction.

Due to the high volume fraction of T' precipitates, many standard quantitative

metallography methods were difficult to implement. 166 Therefore, a method was devised

that involved first counting all the _/precipitates in an SEM micrograph. Next, the total

number of possible "/' precipitates in a micrograph was calculated by using the average size

and assuming a square array of impinging _ precipitates. The actual number of _'

precipitates in the micrograph was divided by the calculated value yielding a volume

62
fraction. Thismethodwasfairly accurate,andthestandarddeviationbetweenmicrographs

of thesamesamplewasvery low.

The importanceof the spacingof the pores and eutecticy/y' in relation to the

deformationandfracturebehaviorwill bediscussedin latersections.Thegeneralequation

for thethreedimensional(3-D) nearestneighbordistance,A3, is:

A 4 N .-1/3 3 = 0.55 (,,) (3-1)

where Nv is the number of particles per unit volume.

can be written in terms of more useful quantifies:

A 3 = 0.89Ro(f) -1/'3

For spheres (pores), this equation

(3-2)

where f is the volume fraction and Ro is the diameter of the object. A similar equation can

be written for oblate spheroids (eutectic 7/"/) with an aspect ratio of two:

A3 = 0.71Ro(f) -1/3 (3-3)

Equations 3-2 and 3-3 are based upon a narrow distribution of feature sizes and

small volume fractions. The size distributions of the pores and eutectic )'/T' are shown in

Figures 3.2 and 3.3. Apparently, these distributions are not narrow enough because A3

changes little after HIP'ing or the alternate heat treatment.

Therefore equation 3-1 must be used in its basic form. Nv can be related to the

number of particles per unit area, NA, and the number of particles per unit length, NL, for

the pores and eutectic y/q/. Both of these quantities are easily measured on micrographs.

63
30

1 3 7 9 11 13 15 17 19 21 23 25 27 29 31 33 35 37 39 41 43 45

size of pores (tim)

Figure 3.2 Size distribution of pores.

30

(.)

e20

O

tt..4

o

._ 10

12 17 22 27 32 37 42 47 52 57 62 67

size of eutectic _'/T'(l.tm)

Figure 3.3 Size distribution of eutectic 3'_.

72 77 82 87

64
Theresultingequationusedfor poresis:

(N 2 _1-1/3

o. oot )

(3-4)

Likewise for the eutectic T/7' (oblate spheroids):

(3-5)

The values obtained from equations 3-4 and 3-5 were very consistent for each sample and

varied as expected for HIP'ed or alternate heat treated samples.

3.3 Mechanical Testing

3.3. I Tensile Testing

It was particularly important in the tensile samples, more so than in the fracture

toughness samples, that the loading axis orientation be as close to [001] as possible to

avoid any variations in the slip behavior. Therefore only single crystal slabs that were

determined to be within 5 ° of [001] were used in tensile testing. These tensile specimens

had a gage section of nominal length of 12.7 mm and diameter of 2.85 mm. The threads

were 1/4-20UNC, and the gage had a machined finished of 16 millionths of an inch, which

provided highly reproducible specimen-to-specimen results. Individual sample dimensions

varied slightly, so each was measured on a travelling microscope. There was little

difference found in the measured total strain whether an extensometer, usually with an 0.5

inch initial span, or the crosshead speed along with the chart speed was used to calculate

65
strain. Testingwasdoneon an Instron model 1125 screw driven machine at a crosshead

speed of 0.2 mm/min, corresponding to a strain rate of 2.6 x 10 4 sec -1.

Samples were pulled either to failure or to small plastic strains, the latter to monitor

the development of the deformation process. Uncharged samples were strained to plastic

strains of approximately either 0.0024 and 0.012. Charged samples could only be strained

to a plastic strain of approximately 0.0015 due to the low strain to failure.

Elevated temperature tensile tests were done using the same Instron machine and a

3-stage Lindberg tube furnace mounted vertically on a portable stand. The temperature was

allowed to equilibrate at least 20 minutes for each test as measured by a thermocouple

placed directly beside the center of the specimen. Inconel pull rods were used, and it was

possible to test up to 800°C.

3.3.2 Compression Testing

Compression tests were conducted on rectangular specimens measuring 5 mm on

each side and 15 mm in length. All sides were flat and parallel to within 0.013 mm over

the specimen length. Both the loading orientation and transverse orientations of the single

crystals tested were determined by Laue back-reflection to assist in determination of

subsequent slip traces. As with the tensile samples, the samples used all had a loading axis

orientation within 5 ° of [001]. Prior to testing, the specimens were electropolished using a

solution of 30% nitric acid in methanol at -25°C and 15 V for 60-90 seconds, sufficient to

remove the machined finish layer. Final dimensions were then measured using a travelling

microscope. Compression tests were performed on an Instron model 1125 screw driven

machine using a crosshead speed of 0.2 mm/min, resulting in an initial strain rate of 2.2 x

10 -4 sec "!. Strains were measured from the chart and crosshead speed, and were within

1% of dimensional changes.

66
Specimens were strained to varying levels of plastic strain to monitor the

developmentof the deformationprocess. Both unchargedand chargedsampleswere

strainedto plastic strainlevelsof 0.005,0.012and0.030.

3.3.3 Kxc Testing

Plane-strainfracturetoughnessvaluesweredeterminedasa functionof orientation

andhydrogenaccordingto ASTM E399-83.17° Compact tension specimens without side

grooves were used as shown in Figure 3.4. Specimens with the load axis near [001] were

tested at CMU, while transverse orientation specimens were tested at Sandia. Prior to

testing, samples were pre-cracked in accordance with ASTM E399-83 at alternating loads

of 1400 lbs and 50 lbs at 3 Hz. This usually required 10,000-30,000 cycles to reach the

desired crack length. After pre-cracking, specimens were hydrogen charged at Sandia in

the manner previously described. Tests on both charged and uncharged samples were

performed on a 20 kip MTS servohydraulic closed loop testing system, and were loaded

under stroke control at a displacement rate of 1.0 mm/min, corresponding to a loading rate

of approximately 100 lbs/sec. Loads and displacements were continuously recorded on an

X-Y recorder and specimen displacement was measured by a clip gage (MTS model

632.03B-20) placed in the integral notch at the specimen mouth. After testing, the fatigue

pre-crack was measured on a travelling microscope, and the load displacement results were

analyzed in accordance with ASTM E399-83. From this data, a KQ value was calculated.

No KIC samples met the thickness criteria of ASTM E399-83, and therefore it was

necessary to perform J integral tests to obtain valid fracture toughness data.

67
Figure 3.4

0.96"

1.20"

.-_ 1.60" ,

2.00" P

(:J

-!

r

F 0.50" dia.

0_0"

1.20"

One-half inch thick compact tension specimen for KIC testing.

3.3.4 JIc Testing

Fracture initiation toughness values were determined by the J-integral fracture toughness

test according to ASTM E813-87.171 Because of material limitations, the single specimen

technique was used with compact tension specimens shown in Figure 3.5. Side grooves

1.25 nun deep were machined in after pre-cracking. All tests were performed on a 20 kip

MTS servohydraulic closed loop testing system. Samples were pre-cracked in accordance

with ASTM E813-87 at alternating loads of 2200 lbs and 50 lbs at 3 I--Iz, which usually

required 10,000-30,000 cycles to reach the desired pre-crack length.

68
L 1.125" i

0.90" _._

oo, g ÷oL.

r _ o.18"

0.20" dia.

Figure 3.5 One-half inch thick compact tension specimen for JIc testing.

After pre-cracking, some specimens were hydrogen charged at Sandia in the

manner previously described. Load-point displacement was measured during the test using

an extended arm clip gage (MTS model 632.03B-30). The clip gage was attached to razor

blades along the load line, held in place by Permabond TM glue. Specimens were loaded in

stroke control at a displacement rate of 2 mm/min, corresponding to a loading rate of 33

lb/sec. During testing, the specimens were repeatedly unloaded to an amount less than

20% of the current load at intervals of approximately 0.04 - 0.09 mm crack mouth opening

displacement. Load and displacement were continuously monitored on an X-Y recorder,

with portions of the unloading and reloading curve magnified -20x using a zerosuppression circuit and recorded on a separate X-Y recorder. The compliance, C, was

obtained from these unloading slopes and the crack length, ai, was calculated using: 171

ai

= 1.00196 - 4.06319gtLL + 11.242B_L- 106.043p3L + 464.335P_L -- 650.677p 5 (3-6)

where:

P-u. (BEC) 1/_ + 1 (3-7)

6g
andW is the lengthfrom load line to thespecimenbackface,B is thethicknessandEis

Young'smodulus. The [001] value of Young'smodulusof 118GPa was used,which

wascalculatedfrom the best literaturevaluesavailableasexplainedin section 2.1.1.3.

Thisvaluecomparedfavorablywith theeffectivemoduluscalculatedfrom ASTM E813-87.

Hysterisisduring unloadingandreloadingwaskeptto a minimumby usingfiat-bottomed

clevisgripsassuggestedin ASTM E813-87.Testswerestoppedafterapproximately1.52

mm of stable crack growth, during which 12-16 unloading-reloading cycles were

accomplished.Specimensweresubsequentlyfatiguedfor a period of time to allow further

crack growth and to determine the amount of crack growth that occurred during J-integral

testing. The area under the load-deflection curves was measured on a Jandel Scientific

digitizer using a program called Sigma-Scan. J-integral values were calculated from the

equations given in ASTM E813-87. JIc values were finally calculated from a J-Aa plot

based on linear regression analysis of J-Aa points within the exclusion lines defined in

ASTM E813-87.

3.4 Deformation and Fractography Characterization

3.4.1 TEM Foil Preparation and TEM Procedure

Transmission electron microscopy (TEM) specimens were cut from tensile and

compression samples normally in the [001] orientation using a slow speed diamond cut-off

wheel. Typically, an 0.4 mm slice was cut, and then thinned to 0.12-0.15 mm by hand

grinding on 600 grit SiC paper. The introduction of dislocations was kept to a minimum

by applying only a small amount of weight on the cut-off wheel and then applying as little

pressure as possible when grinding. The punching of 3 mm discs from the thinned slices

was able to be avoided in the tensile samples as the tensile gage diameter was

:lpproximately 3 ram.

70
Theseground discs were then electropolished in a Fischione dual-jet polisher with a

solution of 55% methanol, 35% butyl cellosolve and 10% perchloric acid at a temperature

of- 20°C and voltage of 13.0 V. The latter corresponded to an amperage of between 8 and

11 mA. After polishing, the foils were rinsed in two separate methanol baths to minimize

contamination. It was found that if the '/' was preferentially polished, then the voltage was

too low and vice versa for the 1' phase. Generally, the success rate for quality foils was

only about 30-40%.

[111] TEM foils were prepared from [001] oriented tensile and compression

specimens as follows: The samples were oriented by the Laue back reflection technique

using a two-tilt axis goniometer. The sample remained attached to the goniometer, while

being cut on a South Bay Technology wire saw using 0.010 inch diameter diamond

impregnated wire blades. The slices were then thinned and polished as described above.

TEM observations were made on a JEOL 120CX electron microscope operating at

120 kV using a two-axis tilt specimen holder. Dislocation types were identified by

standard two-beam dynamical imaging. Most of the pictures were taken with the deviation

vector, sg, slightly greater than zero. Weak-beam dark field images were made using g-3g

diffraction conditions, primarily with g = <200>.

Dislocation densities were measured using the method of Bailey and Hirsch, 172

where the total dislocation line length in a given volume is measured. Due to the

inhomogeneous distribution of dislocations between the phases, this method was preferred

over one using a random line intersection technique. 173 The method used requires the

thickness of the thin foil to be measured, and a simple geometric relation was employed

utilizing the projected width of features such as dislocations or stacking faults, t73

Measuling dislocation densities is a detailed process and care must be taken not to introduce

large errors. An attempt to reduce such errors was made in this study by using an enlarger

to pr_duce images six times larger than the negative. This facilitated measurements by

allowing individual dislocations to be distinguished easier. The total dislocation line length

71
was measured on a digitizer using a program called Sigma Scan. No estimate of the total

error can be made, but the standard deviation was always within 20% of the average value,

which is not bad for these types of measurements.

3.4.2 Slip Band Analysis

Slip band behavior as a function of strain and hydrogen was quantitatively analyzed

using the measurement techniques developed by Mclnteer, Thompson and Bernstein. 174

The average slip band spacing and the degree of waviness were measured, as illustrated in

Figure 3.6. Measurements of these parameters were taken from optical micrographs, and

TEM replicas of the specimen surface were not made. This fact will likely increase the

error associated with the measurements made in this study. It was not the intent to measure

individual slip lines, but rather the collection of slip lines referred to as slip bands. For

each condition of strain and hydrogen, 5-7 photos were used resulting in measurements of

over 100 slip lines.

overaqe spoclnqw• -_ f7 Z

3

n-I

n

l:igurc 3.6 Schematic showing average slip band spacing and the degree of waviness. 174

72
3.4.3 Sectioning Experiments

Sections were made of failed tensile samples and double notch tensile samples in an

attempt to better understand the failure process and mechanisms involved. Longitudinal

sections were prepared from fractured tensile specimens, and Kjc and JIc samples. Prior

to sectioning, the fracture surface was nickel-plated to avoid damage to the surface or

rounding during polishing. Electroless nickel plating was used with a solution of: 1 liter

distilled water; 37.7 g. NiSO4" H20 (nickel sulfate); 26.4 g. NaI-I2PO2" H20 (sodium

hypophosphite); 15.9 g. CH3CO2Na" 3H20 (sodium acetate), and 5-6 drops of sulfuric

acid. The fracture surface and sample were first cleaned in acetone using an ultrasonic

bath, followed by immersion for I-3 hours in the above solution heated to 82°C. This

yielded a deposit of a few micrometers over the entire sample. The samples were then

sectioned on a slow speed diamond cut-off wheel and mounted using the blue diallyl

phthalate, as previously described. Samples were observed on the optical microscope

specifically for the presence of sub-surface cracks.

Double-notch tensile specimens were employed to look for cracks in the material

that may have occurred during deformation and to attempt to observe events just prior to

fracture. Samples were prepared and tested in the same manner as smooth cylindrical

tensiles, except that two 60 ° notches were machined in the tensile gage with a root radius of

0.025 mm. The notches were each 0.25 mm from the specimen center and were 0.70 mm

deep. After fracture, the piece containing the unfractured notch was sectioned

perpendicular to the loading axis using a slow speed diamond cut-off wheel. After

sectioning, the sample was mounted as previously described and examined optically and on

the scanning electron microscope.

73
3.4.4 Pre-Strain and Pre-Charge Tensile Tests

A series of special tensile tests was performed to investigate whether hydrogen was

affecting the deformation process, the fracture process, or both. Two groups of samples

were prepared from the same single crystal and tested as follows:

Group A

Strain to -50% of Ef

Charge with Hydrogen

Pull to Failure

Group B

Charge with Hydrogen

Strain to -50% of _f

Remove all Hydrogen by Outgassing

Pull to Failure

Group A tests were designed to determine if hydrogen affected the fracture process, and

Group B to determine if hydrogen affects the deformation process. Tensile tests were

performed in the same manner as described in section 3.3.1. Hydrogen eharging was

performed using the same conditions as in section 3.1.5, and subsequent outgassing was

done at 300°C for 75 hours, a temperature and time determined to result in less than 5 ppm

hydrogen remaining in the sample.

3.4.5 Quantitative Fractography

Quantitative fractography was performed on tensile samples, as well as on KIc and

JIc samples. Primarily, the area fraction of the eutectic 3'/7' and porosity was measured.

Typically, 15 micrographs from each sample were randomly taken on the SEM at a 400x

magnification. This was higher than the 200x magnification used in quantitative

metallography, because the features were harder to identify on the fracture surface.

Because of the higher magnification used, a grid spacing of 2 cm was employed for the

point counting method of determining area fractions. To check the repeatability of this

method, matching fracture surfaces were analyzed, and the quantitative fractography results

74
for both surfaces was within 5%. Changes in the pore size also were recorded by direct

measurement.

3.4.6 Crystallographic Fracture Orientations

The crystallographic orientation of facets on the fracture surface was determined for

cleavage facets and fracture ridges, as described below. Since all samples were single

crystals loaded in the [001] orientation, analysis of facet orientation was greatly simplified.

Due to the small scale involved, x-ray analysis of the fracture planes could not be used.

While it was not necessary to use the detailed two-trace analysis, 175 as for samples of

unknown orientation, the orientation of some cleavage facets and fracture ridges using twotrace analysis was done as a double-check of the techniques primarily used. As a prelude

to all techniques, the loading axis orientation and, in some cases, the transverse orientation

was determined by Laue back-reflection.

From prior observation, it was known that the cleavage facets were approximately

perpendicular to the [001] loading axis. Therefore, it was assumed that the cleavage facets

were on the {001 } planes, and the goal was to prove this. First, a series of micrographs

were taken of several cleavage facets at different sample angles. The sample was tilted until

the beam direction was parallel to the sample sides to give a starting reference point. From

this point, micrographs were taken at various known angles to the reference point.

The first technique used micrographs of facets at the reference angle and tilt angles of 45-

60 °, and Figure 3.7 shows a schematic of this technique. For demonstration purposes, the

[001] facet in the schematic is shown at a much larger scale than is realistic. Measurements

are made on the cleavage facets at 0 ° tilt and at 60 ° tilt, for instance. If xcos60 ° = x', then it

is concluded that cleavage facet x was perpendicular to the reference point. If this equality

doesn't hold, thcn the angle of cleavage facet x from the reference point can be calculated

75
Electron Beam Direction l

-Jb-,v,

Ixl

Figure 3.7 Schematic showing method to determine cleavage facet orientation, x is the

cleavage facet width and x' is the projected width of the cleavage facet (not

to scale). If xcos60 ° = x', then x was perpendicular to the beam initially.

using the Laue analysis of the sample's orientation. If the angle from the reference point is

equal to the angle that the sample varies from [001], then the cleavage facet is exactly

[001]. Otherwise, it is inclined to [001] by the difference between the two.

Another simple method can be applied to the cleavage facets, due to their apparent

orientation relation with the loading axis. Once again, the beam is aligned perpendicular to

the samples sides, and a series of micrographs taken at various angles. In this ease,

however, the angles are in increments of 2.5 ° up to a total angle of 15 °, and the sample is

rotated 90 ° after each series of micrographs. The projected width of the cleavage facet is

measured, and in this way it can be determined exactly when the facet is perpendicular to

the beam. Once again, this is compared to the Laue data to give the exact orientation of the

cleavage facet.

Lynch reports that crack growth in fcc cleavage sometimes has a preferred

orientation, 176'177 so this possibility was explored. This was done on KIc specimens

where the transverse orientation of the sample was known. By simply turning the

76
specimen and recording the angle to the specimen side, it was possible to determine the line

direction of the river markings and therefore the crack growth direction of the cleavage.

The orientation of the periodic ridges that appear on the fracture surface were

analyzed in a similar simple manner as described above and also by a

stereophotogrammetric analysis given by Wert and Robertson. 178 From the literature, it

was suspected that these fracture ridges would be { 111 } facets. 97'99a°°

As described above, the beam was aligned parallel to the sample sides. If the facets

were of type { 111 }, then there would be a 45 ° angle to the [001] loading axis. Therefore,

the sample was tilted 45 ° to the beam direction to determine if the facet was indeed now

perpendicular to the beam. This was done in the same manner as described above, by

tilting in small increments and measuring the projected width of the facet.

The second technique used was a stereophotogrammetric analysis method

developed by Wert and Robertson, which requires that three variants of the crystallographic

fracture plane exist on the fracture surface. This condition is satisfied with the fracture

ridges. The technique is similar to the one described above, except it is more general

because the specimen loading axis orientation need not be known. First, a suitable area of

the fracture surface must be chosen so that three variants of the ridges are present and can

be tilted through large angles. One of the variants is aligned parallel to the beam, and then

the sample is tilted in increments of 10 ° up to a total tilt of 60 °. From the given geometry

and known orientation relations between { 111 } planes, it can be determined if the fracture

ridges are indeed of the { 111 } type.

3.4.7 Fractographic Analysis Techniques

Various methods were used in an attempt to better understand the relationship

between the fracture and the microstructure. The technique of plateau etching 179 was

employed to directly observe any such relationship in tensile samples, as shown

77
schematically in Figure 3.8. First, the fracture surface was cleaned in acetone using an

ultrasonic bath. This was an important step in order to assure better adhesion of the stopoff lacquer that was applied to part of the fracture surface. The rest of the sample was

electropolished at -25°C using a solution of 30% nitric acid in methanol and applying a

voltage of 15V for 1 minute. This was followed by application of a chemical etch to bring

out the microstructure. Usually this etchant was 33% Nitric acid, 33% acetic acid, 33%

distilled water and a few drops of hydroflouric acid.

Figure 3.8 Schematic showing geometry of fracture surface-microstructure interface

in plateau etching.

After etching, the stop-off lacquer was removed by immersion in an ultrasonic bath of

acetone. Observation of the fracture surface-microstructure interface was done on the

SEM. Typically, the fracture surface was tilted 20-40 ° to the electron beam to yield the best

view of the interface.

Another method used was to etch tensile fracture surfaces in an attempt to better

determine the relationship between the "t/'Y' microstructure and the fracture process. Areas

were chosen and photographed on the SEM before etching to produce a reference

micrograph. Glyceregia applied for 50-60 sec. was generally needed to observe any

differences on the fracture surface. SEM micrographs were taken, and the fracture surface

features were examined for relationships between the ductile voids and the ),/)/, as well as

for crystallographic infonnatioi1.

78
In order to view the fracture surface at higher magnifications, replicas were made

for examination on the TEM. It was determined that two-stage replicas gave the best

results. Cellulose acetate tape was placed on a clean tensile fracture surface after first

softening the contact side with acetone. After the tape hardened, it was carefully stripped

and shadowed at a 30 ° angle with Pt-C on a Veeco evaporator. This was followed by

vertical deposition of carbon onto the tape. The tape was dissolved in successive baths of

acetone, leaving the floating carbon replica. It was important to remove all traces of the

cellulose acetate because of its detrimental effect on the cleanliness of the TEM. The

remaining carbon replica was placed on a copper grid and examined in the JEOL 120CX

TEM.

79
CHAPTER 4

RESULTS

4.1 Microstructural Characterization

PWA 1480 has a dendritic macrostructure with porosity and eutectic _,P/' present in

the interdendritic region. The macrostructure of samples given the normal heat treatment

(sec. 3.1.3) was examined in the [001] and transverse orientations to determine any effect

of growth direction on the size or shape of the porosity or of the eutectic _,/_'. As the

transverse orientation was not controlled during growth of these bars, it can vary randomly

along the <010> - <110> line. The orientation of all bars was given in Table 3.1. All

material was from the same batch provided by NASA, except a small amount received from

Rocketdyne at a later date. Table 4.1 gives the average values for the pertinent

macrostructural features of all of the NASA bars examined; the full table of values can be

found in Appendix A along with the standard deviation of all measurements. A3 is the

three-dimensional nearest neighbor distance as explained in section 3.3.2, and is a useful

quantity in interpreting the effect of microstructure on the deformation and fracture

behavior. The aspect ratio is the largest diameter divided by the smallest and was measured

to assist in subsequent quantitative calculations. There was no effect of orientation on the

macrostructure, except for a slight effect on the average size of the eutectic _,/'y'. This

difference is probably related to the solidification characteristics of the eutectic _,P/'. The

data for the bars from Rocketdyne are given in Table 4.2, and it can be seen that this

material contained less initial porosity and eutectic _,/]t. All future references will be made

to the NASA material unless specifically noted otherwise.

The dendritic macrostructure also was quantitatively analyzed, and the following

average dimensions were found: core diameter - 150 lain; arm diameter - 400-500 l.tm, and

spacing - 200 _ma. Figure 4.1 shows the dendritic structure along with the interdendritic

80
Table4.1

[001]

Trails

Quantitative metallography results of unHIP'ed samples given the normal

heat treatment. [001] orientation and transverse orientation averages shown.

Pores

Average

Size (gm)

14.6

14.9

Volume A 3 Aspect

Fraction (%) (I.tm) Ratio

1.02 165.0 1.13

1.04 152.5 1.12

Eutectic 7/7'

Average

Otm)

32.7

27.4

Volume A 3 Aspect

Fraction (%) (pro) Ratio

5.60 71.9 2.19

5.90 67.3 2.09

Average

Table 4.2 Quantitative metallography results of material received from Rocketdyne.

Average

Size (Bm)

9.1

Pores

Volume

Fraction (%)

0.23 250.0

Eutectic T/7'

Average

Size (grn)

25.9

Volume

Fraction (%)

2.09

A3

(Bm)

150.9

81
ORIGINAL PAGE

BLACK A_D WHIT_ PIiO.T_OGRAPJ_

ORIGINAL PAGE IS

OF POOR QUALITY

Figure 4.1

(a) (b)

(a) Dendritic macrostructure of PWA 1480 and (b) interdendritic porosity and

eutectic 7/Y.

Figure 4.2

1.sum i

(a) (b)

"y/_ microstmcture of PWA 1480; (a) inside dendrites and (b) in the

interdendritic region.

82

1.5 I.tm
porosity andeutectic_'/T'. It wasdifficult to positively ascertainwhethertherewasany

relationbetweenthepositionsof theeutecticT/_/'andtheporosity. It wasdiscussedthatthe

eutecticT/'_mayactasnucleationsitesfor thepores,16andthuseachporewouldbenextto

aneutecticT/_/'.This may havebeenthecase,but it wasnot possibleto determinethis

from a seriesof 2-D micrographs. It was determinedthat becauseof the muchlarger

volume fractionof eutecticT/_/',eacheutecticT/T'wasnot necessarilyassociatedwith a

pore. It alsowasmentionedin section2.1.1.1.that it waspossiblefor theeutecticT/q(to

havea lacelike-appearanceor a more featurelessone. The vastmajority of eutecticT_

observedwereessentiallyfeatureless,and basedon otherstudies15this suggeststhatthe

eutectic_'h/'consistsmostlyof they phase.

Theeffectof thedendriticstructureon theresultingT/q(microstructurewasstudied.

Figure 4.2ashowsa representativeT/_ rnicrostructurein the dendritesand Figure4.2b

showsthe microstructurein the interdendriticregion. Table 4.3 givestheresultsfor the

average"_ precipitate size and volume fraction at various positions relative to the dendrites.

While there was no effect of the dendritic structure on _/' volume fraction, the _/' size was

clearly larger in the interdendritic region. Also, inside the dendrite core and arms, the "{

precipitates were of uniform size, but a bi-modal distribution of T' was found in the

interdendritic region with about 3% of the _' precipitates having an edge length of

Table 4.3 Quantitative metallography results of T/_" microstructure.

Sample

4 hr#1

4 hr#2

4 hr #3

4 hr#4

Location of measurement

interdendritic area

interdendritic area

dendrite core

dendrite arm

Average size (lam)

0.485

0.478

0.416

0.402

Volume fraction

74%

71%

70%

72%

83
approximately1I.tm,andtheremainingy beinglessthan0.5I.tm. However,eventaking

thisbi-modaldistributionintoaccountit cannotexplainthesizedifference,soit is apparent

thatthetypical'y' in the interdendritic region is indeed slightly larger than those found in the

dendrites. No quantitative metaUography was performed on the q( near the eutectic y/q(,

but often there were 1 I.tm q( within a few I.tm of the eutectic y/q(. The volume fraction of

y' precipitates in Table 4.3 is slightly larger than other published values for PWA

1480, 7'76'111 and the average size is slightly less. These variations could be due to

differences in compositions, heat treating practice or measurement techniques. For

completeness, it should also be mentioned that the T/q( misfit has been reported to be

0.28%. 122

4.1.1 Alternate Heat Treatment

It was established that additional time at the solution heat treatment temperature not

only reduced compositional segregation but also dissolved most, if not all of the eutectic

'),/y phase. For the as-received material, it was found that 20 hours at 1288°C was

sufficient to eliminate almost all of the eutectic y/q(. This step was followed by the normal

two-stage aging heat treatment, and this entire heat treatment will be subsequently referred

to as the alternate heat treatment. In fact, this term has been generalized to refer to any heat

treatment that removes the eutectic y_ and is not limited to solutionizing times of 20 hours.

Table 4.4 shows the quantitative metallographic results of samples solutionized for the

normal time of 4 hours and for those solutionized for 20 hours. While no quantitative data

were taken at points between these two times, it was observed that the eutectic y/q( began to

break up after about 8-10 hours. There was no effect of solutionizing time on the 7/q(

microstructure of the samples studied.

A potential drawback to solutionizing for extended periods of time is incipient

melting, which can occur with quite small fluctuations in temperature. In the set of samples

84
studiedin Table4.4, incipientmeltingoccurredin all samplesheld longerthan20 hours,

probablydue to a temperaturesurgethatcanoccurwhenthe furnacedoor is closedafter

removingsomeof thesamples.Otherpossibleexplanationswhich we believelesslikely

involve local concentrationchangesthat could induceincipient melting. The former

rationale could explain why incipient melting is more prevalent in samplesthat have

repeatedlybeenexposedto this temperaturesurge. In supportof this, separatesamples

were held at the solution temperaturefor over 50 hourssuccessfullywith no incipient

melting, solong astheheattreatmentwasnotinterruptedby openingthefurnacedoor. In

futureapplicationsof thealternateheattreatment,carefulexperimentalcontrolwasfollowed

andunlessotherwisestated,sampleswereheldfor 20hoursat 1288°C.

Table4.4 QuantitativemetaUographyshowingresultsof alternateheattreatment.

Pores Eutectic_,/q(

MaterialCondition

NormalHeatTreatment

AlternateHeatTreatment

i

Average

Size 0.tm)

14.6

15.5

Volume

Fraction (%)

Average

Size (l.tm)

31.9

10.0

Volume

Fraction (%)

5.3

0.48

4.1.2 Hot Isostatic Pressing

The removal of porosity by Hot Isostatic Pressing (IMP'ing) has been used by us

and others 23'28 to test the effect of porosity on tensile properties. The difficulties

associated with HIP'ing these alloys is detailed in section 2.1.1.4, however, the cycle used

in this thesis is believed to have overcome these problems. Table 4.5 shows the average

quantitative metallographic results of the first ItlP'ing cycle attempted and a second morc

successful attempt, wflh a hold timc of 8 hours at the final temperature and pressure rather

85
Table4.5 Quantitativemetallographyfor HIP'edsamples.

Average

Size(_m)

no HIP 14.4

Pores

Volume

Fraction(%)

0.99 174.2

Average

Size(_tm)

33.4

Eutectic),/_{

Volume

Fraction(%)

5.25

HIP 1

HIP 2

8.3

4.3

0.57

0.035

206.5

398.4

19.9

15.8

3.04

1.43

A3

(l.tm)

78.0

101.8

147.4

Figure 4.3

(a) (b)

Comparison between the macrostructure of (a) the as-received material and

(b) the HIP'ed material given the alternate heat treatment.

ORIGINAL PAGE IS

OF POOR QUALITY 86

ORIGINAL PAGE

BLACK At'_D WHITF PhCJ JOGRAp_
than 4 hours. In the table, IMP 1 signifies the first IMP cycle used and HIP 2 the second.

The average of the uniMP'ed values is reported in Table 4.5 as no HIP. Figure 4.3

illustrates the differences in porosity content between the as-received material and the

HIP'ed material that was held for 8 hours and subsequently given the alternate heat

treatment. There were no quantitative differences in the _,/y microstructure after either HIP

cycle.

Recrystallization did occur in some of the samples that had been HIP'ed, as shown

in Figure 4.4. This problem was associated with surface damage usually on the ends of the

single crystal slabs, which may have been caused by a sandblasting procedure at TRW.

Fortunately, sections for metallographic specimens had been removed from the ends of the

slabs prior to HIP'ing, and although these sections were sometimes partially recrystallized,

it is not believed that recrystallization affected any test samples taken from the slab interior.

Because HIP'ing is performed at a solution temperature of 1288°C, a decrease in

eutectic 'y/"/' size and volume fraction relative to unHIP'ed material is seen in Table 4.5.

Furthermore, a separate solution heat treatment had to be performed after HIP'ing due to

the slow cooling in the autoclave, thereby further reducing the eutectic 'y/_/' size and volume

fraction. The measurements in Table 4.5 were taken following the full heat treatment after

HIP'ing. Since the amount of eutectic 'y/_' is reduced during HIP'ing, it was difficult to

independently assess the role of pores and eutectic 3'/7'. However, Table 4.5 shows there

still was some eutectic _,/_,' left after HIP'ing, so it was possible to apply the alternate heat

treatment to these samples, and comparisons could be made between HIP'ed material with

and without eutectic _'/_. Table 4.6 gives an overview of the effect of heat treatment and

HIP'ing on the macrostructure.

87
ORIGINAL PAGE

BLACK At_D _HITE PHO.]'_OGRAP.H

ORIGINAL PAGE: IS

O_' POOR QUALITY

Figure 4.4 Recrystallization that occurred in some of the HIP'ed samples.

Table 4.6 Synopsis of the effect of HIP'ing and heat treatment on the macrostructure.

Sample condition

unHIP'ed normal HT

Average

size (_m)

14.6

Volume

fraction (%)

1.02

Eutectic Y/7'

Average

size (I.tm)

unHIP'ed alternate HT

HIP'ed normal HT

HIP'ed alternate HT

15.5

4.3

3.1

1.00

0.035

0.020

33.7

Pores

10.0

15.8

6.1

Volume

fraction (%)

5.6

0.48

1.43

0.30

88
4.2 Hydrogen Trapping Characteristics

Hydrogen trapping was analyzed by performing desorption tests at various

temperatures. To calculate the amount of hydrogen that would be left in a cube after

outgassing and assuming no traps, the following equation was used: 18°

where Co is the initial hydrogen concentration, x is the edge length of the cube, D is the

diffusion constant and t is time. At 0°C there was no measurable loss of hydrogen after 3

months, and at room temperature there was no measurable loss of hydrogen after 14 days.

Figure 4.6 shows the measured results at 150°C for PWA 1480 along with calculated

desorption profiles for similar alloys using the diffusion data described in section 2.3.1.1.

At least two samples were used for each measured point for PWA 1480, and the hydrogen

concentration differences were invariably well within 5%. To assist in interpretation, the

multiplying factor to convert from ppm to appm is 58.6 for PWA 1480. It can be seen that

the MAR-M 246 diffusion data is obviously too high, but the other diffusion data can be

used. This conclusion is supported by recent work that suggested the diffusion constant of

hydrogen in a similar single crystal, CMSX-2,1° was very close to that of hydrogen in IN

903. The desorption profile at 300°C is shown in Figure 4.6 and at 400°C in Figure 4.7.

The measured data of PWA 1480 closely followed the calculated profile obtained using the

diffusion data of IN 903 and assuming no traps.

In order to independently assess the trapping characteristics of the porosity and the

eutectic "f/_/', HIP'ed material was charged and outgassed. This IMP'ed material contained

very close to zero percent porosity, and the results for desorption at 150°C are shown in

Figure 4.8 and at 300°C in Figure 4.9. These figures are almost identical to Figures 4.5

and 4.6 showing that porosity has no effect on the amount of hydrogen that can be charged

into the material or the rate of outgassing. It is important to point out that in samples

without eutcctic T/T', it also was possible to obtain about 300 ppm hydrogen in thc

89
&

e-,

o

"m

o

o

o

0

_0

Figure 4.5

35o 300

250

2O0

150 Ni_

100

50

0

0 25 50 75 100 125

Time at 150°C (hrs)

718

-4- 903

MAR M246

-.o- Ni

1480

Desorption profile at 150°C showing the amount of hydrogen outgassed as a

function of time.

Figure 4.6

350

E

300

&

t_

o 250

.,-4

I_ 2oo

c 150

o

,- 100

._ so

_ o

k k_ "9°3 i-=- Ni

%\\k i-*-,,80

"' i I • I '_ I ' I , .I"

1 0 20 30 40 50 60

Time at 300°C (hrs)

Desorption profile at 300°C showing the amount of hydrogen outgassed as a

function of time.

90
E

&

O

¢)

O

I

30O

250

200

150

100

5O

0

_48_ 71. 8 | 0

i

o lo

! ..o._ 718

903

Ni

1480

m

! ! , •

20 30 40

Time at 400°C (hrs)

Figure 4.7 Desorption profile at 400°C showing the amount of hydrogen outgassed as a

function of time.

E

v

O

or-,I

b

¢..)

O

L)

;:=

2 °t/ \

_oo1/ \

1°°t/

0 25 50 75 100 125

718

•-*- 903

•-_ MAR M 246

--o- Ni

•-_ HIP'ed 1480

Time at 150°C (hrs)

Figure 4.8 Desorption profile at 150°C for HIP'ed samples.

91
._ 3 k' 250 -0" 718 1

,_olhX_,,_---,_,,,o I

0 10 20 30 40 50 60

Tune at 300°C (hrs)

Figure 4.9 Desorption profile at 300°C for HIP'ed samples.

92
microstructure under the same charging conditions. This implies that neither the porosity

nor the eutectic _'/7' are important sinks for hydrogen during charging.

4.3 Mechanical Properties

4.3.1 Tensile Properties

The tensile properties of PWA 1480 were tested as a function of hydrogen content,

eutectic 3'/'f' and porosity. The average results of a dozen tests of uncharged tensile

specimens given the normal heat treatment are shown in Table 4.7. The strength values are

fairly constant from slab to slab, but the strain-to-failure values can vary quite a bit as

documented in the complete results presented in Appendix A. The reduction-in-area (R.A.)

values are reported for some specimens in this Appendix, but accurately measuring this

value was difficult due to the specimen geometry and small changes involved. Appendix A

also gives the standard deviation for all tensile properties. A schematic of a typical stressstrain curve is drawn in Figure 4.10. Generally, after two percent strain the curve began to

rise again reflecting work hardening, which accounts for the slightly higher ultimate tensile

strength than yield strength in most samples. In a few tests, the curve never rose above the

yield strength again, and when this occurred, the reported ultimate tensile strength value is

the yield strength.

4.3.1.1 Effect of hydrogen

Normal charging used a hydrogen pressure of 103.4 MPa at 350°C, which resulted

in uniform total hydrogen contents of around 300 ppm (1.74 at. %). The average results of

room temperature tensile tests on such samples, containing approximately 300 ppm

hydrogen, are given in Table 4.8; the complete results for all samples tested are in

Appendix A. The most striking result compared to uncharged samples was the order of

Inagnitudc decrease in strain to f,filure for the charged specimens, as shown in Figure 4.11.

93
Table4.7 Room temperature tensile test results for samples that were unHIP'ed,

uncharged and given the normal heat treatment.

0.2% Y.S. (MPa)

Average 1112

U.T.S. (MPa)

1153

_(%) R.A. (%)

3.36 3.07

Czh

?=

o

o

t..-

m

1100

! !

0 1.5 3.0

Engineering Strain (%)

h.._

v

Figure 4.10 Typical stress-strain curve for uncharged samples.

94
Table4.8 Roomtemperaturetensiletestresultsfor samplesthatwereunHIP'ed,charged

andgiventhenormalheattreatment.

Average

0.2 % Y.S. (MPa)

1001

U.T.S. (MPa)

1011

8f(%) R.A. (%)

0.26 2.40

¢/J

eu0

r-,

O

O/)

1100

\

Hydrogen charged samples

! r

0 1.'5 3.0

Engineering Strain (%)

Figure 4.11 Typical comparative stxess-strain curves for uncharged and charged specimens.

95
The R.A., however, wasonly slightly affectedby hydrogen,but the yield stresswas

reducedby almost10%.

The hydrogencontentof tensilespecimenswasvariedby changingthe hydrogen

gaspressureduring charging or by desorptionof the samplesafter normal charging.

Different charging pressuresof 13.8MPa, 69.0 MPa and 138MPa were usedto now

examinetheeffectsof differenthydrogenconcentrationson thetensileproperties. Table

4.9 showsthereis little effect on the strength,and Figure 4.12 illustrates the effect of

hydrogenconcentrationonthestrainto failure. Theaveragevaluesfor samplescontaining

nohydrogenand300ppmhydrogenareincludedin Figure4.12.

Thehydrogencontentof tensilesamplesalsowasvariedby degassingsamplesthat

initially containedabout300 ppmhydrogen. Table4.10showsthetensilepropertiesfor

samplesthat were degassedand thentestedat 300°C. Onceagainthereis little effectof

hydrogenconcentrationon the strengthor R.A. Table 4.10 andFigure 4.13 show the

decreasingstrain to failure with increasinghydrogencontent. The end points of 0 ppm

hydrogenand300ppmhydrogenin thefigurearetheaveragevaluesof roomtemperature

tests.Sinceit hasbeenshownthat theductility of unchargedsamplesonly changesslightly

from 0°C to 300°C, 77 the introduction of these points should not introduce large

uncertainties. Comparison of the curves in Figures 4.12 and 4.13 show them to be similar

in shape, except that the minimum ductility occurs around 300 ppm in Figure 4.12 and 200

ppm in Figure 4.13.

96
Table4.9 Roomtemperaturetensiletestresultsfor samples charged under different

hydrogen pressures and tested in air.

Charging Pressure

uncharged

13.8 MPa

69.0 MPa

138 MPa

H Conc.

0

149

202

507

0.2% Y.S. (MPa)

1112

1032

1035

1022

U.T.S. (MPa)

1153

1069

1068

1027

ee (%)

3.4

2.5

2.0

0.26

R.A. (%)

3.1

2.4

2.1

1.9

3

2

¢,o 1

! I " I ! •

0 100 200 300 400

Hydrogen concentration

rl

!

500 600

Figure 4.12 Strain to failure as a function of hydrogen concentration for samples

charged under different hydrogen pressures.

97
Table4.10 300°Ctensiletestresultsfor samplesdesorbedto different hydrogen

concentrations.

Hydrogen Conc.(ppm)

55

135

159

195

0.2%Y.S. (MPa)

1048

1030

1070

1066

U.T.S. (MPa)

1100

1060

1118

1114

ef(%)

2.2

1.2

0.9

0.4

R.A. (%)

2.5

2.3

2.6

2.1

2

! !

o 1oo 200

Hydrogen Cone. (ppm)

300

Figure 4.13 Strain to failure as a function of hydrogen concentration for samples that

were outgassed for various times.

98
4.3.1.2 Effectof eutectic7/7'

After it wasestablishedthat the eutectic7/7'could beeliminatedby anextended

solutionheattreatment,a systematicstudywasundertakento determineitseffectontensile

properties. Tensile sampleswere held for solution timesranging from 4 to 29 hours,

followed by the normaltwo-stageagingtreatment. Theresultsaregraphically shownin

Figure 4.14, which shows that while removal of the eutectic 'Y/T' did not change the

yielding behavior, there was a peak in both the U.T.S. and strain to failure at 20 hours

solution time. The work hardening parameter, n, inexplicably increased from 0.04 at 4

hours to 0.08 at 20 hours.

Metallography was performed subsequently on the tensile thread section of the

specimens, and it revealed that the eutectic 7/7' was broken up at 15 hours and virtually

eliminated at 20 hours. It also was found that incipient melting occurred in the samples

held for 24 and 29 hours but not for shorter times. Therefore, it appears that the increased

ductility and work hardening benefits derived from the removal of the eutectic 7/7' are

offset by the onset of incipient melting above 20 hours. This time is optimum under these

conditions to remove the eutectic 7/7' and improve tensile properties, and while longer

times probably would not be detrimental, proper care must be taken during heat treatment.

Tensile tests also were performed on round bars received from Rocketdyne that

were processed so that they contained less initial porosity and eutectic _,/_, as reported in

Table 4.2. The average values of tensile tests are shown in Table 4.11 for samples given

both the normal heat treatment and those given the alternate heat treatment to remove the

cutectic 7/7'.

99
MPa

1400

1200

1000

8OO

6OO

4OO

2OO

0

I_l_l--l--I

-.......

I i I I

4 15 2O 24 29

Hours at Solution Temperature

8

7

6

5

4

3

2

•o- UTS

a- Yield

4- Strain to Failure

Strain to Failure

Figure 4.14 Effect of time spent at solution temperature on tensile properties.

Table 4.11 Effect of eutectic 7/3/on tensile properties for Rocketdyne material.

Sample

normal heat treatment

alternate heat treatment

0.2% Y.S. (MPa)

1138

1146

U.T.S. (MPa)

1151

1191

el(%) R.A. (%)

100
The tensile curves in this batch were similar to those in Figure 4.11, except that those

samples without eutectic 3'/'_ exhibited higher uniform strains, which produced a higher

UTS. The work hardening parameter, n, was not significantly different. The strain-tofailure differences due to heat treatment were not as dramatic in Table 4.11, as they were in

Figure 4.14, probably because of the lower amount of initial eutectic _,/q( in the Rocketdyne

samples of Table 4.11.

4.3.1.3 Effect of porosity

The effect of porosity on tensile properties was studied at three different porosity

levels: an initial level of about 1.0%; the level after the first HIP of 0.57%; and the level

after the second HIP of 0.035%. Table 4.12 shows the tensile properties for these porosity

levels; the no HIP results are different because the samples came from different single

crystal slabs, and are shown so as to remove any effect of orientation on the tensile

properties. It was found that there was no effect of porosity on any of the tensile

properties, except for an increase in work hardening parameter from 0.04 to 0.07 after the

f'trst HIP attempt, absent in the second HIP attempt.

Table 4.13 gives the results of a series of tensile tests designed to independently

determine the effects of hydrogen, porosity and euteetic T/Y using samples from the same

single crystal bar. The results support all of the observations in this section, namely the

lower yield stress in charged samples and an improved ductility in uncharged and charged

samples without eutectic Y/7'. Further, hydrogen had less of an effect on the R.A. than on

the strain to failure, and there was no effect of porosity on the tensile properties.

lol
no HIP

HIP1

no HIP

HIP2

Table 4.12 Effect of porosity on tensile properties.

0.2 % Y.S. (MPa)

1072

1100

1138

1118

U.T.S. (MPa)

1100

1233

1151

1161

ef (%)

6.8

7.4

3.6

3.9

R.A. (%)

6.4

8.7

3.7

4.2

Table 4.13 Effect of hydrogen, porosity and eutectic T/7' on tensile properties of one bar.

unHIP'ed, norm HT

unHIP'ed, alt HT

HIP'ed, norm HT

HIP'ed, alt HT

unHIP'ed, norm HT, w/H

unHIP'ed, alt ttT, w/H

ttlP'ed, norm HT, w/It

ttlP'ed, ah HT, w/H

0.2 % Y.S. (MPa)

1138

1146

1118

1140

985

1008

1013

1020

U.T.S. (MPa)

1151

1191

1161

1201

985

1008

1046

1051

ef (%)

0.20

1.3

R.A. (%)

102
4.3.2 CompressionProperties

Compressionsampleswere usedto obtain higher amountsof strain in charged

samples,sothat thedeformationbehaviorcouldbestudiedat comparablestrainlevelswith

andwithout hydrogen. Comparativebehaviorof thesesamplesis coveredin section4.4.

Theopportunitywas taken to a/so investigate the reported tension/compression flow stress

asymmetry 88,9°,116 and to extend this as a function of hydrogen. Table 4.14 lists the

average yield strength of compression samples with and without hydrogen, along with the

corresponding tensile yield strengths from the same slab, #16. The observed asymmetry

supports previous claims s8.9°.]]6 that the tensile yield stress is slightly higher than

compression in the [001] orientation. Further, it is now established that the presence of

internal hydrogen does not affect the magnitude of the asynm'_try.

Table 4.14 Effect of hydrogen on the tension/compression asymmetry.

Oys (tension)

Oys (comp.)

T C

(_ys/(_ys

no hydrogen

1135

1076

1.05

with hydrogen

1100

1060

1.04

4.3.3 Fracture Toughness

4.3.3.1 Ktc results

The effect of hydrogen on fracture toughness was studied using KIc tests on

compact tension specimens. [001] uncharged samples were found to have a fracture

toughness value of 122 MPa',/m, while samples with about 300 ppm hydrogen had a

toughness of 103 MPa',]m, a decrease of 16%. Each of these values is the average of two

tt'SlS;, all(] the fracture totJghness values ;_rc similar to those reported in other nickel-base

103
alloys and iron-base superalloys (sec. 2.4). The fracture toughness tests passed all validity

criteria in ASTM E-399, except for the thickness criterion. In order to be considered valid,

the 1.27 cm-thick specimens would had to have been just over 2 cm thick. However,

based upon valid short rod fracture toughness tests reported in Appendix C, it was believed

that the obtained KQ values were close to the true fracture toughness.

4.4.3.2 hc results

JIc tests were subsequently conducted to obtain valid fracture toughness data as a

function of hydrogen, porosity and eutectic T[_, as shown in Table 4.15. All of the values

in Table 4.15 are the average of two tests, and all samples were taken from the same slab.

The hydrogen content of charged samples was approximately 300 ppm, and the initial

porosity and eutectic 'y/_ volume fractions were 1.07°)'o and 5.24%, respectively. As a

result of HIP'ing and the alternate heat treatment, these values were then reduced to

essentially zero and these changes were assessed. Table 4.15 shows that introduction of

hydrogen had a minimal effect on JIc, in contrast to the definite decrease observed in

toughness in charged KQ samples. This point will be discussed later in terms of the

differences in testing techniques. There does not appear to be any effect of eutectic T/'_/'on

the toughness, but eliminating the porosity did appear to slightly increase the toughness.

Due to the presence of stable crack growth in this material, the tearing modulus was

calculated using: 181

T = _-_ (4-2)

where dJ/da is the slope of the J-Aa curve using points within Aa < 0.06bo, and bo is the

initial remaining ligament length. Calculation of the tearing modulus was restricted to this

rcgion of Aa, since it has been reported that J-controlled crack growth does not extend

104
Table 4.15 J-integral fracture toughness values (kPa.m) as a function of hydrogen,

porosity and eutectic 7/'Y'-

normal heat treatment

(with eutectic T/"/')

alternate heat treatment

(without eutectic T/Y)

with pores

w/hydrogen

129

w/o hydrogen

128

no pores

w/hydrogen w/o hydrogen

130 134 140 148

Table 4.16 Tearing modulus values as a function of hydrogen, porosity and eutectic y/_'.

normal heat treatment

(with eutectic T/7')

alternate heat treatment

(without eutectic T/Y')

with pores

w/hydrogen

9.6

12.4

w/ohydrogen

7.2

8.9

no pores

w/hydrogen

13.0

w/o hydrogen

17.8

105
beyondthisregion.181Table4.16showsthe results of the tearing modulus calculations in

dimensionless units. Once again the effect of hydrogen is inconclusive, there was no effect

of eutectic T/7', and the elimination of porosity improves this measure of toughness.

4.4 Deformation Behavior in Tension and Compression

4.4.1 Development of Deformation Structure

Room temperature tensile tests were interrupted at various plastic strain levels to

monitor the development of the deformation structure. The as-received microstructure

contained very few dislocations, as expected, likely because of the low reported '//'¢

misfit. 122 Micrographs of [001] foils in Figure 4.15 illustrate the development of

deformation behavior in uncharged samples. At the smallest plastic strain level examined,

Ep = 0.24%, loosely associated a/2<110> dislocation pairs and single dislocations are seen

in the Y matrix, as determined by conventional g-b anaiysis. 182 At this strain, the

deformation was inhomogeneous across the foil, and isolated deformations bands appeared

to be the primary deformation mechanism. In these narrow bands, there was shearing of

the 7' precipitates by a/2<110> dislocation pairs. Densities were not measured in these

bands because it was difficult to obtain accurate results due to the heavy activity in this

region. However, the activity outside of these bands and in other regions consisted

primarily of dislocations in the 3' matrix with little activity in the 7' precipitates.

The eutectic y/7' was rarely observed in the thin foils, but when it was present, little

dislocation activity inside the constituent was observed. The eutectic T/7' phase appeared

featureless, without even the expected lamellae-type structure. Figure 4.16 shows the

interface between the y/y microstructure and the eutectic 3'/7' in an uncharged tensile sample

with 0.24% strain. There is a layer of 3' matrix surrounding the eutectic 3'/7' that contains a

measured dislocation density of 5.2 x 10 )° cm 2. This high density of dislocations at the

106
ORIGINAL PAGE

BLACK. AND WHITF. PHOTOGRAPH

0.15 t.tm

(a) (b)

Figure 4.15

0.30 I.tm

(c)

[001] foils showing deformation structure as a function of plastic strain in

uncharged tensile samples. (a) _?.p= 0.24%. (b) gp = 1.2%. (c) gf = 3.0%.

107
eutectic"y/7' interface was commonly observed in all samples.

At a higher plastic strain of 1.2%, as in Figure 4.15b, the deformation became

much more uniform across the entire foil, although concentrated to the 7 matrix, with very

few dislocations in the 7' precipitates. At this strain, there is no longer any evidence of the

local deformation bands seen at the smaller strain. Dislocation densities were measured in

each phase, and it was found that the "/matrix contained 4.5 x 10 l° dislocations/cm 2 and

the 7' precipitates had 4.0 x 109 dislocations/cm 2. Figure 4.17 is a <111> foil at Ep =

1.2%, suitably oriented to show dislocations on the slip plane, and it can be seen that the

dislocations are bowing around and between the 7' precipitates, while attempting to

maintain a screw orientation. There was no evidence for shearing found in these <111>

foils implying that the deformation mechanisms in PWA 1480 are different than in most

superalloys. It was previously mentioned that the 7' precipitates were slightly larger in the

interdendritic areas with some precipitates as large as 1 I.tm; however, there did not appear

to be any effect of the 7' size on the deformation behavior at any strain level.

At the failure strain of 3.0%, Figure 4.15c, the dislocation density in the y matrix

was 1.0 x 1011,with a slight amount of ),' shearing observed. The density in the

precipitates at this strain was 4.9 x 109. The trapping of dislocations in the y matrix was

still uniform across all of the foils. The dislocation density measurements show that the

number of dislocations continue to increase steadily in the matrix with strain, but the

dislocation activity in the precipitates does not increase as much with strain.

108
ORIGINAL PAGE

BLACK AND WHITE PHOTOGRAPI.t

Figure 4.16

0.50 izm 0.50 I_m

(a) (b)

Interface between eutectic T/_' and ),/_ microstructure showing dislocation

tangles at interface. _ = 0.24%.

Figure 4.17

0.50/_m

<1 11> foil in uncharged tensile sample showing dislocations bending

around f precipitates, _'_p= 1.2%.

109
Severaldifferenta/2<110>Burgersvectorswerefoundindicatingthatmultipleslip

wasoccurring,asexpected.Occasionally,superlatticeintrinsicstackingfaults(SISF)were

observedin theT' precipitate,asseenin Figure 4.18 anddeterminedby comparisonsof

bright field anddark field TEM micrographs.182Thesewere not observedto the same

extentin eachfoil, andtheyneverappearedto play amajorrole in theoveralldeformation

process.

The deformationbehaviorof compressionsamplesalsowasexaminedat various

plastic strains,and Figure4.19 showstheprogressionof deformationin thesesamples.

Therewasno significantdifferencesobservedin thedeformationbehaviorbetweentension

and compression samples. In both at low strains, Figure 4.19a, inhomogeneous

deformationoccurredin narrowbands,while athigherstrainsthedeformationwasuniform

acrossthefoil. As in tension,dislocationmotionwasprimarily restrictedtothe"1'matrix at

higher strains, and this appeared to be the main deformation mechanism.

At a plastic strain of 3.0%, Figure 4.19c and 4.19d, a slight amount of activity was

observed in the 'f precipitate with a very high dislocation density in y matrix. It was also

observed, Figure 4.19c, that some 'y' precipitates contained a moderate density of

dislocations, while others contained very few, if any, dislocations. This observation also

was made in tensile samples at 13f= 3.0%. Various a/2<110> single dislocations and paired

dislocations were analyzed, again indicating multiple slip, and it appeared that SISF's in the

_' precipitate occurred to even a lesser extent than in tensile samples.

110
ORIGINAL PAGE

BLACK AND WHITE PHOT.OGRAP.H

If

Figure 4.18 Superlattice intrinsic stacking faults occasionally observed in y precipitate.

111
ORIGYNAL PAGE

BLACK AI'_D WHJTE PJ-_OI.OGRARH

(a) (b)

Figure 4.19

0.25 _tm

(c) (d)

Deformation structure of uncharged compression samples at various

strains. (a) ep = 0.5%. (b) ep = 1,3%. (c & d) gp = 3.0%.

112
4.4.1.1 Effect of hydrogen

The introduction of internal hydrogen per se did not visibly affect the microstructure

or introduce any dislocations. Tensile samples containing approximately 300 ppm

hydrogen were pulled to plastic strain levels of 0.15% and to the failure strain of 0.30%.

Figure 4.20 shows the deformation behavior for both strains in [001] oriented foils. At Ep

= 0.15%, there was inhomogeneous deformation, as shown in Figure 4.20a and 4.20b,

similar in appearance to the uncharged samples at Ep = 0.24%. In regions outside these

narrow bands of deformation, single dislocations and paired dislocations were observed

exclusively in the ymatrix. At the failure strain of 0.30%, there was slightly more uniform

dislocation activity, Figure 4.20c, but it appeared that the majority of deformation still was

occurring in these thin bands, as shown in Figure 4.20d. These bands still were the only

area in which sheafing of the y precipitates occurred. Dislocation densities were measured

at this strain, but the average values are misleading because of the inhomogeneous

deformation. The density in the matrix was 8.9 x 109 cm -2 and in the precipitate it was 4.1

x 109 cm -2. These measurements do show that the average dislocation density in the y

matrix of the failed, charged tensile samples was very much lower than was observed in

failed, uncharged tensile samples pulled to higher strains, and this can be seen by

comparing Figures 4.15c and 4.20b.

Burgers vectors of various slip systems were analyzed showing that multiple slip

also occurred in charged samples. SISF's were seen on occasion, and it did not appear that

the presence of hydrogen had any effect on their appearance or number.

The primary purpose of the compression samples was to permit equivalent plastic strain

levels in samples with and without hydrogen to be reached. Figure 4.21 shows the

compression deformation behavior of charged samples at strains up to the uncharged tensile

failure strain, which can be compared directly to Figure 4.19.

113
ORIGINAL PAGE

8LACK AND WHITE PHOT.OGRA_.H

(a) (b)

Figure 4.20

(c)

Deformation structure of tensile samples with hydrogen at various strains.

(a) Ep = 0.15%. (b &c) Ef = 0.3%.

114
Figure4.21

(a)

- -'_ t'b ! "

_,L.A,CK A['-;D WHITE pt-tOIo_RAPH

0.50 l,tm

(c) (d)

(b)

Defomaation structure of compression samples with hydrogen at various

strains. (a) _ = 0.5%. (b &c) gp = 1.3%. (d) gp = 3.0%.

115
At the smallest strain, Ep = 0.5%, again there is evidence of local deformation bands,

with little dislocation activity elsewhere. At 1.3%, there now is a def'mite pattern of localized

bands of deformation, in marked contrast to the uniform deformation in uncharged samples at

this strain. Figure 4.21b shows examples of several of these bands, and Figure 4.21c

provides a close-up view of the edge of one of these bands showing the severe shearing of the

_/' precipitates inside the bands and dislocations trapped in the _' matrix just outside the bands.

At ep = 3.0% in charged compression samples, Figure 4.21d reveals that the deformation is

now uniform across the foils, and the presence of dislocations in the _/matrix is the

predominant observation. There still was some evidence of the intense band activity, but the

deformation had now spread out more uniformly. Dislocation densities were measured at this

strain, and it was found that there were 1.7 x 1011 dislocations/cm 2 in the matrix and 3.9 x

109 dislocations/cm 2 in the precipitates.

Figure 4.22 is a <111> foil of a charged sample at Ev = 1.2% showing an area where

the dislocations are bowing around and between the y precipitates. This can be compared to

the uncharged sample at Ep = 1.2% in Figure 4.16, and it can be seen that they both display a

similar deformation appearance. It is not inconsistent that the observations in the charged

<111> foil show dislocations primarily in the 7 matrix, while the [001] foils show bands of

intense y shearing. The areas outside these bands do contain dislocations that are trapped in

the _, matrix, and since the bands can be widely spaced, it is believed they just were not

observed in the limited viewing area afforded in the <111> foils. Since many more [001]

foils were observed, it is felt that these bands of intense activity are representative of

deformation behavior of charged samples at strains up to at least 1.2%.

116
ORIGINAL PAGE

BLACK AND WHITE PHO[OGRAP.H

Figure 4.22 <111> foil showing deformation structure with hydrogen. Ep = 1.2%.

;0.50 I.tm

Figure 4.23

(a) (b)

Deformation structure of tensile samples without eutectic T/Y'. (a) without

hydrogen, Ef = 5.3%. (b) with hydrogen, Ef = 1.2%.

117
4.4.1.2 Effectof eutectic"//7'

Eliminating the eutectic y/y resulted in higher su'ains to failure both in uncharged and

charged tensile samples (Table 4.13). [001] thin foils were examined from such samples, and

the observations are shown in Figures 4.23a and 4.23b. The uncharged samples in Figure

4.23a showed a much higher degree ofT' shearing than previously had been observed in any

samples. There was now a high dislocation density in the 3' matrix, but with much more

activity in the 7' precipitate.

In the charged samples without eutectic T/7' shown in Figure 4.23b, the failure

strain was about 1.2%, which was comparable to the strain level in the charged

compression foils examined in Figure 4.21b. The deformation behavior of these two

samples was similar, and the primary mode of deformation in both cases occurred by local

intense bands of 7' shearing, with some activity in the rest of the foil in the form of

dislocations trapped in the T matrix.

In summary of the overall deformation behavior, there was no qualitative

differences observed in the deformation behavior of samples tested in tension or

compression. In uncharged samples deformation began in isolated bands consisting of

intense shearing of the 7' precipitates along a single slip system. These bands eventually

hardened and the deformation structure became uniform across the sample consisting

primarily of dislocations trapped in the T matrix. This activity persisted up to failure in

tensile samples in contrast to the primary deformation mechanism of shearing in most

superalloys. 51.52.53.77 The disproportionate trapping of dislocations in the matrix is

supported by the dislocation density measurements summarized in Table 4.17. The

standard deviation on all measurements was within 20% of the reported average value.

Hydrogen-charged tensile and compression samples initiated deformation in the same

manner as uncharged samples with the isolated deformation bands. This activity continued

up to failure in charged tensile samples and to 1.2% in compression samples. Densities

,,,.'ere not measured in samples where deformation was accomplished primarily by the

118
narrowbandspreviouslydescribed,with theexceptionof the failed charged tensiles at Ep =

0.3%. This continuation of localized deformation to higher strains than in uncharged

samples represents the major effect of hydrogen on the deformation behavior and is

consistent with hydrogen causing localized plasticity in other alloys. 126'127'148'149'150 At

3.0% strain in charged compression samples, the bands had hardened, and the deformation

structure resembled uncharged samples with dislocations primarily trapped in the y matrix.

Eliminating the eutectic yP/' allowed higher strains to be achieved, and while dislocations

were still trapped in the y matrix, shearing of the _ precipitates also was observed in

uncharged samples.

Table 4.17 Summary of all measured dislocation densities in cm 2.

Sample

Condition

= 1.2%, no H

Location of Measurement

Y

4.5 x 101° 4.0 x 10 9

I Eutectic interface Y/7'

Ep = 3.0%, no H

= 0.3%, w/H

ep = 0.24%, w/H

£p = 0.24%, no H

ew = 0.15%, w/H

ii

1.0 x 1011

8.9 x 109

1.7 x 10 it

4.9 x 10 9

4.1 x 109

3,9 x 109

5.2 x 101°

3.9 x l0 t 1

4.4.2 Slip Character

The slip character was characterized on compression samples at plastic strains of

0.5%, 1.3% and 3.0%, following the guidelines proposed by Williams, Thompson and

Baggerly. 183 The spacing of the slip bands and their degree of waviness was quantitatively

mcasured following the methods used by McInteer, Thompson and Bernstein 17'* and

described in section 3.3.2. Measurements were made directly from optical micrographs of

119
the sample and may not be as accurate as if replicas of the slip bands were measured. At

0.5% strain, there was coarse planar slip with large areas that did not contain any slip

traces. Figure 4.24 shows the slip bands present at Ev = 1.3%, and it can be seen that there

is evidence for considerable cross-slip. The dark bands associated with the slip lines have

been analyzed by stereo microscopy and found to simply be inclinations on the surface.

At Ep = 3.0%, the slip activity was more homogeneously spread across the

samples, but the slip character could still be classified as coarse planar slip as illustrated in

Figure 4.25. At this strain, cracks that opened up along slip bands were occasionally

observed, in agreement with observations of slip band fracture discussed later. Nomarski

contrast was used in Figure 4.26, and in this figure the cracks appear as black lines and the

slip bands as white lines. This figure also shows the presence of some coarse wavy slip

that was only infrequently observed at Ep = 3.0%. Figure 4.27 graphically illustrates the

results of the quantitative measurements for samples with and without hydrogen. It was

found that the slip band spacing decreased continually with strain implying that new slip

sources were continually operating. The degree of waviness also increased steadily with

strain in uncharged samples suggesting that cross-slip was occurring to a larger extent.

Laue analysis of the specimen orientation in conjunction with slip trace analysis revealed

that all of the planar slip bands were of <11 l>-type, as expected, and all four <111> planes

were active. There did not seem to be any predominant slip system, and all {111 } planes

were observed to occur approximately to the same extent. It also was consistently seen that

the slip bands did not cut across the eutectic 7/7' as seen in Figures 4.24, 4.25 and 4.28.

This is consistent with the TEM observations in Figure 4.18 that show high densities of

dislocations at the interface but little activity inside the eutectic 7/7'.

120
ORIGINAL PAGE

BLACK AND _HIT_ P_HQ_T_QGRJAB_I_

Figure 4.24

501am

Compression slip bands at e_,p= 1.3%.

Figure 4.25

50 gm

Compression slip bands at ,C_,p= 3.0%.

121
ORIGINAL PAG_

BLACK AND WHITE. PHOTOGRAPH

Figure 4.26 Compression slip bands at Ep = 3.0% showing cracks along slip bands and

some wavy slip.

122
:=k

d_

140

120

100

80

60

40

20

C

no hydrogen

I " I ' "" ' I

I 2 3

plastic strain (%)

(a)

Figure 4.27

14

t_

._,.<

12

10

8

E

4

2

gen

C 1 2 3 4

plastic strain (%)

(b)

Quantitative results of slip band measurements as a function of strain and

hydrogen showing (a) slip band spacing and (b) degree of waviness.

123
ORIGINAL PAGE

BLACK AND WHJ.TF., pHO.11OGRAP-H

.......

i.tm

Figure 4.28 Note that slip bands do not cut through the eutectic y/_/.

124
The presenceof 300 ppm internal hydrogencausedthe slip to be confined to

slightly fewer,but coarserslip bands,asshownby thelargeraveragespacingfor charged

samplesin Figure4.27. Theslipcharacterwasverysimilarto theunchargedsamplesatEp

= 0.5%,but atEp= 1.3%theslipbandsin thechargedsamplesremainedmoreplanarthan

in unchargedsamples.This isin agreementwith theTEM resultsshowingthat hydrogen

causeslocalized planar slip up to higher strainsthanin unchargedsamples. The same

comparativetrendscontinuedto 3.0%,but thedifferenceswith unchargedsamplesbecame

lessobvious. Crackswereobservedalongsomeslip bandsatEp= 3.0%,andcoarsewavy

slip was seenona coupleof occasions.

4.4.3 DoubleNotchTensilesandSectionedTensiles

In orderto determinethefractureinitiation site,doublenotchtensilesampleswere

tested. Due to materialconstraints,it wasonly possibleto testsampleswithout hydrogen.

Sectionswere takenon theunfracturednotch,parallelto andperpendicularto the stress

axis,butcrackswereneverobservedin anyof thesections.

Next, previously testedandfracturedtensileswerenickel-plated,andlongitudinal

sectionswereprepared. In all cases,therewere somecrackedeutectic3'/7'within a few

micrometersof the fracturesurface. Therewasa slight tendencyfor thereto be more

crackedeutectic3'/_'in chargedsamples.Failedtensilesamplesalso were mounted,and

the fracturesurfacewaspreparedby grinding sosectionsperpendicularto thestressaxis

could beobserved. In thesetensiles,somecrackedeutectic3'/7'were seenasin Figure

4.29. All of theobservationssupporttheideathatcrackinitiation beganin theeutectic7/7,

or at theeutectic3'/7'interface.

125
ORIGINAL PAGE

BLACK AND WHIT£ PHOT_OGRAP.H

_q

Figure 4.29

d

Failed tensile sample sectioned perpendicular to the tensile axis showing

cracks running the eutectic T/Y.

126
4.4.4 Pre-StrainandPre-ChargeTensileTests

A study was initiated to determine if hydrogen affected either or both the

deformation or fracture behavior. The details and rationale of the testing procedure are

given in section 3.4.4, and Table 4.18 gives the conditions of each specimen along with the

associated tensile properties. The pre-strained samples were intended to help determine

whether hydrogen primarily affected the fracture behavior, while the pre-charged samples

were designed to determine hydrogen's effect on the deformation behavior. The yield

strength of all samples during the "straining only" portion of the test was within 3% of the

yield strength obtained when the same samples were pulled again to failure.

The first two samples in Table 4.18 were used as reference points. The next two

samples in Table 4.18 showed that hydrogen, as expected, had a dramatic effect on the

fracture behavior. This can be argued since without the subsequent charging, the prestrained samples would have attained a failure strain on the order of 3%. However, the

hydrogen somehow initiated the fracture process at a much earlier strain and additional

strains of around 0.2% were only possible. The last set of samples that were pre-charged

suggested that hydrogen also affected the deformation behavior. Similar to the argument

above, if hydrogen did not affect the deformation structure, then upon outgassing the

sample would have attained a failure strain of around 3% in later testing.

This large effect of hydrogen on the deformation behavior and subsequent failure

strain was not expected based on TEM studies, so the effect of hydrogen on these was

analyzed using longitudinal sections of samples that had only been strained and not broken.

It was found that in the sample that was pre-strained 2% without hydrogen, there were no

cracks to be found in the gage section. However, in samples that contained hydrogen and

had been strained to 0.2%, there were several cracked eutectic y/_/' already present. This

shows that hydrogen causes the eutectic T/Y to crack at a much earlier strain.

127
Table4.18

| i

Sample Condition

pre-charge

no hydrogen

pre-strain 1%, charge

pre-strain 2%, charge

pre-charge, strain 0.1%, outgas

pre-charge, strain 0.2%, outgas

Pre-strain, pre-charge tensile results.

0.2% Y.S. (MPa)

1100

1145

1107

1093

1122

1129

U.T.S. (MPa)

1108

1189

1107

1093

1136

1143

0.33

3.9

0.16

0.26

R.A. (%)

0.57 2.2

0.46 2.1

4.5 Tensile Fracture Behavior

4.5.1 General Fracture Characteristics

The majority of the tensile fracture surface appeared ductile, as shown in Figure

4.30a. At higher magnification, Figure 4.30b, the ductile "voids" are seen to consist of a

flat region surrounded by a thin rim of ductile tearing. These "voids" are about 0.75 I.tm

across and are essentially the same size throughout the fracture. There also were areas that

have a similar appearance at high magnification, but were much more brittle-appearing at

lower magnification. This type of fracture is illustrated in Figure 4.31; the facets have been

identified as { 111 } planes using techniques described in section 3.4.6. The more ductileappearing fracture of Figure 4.30 also has been identified as occurring macroscopically on

{ 111 } planes. Therefore there appear to be two types of fracture surfaces when viewed

macroscopically, designated as Type A (rough) and Type B (smooth), for convenience.

Type A fracture was more commonly observed in all tensile samples studied.

128
Crystallographiccleavagefacetswerealsofoundonthefracturesurfacethatvaried

in sizefrom 201.tmto 501.tm.Figure4.32ashowstheirrelationto thegeneralfracture,and

Figure4.32b is a highermagnificationview of oneof thecleavagefacets. Typical river

linesindicativeof cleavagewereseenonall facets,andtheinitiation sitewasalwaysat the

surfaceof the facet, which will be important in later discussions. Sometimes there is a pore

at the initiation site, but the majority of cleavage facets had no identifiable microstructural

feature at the initiation site. The orientation of the river lines was checked in view of

Lynch's observations of fee cleavage, 176,177 but no consistent orientation of these features

was found. Laue analysis of the specimen orientation and two-trace analysis have verified

that these cleavage facets occurred on macroscopic {001} planes. Figure 4.33 shows a

cleavage facet on an {001 } plane with several { 111 } fracture planes surrounding it.

129
ORIGINAL PAGE

BLACK AND WHITE PHO[OGRAPi-t

Figure 4.30

(a) (b)

(a) Ductile fracture appearing over most of the tensile fracture surface.

(b) High magnification view of { 111 } facet.

Figure 4.31

10 lam

(a) (b)

(a) Smooth { 111 } ductile fracture seen on parts of the tensile fracture

surface. (b) High magnification view of { 111 } facet showing similar

appearance to the fracture in Figure 4.30b.

130
ORIGINAl..; PAGE

BLACK AND WHITE PHOTOGRAPH

Figure 4.32

100 p.m I0 um

(a) (b)

(a) Cleavage facets seen on the tensile fracture surface. (b) Single cleavage

facet showing initiation site at the edge opposite of the pore.

201am

Figure 4.33 {001 } cleavage facet surrounded by { 111 } ductile fracture planes.

131
Plateau etching was used to determine the relation of the microstructure to the

ductile fracture of Figures 4.30 and 4.31 and to the cleavage facets. Figure 4.34

convincingly shows that the cleavage facets are the eutectic _,/y in the microstructure.

It was more difficult to make observations utilizing the plateau etching technique at

higher magnifications, because the interface between the microstructure and fracture surface

was not very clear. The etchant used to polish the sample was able to slightly penetrate

under the stop-off lacquer, resulting in a narrow region of the fracture surface being

slightly etched. But this turned out to also be beneficial, as Figure 4.35 shows a partially

etched fracture surface illustrating the relation between the _,/?' microstructure and the small

ductile fracture areas. It appears that the flat fracture feature is the y precipitate and that the

ductile tearing is occurring through the _ matrix. TEM replieas were made to try to

determine if any river lines or other evidence of cleavage existed in the flat region identified

as the 3/precipitate. The results of this attempt were not convincing one way or the other;,

this fracture will simply be referred to as "brittle".

Fracture surfaces were also heavily etched to bring out the _'/'y' microstrueture, and

Figure 4.36 shows the white tear ridges occurring along the _ phase as indicated by the

arrows. Further evidence is shown in Figure 4.37 where microcracks were seen on a few

occasions running along what appeared to be the 7 matrix phase.

132
ORIGINAL PAGE

BLACK AND WHII£ PHOTOGRAI_.H

Figure 4.34

1 0 lain

Plateau etch showing that the cleavage facet in the top half of the micrograph

corresponds to the eutectic y/y indicated by the arrow in the bottom half of

the micrograph.

Figure 4.35

0.50 I.tm

Plateau etch showing relationship between ductile fracture and y/y

microstructure.

133
OIRfGfNAL PAGE'

BLACK AND WHIT_ P_OTOGRAPH

Figure 4.36 Type A {111 } fracture showing triangular regions corresponding to 7'

precipitates. The white tear ridges are occurring along the _ phase.

Figure 4.37

2 iJm

Microcracks on tensile fracture surface showing outline of 7/3/ microstructure.

134
When there is less ductile tearing as previously shown in Figure 4.31 (Type B), an

outline of the T/_ microstructure can actually be seen on the fracture surface. Figure 4.38

shows a { 111 } fracture facet with the outline of many triangular _/' precipitates evident.

The cuboidal 'y' faces are aligned along {001 } planes, and thus a { 111 } section through the

precipitate would appear triangular.

4.5.2 Effect of Hydrogen

The presence of internal hydrogen did not have any apparent effect the type of

fracture features seen. The area surrounding the pores appeared the same with and without

hydrogen, as did the ductile areas and cleavage facets. Type A { 111 } fracture remained the

predominant type of slip band fracture in tensile samples.' The major effect of hydrogen on

the fracture behavior was the large increase in the number of cleavage facets now observed

on the fracture surface. Without hydrogen, the area fraction of cleavage on the fracture

surface was consistently 5% to 6%, while the charged samples showed a larger and wider

range of between 11% and 15%. No significant effect of hydrogen on the number of pores

on the fracture surface was observed, suggesting that large pressures do not build up in the

pores after charging that could lead to crack initiation, as suggested by others. 9

Surface cracking was seen along the gage length only in tensile samples containing

hydrogen. The cracks were comprised of segments along the (001) plane and { 111 }

planes. Samples were subsequently heavily etched, and Figure 4.39 shows that the (001)

segment was always through a eutectic _'/_', while the {111 } segment was through the _'/'y'

microstructure. There were numerous cases of isolated cracked eutectic "t/q( on the surface,

but only a few cracks were seen on { 111 } planes that were not linked at some point to a

eutectic 7/Y.

135
ORIGINAL PAGE

BLACK AND WHITE; PHO.'[.OGRAP_H

Figure 4.38 Relatively smooth {111 } fracture facet showing the triangular outline of

many "y'precipitates.

Figure 4.39

(a) (b)

Surface cracking on charged tensile samples showing the eutectic "{/y

cracked on the (001) plane and the crack propagating on a { 111 } plane.

136
4.5.3 Effect of Alternate Heat Treatment and ltlP'ing

The effect of removing the eutectic 7/3" and/or porosity did little to change the

fracture surface other than eliminating these features. There was, however, some tendency

to observe more of the smooth, type B { 111 } fracture in samples with less eutectic "//_'.

The amount of cleavage seen on the fracture surface was very close to zero in the alternate

heat treatment samples and also was reduced in the HIP'ed samples. The area fraction of

cleavage approximately doubled in the charged HIP'ed samples compared to the uncharged

I-IIP'ed samples, as it did in the unHIP'ed samples. These observations are all consistent

with the metallographic observations and previous tendencies noted in uncharged samples.

4.6 Fracture Toughness

4.6.1 Fracture of KIc Samples

The fracture appearance of KIC fracture toughness samples was very similar to

tensile samples in the type of features seen, but there were major differences found between

the fractography of the KIc samples and the JIc samples. All four { 111 } planes were

active in the fracture of both types of fracture toughness samples. The KQ samples had a

macroscopically rough fracture surface, and slightly more of the smooth, type B { 111 }

fracture was seen than type A, as in Figure 4.40. Secondary cracks were seen emanating

from many pores, and Figure 4.41 shows one such case.

Quantitative fractography of the observed cleavage facets revealed that the area

fraction of cleavage facets on the fracture surface was about equal to the volume fraction of

eutectic "f/_ in the microstructure. This means that without hydrogen the crack path is

moving randomly with respect to the eutectic _'/"/'. This is consistent with the observations

in uncharged tensile samples. Longitudinal sections of the KIC fracture surface revealed

many cracked eutectic q,/_, as seen in Figure 4.42. There were even some cases where a

few cracked eutectic TJ_" linked up through the 7/Y' microstructure, as shown in Figure

4.43.

137
Thepresenceof 300ppmhydrogencausedthefractureto bemuchmoretortuous,

Figure 4.44, and type A { 111} fracture predominated. Onceagain, therewere many

secondarycracksemanatingfrom poresand a large numberof cleavagefacets. It was

determinedthaton thefracturesurfaceof chargedKIc samplestherewasanaveragearea

fractionof cleavagefacetsof 21.5%. Thisrepresentsa four-fold increaseovertheamount

of eutecticT/_'in themicrostructure,andit canbeconcludedthatthecrackpathis seeking

out the eutecticT/_' in chargedsamples. This is analogousto the increasedcleavage

observedin chargedtensilesamples.

Longitudinalsections of the charged KIC fracture surface showed a large number of

cracked eutectic T/"/'. The plastic zone size, ry, was estimated using: 184

(4-3)

where K1 is the stress intensity and Oys is the yield strength. It was found that all of the

cracked eutectic T/_' were within the plastic zone, and it was calculated that approximately

12% of the eutectic T/T' in the plastic zone were cracked. There did not appear to be any

relation between the eutectic T/q/size and whether or not the eutectic T/'Y'was cracked.

138
ORIGfNAL PAGE

BLACK AND WHITE PHO[OGRA__H

50

Figure 4.40 Fracture surface of uncharged KIC samples.

20 I.tm

Figure 4.41 Secondary crack emanating from a pore on KIc fracture surface.

139
ORIGINAL PAGE

BLACK Abi_ WHITE, PHOT__OGRAPH

Figure 4.42 Longitudinal section of uncharged, failed KIC sample showing several

cracked eutectic _,/'_ beneath the fracture surface.

Figure 4.43 Longitudinal section of fractured KIC sample showing linked cracked

eutectic T/_

140
ORIG_N,_L PAGE

8LACK AriD. WHITE P,"'IO]OCRAPI.-I

401._ I

Figure 4.44 Fracture surface of charged KIC sample.

141
4.6.2 Fractureof JIcSamples

The fracture surfaceof the hc samples was macroscopically much flatter than the KIC

samples and consisted almost exclusively of smooth, type B {111} fracture, as seen in

Figure 4.45. Once again the distance traversed on one {111 } plane before switching to

another { 111 } plane was around 5 I.tm, as found in the tensile samples and KIc samples.

There were some secondary cracks observed, Figure 4.46, and it is possible that these

cracks occurred during the unloading-loading cycle in the J integral test. The spacing

between cracks of 100-150 I.tm corresponds to approximately the length of crack growth

that occurred in each loading cycle.

A crack also was consistently observed running between the end of the fatigue precrack and the beginning of the J integral test region, as shown in Figure 4.47. It was noted

that the fracture in the fatigue pre-crack region was slightly different, having a smaller

{111 } facet size. It is important to note that in the final overload region after the second

fatigue cycling, the fracture was much more macroscopically rough and resembled the KIc

fractography, as seen in Figure 4.48. The bottom of Figure 4.48a is the fatigue area used

to identify the end of the J integral test region, and the upper portion of the photo is the

overload region that occurred as a result of failure due to the fatigue conditions employed.

The rate of loading was much faster in this overload region than in the J integral region.

142
ORIGINAL PAGE

BLACK AND WHITE PPIOT.OGRAI_.H

Figure 4.45

5O Ima

Fracture surface of uncharged JIc sample.

Figure 4.46

5O _m

Secondary cracks on JIc fracture surface,

143
ORIGINAL PAGE

BI.,A.CK A_D. _Hi ]'_" PHO[OGRA£H

Figure 4.47 Crack running between fatigue pre-crack and beginning of J integral test.

Figure 4.48

600gin

(a) (b)

(a) Bottom half of the photo shows the J-controlled crack growth region,

and the top half and (b) show the fracture behavior in overload region

after completion of J integral test.

144
Theporeson the JIc fracture surface rarely had any secondary cracks around them,

and their main effect was to slightly alter the crack path, as shown in Figure 4.49. The

crack growth in this figure is occurring from the bottom of the photo to the top. The area

fraction of porosity and eutectic y/_/on the fracture surface was approximately equal to the

same quantities found in the microstructure, in agreement with uncharged tensile and KIc

samples. Longitudinal sections of interrupted JIc tests supported the idea that the crack

path was not affected by the presence of eutectic T/T'- In contrast to similar sections of KIC

samples, there were rarely any cracked eutectic T/?' in the plastic zone of JIc samples, as

shown in Figure 4.50.

Fortunately, the transverse orientation of the JIc samples tested was close to <110>

so that the crack growth could be easily observed, as in Figure 4.51. These sections also

showed that crack growth occurred across a single { 111 } facet and not along the sides of

two { 111 } planes. It appeared that in many cases cracks were opening up along { 111 }

planes slightly ahead of the main crack.

The presence of 300 ppm internal hydrogen did not alter the fracture surface

appearance or the crack growth mechanisms, in contrast to KI¢ samples. In JI¢ samples,

there were very few eutectic T/I/cracked in the plastic zone, and there was no quantitative

increase of cleavage on the fracture surface. These fractographic differences between KIC

and JI¢ samples were surprising and will be discussed further in the next chapter.

145
ORIGfNAL PAGE"

81JE,K _ WBITE, £HO;OGR!__H

Figure 4.49 Effect of pores on the JIc fracture behavior. Direction of crack growth is

towards the top of the photo.

125 lain

(a) (b)

Figure 4.50 Sectioned JIc sample showing little effect of porosity or eutectic "[/y on

crack growth.

146
There was no apparent effect of the ahemate heat treatment on the JIc fractography,

except for eliminating the presence of cleavage, of course. HIP'ing the material eliminated

the alteration of the crack path by the porosity, but there were no other effects on the

fractography. Figure 4.52 is a longitudinal section of a HIP'ed JIc sample given the

alternate heat treatment, showing the crack growth mechanism to be identical to the sample

in Figure 4.51, which contained porosity and eutectic T_-

147
ORIGINAL PAGE

BLACK AND WHITE PHOT_OGRAP_H

Figure 4.51

•_ .............. 10 pm

(a) (b)

Crack growth along { 111 } planes shown in sectioned JIc samples.

(a) (b)

IOwa

Figure 4.52 Crack growth behavior in a HIP'ed JIc sectioned sample.

148
CHAPTER5

DISCUSSION

The previouschaptershowedtherearemanysimilarities anddifferencesbetween

PWA 1480andother superaUoys,andthis chapterwill attemptto explainmanyof these

observations.Oneof themoresurprising results was that the porosity and eutectic 7/_/' did

not affect the amount of hydrogen that could be charged into the material, nor did they

appear to act as irreversible trapping sites. Considering these observations, an ahernative

explanation is needed for the high hydrogen content found after charging, and possibilities

will be offered.

The deformation behavior of uncharged samples differed from that of most

superalloys 11.-sl-53 because precipitate shearing was not observed at all strains. The factors

that affect shearing will be covered in terms of how PWA 1480 may differ from similar

alloys. The relationship between the deformation behavior and possible failure

mechanisms also will be covered, as well as the role of the eutectic 7/_ on both of these

processes. Cleavage of the eutectic y/y is one of the unique attributes of PWA 1480, and

its origins and effects will be examined in detail. Finally, it was found that fracture occurs

on {1 11 } slip planes, in agreement with fracture of many other single crystal

superalloys 42.95.96,100 and fcc single crystals. 91'92'94

The effect of hydrogen on the tensile properties was similar to other single crystal

superalloys, 25 in that the failure strain was reduced to a greater extent than the R.A. This

suggests that hydrogen is localizing the plasticity, and this will be discussed in view of

available models of hydrogen embrittlement processes. 126'127'131'132'185 This observation

is supported by analysis of the slip line characteristics, and agrees with observations in

other alloys of increased slip planarity due to hydrogen. 25'174'186'187 Again, the

149
relationshipbetweenthe deformationstructureandpossiblefailure mechanismswill be

explored.

Thefracturetoughnessresults were unusual because the Kic and JIc values did not

agree as expected. Further, the fractography and crack growth behavior of these two types

of samples was quite different, and this must be carefully examined. It appears that the role

of the eutecfic _,/'/' was not the same in each test, and this will be discussed in terms of

possible differences in the testing techniques, such as the unloading involved in the Jxc test.

5.1 Hydrogen Solubility and Trapping

The desorption results suggested that no hydrogen was trapped at irreversible traps

after gas-phase charging at 1020 atmospheres and 350°C. The measured rate of hydrogen

outgassing for PWA 1480 closely followed the calculated profiles for IN 718 and IN 903 at

all temperatures. This suggests that hydrogen is freely diffusing out of the material, and

there is no hydrogen remaining behind at irreversible trap sites. Recent results for a similar

alloy, CMSX-2,1° showed the estimated diffusion constant of hydrogen in this alloy to be

very close to that of hydrogen in IN 718 and IN 903.14° These results were obtained

through a combination of outgassing experiments and permeation tests. Therefore, the use

of diffusion data from IN 718 and IN 903 appears appropriate for single crystal nickel-base

superalloys.

Because the absence of irreversible traps is somewhat surprising, the results

warrant critical examination. It is possible that the hydrogen analysis technique of vacuum

hot extraction performed at 900°C did not detect hydrogen at strongly trapped sites. The

resultant binding energy of the traps for such an occurrence can be calculated by employing

equation 2-20 that relates outgassing data and binding enthalpy. The subsequent binding

energy is almost 2 eV, which is more than twice any previously reported binding

energies. 9,135,136,141 It is unlikely that traps of this strength would exist in PWA 1480,

150
and thus it is probablethat the vacuum hot extraction techniquedid detectall of the

hydrogen.

One should also consider the possibility of hydrogen rearrangementduring

desorption,which mayaffectinterpretationof theresults.This wouldinvolvediffusion of

hydrogenfrom eitherthelatticeor reversibletrapsintoirreversibletrapsduringdesorption

at agiven temperature.It is not believedthatthisoccurredduringdesorptionat300°Cor

higher basedon the work of Baker, et at.,9and Roux, et at., on CMSX-2,1° which

showedtherewere noirreversible trapsoperatingat thesetemperatures.However,it is

possiblethatduringdesorptionat 150°Cthehydrogencoulddiffuseinto irreversibletraps

andnot be ableto overcomethe binding energyof the trap to diffuse out. Subsequent

desorptiondatawould suggestapparenttrapping,butthiseffectwasnot seenin this study

for the desorptiontimes usedat 150°C. This type of secondarytrapping would be an

artifact of thedesorptiontestmethodand wouldnot representany trappingthat occurred

during charging. All room temperature mechanical testing was completed in times that

were far shorter than required for significant diffusion and subsequent possible secondary

trapping to occur. Thus, it is believed that the distribution of hydrogen after charging did

not change before or during testing.

In other studies concerned with hydrogen trapping in similar alloys, irreversible

trapping was observed in CMSX-2, 9'1° but this occurred after cathodic charging at 150°C.

The binding energy calculated by Baker, et al., 9 for the pores was such that trapping could

occur during charging at 150°C but not during charging at higher temperatures. Roux, et

al., report similar results of irreversible trapping in CMSX-2 after cathodic charging at

150°C but did not observe irreversible trapping after gas-phase charging at 350°C. l° These

observations are consistent with the absence of irreversible trapping observed in this study

after gas-phase charging at 350°C. Roux, et al., also performed tritium autoradiography

tests that suggested a higher concentration of hydrogen can exist in the eutectic 7/7' than in

the surrounding microstructure. 1° In these tests, tritium was charged into the material at

151
150°C;it is possiblethattheeutectic7/Ybehavesasanirreversibletrap at this temperature

but not at higher temperatures, such as 350°C.

In order to better evaluate the distribution of hydrogen after charging, it is important

to analyze the solubility of hydrogen in the microstructure. In this study, analysis of the

hydrogen content after charging showed that the amount of hydrogen in the material was

essentially independent of the amount of porosity or eutectic 7/7' in the microstructure.

This in itself suggests that these features were not acting as large sinks for hydrogen during

charging. Thus, the relatively high concentration of 300 ppm hydrogen must be due to

other factors, and calculating the equilibrium solubility for the charging conditions will

assist in determining if hydrogen exists in an equilibrium state inside the material or if it is

reversibly trapped in the microstructure. The following expression can be used to

approximate the hydrogen solubility: 14°

It2 -Hs

S= SoP exp(-R--_- ]

(5-1)

where S is the solubility in cm 3 H/cm 3, p is the pressure in atmospheres, So is the

solubility constant, Hs is the heat of solution, R is the gas constant and T is the

temperature. Table 5.1 shows the available solubility data for similar alloys. The data for

each of these alloys reveals that the equilibrium solubility at 1 atm. and 25°C is less than 1

ppm. The gas-phase charging conditions of 1020 atmospheres and 350°C yield hydrogen

solubilities ranging from 25-62 ppm. This is well below the average value of 300 ppm

found in PWA 1480 after charging.

Thus, it is apparent that since neither the eutectic 7/7' nor the porosity act as sinks,

the 7/7' microstructure must either have a much larger solubility than the alloys in Table

5.1, or there exists a large amount of hydrogen reversibly trapped in the lattice. The issue

of higher hydrogen solubilities in single crystal superalloys compared to alloys such as IN

903 has been raised by Roux, et al., 1° in their study of CMSX-2. It was suggested that the

152
Table5.1 Solubility datafor nickel-basesuperalloys.

Alloy

IN903

So(cm3/cm3'atm'a)

0.166

Hs(cal/mole)

1120

IN718

HasmHoy

0.89

0.53

1930

1240

Reference

140

140

188

compositional differences betweenIN 903 and CMSX-2 could possibly explain the

apparentsolubility differences;in particular,it wasthoughtthatTacould stronglyincrease

hydrogensolubility. In their study,1°CMSX-2 waschargedin theexactmannerasin this

work, and a hydrogenconcentrationof approximately 350 ppm was reported- only

slightly higherthantheamountof hydrogenin PWA 1480foundin this study. Sincethese

values are similar, and PWA 1480containsabout twice as much Ta as CMSX-2, the

effectivenessof Ta in increasinghydrogensolubility appearsto beminor.

Suffice it to say that the solubility questionremainsopen, and further work is

neededin this areaon single crystal nickel-basesuperalloys. If solubility differences

cannotsatisfactorilyexplainthediscrepancybetweenequilibrium solubility andhydrogen

content,it is possiblethatreversiblelatticetrappingcan. Reversibletrapsaredefinedas

microstructuralfeaturesthat cancontain hydrogenand,undergiven testconditions,act

either assourcesor sinksfor hydrogen.132This is in contrastto irreversibletraps,which

alwaysact assinksuntil theyaresaturated.It is believedthat duringgas-phasecharging,

hydrogen is forced into lattice sites by the hydrostatic pressureand remainsin this

supersaturatedstateafter cooling from 350°Cto room temperatureunderpressure.This

argumentis supportedby charging experimentsconductedat different pressuresthat

showedthehydrogencontentto beproportionalto thechargingpressure.

153
The argumenthathydrogenis only at latticetrappingsitesis supportedby thefact

that CMSX-2 containedaboutthesamehydrogenconcentrationafterequivalentcharging

conditions. CMSX-2containsnoeutecticT/q( and less initial porosity than PWA 1480, but

otherwise the T/Y' microstructure is very similar. Thus, it makes sense that this similar

microstructure contains about the same amount of hydrogen as PWA 1480. This also

implies that small compositional differences are not important in determining the amount of

hydrogen that can be charged into these types of alloys. Based on the similarities found

between CMSX-2 and PWA 1480, there is no reason to believe that PWA 1480 would not

respond similarly to cathodic charging at 150°C. In particular, it is probable that the pores

and eutectic 7/7' would act as irreversible trap sites during such charging.

While it appears that hydrogen is primarily at lattice sites in the 7/Y' microstructure,

it is possible that hydrogen is not distributed uniformly between the _ and 7' phases

because of different interaction parameters. But this is unlikely based on preliminary

results of Chtne and co-workers on CMSX-2.1°8 In their study, an atom microprobe with

a deuterium tracer was used, and little, if any, partitioning to either of the phases was

observed. A slight tendency for hydrogen trapping at the interface may have been

observed, but this is not believed to be a strong trap because the y/_/' mismatch in these

alloys is very small (0.1-0.5%).

Although hydrogen appears to be "trapped" at lattice sites in PWA 1480, it is free to

diffuse through the lattice at room temperature and higher, based on desorption results.

Thus, it is believed that hydrogen is distributed uniformly through the microstructure and is

probably located at tetrahedral or octahedral sites of the fcc lattice in the _, phase and in the

same sites in the ordered LI2 lattice of the _ phase. Therefore, the traps can be defined as

reversible under the testing conditions of interest; however, in actuality the diffusion

constant at room temperature is so low that substantial diffusion is not likely. Stressassisted diffusion should also be considered during testing, but Baskes 189 has suggested

that this does not significantly aid the hydrogen diffusion process in fcc metals. Thus,

154
althoughhydrogenis ableto diffuseduring testing,it is not likely due to the shorttimes

involvedin thetestsof thisstudy.

In summary,it appearsthatthepresenceof porosityandeutecticT/T' do not affect

either the amount of hydrogen that was charged into PWA 1480 or the trapping

characteristics, and based on desorption results, all hydrogen in the material is reversibly

trapped. Further, the atom microprobe results of Chdne and co-workers :°8 suggest that the

hydrogen is uniformly distributed with respect to the y/7' microstructure. These results

have major implications on how hydrogen affects the subsequent tensile behavior. Since

the hydrogen concentration apparently is evenly distributed, dislocations will generally

encounter the same hydrogen atmosphere throughout the material. This implies that

whether hydrogdn assists or hinders the motion of dislocations, the effect will be uniform

through the material. The role of strain rate in hydrogen embrittlement is also minimized,

because it is not as important for the hydrogen atmosphere to move with the dislocations

since there will be a high, uniform concentration of hydrogen wherever the dislocations

travel. This analysis is consistent with the results of Roux, et al.,10 that showed no effect

of tensile strain rate on the degree of hydrogen embrittlement in gas-phase charged CMSX2. Finally, any role of irreversible trapping sites that could contain large amounts, of

hydrogen and possibly lead to cracking is eliminated.

5.2 Tension and Compression Properties

The tensile properties of PWA 1480 were consistent with those found by other

investigators. 9°,111,116 The yield stress was slightly higher than in some

studies,23,76.90,1 ll.116 and this can be correlated with the smaller 7' size reported in this

study, as well as the use of tensile samples within 5 ° of [001]. At constant volume fraction

ofT, precipitates, it has been shown that strength increases with smaller precipitate size. 9°

155
To betterillustratetheeffectof orientation,oneof theslabstested (#134) was a few more

degrees off of [001] than the other slabs, and consequently the properties were between the

average for the [001] orientation and the single slip <130> orientation discussed in

Appendix A. This indicates that multiple slip is being delayed longer in slab #134 than in

other samples; this has been observed by Shah and Duhl in a careful study of orientation

effects. 9° In another study, Fritzemeier 23 reported some exceptionally high strain-tofailure values. It is believed that these orientation effects, along with differences in the

eutectic 7/7' volume fraction of their material and that of this study, can explain these

results.

The presence or absence of work hardening in these types of alloys is not well

understood. Most investigations only report the tensile properties, with no analysis of

work hardening behavior. In studies that have reported the stress-strain

behavior, l°AI.52a9° two types of curves have been observed for the [001] orientation at

room temperature, as illustrated in Figure 5.1. Type I, Figure 5.1a, is typical of many

polycrystalline metals and also has been observed in hydrogen charged CMSX-2 l°ai and

single crystal MAR-M 200.19° Type II stress-strain curves have been observed in

uncharged CMSX-2,11 Rent N452 and the PWA 1480 of this study. Type I behavior is

indicative of simultaneous multiple slip, while Type II behavior is observed in single slip

orientations. Type I behavior would occur if multiple slip is initiated uniformly across the

sample, and this is observed in the two single crystals cited above that exhibit this

behavior. Multiple slip is favored in single crystals with the [001] orientation, but it is

possible to have multiple slip occur in the crystal and still observe Type II behavior. This

could occur if slip initiated on isolated slip systems and continued uninterrupted along these

systems. Narrow bands of intense dislocation activity would form, which would

evcntually lead to a more uniform slip across the sample as dislocation interactions and

tangles occurred in these bands. This in turn would lead to an increased work hardening

rate, :is observed in l:igure 5. lb. This scenario is exactly what is obserwed in this study

156
03

etO

_0

Type I

Engineering Strain

(a)

03

v v

Engineering Strain

(b)

Figure 5.1 Two types of stress-strain curves observed in [001] single crystal nickel-base

superalloys.

and will be discussed later. Tensile samples without eutectic 7/3e showed the same stressstrain behavior as in Figure 5.1b, except that the curve was extended to higher stresses and

strains. There was no effect of porosity on the tensile properties, which is consistent with

other studies. 23,28

The tension/compression asymmetry found in the [001] orientation is similar to

other investigations showing the critical resolved shear stress (CRSS) in tension to be

about 5% greater than in compression, s8,9°,116 The models used to explain the

tension/compression asymmetry are based on changes in the dislocation core s6 or the

behavior of the partial dislocations with stress state. 9° In this study, hydrogen lowered

both the tension and compression yield strength equally, and subsequently there was no

effect of hydrogen on the tension/compression asymmetry. This may be used to argue that

hydrogen does not substantially affect either of the above two factors. This is indirect

evidence, and there have not been any studies done by others to substantiate this.

The effect of hydrogen on the tensile properties will now be discussed in terms of

available models of hydrogen embrittlement processes. The presence of 300 ppm internal

157
hydrogen lowered the flow stress by about 10%, decreased the strain to failure by an order

of magnitude and had little effect on the reduction-in-area (R.A.). All of these observations

are consistent with the hydrogen-enhanced localized plasticity (H.E.L.P.) model, 126'146

which argues that hydrogen enhances dislocation generation and motion. The large

decrease in failure strain accompanied by little change in R.A. signifies that the majority of

deformation is concentrated in a narrow region near the resultant fracture surface. This

plastic strain localization has been observed in other single crystal superalloys tested in

hydrogen. 25 Other experiments supporting the H.E.L.P. model have shown very localized

deformation near crack tips in thin foils that are strained in a hydrogen

atmosphere, 142.t43,145 which could be manifested in bulk samples by the strain localization

described above. Another manifestation of H.E.L.P. observed in PWA 1480 is the lower

flow stress observed in charged samples, which could be caused by enhanced dislocation

generation.

There are alternative explanations for hydrogen embrittlement, but they do not

satisfactorily explain the results of this study. The trap theory of hydrogen embrittlement

argues that a critical concentration of hydrogen will reach a microstructural feature and

possibly lead to cracking, t32't33 This is unlikely to occur in PWA 1480 for a few reasons.

First, the previous section demonstrated that there were no sites containing uniquely large

concentrations of hydrogen, and hydrogen existed more or less uniformly across the

microstructure. Therefore, hydrogen must diffuse to a site in order to attain the "critical

concentration" required for hydrogen embrittlement. The issue of diffusion also was

discussed previously, and it was believed that in room temperature testing, long-range

diffusion would be insignificant. Further, the presence of stress fields would not likely aid

diffusion very much. 189 Thus, the trap theory of hydrogen embrittlement is not applicable

to this study, because there does not appear to be any trapping sites that contain

significantly more hydrogen th:_n the surrounding microstructure.

158
Thedecohesionmodelof hydrogenembrittlementcould alsobeusedto explainthe

degradationin tensileproperties,butexperimentalprooffor thismodelisdifficult to obtain.

The decohesionmodelarguesthatthepresenceof hydrogensolutesdecreasestheatomic

bondstren_hof themetalcausingabrittlefracturewhenthelocalstressexceedstheatomic

bondstrength,t3°,t91The factthatthetensileR.A. is largelyunaffectedby thepresenceof

hydrogensuggeststhat this modelis notapplicable. Having similar R.A.'sfor uncharged

and chargedsamplesimplies that thefracture mechanismis not beinggreatly changed,

which is not theeffectthatonewould generallyassociatewith areductionin atomicbond

strength. Further,thereis quite a bit of fractographicplasticity on thefracturesurface,

especiallyin thematrix phase.Again, thisdoesnotcorrelatewith a brittlefracturecaused

by weakenedatomicbonds.

Thus,it appearsthattheconceptof hydrogen-enhancedlocalizedplasticity canbe

used to explain much of the observed tensile behavior of charged samples. The

observationthat hydrogencausessubstantialembrittlement,asmanifestedby a four-fold

decreasein failure strain in sampleswithout eutectic 7/T', will now be examined.

Previously, only embrittlement of samples with eutectic Y/T' has been discussed, and in

these samples a ten-fold decrease in failure strain was observed. Table 4.13 shows that in

samples without eutectic T/T', hydrogen still causes a decrease in the yield stress of about

10%. The failure strain is not decreased as much as in samples with eutectic Tlq/and the

R.A. is reduced about 50% instead of having little change. These are still manifestations of

H.E.L.P., but the R.A. differences suggest that the plasticity is not as localized. For

reasons discussed above, it is not felt that the trap theory or decohesion model are

applicable. In charged samples containing eutectic T/q/, this constituent cracked at a very

early stage of deformation, and it is possible that this crack could cause intense localized

deformation, compounded by the effect of hydrogen on localizing the deformation. This

scen_wio may explain why the deformation is more localized in charged samples that contain

eutectic y/q/. In samples without eutectic y/I/, it is possible that only the effect of hydrogen

159
on localizing deformation is operating and not the effect of cracked eutectic T/T'. As a

result, there is a decreased localization of plasticity and decreased embrittlement due to

hydrogen, as measured by change in failure strain. All of the evidence used for hydrogenenhanced localized plasticity in this study are bulk observations, and to definitively

determine that hydrogen is enhancing dislocation motion and generation, in situ straining

experiments on the TEM must be performed. 142,143,145 This was not possible in this

study, so the application of the H.E.L.P. embrittlement process remains based on bulk

observations.

The effect of hydrogen concentration on tensile properties was studied in two sets

of experiments. In both, a steady decrease in strain-to-failure was observed until a

minimum value of about 0.3% was reached. The hydrogen concentration at which the

minimum failure strain was achieved was different depending on the method used to vary

the hydrogen concentration, and this discussion will be aimed at understanding this

difference. One method used to vary the hydrogen content was gas-phase charging at

different hydrogen pressures, which should result in a uniform concentration of hydrogen

across the sample because of the long hold times during charging at 350°C. The second

method used was to gas-phase charge the samples at a single hydrogen pressure and

subsequently outgas them to different hydrogen concentrations. It is probable that this

method would result in hydrogen concentration gradients. It only took 200 ppm total

hydrogen content to fully embrittle these samples but 300 ppm to fully embrittle the

samples charged at different pressures. It is possible that the lower total concentration of

hydrogen required to obtain the minimum failure strain in outgassed samples is related to

the presence of hydrogen concentration gradients. The outgassing was performed at

300°C, so long-range diffusion is possible, but any subsequent secondary trapping would

not be expected because the binding energy of any irreversible traps is likely overcome at

this temperature. To understand the diffcrences in hydrogen distribution in the two sets of

s:_mples, consider samples theft contain the same total amount of hydrogen. The one

160
charged at a lower pressure would have the hydrogen evenly distributed, whereas the

outgassed sample would have the hydrogen concentrated more toward the center of the

sample. Thus, the outgassed sample could have a much higher, local concentration at the

center. This effect could lead to a degradation in properties at a lower total amount of

hydrogen than in the ones charged at different pressures.

In analyzing these results, it should be noted that the outgassed samples were tested

at 300°C, but the results for samples with close to 300 ppm hydrogen were very close to

the results of tests conducted at room temperature. The fact that there was little difference

in tensile properties of charged samples tested at room temperature and at 300°C brings up

this question: At what temperature does hydrogen embrittlement cease? Higher

temperatures were not tested due to material constraints, but the behavior at these'

temperatures can be surmised. Of course, this question cannot be answered without

incorporating time, and thus 20 minutes will be used in this thought experiment, since this

was the time allowed for the temperature to equilibrate before testing. Hydrogen will be

able to diffuse out of the material, but based on the desorption results, this would only

effect a narrow surface region of about 200 I.tm when held for 20 minutes at 300°C. The

bulk of the sample should still contain a uniform, untrapped concentration of approximately

300 ppm. At higher testing temperatures, the depleted surface layer would become larger

and more important, and eventually the properties would approach those of the uncharged

material. Using diffusion data of IN 903,14° it can be calculated that at 400°C after 20

minutes, there would still be over 290 ppm hydrogen in the sample, while at 500°C there

would be 190 ppm, and at 600°C there would be 80 ppm left. Also at 600°C, about twothirds of the specimen surface would be almost hydrogen-free after 20 minutes using the

simple approximation of x = _Dt, where x is hydrogen-free surface layer.18°

Using thesc data and the results of tensile tests at 300°C, it appears that the

unifomlly charged samples would have to be tested at 600°C in order to obtain properties

approaching those of uncharged samples. Further, the above discussion did not account

161
for therole of temperature on the deformation behavior. It has been shown by Dollar and

Bernstein 77 that the deformation structure does not appreciably change from room

temperature to 300°C. Above 300°C, there was more shearing of the 7' precipitates,

however the ductility only increased slightly from room temperature to 600°C. 192

Considering the slight increase in ductility with temperature, as well as the role of the

hydrogen-free surface layer, it is possible that testing at 400°C-500°C could result in a large

improvement in properties. This is important because hydrogen embrittlement is typically

thought to be a low temperature phenomenon, and the possibility of hydrogen

embritflement at temperatures up to 500°C is a concern for service applications. Of course,

this analysis applies to the specific case of an initially uniform hydrogen concentration, and

any applications to situations involving surface gradients or different hold times may result

in less severe conclusions.

5.3 Deformation and Fracture Behavior without Hydrogen

At the earliest strains studied in tension and compression samples, deformation was

accomplished by narrow, widely-spaced, intense bands of dislocations shearing "y

precipitates. There was limited activity elsewhere in the form of single a/2<l10>

dislocations and a/2<110> pairs in the 7 matrix. This type of dislocation procession has

previously been observed in other Ni-base superalloys 193196 and aluminum alloys

containing coherent, ordered particles. 197 It has been argued in these cases that a

dislocation cutting an ordered panicle will reduce its effective cross-sectional area, thereby

weakening it by causing a region of decreased order. 35 Hornbogen 198 has quantified this

process by modifying an existing low volume fraction 'y' CRSS model, 199 now

incorporating the decrease in CRSS due to the passage of dislocations along the same slip

plane. The expression for the CRSS due to coherent, ordered, cuboidal panicles, A't,

162
accounting for the passage of n dislocations was written as: 198

A'c(n)- 4 bD\ 1 -

(5-2)

wherc yis the APB energy, d is the effective particle diameter, b is the Burgers vector and

D is the spacing between particles. This equation shows that the CRSS on a given slip

plane is decreased for every dislocation that passes along the same slip plane. Therefore, at

small strains it appears easier to move dislocations on isolated slip planes where shearing

has already occurred, than it is to initiate slip uniformly across the sample.

At the next plastic strain level examined- 1.2% in tension and 1.3% in

compression - narrow deformation bands were no longer observed, and more uniform

deformation occurred across the samples. The vast majority of dislocations were in the ),

matrix, and there was little evidence of precipitate shearing. The previous bands were not

observed, probably because of their wide spacing and the fact that the heavy activity now

occurring over the entire sample made their observation more difficult. This is not to say

that they do not exist anymore, but that their relative importance in the deformation process

has become minor. Thus, it appears that the isolated slip bands active at smaller strains

have hardened, either by generation of multiple slip within the bands or by interaction with

accumulated debris. The work hardening rate increases at this point because of more

dislocation interactions and the trapping of dislocations in the 3, matrix. This, along with

the work softening described above, provide a possible explanation for the stress-strain

behavior observed in Figure 5. lb.

At 3.0% strain, there was a very high dislocation density in the 3' matrix and little

activity in the "y' phase. Normally, in high volume fraction y' superalloys, the deformation

mechanism at these strains is thought to be continued shearing of the precipitates. 11,51-53

To assist in interpreting the differences between PWA 1480 and othcr similar superalloys,

the Copley and Kear approach to modeling the critical resolved shear stress (CRSS) can bc

163
employed.Themodelis basedona forcebalanceof apair of dislocationstryingto entera

7'precipitate.TheCRSScanbeexpressedas:2

7 T

zc- 2b br + 0"5(Zm + Xp) (5-3)

where T is the dislocation line tension, r is the particle radius on the slip plane, and XM and

Xp are the shear stresses of the matrix and precipitate, respectively. The APB energy of

PWA 1480 is close to that of similar alloys, 11,2°°,2°1 as is the precipitate size 23.76.9° and

shear modulus. 4345 Of the parameters in equation 5-3, only the matrix and precipitate

flow stress of PWA 1480 may differ greatly from similar alloys. The matrix flow stress

used was the one employed by Copley and Kear in their study of a similar alloy and should

be fairly accurate, because the composition of the )'phase of PWA 1480 is not too different

from their alloy. 2,zl It was pointed out in section 2.1.1.2 that the 7' flow stress in PWA

1480 may be higher than in most other superalloys because of the larger amount of Ta in

PWA 1480. "/'he cutting stress is proportional to the precipitate flow stress, so naturally a

higher precipitate flow stress would make it more difficult for a pair of dislocations to enter

the precipitate. Thus, it is possible that the 7' flow stress could explain the differences in

deformation structure between PWA 1480 and other superalloys.

As previously mentioned, there was little dislocation activity in the 7' precipitates,

but occasionally Superlattice Intrinsic Stacking Faults (SISF's) were observed in this

phase. Theft appearance was not consistent from foil to foil within the same sample, and

the separation of the a/3<112> partials varied greatly with each SISF. These observations,

along with the overall sparse occurrence of SISF's, lead to the conclusion that their role in

the deformation process is minor. Further, Dimiduk suggested that the presence of SISF's

in Ni3AI was directly related to the thin foil preparation technique, 2°2 casting possible

doubt on their reproducibility and meaningfulness. SISF's were seen to a lesser extent in

compression samples in this study, but again this simply may be due to variations in

164
specimenpreparation, since SISF's have previously been observedmany times in

compression.57,203

Theseobservationscontradict thoseof Milligan andAntolovich, 123 who report a

high density of SISF's within the 7' precipitates after room temperature deformation. The

density of SISF's shown in their micrographs is much higher than was ever observed in

this study. Similar local areas may have been observed in this study, but the representative

density of SISF's was much lower than that shown in their mierographs. These

differences could be due to many factors, such as orientation, specimen preparation or

testing conditions, but the most likely cause is slight compositional differences in the

separate batches of material used. These compositional differences are only a possibility,

and a definitive explanation for the differences observed in these two studies cannot be

given.

The tensile fracture behavior of uncharged samples will now be discussed relating

the deformation behavior and fractography to possible fracture mechanisms. The tensile

fracture consisted of ductile separation on { 111 } planes combined with cleavage facets

dispersed across the fracture surface. Sectioned hydrogen-free tensiles revealed that there

were no cracked eutectic 7/7' observed at 2.0% strain, which means that the eutectic 7/7'

cleaved very close to the fracture strain. Further, only a few cracked eutectic 7/7' were

observed on sectioned fractured tensiles, and these were very close to the fracture surface.

Thus, it seems that cleavage of the eutectic Y/7' was confined to the fracture zone and

occurred very late in the deformation process. Although, this is not to say that the eutectic

"//3" did not play a role in the fracture process. In fact, samples without eutectic y/y'

achieved failure strains twice as high as in samples with eutectic "f/7'. Further, there was

no effect of eutectic y/7' on the deformation process per se, and these observations suggest

that cleavage of the eutectic q,/y' initiated the fracture process.

In analyzing the effect of cleavage on the fracture process, it must be recognized

that the edge of the cracked eutectic 3'/7' can act as a local stress riser. This is analogous to

165
pre-cracked carbides in MAR-M 200, 2o4 which also were cracked normal to the stress axis.

Observations in that study showed localized slip along { 111 } planes emanating from the

cracked carbides. A similar situation could occur for the eutectic T/7' in PWA 1480, thus

creating shear localization in this manner. Continued shear localization across the samples

could be aided by the stresses in the T matrix created by the large density of dislocations

there. It was found in another study that when only the matrix deforms, large stresses can

build up within the precipitates. 2°5 Shear localization caused by the cracked eutectic T/7'

and exclusive deformation in the T matrix could be manifested by rapid shearing across the

7' precipitates. If such shearing did occur, then it is possible that this could lead direcdy to

failure. For instance, in another superalloy, Gell and Leverant 2o4 argue that localized slip

led to weakening of the atomic bonds across the slip plane. Then the normal stresses,

which are not relaxed in these alloys, 1°1A°3 could separate the bonds and create a britdeappearing fracture. This so-called "glide band decohesion" is believed to be a fairly

common fracture mechanism in alloys that deform by shear of coherent particles. 2°6

Completing the story, the ductile T phase is the last to undergo fmal separation and thus

necks down. This scenario is in agreement with the fractographic evidence showing that

the eutectic T/T' cracked fin-st, the 7' precipitates had a britde appearance and the T phase

appeared ductile. It is also consistent with the TEM observations of high dislocation

densities in the T phase with little activity in the 7' precipitates. Intense shear on subsequent

fracture planes would not be observed on TEM foils because the foils were taken from

sections below the fracture surface.

In samples that do not contain eutectic T/T', cleavage of this constituent cannot

initiate fracture, but the fracture sequence could be much the same. In these samples,

continued deformation leads to the initiation of significant precipitate shearing, while the y

phase continues to contain high dislocation densities. Previously, it was argued that

precipitate shearing is difficult in PWA 1480, but it appears that the stresses created by the

hi eh dislocation density in the I'natrix may tk_rcc initiation of shem-ing to partially relieve

166
these stresses. It is possible that this does not occur uniformly across the sample and in

some areas large stresses continue to build up within the precipitates. 2°5 This could, in

turn, lead to a similar shear localization phenomenon as described above. This would

explain the similar fracture surface appearance between samples with and without eutectic

y/3/, even though the deformation behavior is slightly different. Thus, it appears that shear

localization along slip planes offers a bridge between the observed deformation behavior

and fractography.

In comparing the observed fractography to that of other alloys, it is found that this

type of { 111 ] crystallographic fracture with ductile tearing of the T matrix has been

observed in many other single crystal superalloys. 95'96'97'100 So it appears that although

the deformation behavior of PWA 1480 differs from_most similar alloys, the fracture

behavior does not. Further, shear localization leading to slip band fracture is apparently

common in superalloys and other systems with coherent particles. 204"206'207

Another type of fracture has been observed in single crystal superaUoys and will be

discussed now. Crystallographic fracture along apparent {001 } planes has been observed

in another study of PWA 1480 under certain cyclic loading conditions. 95 This type of

fracture was never observed in this study, but it is useful to examine its origin in the other

study. A similar macroscopic type of fracture also has been observed in hydrogen-charged

CMSX-2 tensile samples. 9 Figure 5.2 is a schematic of a profile of this type of {001 }

fracture surface compared with the {111 }-type fracture more commonly observed in these

alloys. In both studies reporting the {001 }-type fracture, it was suggested that the failure

was confined to the matrix phase, and that the precipitates had not failed. Analysis of the

deformation structure in CMSX-2 revealed cracking along { 111} planes inside the T

matrix, 11 so it is believed that failure occurs inside this phase rather than at the 7/Y'

interface. Cuboidal images of the precipitates were still observed on the fracture surface,

because the matrix phase is so thin, causing the 7/7' outline to be exposed. In the fatigue

167
{_1 } ma_oscopic _cmre plane

naunaauna'iaamuauua

anaanaannanoaanaan

(a)

{ 111 } failure of ? matrix

gmug

{111 } macroscopic

_ fracture plane

(b)

ductile tearing of

K T matrix

Figure 5.2 Profile of (a) the {001 }-type fracture sometimes observed in single crystal

superalloys at low and high magnifications and (b) {111 }-type fracture

observed in this study at low and high magnifications.

study of PWA 1480, 95 {001 } fracture was confined to the lowest AK values tested, and

{ I 11 } fracture was observed at all higher AK values. The {001 } fracture appearance and

limitation of fracture to the "/matrix can be explained by considering the dislocation

structure. It was argued that at low AK values, the cutting stress for the precipitates had

not been overcome, and thus deformation was limited to the matrix phase. 95 No relation

between this behavior and the fracture was given, but in the hydrogen-charged CMSX-2

tensile samples, it was observed that the macroscopic {001} fracture appearance was

caused by a high dislocation density in the "/matrix that led to strain exhaustion and failure

through this phase along { 111 } planes, l_ This is illustrated in Figure 5.2a showing that

macroscopically the fracture was {001 }, but microscopically it was actually along { 111 }

pl,_ncs in the 7 rn;Hrix. This is likely what happened in Tclcsnmn and Ghosn's fatigue

san_plcs a! l_w AK. A {()()1 }-lypc flacture was not seen in oi]r stl]dy, because il is lholl_,.bl

168
that the high density of dislocations in the y matrix caused shear localization across the 7'

precipitates. A { 111 }-type fracture was reported at higher AK values in the Telesman and

Ghosn work, but there was no TEM evidence to determine if a similar scenario occurred as

in the tensiles of this study.

Now that the ductile fracture observed in this and other studies has been discussed,

the cleavage facets and their origin will be covered. Cleavage in an fcc material is unusual,

and thus the observed cleavage of the eutectic T/7' warrants further examination. It is

believed that this is crystallographic cleavage, which is supported by evidence that includes

the flat fracture surface of the eutectic _,/'/', the river markings present and the facets being

parallel to a low index crystallographic plane. A theoretical analysis of fcc cleavage by

Gilman 2°8 suggests that the cleavage plane would be {001}, as found in this study. In

practice, most fcc cleavage occurs in the presence of an aggressive environment, 176 so it is

unusual that the eutectic '//'/' cleaved in uncharged samples tested in air at room

temperature. To understand the origins of this cleavage, the nature of the eutectic y/y must

be understood.

The modulus difference between the eutectic '//_' and the '//7' microstructure is not

believed to be large, so this does not appear to offer an explanation for cleavage. The vast

majority of eutectic '//7' showed little evidence of a lamellae structure and it is believed to

consist mostly of the 7' Phase-15 Therefore, the behavior of the eutectic '//7' phase should

be very similar to that of the 7' precipitates. It has already been noted that dislocations did

not easily penetrate the 7' precipitates, and likewise, it was shown that surface slip lines did

not cut through the eutectic '//7' phase. Further, the dislocation densities were higher in the

immediate region of the eutectic '//7' and at the interface than in the surrounding '//7'

microstructure. This leads to the likelihood of large stresses existing at the interface, which

in turn could initiate cleavage. 2°9.21° Also, it is probable that the eutectic y/7' is inherently

more brittle than the _/' precipitates, because the precipitates apparently undcrwent some

shear bel\_re failure, but the eutectie y/7' seemingly did not. The impingement of slip bands

169
ontheeutectic_,/_'is only apossibilityfor why theeutectic_'/,/appearsto cleaveandis not

claimed to be definitive. However, it is consistent with the experimentalevidence,

including the observationthat all cleavageinitiation sitesoccurredat the edgeof the

cleavagefacet. A similar finding wasreportedin MAR-M 200,wherecrackinitiation of

carbideswasbelievedto becausedbyimpinging slip bands,but noexperimentalevidence

was provided.2°4

It also hasbeenobservedthatporesinitiate cleavagein somecases,but thereare

twoexperimentalobservationsthatdemonstratethattheporescannotinitiatethemajorityof

cleavagefacets. Perhapsthestrongestevidenceis thatthevastmajorityof cleavagefacets

arenot seenadjacentto pores. It is quite possiblethata poremay abutthe eutecticqt/7'

below the fracture surface,but this pore cannotinitiate cleavagebecausecleavageis a

normal stress-controlledprocess,so the initiation site must be on the fracturesurface.

Secondly,in HIP'edsamplescontainingessentiallyzeroporosity,cleavagewasobserved

to thesamerelativeextentasin unHIP'edsamples.

5.4 DeformationandFractureBehaviorwith Hydrogen

Thedeformationstructureof hydrogen-charged samples was similar to uncharged

samples at strains up to the failure strain of 0.3% in tension and 0.5% in compression. At

these strains, narrow, widely-spaced bands of dislocations shearing the 7' precipitates

dominated the deformation process. The deformation bands persisted in charged

compression samples up to at least Ep = 1.3%, whereas in uncharged samples uniform

deformation was seen at ep -- 1.3%. It has been observed in pure nickel and nickel-base

superalloys that hydrogen promotes planar slip, 25'174'187 causes localized

deformation 25,186.187 and possibly lowers the stacking fault energy. 211 These

observations are consistent with the persistence of the deformation bands to higher strain,s

170
in charged compression samples and the increased planarity of the slip lines shown in

Figure 4.27. As in the uncharged samples, it is probable that the dislocation structure in

these deformation bands eventually will become tangled because of initiation of multiple

slip and debris accumulation. Thus, slip will initiate uniformly across the sample, as

further slip in the local deformation bands becomes unfavorable. This was observed at a

strain of 3.0%, resulting in little difference between the deformation structures of

compression samples with and without hydrogen. Both consisted of uniform, multiple slip

with the vast majority of activity in the ymatrix.

Dollar and Bernstein 11 have performed the only other study that has examined the

effect of a uniform concentration of hydrogen on the deformation structure of a single

crystal superalloy. Their study on CMSX-2 found that the presence of hydrogen caused

the dislocations to be confined to the matrix phase instead of shearing the precipitates, as

occurred in uncharged samples. This change in deformation behavior was believed to

account for much of the decrease in ductility of charged samples. This type of matrix

trapping behavior occurred in PWA 1480 without hydrogen, so such a dramatic change due

to hydrogen was not possible, although the scale and extent of trapping could be affected

by the presence of hydrogen. It was argued in CMSX-2 that the accumulation of

dislocations in the 3' matrix at small swains led to an increased local hydrogen concentration.

This was followed by enhanced dislocation generation and a decreased mean free path in

the T matrix, which persisted up to failure. This scenario differs significantly from what is

believed to occur in PWA 1480, where the gapping of dislocations in the matrix is thought

to be due to the large 7' flow stress, as previously explained. Dollar and Bernstein also

ob_el-v6d room temperature hydrogen-induced cross slip of dislocations onto {001} cube

planes_. 2z No Such observations were made in the PWA" 1480 of this study, althou___ti the

evidence may have been overlooked, because even in CMSX-2, cross-slip segments were

not frequently observed.

171
Thefracturebehaviorof chargedtensilesampleswill now beexaminedin termsof

the deformation structure, the role of the eutecticT/?'and the H.E.L.P. embrittlement

process.The main differencebetweentheunchargedandchargedfracturesurfacesis the

drastic increasein the amountof cleavageon the fracturesurfaceof chargedsamples.

Therewasnoqualitativedifferenceobservedin theductilefracturealong{ 111 } planes, and

possible explanations of these two observations follows.

Sectioned tensiles revealed that the eutectic T/_/' cracked at a much earlier strain in

charged samples (0.2%) and to a greater extent than in uncharged samples. In some

manner, the presence of hydrogen is initiating cleavage at an earlier strain. It was

suggested in the previous section that cleavage may be initiated by the stress created from

impinging of slip bands on the eutectic T/l(. If this does occur, then the H.E.L.P. model

may offer an explanation for the increased cleavage in charged samples. At 0.2% strain,

the deformation structure of uncharged and charged tensiles consisted of the narrow slip

bands previously described. The H.E.L.P. model contends that hydrogen can enhance

dislocation generation and motion, so one might expect higher densities of dislocations to

be piled up at the eutectic y/1e interfaces in charged samples. In fact, the density in the

matrix layer surrounding the eutectic _,/_/' was substantially higher in samples with

hydrogen compared to uncharged samples at equivalent strains, as seen in Table 4.17.

This supports the possibility that in charged samples higher stresses exist at the eutectic T/Y

at a given strain. This could lead to cracking of the eutectic T/_ at earlier strains, as

observed. Supposing hydrogen is enhancing the generation and motion of dislocations, it

is likely that the critical stress for cracking would be reached at more eutectic T/T' before

final fracture. This was observed, and thus it appears that the H.E.L.P. embrittlement

process can be used to offer explanations for the increased cleavage in charged samples.

It should not be overlooked that the decohesion model of hydrogen

cmbrittlement 131,133 cotll<l _flmobc used to explain the accelerated cleavage by arguing that

the presence of hydrogen solutes decreases the atomic bond strength in the eutectic y/y.

172
This model is usually applied to systems that change from a ductile to brittle fracture, but

this is not necessary. It is not believed that the decohesion model can explain the increased

propensity for cleavage in this study because the eutectic y/y already cleaves in the absence

of hydrogen. However, it cannot be proven that the decohesion model does not offer the

means to enhance cleavage of the eutectic 3'/q(, and thus it must remain a viable alternative.

It is not likely that the trap theory of embrittlement 132'1_ is applicable, since it was found

that the eutectic y/_' was not a sink for hydrogen. Nor did it appear that substantial

amounts of hydrogen would be able to diffuse to the eutectic y/q( during testing. Thus, it

does not seem that a "critical concentration" of hydrogen necessary to cause

embrittlement 132,134 can be reached within the eutectic %/q(.

Turning now to the effect of hydrogen on ductile fracture, it is surprising that

despite the increased propensity for cleavage and much smaller failure strains in charged

samples, the ductile fracture regions are not affected. The fracture surface is generally

rougher because the crack path is now seeking out pre-cracked eutectic y/l/, but the

microscopic fracture still consists of the ductile "voids" on {111} planes. It was

previously suggested that these ductile areas were a result of shear localization in the

precipitates followed by a brittle-appearing fracture of the precipitates along { 1 11 } planes

and ductile tearing of the matrix. This was supported by a large density of dislocations in

the matrix that could create stresses sufficient to initiate shear localization. At first glance,

this scenario does not appear possible in charged samples, because at the failure strain there

was a relatively low density of dislocations in the matrix.

Therefore, it remains to explain why the ductile fracture behavior did not change

with hydrogen, although the deformation behavior at fracture was apparendy quite different

than in uncharged samples. To approach this problem, the effect of the eutectic y/q( on the

tensile behavior must be separated out. In charged samples with eutectic T/q(, the eutectic

y/y' cracked at a small strain resulting in high subsequent stresses at the tip of the cracked

eutectic y/q(. It was already suggested that this could lead to shear localization, but now

173
incorporating the ideas of H.E.L.P., it is possible that shear localization is directly

enhanced by the presence of hydrogen. Independent experiments supporting this model

show a large increase in the local dislocation activity surrounding a crack in charged

samples compared to uncharged samples. 142147 This could lead to intense shear

localization emanating from the cracked eutectic T/'/and eventually fracture along this slip

band. This would likely occur in a narrow region encompassing the resultant fracture

surface, and thus this activity would not be observed on the TEM thin foils. It was already

suggested in uncharged samples that the final shear across the precipitates may not have

been observed because it only occurred at the resultant fracture surface, and the same

phenomenon may have occurred in charged samples.

In tensile samples without eutectic _'/7', there were no pre-existing cracks, so

another mode of explanation must be used. In this case, the failure strain is increased to

just over 1%, and the deformation structure still consists of the narrow slip bands. Due to

this deformation structure, it cannot be argued, as in uncharged samples, that the stresses

caused by the high dislocation density in the matrix promote strain localization and

subsequent fracture in the precipitate along the slip band. Accounting for this difference

and absence of pre-cracked eutectic 3'/7', it appears that the most plausible explanation

involves the observed narrow deformation bands. At 1.2% strain, these bands are much

more developed than those observed at 0.2% strain, and the precipitates along this band

have undergone massive shearing. Thus, it seems appropriate to cite the GeU and Leverant

work 2°4 that suggests the precipitates are greatly weakened after this massive shearing.

Subsequently, the fracture occurs along these bands producing a similar fracture

appearance to all other types of samples since ultimately the same shear localization

mechanism was operative. So it appears possible that two different deformation processes

occur in charged samples depending on the amount of eutectic 7/'/present. When there is a

large amount of this constituent, it cracks at an early stage of deformation, and this fact,

along with the H.E.L.P. model, provide a means to shear localization and subsequent

174
fracture. In sampleswithout anyeutectic3,/_/,thedeformationstructureconsistsof bands

of intensesheafingof theprecipitates.Thiscoulddirectly leadto weakeningalongthisslip

planeandsubsequentfracturethere. Bothof thesepossiblescenarioswouldcreatesimilar

appearingfracturesurfaces.Further,thefractureis similar to unchargedsamplesbecause

in the end,all processesculminatein slip bandfracture of the precipitatesandductile

separationof thematrix.

5.5 Analysisof Strengthening

In alloys containing a high volume fraction of coherent, ordered particles, the main

strengthening contributions are from order strengthening and solid solution strengthening. 3

There are two models that are commonly used in such alloys to predict the critical resolved

shear stress (CRSS) and study the effect of various parameters on the CRSS. Reppich and

co-workers 73.8°,81 have developed one of the models that analyzes the flow stress due to

the presence of these particles as a function of their size and volume fraction. The

quantitative application of this model to the PWA 1480 of this study is left to Appendix F

because of its length and inconclusive results. In summary, the total CRSS predicted by

this model is 1143 MPa, which is considerably higher than the 450 MPa obtained for PWA

1480 in this study. There are two reasons why the model did not predict the observed

CRSS more closely. Perhaps the largest inaccuracy is that the Huther-Reppich model only

considers the leading dislocation and does not account for the interaction term between the

pair of dislocations. This would have reduced the CRSS by almost half, resulting in a

reasonable value. Further, Ardel169 points out that there is a large spread in the values of

some of the constants obtained by different investigators, 75'8285 adding unreliability to the

model.

The other model generally used was developed by Copley and Kear to predict the

CRSS of alloys with a high volume fraction of coherent, ordered particles. 2 The model i_

175
basedon a force balanceof a pair of dislocationstrying to enteraT' precipitate,and the

CRSScanbeexpressedas:2

7 T

"_c- 2b br + 0"5('_m + "tp) (5-4)

where r is the particle radius on the slip plane, and '17M and Xp are the shear stresses of the

matrix and precipitate, respectively. Using 1_M = 115 MPa and Xp = 180 MPa from Copley

and Kear, 2 G = 54.5 GPa and r = 0.32 I.tm, a CRSS of 370 MPa was calculated. This is

slightly lower than the CRSS of 450 MPa found in PWA 1480, but the agreement is not

bad. Dollar and Bernstein used a similar approach and credited such a difference to work

hardening that resulted from a high density of dislocations in the "t matrix.77 However,

Copley and Kear's model is based on events at yielding, and it is not believed that the

contribution to strengthening due to work hardening would be very large at the microstrains

involved.

Although the comparison between the predicted CRSS and the measured value for

PWA 1480 in this study is relatively good, it is believed that better agreement is possible.

It could be that the flow stresses used in equation 5-4 for the matrix and precipitates may be

inaccurate. The APB energy of PWA 1480 is close to that of similar alloys, 11.2°°.2°1 as is

the precipitate size 23°76'90 and shear modulus. 43-45 The matrix flow stress used was the

one employed by Copley and Kear in their study of a similar alloy and should be fairly

accurate because the composition of the 7 phase of PWA 1480 is not very different from

their alloy. 2,21 It was pointed out in section 2.1.1.2 that the T' flow stress in PWA 1480

may be higher than in most other superalloys because of the larger amount of Ta in PWA

1480. Assuming that the flow stress of the 7' is totally responsible for the difference in

CRSS values, then a value of xp = 340 MPa would have to be used to eliminate the

diffcrence. This represents almost a 90% increase over the previous "t_,value, but this may

not be unreasonable considering the large solid solution strengthening effect of Ta 4'3739

176
andthefact that CopleyandKear's7' alloy did notcontainanyTa.2 This analysisis not

meantto beexact,but it doespoint out thatif moreaccuratevaluesof theflow stressesof

eachphaseare used,rather thansimply using the valuesquotedby Copley andKear,2

betteragreementcanbeobtained.

An alternate explanation for the effect of temperature on the CRSS of PWA 1480,

independent of the Copley and Kear model, was provided by Milligan and Antolovich 123

incorporating the effect of Superlattice Intrinsic Stacking Faults (SISFs). They observed a

high density of SISF's in the 7' precipitates at room temperature, and proposed that the

faults act as barriers to glide. They extend this idea to explain the effect of temperature on

the CRSS and convincingly con'elate their observations with the CRSS results. However,

in this study there was not a consistent or large density of SISF's observed, so this

possible strengthening effect cannot be assessed.

5.6 Fracture Toughness

The fracture toughness of PWA 1480 without hydrogen was similar to other nickelbased alloys shown in Table 2.4. Direct comparisons cannot be made though, because to

the author's knowledge there have not been any other studies on high q( volume fraction

superalloys, let alone single crystals. The fracture mechanisms also are quite different

because of the generally crystallographic fracture of single crystals and the different

microstructural features involved. Further, direct comparisons of charged samples cannot

be made unless the same hydrogen-charging conditions are used.

There are several points to discuss concerning the differences observed in KIc and

JIc samples of PWA 1480. To begin with, JIC values can be converted to KIC values

using: 158

= _ (5-5)

JIc E

177
where E is Young's modulus in plane strain. It was found that the uncharged KIC samples

had a toughness value that was 10% less than in comparable Jic samples, and the charged

KIc value was 20% lower. Normally, KIc values would be higher than comparable Jxc

values, as shown in Figure 5.3. This difference in values is believed to arise from the fact

that JIc measurements are taken at minimal (0.2 mm) crack extension, while KIc is

measured at 2% apparent crack growth. 212 The agreement is a function of the slope of the

R curve, with steeper R curves resulting in worse agreement. PWA 1480 has an

intermediate R curve slope in comparison to those shown in Figure 5.3, and therefore KIC

should be slightly higher than JIo Further, the KIc values used for comparison in this

study were KQ values and may have been slightly higher than true KIC values because of

thickness effects. 181 The discrepancy in the JIc and K]c values can be better understood

after examining the differences observed in the fracture surfaces.

JoK

Figure 5.3

KIc

KIC

JiC I

I

I

I

I

1

2%

Comparison between JIc and KIC values based on the slope of the resistance

curve. 212

Perhaps the most striking difference in the two tests is the role of the eutectic TH'.

In Klc samples, hydrogen caused the crack path to seek out this constituent in a similar

manner to what occurred in tensile samples. But in Jlc samples, hydrogen does not

178
promote cleavage, nor does the eutectic _,/_ appear to have an influence on crack growth in

uncharged or charged samples. This, in itself, may offer an explanation for the differences

in the toughness values of the two tests. Since many eutectic _,/y are cracked in the plastic

zone of uncharged Km (-5% of all eutectic _,/',/) and charged Kic (-12%) samples, this

would probably lower the toughness of the material. However, in JIc samples, there were

virtually no eutectic 7/'1/cracked in the plastic zone, and intuitively this would lead to a

higher JIc toughness value than KI¢ value.

There is apparently a difference in the two tests that is affecting cleavage of the

eutectic _'/_/'. One would not expect to observe such marked differences in the fracture

behavior and appearance of Km and Jm samples when similar stress intensities are

involved, as was the case. The stress intensity determines the plastic zone size, and

therefore both types of samples should have comparable plastic zones. The loading rate of

the KI¢ samples was only slightly higher than in JIc tests, but this difference is not believed

to be a significant factor. Any orientation effects on the crack growth are not likely because

the charged KIC samples had a transverse orientation very close to the JIC samples. The

primary difference is the unloading-reloading cycles that occur in a JIc test compared to the

constant rising load in a KIC test. This cycling can affect the strain energy stored in the

plastic zone, as well as the deformation structure. It may be possible that the effects of

unloading in the JIc test could explain why crack growth is not affected by the presence of

the eutectic _,/y.

In tensile samples, the possibility of cleavage being initiated by the stress created at

impinging slip bands on the eutectic _,/_/was raised. It seems likely that a similar process

could operate in the plastic zone of the fracture toughness samples. This offers an

explanation for the cracked eutectic _'/y in the KIC samples, and it could be possible that the

slip bands impinging on the eutectic T/Y in the Jlc samples did not provide a large enough

stress to initiate cleavage. This may have occurred if the unloading cycle in the JIC tests

relaxed the deformation structure ahead of the crack tip in some manner. The observation

179
that the fractography in the final overload region of the JIC samples resembled that of Kzc

samples supports the idea of a load cycling effect because this fracture region was created

by a more rapid rising load similar to the conditions of a KIC test. Unfortunately, the

loading rate or stress intensities involved in this final portion of the test were not

monitored.

This above scenario incorporating the effect of unloading is only a possible

explanation for the fractography differences, and it should be noted that generally the effect

of small amounts of unloading on crack growth behavior is minimal. 213'214 Thus, more

work is needed in this area to either conftrm the above scenario or to find a new explanation

for the differences observed in KIC and JIc samples.

There were only minimal improvements in JIc when the eutecdc 7P_/was removed,

and these differences are probably within experimental error. Based on the sectioned JIc

samples that showed no effect of the eutectic ),/q( on the crack path, it is not surprising that

the eutectic 3,/3e did not affect the JIc values. There was an improvement in JIc for HIP'ed

samples, and this can be related to the fractography shown in Figure 4.48. The pores

altered the crack path, and it is possible that the crack growth was locally accelerated.

Elimination of the pores allows the regular, periodic crack growth along { 111 } planes,

which could be responsible for the higher toughness of the HIP'ed samples. The

unloading-reloading cycles also may affect the behavior of the porosity, because in KIC

samples most of the pores were associated with secondary cracks, but few were in the Jlc

samples.

The crystallographic fracture along { 1 11 } planes has been explained by Chan and

co-workers 42,101.103 in terms of the stresses present at the crack tip. Fracture along

coplanar slip bands is caused by the buildup of large normal stresses along these slip

bands, which can explain why the crack does not deviate from the coplanar slip band.

Simultaneous fracture along a pair of { 111 } planes can occur because similar normal

stresses build up on the cross-slip planes. 1°3 t:racture along these two planes give the

180
characteristic ridge pattern observed on fracture surfaces of PWA 1480 and other single

crystal superalloys, 4%95'96'1°° as shown schematically in Figure 5.4. Fracture along pairs

of cross slip systems implies that crack growth occurs along <110> directions. Sectioned

JIc samples clearly showed that the crack path was along a <110> direction, such that the

crack travelled along a single { 111 } plane and then changed angles to travel along another

{111 } plane. Referring to Figure 5.4, this corresponds to growth in the [101] direction

along the (111) and (111) planes. However, as seen in Figure 4.45, the fracture ridges

also could be oriented parallel to the direction of crack growth, and in this case the crack

path travels simultaneously along two { 111 } planes. This would still correspond to crack

growth in the [101] direction in Figure 5.4, but now the crack grows simultaneously along

the (I 11) and (111) planes. This second type of crack path is the one used by Chan in

explaining the presence of ridges on the fracture surface, x°3 It is not known how the crack

would grow when the <110> directions are at an angle other than 90 ° to the crack front, but

it is possible that a combination of the above described paths could occur.

It would be useful to extend Chan's results to attempt to explain the repeatable,

periodic nature of the fracture ridges in all samples. The distance that the crack traversed

on one { 111 } plane before switching to the corresponding cross-slip plane was typically

around 5 l-tm, as shown in Figure 5.4. Chan does not incorporate any such distance in his

model, 1°3 and it does not appear possible to modify the model to derive this distance. It

can only be assumed that when the crack is growing along a single {111 } plane, the normal

and shear stresses on the associated cross slip plane are increasing until they reach a critical

value. At this critical value, the crack turns and begins growth on the cross-slip system.

The interesting point is that the ridge length was about 5 i.tm in both tensile samples and

fracture toughness samples. The stress states in these two samples are different, but

perhaps on a local scale the stresses on the { 111 } cross-slip system are comparable.

181
[0011

original crack plane

[110]

[1101

Figure 5.4 Schematic of the crystallographic fracture predominant in JIc samples.

It was the original intent of this work to model the effects of microstructure on the

fracture toughness using micromechanistic fracture initiation models. 16116a In order to

apply these models, there needs to be a clear understanding of the fracture initiation site,

and the spacing of these sites must be small enough to affect the local fracture process. It is

clear in KIc samples that many eutectic Y/7' are cracking in the plastic zone, and this

constituent or the porosity may be a fracture initiation site. But in the JIc samples, there is

no indication of any fracture initiation sites ahead of the crack tip. Further, the threedimensional nearest neighbor distances of the porosity and eutectic 7/Y' appear to be too

large to affect the local fracture process based on other applications of these

models. 160'169,215-217 Another problem encountered is that the failure strain in many of

these models158'16n'168 is estimated by incorporating the growth of the initiation site during

the deformation process. But the eutectic 3'/Y' does not change size, and the increase in the

pore size is minimal and not enough to give reasonable failure strain values using the

182
methods of these models. 158'164'168 Therefore, using any of the available

micromechanistic fracture initiation models to correlate fracture toughness and

microstructure in PWA 1480 does not appear to be possible.

However, based on the apparent shear localization in tensile samples, it may be

possible to model the fracture process in these terms. A problem arises in applying this to

the JIc samples, though. The eutectic T/l/does not appear to be a viable site for initiation

of shear localization, as in the tensile samples, nor is the deformation structure known.

Closer examination of these two factors in future work may allow modeling of the fracture

process as a shear localization phenomenon. Using this approach, it may even be possible

to model the apparently constant 5 I.tm distance over which crack growth occurs on a single

{ 111 } plane before switching to another {111 } plane.

183
CHAPTER6

CONCLUSIONS

,

.

.

Hydrogen desorption experiments revealed that hydrogen was not irreversibly trapped

at microstructural heterogeneities after gas-phase charging at 350°C. Estimates of the

hydrogen solubility showed that PWA 1480 contained much more hydrogen after

charging than could be explained by available solubility data of similar alloys. It is

possible that this data is not applicable to single crystal superalloys, but this was not

proven and this area is in need of future research. It is more likely that reversible

lattice trapping can explain the larger than expected hydrogen content after charging.

Further, it is believed that the hydrogen exists uniformly throughout the microstructure

with respect to the T and q/phases, as well as the eutectic _//'y', although the role of the

T/Y interface is not presently well understood.

Uncharged samples initially deformed by slip on isolated single-slip bands resulting in

work softening. This phenomenon was shown to be in common with other

superalloys and precipitation hardened alloys. These bands eventually hardened and

deformation became more uniform across the sample. Subsequent deformation was

primarily confined to the T matrix, which led to a higher work hardening rate. It is

believed that large scale precipitate shearing did not occur, as with most other

superalloys, because the T' precipitates in the PWA 1480 of this study had a higher

flow stress than precipitates of other superalloys.

Superlattice intrinsic stacking faults were seen only occasionally and never in large

enough densities to play a role in the deformation process at room temperature, in

contrast to observations by othcrs in the same system.

184
°

.

The Copley and Kear flow stress model for high volume fractions of coherent, ordered

particles was used to rationalize the CRSS of PWA 1480 at room temperature. Initial

agreement was good, but it was shown that better results would be obtained if a more

accurate value for the precipitate flow stress was available.

It was found that porosity had no effect on uncharged tensile properties and that

eliminating the eutecfic _//_/' did not effect the strength, but did increase the ductility.

The observed tension/compression flow stress asymmetry was in agreement with other

studies, and it was found that hydrogen had no effect on this phenomenon.

. In uncharged samples, it was shown that the the fracture process was initiated by

cleavage of the eutectic 7/7' along [001 } planes, which may have been caused by the

stress created from impinging slip bands upon the eutectic 7/?'. The cracked eutectic

),/y, in conjunction with the stresses caused by the high dislocation density in the

matrix, likely causes shear localization in the precipitates. This localization of slip

could then lead directly to fracture along this slip band, as suggested in other

superalloys and precipitation-hardened alloys. In specially heat treated samples

without eutectic _/'1/, it was discussed how the same process can occur, but now

without the assistance of the eutectic ),/y in initiating shear localization.

, The tensile strength only changed slightly as a function of internal hydrogen content,

but ductility continually decreased with larger concentrations until a threshold level of

about 0.3% strain to failure was reached. It was shown that hydrogen may degrade

properties up to testing temperatures of 500°C, when samples contain an initially high,

uniform concentration of hydrogen. Many of the effects of hydrogen on the tensile

properties and deformation behavior are explainable by the hydrogen-enhanced

localized plasticity (H.E.L.P.) model. It was not believed that tile decohesion model,

185
trap theory or any other embrittlement processes were especially applicable to this

study.

.

It was found that porosity had no effect on charged tensile properties and that the

material was more resistant to hydrogen embrittlement when the eutectic T/_' was

eliminated. However, there was still a significant degradation in properties due to

hydrogen in "clean" material without porosity or eutectic T/]/. It is possible that the

ideas involved in the H.E.L.P. model can help to explain this, but a definitive

explanation for this embrittlement remains elusive.

. The presence of 300 ppm internal hydrogen caused the deformation on isolated singleslip bands to persist to much higher strains. This continuation of localized slip is

consistent with the effect of hydrogen on many other alloys. This observation, along

with a lower flow stress and localized plastic deformation in charged samples, can be

satisfactorily explained by the H.E.L.P. model.

10. The primary difference in the fracture behavior of charged samples was that the

eutectic T/It' cleaved at a much smaller tensile strain and to a greater extent. An

explanation for these observations was offered, involving the impinging slip bands on

the eutectic T/T' and the premises of the H.E.L.P. model. It was believed that the

impinging slip bands would be more intense due to hydrogen's effect on the

dislocation behavior, and thus create higher stresses at the eutectic T/'t" at smaller

strains than in uncharged samples. Experimental evidence for this suggestion was

obtained in the form of higher densities at the eutectic T/7' interface in charged samples

than uncharged samples at equivalent swains.

I 1. The failure process in charged samples was explained in a similar manner to uncharged

samples using the concepts of shear localization and slip band fracture. It was

suggested that samples containing eutectic T/Y' initiated shear localization in the same

186
12.

13.

14.

mannerasunchargedsamples,but now the shearlocalizationwasenhancedby the

presenceof hydrogen.Again, the shear localization is believed to lead directly to slip

band fracture. More strain was obtained in tensile samples without eutectic ),/y, and

thus it was thought that the observed deformation bands were sufficiently intense to

weaken the slip plane and initiate fracture.

The fracture toughness of PWA 1480 was similar to other nickel-based alloys,

although direct comparisons to similar alloys was not possible because no other work

was found on high volume fraction y superalloys. In JIc samples, the eutectic ")'/_' did

not affect the fracture toughness, but eliminating the porosity slightly improved the

toughness. Further, the presence of 300 ppm internal hydrogen did not influence the

crack growth behavior or fracture toughness.

KQ values were relatively lower than comparative JIc values, and there was a large

influence of porosity and eutectic ),/q( on the KQ fractography, as compared to the JIc

fractography. It appeared that differences in the amount of cleavage of the eutectic _,/y'

could be used to explain the differences in the toughness values, but the reasons for

the different cleavage amounts was unknown. It is possible that the effect of

unloading in the JIc tests could explain these observations, but any definite

conclusions were not reached.

Failure occurred in JIc samples by fracture on pairs of {111} planes resulting in

periodic ridges on the fracture surface. Crack growth occurred in the <I 10> direction

and could occur simultaneously along two { 111 } planes or traverse a single { 111 }

plane at a time depending on the initial orientation of the sample. The eutectic ),/',/did

not influence the crack growth behavior or toughness value, but the pores apparently

accelerated the crack growth and slightly reduced toughness

187
15. In summary,it wasshownthatinternalhydrogenseverelydegradesthepropertiesof

PWA 1480,not only at room temperaturebut likely at temperaturesup to at least

500°C. Further,it wasshownthattheresistanceto room temperaturecrackgrowthin

a gaseoushydrogenatmosphereis very low. It wasalsofound thatthepresenceof a

largeamountof eutecticT/7'enhancestheembrittlingaffectof hydrogenandshouldbe

eliminatedwheneverpossible.Thus,it doesnotappearthatPWA 1480is anattractive

candidatealloy for usein situationsthatmaysimulatethetestingconditionsexamined

in this work. More specifically, this would involve room temperatureexposureof

crackedpartsto gaseoushydrogenor low temperature(<500°C) situationswherethere

is a possibility of an uptakeof large quantitiesof internal hydrogen followed by

applicationsof stresses.

188
CHAPTER 7

SUGGESTIONS FOR FUTURE RESEARCH

The present study has pointed out several areas of study on hydrogen embrittlement

of single crystal nickel-base superalloys that are in need of future research. Some of the

areas listed below would best be served by further work on PWA 1480, but many are

research topics applicable to any single crystal superaUoy. Unless specifically necessary as

an experimental variable, the presence of porosity and eutectic 7/7' should be avoided, so

results can be more easily compared to other alloys.

. The results of this study suggested that the role of porosity, eutectic 3'/3" and

composition are minor in determining the amount of hydrogen that can be gas-phase

charged into these types of alloys. In order to learn what factors are important, there

are two areas that should be examined. First, the role of the 7/_' interface in hydrogen

trapping must be better defined possibly through either atom microprobe studies or

high resolution tritium autoradiography. Secondly, further research is needed on

compositional effects on hydrogen solubility and distribution. Both of these areas are

very important in terms of being able to fundamentally design a hydrogen-resistant

single crystal superalloy. Also, to better determine the hydrogen trapping behavior of

a given alloy, a wide range of charging methods should be used including gas-phase

charging at different pressures and cathodic charging.

. There were several indications in this study that hydrogen was enhancing the

occurrence of localized plasticity. Similar observations have been made on other

single crystal superalloys, providing an impetus for more work in this area. Definitive

proof for hydrogen enhancement of dislocation generation and motion and localized

deformation can best be obtained by performing in situ straining experiments on a

189

C-3
.

.

TEM hydrogen charging capabilities. In this way, bulk observations from this study

can be linked to disk_cation behavior and comparisons made on similar in situ tests that

have already been performed on other alloys.

Countless studies have been performed on superalloys in an attempt to understand the

effect of composition on tensile properties, but this area remains in need of more

fundamental-type research. Comparisons of the PWA 1480 of this study with alloys

of similar microstructures, such as CMSX-2 and MAR-M 200, have demonstrated that

the tensile properties and deformation behavior can dramatically vary. It is likely that

composition effects could explain some of these differences, and detailed studies are

needed in areas such as the effect of composition on fault energy, _'/T' misfit and the ),

and _' flow stresses.

It was found that strain localization played a major role in the tensile behavior of

charged and uncharged samples. It is known that once shearing has begun, it is easier

for dislocations to continue slip through ordered precipitates, but litre is understood

about the origins or results of this localization. This study has suggested links

between the macroscopic stress state, stress risers and shear localization on slip

planes, but more evidence for such links is necessary. Observation of slip traces

emanating from eutectic _,/'_' and correlation with localized slip is needed, and may be

accomplished by performing in situ SEM straining experiments in conjunction with

TEM studies of the deformation structure. Finite element modeling of the "t/'Y'

microstructure could assist in interpretation of the development of the deformation

structure and possibly suggest origins for the occurrence of localized slip. Finally,

correlation of these microscopic events to the stress state could be aided by the use of

plane strain tensile tests.

190
. Differences in the fracture behavior of KI¢ and JIc samples remains an unresolved

issue. It is possible that more sectioning experiments on both types of samples at

different stages of testing could be useful. Thin foils taken from the plastic zone of

each sample would also assist in determining any effect of unloading on the

deformation structure.

191
CHAPTER 8

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Effect of Hydrogen on the Fracture and Slip Behavior of Nickel," 1981.

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2OO
196. H.F. Merrick, Met. Trans., 5, (1974), 891-897.

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205. T.M. Pollock and A.S. Argon, Superalloys 1988, eds. D.N. Duhl, et al., TMSAIME, 1988, 285-294.

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April 1976, 180-189.

207. J.P. Tanaka, C.A. Pampillo and J.R. Low, ASTM STP 463, 1970, 191-215.

208. J.J. Gilman, "Fracture," Wiley, New York, 1959, 193-222.

209. A.H. Cottrell, Trans. TMS-AIME, 211, (1958), 192-203.

210. A.N. Stroh, Adv. in Physics, 6, (1957), 418-465.

211. A.H. Windle and G.C. Smith, J. Metals Sci., 2, (1968), 187-191.

212. R.W. Hertzberg, "Deformation and Fracture Mechanics of Engineering Materials,"

Wiley and Sons, New York, 1989, p. 320.

213. N.E. Dowling and J.A. Begley, ASTM STP 590, "Mechanics of Crack Growth",

1976, 82-103.

214. J.W. Hutchinson and P.C. Paris, ASTM STP 668, 1979, 37-50.

215. W.M. Garrison, Met. Trans., 17A, (1986), 669-678.

216. S. Lee, L. Majno and R.J. Asaro, Met. Trans., 16A, (1985), 1633-1648.

217. R.E. Stohz, N.R. Moody and M.W. Perra, Met. Trans., 14A, (1983), 1528-1531.

201
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APPENDIX B

DEFORMATION AND FRACTURE BEHAVIOR OF <130> SAMPLES

Tensile tests were performed to determine the effect of hydrogen on the deformation

and fracture behavior of samples oriented for single slip, in contrast to the [001] samples

that undergo multiple slip. All experimental procedures were the same as for [001] samples

described in section 3.3.1. Single slip oriented samples were obtained by machining slabs

perpendicular to the [001] crystal growth axis, as shown in Figure B.la. These transverse

oriented samples were all taken from the same single crystal slab and were 2 ° off <130>, as

determined by the Laue back reflection technique and shown in Figure B. 1b.

t [001]

I

i

I

<130>

I

I

I

[ira] [] 1l]

[001] [013] [011]

(a)

Figure B. 1

(b)

(a) Orientation of <130> tensile samples relative to [001] single crystal

slab. (b) Equivalent orientation in stereographic triangle at point x.

207
The average tensile properties as a function of hydrogen are shown in Table B. 1,

and a typical stress-strain curve is shown in Figure B.2. Comparative tensile properties

and stress-strain curves for [001] oriented samples are also shown in Table B.1 and Figure

B.2, respectively. These tensile properties are consistent with other investigations that

have studied the effect of orientation on PWA 1480 and found a lower yield strength for

single slip orientations. 1'2 It is believed that this is purely an orientation effect since the

CRSS for both the [001] and <130> oriented samples is close to 450 MPa. The tensile

stress-strain behavior for the two orientations is qualitatively similar, and this can be related

to the single slip behavior that the [001] orientation exhibits at low strains, as discussed in

section 5.3. In the [001] orientation, deformation initiates and proceeds in narrow bands of

single slip until these bands harden either by multiple slip initiation or debris hardening.

Similarly, the <130> orientation stress-strain behavior can be explained by initiation of

single slip resulting in work softening. As the tensile axis rotates during deformation, a

duplex slip orientation is reached around [I 14] in Figure B.lb, and the work hardening rate

increases up to failure. Failed <130> samples were sectioned and examined on the TEM,

Table B. 1 Tensile properties for the <130> and [001] orientations as a function of

hydrogen.

[001 ] uncharged

[001] charged

< 130> uncharged

<130> charged

(Iys (MPa)

1112

1001

929

906

Otrrs (MPa)

1153

1011

929

906

ef (%)

24.3

6.9

R.A. (%)

25.3

21.2

208
I100

vi

II900

r_

.E

tt_

Engineering Strain (%) Engineering Strain (%)

(a) (b)

Figure B,2 Typical stress-strain curves for (a) <130> samples and (b) [001] samples.

but the dislocation density was too high to determine the active slip systems. It was

possible, however, to determine the f'mal specimen orientation of <112>, as predicted from

typical fcc single crystal deformation. 3 The rotation of the sample during deformation is

supported by fractography observations, which showed that the initial circular crosssection had become oval, Figure B.3, as observed in other samples oriented for initial

single slip. 3

The fracture surface was macroscopically much flatter in the <130> samples, and

fracture occurred primarily on a single { 111 } plane, as opposed to multiple { 111 } planes in

[001 ] samples. The microscopic ductile fracture of the Y/'/microstructure was not affected

by orientation, as seen by comparing Figures B.4 and 4.30b.

209
ORIGINAL PAGE

BLACK AND WHITE PHOT.OGP,.AP_H

3(X) l_m

Figure B.3 Oval cross-section of uncharged <130> tensile sample.

Figure B.4 Ductile fracture is not affected by orientation. Compare with Figure 4.30b.

210
A significant difference in the fracture behavior of the two orientations was that the

eutectic y/y in the <130> oriented samples did not fail by crystallographic cleavage. Figure

B.5 shows that the eutectic y/y' failed in a more brittle manner than the surrounding y/y

microstructure, but the characteristics of crystallographic cleavage are not present, such as

river fines or a definitive fracture plane. It is possible that the {001 } cleavage plane of the

eutectic y/'y' is not favorably oriented for cleavage in these samples. Assuming that the

impingement of slip bands on the eutectic y/_/initiates cleavage, it can be argued that the

normal stress created by the slip bands is not as high on the {001 } cleavage plane in the

<130> oriented samples, as in the [001] samples.

Figure B.5

_, 5 lam

Brittle fracture of eutectic y/y in < 130> oriented samples.

BLACK

The effect of hydrogen on the tensile behavior of the <130> samples was similar,

but more pronounced, compared to the [001] samples. For instance, the strain to failure

was greatly reduced with little change in the R.A., as previously observed, but the

magnitude of the values was much larger in the <130> samples. This provides assurance

211

ORIGIHAE PAGE

AND WHITE PHOTOGRAPH
that this effect was not merely an artifact of experimental error in the [001 ] samples and

gives further support to the idea that hydrogen localized the plastic flow in tensile samples.

Previously, it was shown that hydrogen promoted cleavage of the eutectic T/q(, and there

was evidence in the fractography of these samples that hydrogen embrittles this constituent.

The type of fracture shown in Figure B.5 was observed to a much greater extent in

hydrogen-charged samples. This may be because the fracture has a more brittle appearance

in charged samples and was easier to observe, rather than an actual increase in the area

fraction of eutectic TP/' on the fracture surface, as in the [001] samples. As in the [001]

samples, the presence of hydrogen did not change the microscopic ductile fracture process.

References

.

.

.

D.M. Shah and D.N. Duhl, Superalloys 1984, eds. M. Gell, et al., TMS-AIME,

1984, 105-114.

T.G. Gabb, NASA TM 100269, "The Cyclic Stress-Strain Behavior of a Single

Crystal Nickel-Base Superalloy," 1988.

R.W.K. Honeycombe, "The Plastic Deformation of Metals, 2nd ed.," Edward

Arnold Publ., London, 1984, pp. 10-32.

212
APPENDIX C

SHORT ROD Klc TESTING

C. 1 Background

There are a growing number of researchers studying various smaller, alternative

fracture toughness specimens in order to make the determination of KIC a simpler, more

economical exercise. The size of the current KIC specimens has always been a drawback,

as well as the tedious testing method, especially fatigue pre-cracking. JIc testing offers a

smaller specimen size, but the testing procedure is even more complicated. The most

common alternative candidate to replace the current specimens is the short rod or short bar

specimen. To add to the knowledge base in this field, these specimens will be used to

measure the fracture toughness of PWA 1480 and then compared to KQ values and valid

ASTM E 813-871 Jic values. This work primarily will be done by Jim Lucas at Sandia

National Labs in Livermore, Calif. This literature review will examine the motivation and

background behind these alternative specimens, and then look at experimental studies that

have compared the KIC values obtained in short rod specimens to valid ASTM values.

L.M. Barker is one of the leading proponents of short rod or short bar specimens and has

written many articles on the theory and testing methods involved. 2-4 A drawing of a short

rod specimen is shown in Figure C.i. One of the primary advantages of this design is that

fatigue pre-cracking is not necessary because of the chevron notch design, which allows

initially stable crack growth. When considering the use of small specimens to measure

plane strain fracture toughness, there are two problems that must be overcome. 2 First, the

condition of plane strain along a vast majority of the crack must be maintained. This

problem is overcome by the chevron notch design, which forces the crack tip to be wellconstrained. The second dilemma is that the material response changes from an elastic

nature to an elastic-plastic one.

213
y

I,

Figure C. 1

1

l

re) y - 0 olone.

Knife-edge loaded chevron notched short rod specimen. 5

Barker's rederivation of the compliance equation to include plasticity will now be

considered. 2 The short rod testing method provides a measure of the amount of energy per

unit crack area that is required to advance a steady state crack under plane strain conditions.

At first it appears that this definition violates the premise upon which KIC or JIc is

calculated; that is measuring initial crack growth, and not steady state crack growth.

However, Barker 2 goes through a lengthy argument stating that KIC and hc testing is

actually measuring steady state crack growth. Accepting this, now consider the effect of

plasticity on the specimen behavior. First, the stress intensity parameter must be derived

for the short rod geometry: 2

Prnax

KIC = B3/2(I _ v2)1/2 f(a/b) (C-l)

where P is the maximum load applied, B is the specimen thickness and f(a/b) can be

replaced by a constant for each specimen geometry. There are other expressions used in

tile literature that are very similar, 6.7 but equation C- 1 is generally accepted. There also is

some debate over the correct v:due of f(a/b), but after a comprehensive revicw. Newnlan 7

214
gives the currently accepted values for many specimen geometries. To correct for the

plasticity, Barker 2 uses a parameter, p, which can be calculated from the load-displacement

chart upon unloading and reloading the specimen during the test. The modified KIC

equation becomes: 2

( 1 + p-_U2

KIC = k']-'__ p j KQ

(C-2)

where KQ is the previously calculated KIC (C-l) from LEFM principles.

An important point to note is that the load used to calculate KIC in (C-l) and (C-2) is the

maximum load, as seen in Figure C.2. 7 The dashed lines represent the stress intensity

factors as a function of crack growth at various applied loads. The fact that they drop and

form a minimum with increasing crack growth is important. Figure C.3 shows that the

stress intensity for the short rod specimen (solid line) is very large when a=ao. 7 When the

crack begins to grow, the stress intensity drops rapidly and reaches a minimum between ao

and at. This minimum, ac, is constant for a given specimen geometry. As the crack

reaches al, the stress intensity rises and meets that for a normal through cracked specimen

(dashed line). Looking again at Figure C.2, it is seen that a small load, Pt, can initiate a

crack. This crack grows until it intersects the "ideal" crack growth resistance curve (solid

line) at point A. This solid line represents the curve for brittle materials. The load must

now be increased to get crack growth to points B and C. When the maximum load is

reached the crack driving force curve is tangent to the crack resistance curve at point D.

Thus, the K value obtained is KIc and the only load of concern is Pmax.

Many researchers have performed tests on short rod samples analyzing the effect of

specimen size, equation parameters and plasticity. 4'6'8-12 It seems that specimen thickness

may be as small as 1.25(Kic/_ys) 2, although this is still being studied. 4 The other

215
I : Pax}. I_

It_ ,

II| \ '\l'llnlmum J.

A,_ \ _- \ /%"

I Pl " P2 4 P3 • PmO_

oo

CrOCk lenqth. 0

Figure C.2 Stress intensity vs. crack length. 7

ORIGINAL PAGE IS

OF POOR QUALITY

Figure C.3

P

./_ / I L'--o ---__1

1 I

oo 01

_.

CroEk-len9tn-to-wlOtn rotlo, O/N

Comparison of normalized stress-intensity factors for chevron-notch and

straight-edge, through-cracked specimens. 7

specimen dimensions are scaled with the thickness and initial crack length, as in normal

fracture toughness testing, but standard ratios have yet to be established. As previously

mentioned, Newman has compiled the data on various equation parameters and

recommended the best values. 7

216
In comparing short rod Kzc values to ASTM E 399-8313 Kic values, there are many

instances where comparative values are well within 10% of each other for a wide variety of

materials, such as high Cr cast irons, 6 various A1 alloys 2.8 and Ti-6A1-4V. 4 However, in

high toughness materials, excessive plasticity is a problem, and the comparison becomes

somewhat worse. 7d° Many of these materials have rising resistance curves, which causes

the estimate of KIC to be too high, as seen in Figure C.4. 7 The minimum stress intensity

curve for Pmax no longer intersects the resistance curve at a tangent, as in Figure C.2. This

leads to overestimation of KIC, as well as a new specimen width effect.

ORIGINAL PAGE !$

OF POOR QUALITY

Figure C.4

I

t

t

II

I_ _, %_

I\ \'-..&/.(_ ,_

I \\/--'_// P_

o o

Cr_k length, o

Fracture of "rising resistance curve" material using a chevron-notched

specimen. 7

The problem of the rising resistance curve for some materials is one of the major

drawbacks still associated with the short rod test. It seems that this method works well for

relatively brittle materials, but its future for highly ductile materials is still being

researched. 7 ASTM is currently reviewing standardization, and more research can only

lead to more accurate KIc values along with a better understanding of the phenomena

involved. It is not obvious at this point whether the short bar testing method will provide

valid Kit" values for PWA 1480, but that is not the only goal of this research. The object is

217
to understand and apply the test method, and examine the results in light of the current

knowledge in the area.

C.2 Experimental Procedures

Chevron-notched, short rod fracture toughness specimens were tested as a function

of microstructure and hydrogen at Sandia National Laboratories by Jim Lucas. Specimens

were machined from fractured halves of KQ and JIc specimens in the geometry shown in

Figure C. 1. Because of the configuration of the chevron notch, no pre-craeking was

necessary, and loading was accomplished by special knife-edge grips at point P in Figure

C. 1. Short rod specimens that were to contain hydrogen were machined from charged K1c

specimens, and subsequently re-charged in the manner given in section 3.1.5. This led to

higher hydrogen concentrations in the short rod specimens than in the KQ specimens. As a

result, the toughness values did not agree well, and subsequent samples machined from

charged J1c samples were outgassed prior to machining into short rod samples. The KIc

values and the fractography of the short rod specimens were compared to full size KQ and

JIc specimens.

C.3 Results and Discussion

The results of short rod KIC tests compared to ASTM E 399-83 KQ tests are shown

in Table C. 1 for two orientations. For uncharged specimens the agreement is very good,

but the charged short rod values are consistently lower than the ASTM E 399-83 values.

This is probably due to higher hydrogen concentrations in the short rod samples, as

explained above. Short rod tests are currently underway on samples machined from

broken [001l J)c samples and will be performed as a function of porosity, eutectic T/Y' and

hydrogen.

218
TableC.1

Orientation

[001]

[001]

transverse<130>

transverse<130>

ShortRodKI¢ values as a function of orientation and hydrogen.

I

Condition

no hydrogen

with hydrogen

no hydrogen

with hydrogen

ASTM E 399-83 KQ

122 MPaqm

103

133

100

Short Rod KIc

121 MPa4m

82

129

83

The close agreement in uncharged samples is further support for the short rod

testing technique. It also shows that rising resistance curyes are apparently not a problem

in PWA 1480, and completed JIc tests show that the resistance curve is not as steep as in

Figure C.4. Based on this, the short rod toughness values obtained from the J1c samples

should correlate well. However, the discussion in section 5.6 concerning the discrepancy

in JIc versus KIe values must be kept in mind. It is hoped that the results of the short rod

tests will help in understanding these differences.

There were no differences in the short rod fractography compared to the full size KQ

fractography. The fracture occurred along { 111 } planes, and the roughness of the surface

appeared similar to the full size KQ samples. It will be interesting to see if the subsequent

short rod tests show the same fractography differences that the full size KQ and JI¢ samples

did. Like the tensile and full size KQ samples, cleavage of the eutectic 7/7' was seen in the

[001] orientation but not in the transverse <130> orientation. It is believed that this is

because the eutectic 7/Y is no longer favorably oriented for cleavage along the {001 } planes

in the transverse oriented samples.

219
ORIGINAL PAGE

BLAC.K KI_,£). _t:-IL]'.F., P_HO.T_OGRAPJ-J

(a)

Figure C.5

25 I.tm

(b)

Fractography of short rod fracture toughness specimens at (a) low

magnification and (b) a higher magnification.

220
References

°

,

.

,

.

.

7.

,

.

12.

13.

ASTM E813-87, "1987 Annual Book of ASTM Standards", ASTM, Phila, Pa.,

1987, pp. 686-700.

L.M. Barker, Terra Tek Report RT 78-6R, "Theory for Determining Kic from Small,

non-LEFM Specimens Supported by Experiments on Aluminum," 1978.

L.M. Barker, Fracture Mechanics: 13th Conference, ASTM STP 743, eds. R.

Roberts, ASTM, 1981, 456-475.

L.M. Barker, Chevron Notched Specimens, ASTM STP 855,Chevron-Notched

Specimens: Testing and Stress Analysis", eds. J.H. Underwood, S.W. Freiman and

F.I. Baratta, ASTM, 1984, 117-133.

I.S. Raju and J.C. Newman, Jr., ASTM STP 855, "Chevron-Notched Specimens:

Testing and Stress Analysis", eds. J.H. Underwood, S.W. Freiman and F.I. Baratta,

ASTM, Phila., Pa., 1984, pp. 32-48.

S.B. Biner, J.T. Barnby and D.W.J. Elwell, Int. J. Frac., 26, (1984), 3-16.

J.C. Newman, Jr., Chevron Notch Specimens, eds., J.H. Underwood, S.W.

Freiman and F.I. Baratta, ASTM, Phila., 1984, 5-31.

K.R. Brown, ASTM STP 855,"Chevron-Notched Specimens: Testing and Stress

Analysis", eds. J.H. Underwood, S.W. Freiman and F.I. Baratta, ASTM,

Philadelphia, 1984, 237-254.

J.L. Shannon, Jr. and D.G. Munz, ASTM STP 855, "Chevron-Notched Specimens."

Testing and Stress Analysis", eds. J.H. Underwood, S.W. Freiman and F.I. Baratta,

ASTM, Philadelphia, 1984, 270-280.

J. Eschweiler and D. Munz, lnt. J. Frac., 20, (1982), R131-R133.

R.T. Bubsey, D. Munz, W.S. Pierce and J. J.L. Shannon, Int. J. Frac., 18, (1982),

125-132.

J. Eschweiler, G. Marci and D.G. Munz, ASTM STP 855,"Chevron-Notched

Specimens." Testing and Stress Analysis", eds. J.H. Underwood, S.W. Freiman and

F.I. Baratta, ASTM, Philadelphia, 1984, 255-269.

ASTM E399-83, "Annual Book of ASTM Standards", ASTM, Phila., Pa., 1983, pp.

518-553.

221
APPENDIX D

HYDROGEN-INDUCED CRACK GROWTH

D. 1 Background

The determination of a material's susceptibility to sustained-load cracking in a

hydrogen atmosphere is very important to design considerations. Through the use of

wedge-opening-loading (WOL) specimens, it is possible to measure the stress intensity

threshold (KTH or KIscc) below which sub-critical cracks will not grow to a critical size in

a hydrogen atmosphere. 1 Most investigations and modeling attempts have focused on

crack growth in aqueous solutions and not gaseous hydrogen. However, the mechanisms

of hydrogen-induced crack growth will be similar, because in both cases the presence of

hydrogen at the crack tip will be the driving force for crack propagation. This section will

first describe the type of specimen used in hydrogen-induced crack growth tests, along

with the corresponding linear elastic fracture mechanics equations. Next, typical crack

growth behavior is discussed and models that relate the stress intensity threshold to the

microstructure and hydrogen concentration are covered. Finally, experiments performed

on fcc alloys will be discussed in an attempt to better understand the behavior observed in

PWA 1480 in this study.

The specimens typically used for hydrogen-induced crack growth are the modified

WOL specimens suggested by Novak and Rolfe for environmental testing, t The specimen

is self-stressed with a bolt, and there is constant displacement loading so that the initial K1

value decreases to a threshold value, Kth. as the crack propagates. Currently, there is no

ASTM standard concerning WOL specimen dimensions or test procedures for determining

crack growth threshold values, although standardization efforts are underway. _ The

222
specimen dimensions proposed by Novak and Rolfe and commonly used for a one-inch

thick specimen are shown in Figure D. 1.

-_ W t

w

i

IX

'

z

r-- ----_

---_ os i_--

Geometry in inches: B = 1.00 W = 2.55 (H/W) = 0.486

W1 = 3.2 Hp = 1.00

Ds = 0.625 Dp = 0.70 (a/W) = 0.302

a = 0.77 N = 0.094

Figure D.1 One-inch thick modified WOL specimen proposed by Novak and Rolfe. 1

Slight modifications can be made to the dimensions shown in Figure D.1; Hudak and

Saxena 3 have given compliance expressions for modified specimens. Stress intensity

values can be calculated from: 1

PC3(a/w)

K = (D-l)

(BBN)I/2a 1/2

223
where P is the load, C3(a/w) is a function of (a/w), a is crack length, W is specimen depth,

B is specimen thickness and BN is specimen thickness with side grooves. C3(a/w) is

analogous to f(a/w) in ASTM E 399-834 and is given by: 1

C3(a/w ) = [ 30.96(a/w) - 195.8(a/w) 2 + 730.6(a/w) 3 - 1186.3(a/w) 4 + 754.6(a/w) 5] (D-2)

for samples with a height to depth ratio (H/W) of 0.486.

independently determined from the compliance: 3

BEVEL 16 24

P [ 1 - (a/w)] 2 [ 1 - (a/w)]

The crack length can be

4.5 In[ 1 - (a/w)] + 32.12 (D-3)

where E is Young's modulus and VLL is the displacement measured along the load line.

Crack growth specimens tested in a gaseous hydrogen atmosphere generally obey a

three-stage crack growth rate versus stress intensity curve, as shown in Figure D.2. At

low K values, the crack growth rate is dependent on the stress intensity, as well as

temperature, pressure and environmental conditions. In stage II, da/dt is independent of K,

and crack growth is associated with the diffusion of hydrogen through the lattice. The last

stage occurs at K values approaching KIc, where crack growth primarily is controlled by

stress intensity effects.

Fatigue pre-cracked WOL specimens axe loaded in a non-aggressive atmosphere to

a stress intensity value ranging from 30% to 90% of the KIC value. 1 The specimens are

then immersed in the hydrogen atmosphere, and hydrogen-induced crack growth occurs

causing both the stress intensity and crack growth rate to decrease. The stress intensity at

which crack growth stops (KTH) is considered a material property for the specimen

conditions and hydrogen pressure used. 6

224
Figure D.2

g

K_A

t

KIc

K

Three stages of hydrogen-induced cracking under sustained load. 5

There are several models that argue when a critical stress is exceeded ahead of the

crack tip, fracture will occur. 7q° For samples tested in hydrogen, the stress ahead of the

crack tip can lead to a higher local hydrogen concentration: ll

[ oN. ]

c =Co exp[_3- j (19-4)

where Co is the equilibrium concentration in the absence of stress, VR is the partial molar

volume of hydrogen, R is the gas constant and (I i is the summation of the principal stresses

such that _r-d3 is the hydrostatic component of the applied stresses. The high concentration

of hydrogen will lower the critical stress needed for fracture 7"9 and can be illustrated by the

simple relation: 9,12

(_" = (3: - AC H (D-5)

where A and Co ,are material dependent constants and CH is the hydrogen concentration.

Consequently, the material in hydrogen fractures at a lower critical stress and KTIt is much

less than KIc. In some cases a critical distance is introduced into the model such thai the

225
critical stress must be exceeded over this distance. 6'8'13'14 This relates KTH tO

microstructural features, such as grain size or inclusion spacing.

The vast majority of hydrogen-induced crack growth studies have been performed

on bcc ferrous alloys 2'6,9 and little work has been done on fcc or nickel-base alloys. 1318

Table D. I lists some results for two 7' strengthened fcc superalloys (IN 903 and JBK 75)

and a nickel-base superalloy (IN 718) tested in a gaseous hydrogen atmosphere. The

results for IN 903 and IN 718 are especially applicable to this study since it was shown in

section 5.1 that the diffusion constants for hydrogen in these materials closely matches that

of hydrogen in PWA 1480. However, the microstructures are significantly different, and

to the author's knowledge there have been no other studies on high volume fraction 7'

superalloys. The ratio of the fracture toughness in air (KIc) to the threshold in hydrogen

(KTH) is typically about 3-5.5. The KTH dependence on hydrogen pressure is illustrated in

the IN 718 series B alloys in Table D. 1. KTH reaches a plateau value at a certain pressure,

and this type of behavior also has been observed in IN 903.13

Alloy

IN 903

JBK-75

IN 718 L

IN 718 A

IN 718 B1

IN 718 B2

IN 718 B3

IN 718 B4

Table D. 1

0.2% Y.S.

1080 MPa

Available KTrl values of similar alloys.

KIC

90 MPa_/m

KTH

30 MPaqm

Pressure

207 MPa

924

1331

1317

1317

1317

1317

m

95

78

119

119

119

119

65

21

14

80

64

42

42

207

34.5

34.5

5

8

34.5

72

Reference

19

18

16

15, 17

15, 17

15, 17

15, 17

15, 17

226
D.2 Experimental Procedures

The dimensions for the WOL specimens used in hydrogen-induced crack growth

tests are shown in Figure D.3. The important ratios for compliance calculations were

preserved from Novak and Rolfe's original specimen dimension suggestions, l It was

extremely difficult to tap this material, so electrodischarge machining (EDM) of the two

holes was done. Specimen preparation was done at Carnegie Mellon and testing was done

at Sandia National Labs by Neville Moody. This procedure has been developed

extensively there, and Sandia possesses capabilities for testing in high-pressure hydrogen

that are not available at Carnegie Mellon.

75" 5

tl

Counterbore

_------. 0.375" - 24UNF-2B

'

0.625" dia. _d

1.40"_ 4" 0.425_

l

-'_ 1.75"

Figure D.3 Actual mcasuren_cnts of WOL specimens tl,,;c_f.

227
The samples were fatigue pre-cracked, and subsequently a torque-wrench was used

to apply a KI value to the specimen via the bolt. Crack growth was monitored with a

KRAK-GAGE TM mounted on the surface of each sample. All tests were performed at

room temperature using A-286 stainless steel primary vessels that were stored within

secondary safety vessels. The primary vessel could hold multiple specimens and was

capable of withstanding hydrogen pressures of up to 200 MPa. Hydrogen

dehumidification and purification was accomplished with a combination of sieve beds and

cold traps, and a series of pressure-vacuum cycles was run to purge the gas lines of

impurities. Each specimen was equipped with individually calibrated load cells similar to

ones used by Loginow and Phelps. 2° The loaded samples were placed in the test cell, and

the loads and crack lengths were continuously monitored. Average crack velocities as

small as 10 dl m/see could be resolved and crack arrest was defined as no detectable crack

growth for four months. The threshold stress intensities were determined from the final

load and crack length using the assumption of a rigid bolt and equations D-1 to D-3.

D.3 Results and Discussion

Based on the data in Table D.1 and the fracture toughness of PWA 1480, a KTH

value of between 25 MPa'4m and 40 MPa4m was expected. The first set of samples

immediately fractured when hydrogen gas was introduced into the system. The initial K

value was about 75 MPa4m, so this was reduced on the second set of samples to 50

MPax/m and 35 MPa_/m. These samples also fractured quickly, but some data was

obtained. Figure D.4 shows a typical profile of crack growth rate versus stress intensity.

The crack growth rate increased in the beginning of the test and leveled off just before

fracture. Normally, in a hydrogen-induced crack growth test, the crack begins growing

rapidly at the onset but quickly slows down as the stress intensity decreases, as shown in

l=igure D.2. This is followed by a long period of steady crack growth before the crack

228
finally slows and stops. These stages of crack growth were never observed, and the crack

_ew rapidly and independently of the stress intensity level.

-4.4

-4.6

-4.8

"1:3

o -5.0

-5.2

-5.4

0 [] 0 O0

0

0

[]

o 50

o

beginning of test _ [] []

I I I

20 30 40

Stress Intensity (MPaqm)

Figure D.4 Typical crack growth rate versus stress intensity in PWA 1480.

Figure D.5 shows the fractography of one of the samples. It is evident from the

7/7' outline in Figure D.5c that the crack is growing macroscopically along {001 } planes.

This type of fracture previously has been observed in these types of alloys in hydrogen

charged 21 and uncharged samples. 22 Fracture along {111} planes was observed, similar

to the fracture in tensile and fracture toughness samples, however the {001 } fracture was

predominant. Cube slip has not been reported at room temperature in these alloys, z3 so all

of the deformation must be occurring by { 111 } slip in the y matrix. This is similar to

tensile samples where the majority of deformation was in the y matrix, but now the final

fracture is probably occurring through the 7 matrix on the {001 } planes instead of through

both phases on { 111 } planes.

229
ORIGINAL PAGE IS

OF POOR QUALITY

ORIGINAL PAGE

BLACK A!XlD WHITE PHOT_OGRAP_H

25 gm

(a) (b)

25 lain

(c)

Figure D.5 Fractography of hydrogen-induced crack growth samples. (a) 100x (b)

500x (c) 5000x, showing {001} faces ofy precipitates.

230
Thecausefor thisrapidfailureis unknown,but someaspectsbecomemoreclearby

examiningtherole of hydrogendiffusion during thetest. The final fractureis occurring

within afew minutesof theintroductionof hydrogengas,andit is not believedthat long

rangediffusion of hydrogeninto the alloy is occurring. An approachsimilar to the one

usedby Johnson24canestimatethedistancethathydrogenhaspenetrated.Using theIN

903 diffusion constantat room temperatureof 1.4x 10-I1 cm2/sand an averagecrack

velocity of 2 x 10"5m/s, Johnson'sanalysisindicates that hydrogencanonly penetrate

about 10A(aboutfour atomic spacings). This is in close agreementwith the resultsof

PerraandStoltzz5 using the same approach for JBK-75. It also has been shown that the

maximum local stresses are reached ahead of a crack tip, z6 and this has been suggested as

the site of embrittlement. 12'27 For a stress intensity of 35 MPa_/m and a flow stress of

1100 MPa, the peak stress occurs approximately 200 Ixm ahead of the crack tip. This is

five orders of magnitude farther from the crack ill) than the distance hydrogen can diffuse

by bulk diffusion. This analysis suggests either that embritflement is occurring near the

surface or that bulk diffusion is not the controlling parameter.

The diffusion analysis ignored the effect of stress on diffusion, but Baskes and

Melius 28 have shown that hydrogen diffusivity is not greatly enhanced in the presence of

crack tip stress fields in fcc metals. Hydrogen transport by dislocations is another possible

transport mode, 29 but given the almost immediate rapid crack growth that occurred, this is

not likely to be an applicable mechanism. It remains that some type of surface effect is

controlling crack initiation and propagation, and further studies are needed to examine this

possibility. Experiments are planned using the same experimental setup but with an initial

applied stress intensity of about 20 MPa_/m.

231
References

1. S.R. Novak and S.T. Rolfe, J. Materials, 4, (1969), (1969).

2. R.P. Wei and S.R. Novak, J. Test. Eval., 15, (1987), 37-75.

3. A. Saxena and J. S.J. Hudak, Int. J. Frac., 14, (1978), 453-467.

4. ASTM E399-83,1983 Annual Book of ASTM Standards, Phila., Pa., 1983, pp.

518-553.

5. R.W. Hertzberg, "Deformation and Fracture Mechanics of Engineering Materials,"

Wiley and Sons, New York, 1989, 431-453.

6. K.N. Akhurst and T.J. Baker, Met. Trans., 12A, (1981), 1059-1070.

7. R.A. Oriani and P.H. Josephic, Acta Met., 25, (1977), 979-988.

8. P. Doig and G.T. Jones, Met Trans., 8A, (1977), 1993-1998.

9. W.W. Gerberich and A.G. Wright, Environmental Degradation of Eng. Mat. in

Hydrogen, eds. M.R. Louthan, Jr., R.P. McNitt, R.D. Sisson, Jr., Virginia Tech.

Publ., Blacksburg, Va., 1981, 183-206.

10. S.V. Nair and J.K. Tien, Met. Trans., 16A, (1985), 2333-2340.

11. J.C.M. Li, R. Oriani and L.S. Darken, Z. Phys. Chem. Neue Folge, 49, (1966),

271-290.

12. C.L. Briant, H.C. Feng and J. S.J. McMahon, Met. Trans., 9A, (1978), 625-?

13. N.R. Moody, R.E. Stoltz and M.W. Perra, Scripta Met., 20, (1985), 119-123.

14. R.E. Stoltz, N.R. Moody and M.W. Perra, Met. Trans., 14A, (1983), 1528-1531.

15. R.J. Walter and W.T. Chandler, Environmental Degradation of Eng. Mat. in

Hydrogen, eds. M.R. Louthan, Jr., R.P. McNitt and R.D. Sisson, Jr., Virginia Tech

Publ., Blacksburg, Va., 1981, 513-522.

16. P.M. Lorenz, Boeing Document D2-114417-1, The Boeing Co., Seattle, Wash.,

1969.

17. R.P. Jewett, R.J. Walter, W.T. Chandler and R.P. Frohmberg, NASA CR-2163,

"ttydrogen Environment Embrittlement of Metals," 1973.

18. R.E. Stoltz and A.J. West, Hydrogen in Metals, eds. I.M. Bemstcin and A.W.

Thompson, TMS-AIME, 1980, 541-553.

1c). N.R. Moody, M.W. Pcrra and S.L. Robinson, Scripta Met.. 22, (1988), 1261-

1266.

232
20. A.W. Loginow and E.H, Phelps, Corrosion, 31, (1975), 404-412.

21. C.L. Baker, J. Chene, I.M. Bernstein and J.C. Williams, Met. Trans., 19A, (1988),

73-82.

22. J. Telesman and L.J. Ghosn, Superalloys 1988, eds. D.N. Duhl, et al., AIME,

1988, 615-624.

23. R.V. Miner, R.C. Voigt, J. Gayda and T.P. Gabb, Met. Trans., 17A, (1986), 491-

505.

24. H.H. Johnson, Hydrogen in Metals, eds. I.M. Bernstein and A.W. Thompson,

TMS-AIME, 1974, 34-49.

25. M.W. Perra and R.E. Stoltz, Hydrogen in Metals, eds. I.M. Bernstein and A.W.

Thompson, TMS-AIME, 1981, 645-653.

26. J.R. Rice and M.A. Johnson, Inelastic Behavior of Solids, eds. M.F. Kanninen, et

al., McGraw Hill, New York, 1970, 641-672.

27. A.R. Troiano, Trans. ASM, 52, (1960), 54-80.

28. M.I. Baskes and C.F. Melius, Z. Phys. Chem. Neue Folge, 116, (1979), 289-299.

29. J.K. Tien, S.V. Nair and R.R. Jensen, Hydrogen in Metals, eds. I.M. Bernstein and

A.W. Thompson, TMS-AIME, 1980, 37-56.

233
APPENDIXE

TRITIUMAUTORADIOGRAPHY

E.1 Background

The distribution of hydrogenin metalscanbe studiedby autoradiographyusing

tritium asa tracer. Tritium autoradiographyfor metallurgicalpurposeshasbeendoneona

limited scale for more than20 years.1'2 The techniqueis not widely usedin hydrogen

embrittlement studies possibly becauseof the complicated experimental procedures

involved andthe considerableeffort requiredto perfectthe technique. Thereis also the

safetyaspectto consider,althoughtritium doesnotposeseriousproblemsunlessreleased

in largequantities. Thevastmajorityof studieshavebeenperformedonferrousalloys,3-11

andonly two haveexaminednickel-basealloys.12,13

Tritium autoradiographyinvolvesplacingtritium-chargedsamplesin contactwith a

photographicfilm so thatuponthedecayof tritium to helium-3,theemitted13particlecan

interactwith the silver halidegrainsin theemulsion.1 During subsequentphotographic

development,thosegrainsthat havebeenstruckby 13particlesareconvertedinto silver

Filaments, and the unexposed grains are dissolved. The silver grains can then be examined

on the microscope and correlated to the underlying microstructure.

The introduction of tritium into the samples is similar to hydrogen charging and can

be accomplished by either cathodic charging 36'8'11 or gas-phase charging. 7'9.1° In

cathodic charging, tritiated water is simply added to the molten salt, and the same

conditions are used as in hydrogen charging. 1'_ Likewise, gas-phase tritium charging

involves the same pressures and temperatures as in gas-phase hydrogen charging. 15

After tritium charging, the photographic emulsion is applied; there are a variety of

emulsions that can be used depending on the results desired. For observations wherc sub234
micron resolution is not required, stripping films such as AR-10 may be used. 7'16.17 The

advantages of the stripping film are that it provides a uniformly thick emulsion and gelatin

layer, and is relatively easy to use. Care must be taken, however, not to introduce artifacts

caused by mishandling the film. This technique has been successful when used in

conjunction with optical microscopy 7.16.17 but not with scanning electron microscopy

(SEM). 7 This stems from the thick gelatin layer (-4p.m) that remains on the sample after

photographic developing and results in charging effects and foggy pictures on the SEM.

The other type of emulsion commonly used is a liquid emulsion; Ilford LA or

Kodak NTB-2 are the two most popular brands. 3-6,81°,12A3 The emulsion comes in the

form of a gel, and there are many procedures that must be followed to prepare the emulsion

for application and ensure a uniform coating thickness. There also is a gelatin layer in these

emulsions, but with careful preparation the thickness of this layer can be minimized. 7.18

This allows easier observation on the SEM or TEM, and resolutions approaching 0.1 lam

have been reported.18

To avoid diffusion of tritium during exposure that would result in ambiguous

results, 19 exposure usually takes place at liquid nitrogen temperatures. 1,18 The length of

exposure and developing conditions are among the variables that must be adjusted to

achieve the best results.

Tritium autoradiography has been used on ferrous alloys to determine trapping at

heterogeneities, such as particle interfaces, grain boundaries and dislocations. 3-11 Attempts

have been made to quantify the amount of tritium in the samples based on the appearance of

the autoradiograph, 2,7,19 but currently there are too many unknown variables to provide

reliable data. Brass, et al., 12 have found that in pure nickel, trapping of hydrogen occurs at

grain boundaries. In the same laboratory, Roux, et al., 13 studied a single crystal alloy

based on the _ phase of CMSX-2 and found that hydrogen strongly segregated to the

eutectic 7?'{. Since the silver grains were observed uniformly through the eutectic TIT', it is

believed that a higher tritium solubility in this constituent compared to the surrounding

235
microstructure can explain the enhanced tritium concentration. 2° This is especially relevant

to this work because of the eutectic 7/7' present in PWA 1480. There have not been any

efforts to study the hydrogen solubility or trapping characteristics of the 7/7' microstructure

using tritium autoradiography. Since the best resolution attainable is only about 0.1 l.tm,

the results would be ambiguous because the _, matrix channels in the _'/7' microstructure are

about 0. l_m, and it would not be possible to separate trapping at the 7/Y interface from

lattice trapping or solubility differences.

E.2 Experimental Procedures

Samples were cut to approximately 1 cm x 2 cm x 0.2 cm on a slow speed diamond

cut-off wheel. The dimensions were chosen to give both a proper current density during

charging and optimum resolution from the emulsion. The samples were hand-ground to

600 grit silicon carbide paper and then polished with 1 lam A1203 slurry. The tritium

charging facility is housed in a self-contained glove box, and tritium levels were measured

in the glove box and an interchange container by a Betatec S/N 611 ionization chamber

manufactured by Overhoff Technology. Figure E. 1 is a schematic of the tritium charging

facility. The fume hood has a nominal flow rate of 600 cfm and exhausts to a roof vent. A

small negative pressure was maintained in the glove box and monitored by a magnehelic

gauge. Periodic smear samples were taken from the glove box, and continuous room air

monitoring occurred during the test.

Tritium was introduced into the sample by cathodic charging in a witiated molten

salt solution. The initial amount of tritium was 10 Curies and a specific activity of 0.2

Ci/cm 3 was used. By-products of the tritiated solution were channeled through a series of

condensers and traps to minimize the amount of released tritium and tritiated water vapor.

The molten salt consists of 57% by weight NaHSO4-H20 (Sodium Bisulfate) and 43%

KI_tSO4 (Potassium Bisulfate). Cathodic charging took place at 150°C with the sample

236
Electrolyte: Eutectic mixture.

NoHSO 4.HzO 57%

KHSO 4 4;5%

A di° ,,.oc.io°:.,o ° • 3so.'-- 2. - 3.so; * 1⁄2o;"

Cathodic reaction : HSO 4- • • _ H • S04 z"

love box

• rtf. llllle.

/1

Potent io_tot

Figure E. 1 Schematic of the tritium charging apparatus.

acting as the cathode. The surface of the sample was fLrst etched by applying -25 mV and

10 mA for four minutes. Charging was done at -800 mV and 10 mA for 2-4 hours, which

resulted in a low current density of 5 mA/cm 2. After charging, the sample was removed

from the pyrex kettle and rinsed with distilled water. The tritiated absorbed layer is then

removed by hand-grinding with 600 grit paper. The sample remained in the glove box at

room temperature for 24 hours to allow outgassing of the "diffusible" tritium.

Subsequently, it was transferred into the interchange chamber and removed after significant

tritium outgassing ceased. Further sample manipulations were made using only disposable

rubber gloves since the "diffusible" tritium had been removed. The sample was further

polished under a fume hood using 1 I-tm diamond paste. After polishing, the tritium

concentration was measured by a liquid scintillation counter. Typical concentrations ranged

from 6,000 to 20,000 disintegrations per minute (dpm). Polishing was continued to 0.05

_lm Ai203 and samples were etched using one of the etches previously mentioned in section

3.3.1.

237
Typically, a liquid emulsion then would be applied to the sample, however in this

study a stripping film containing an emulsion layer was used instead. The liquid emulsions

are much harder to work with and, more importantly, only have a shelf life of two

months. 21 The liquid emulsion does have the advantage of a better resolution of slightly

less than 0.1 I.tm, 18 however this is not sufficient to observe trapping at the T/q( interface.

The primary concern of this study was trapping at the dendritic macrostructure and eutectic

y/q/, so such a fine resolution was not needed. Therefore AR-10 stripping film from

Kodak was used, which has a resolution of 2.0 I.tm and a shelf life of one year. 22 The

stripping layer, mounted on a glass plate, consists of a 5 l.tm thick emulsion coating on a

gelatin layer 10 p.m thick. In order to transfer the emulsion to the sample, squares slightly

larger than the samples were cut and placed emulsion side down in a sugar solution of

controlled bromide concentration. It is claimed that this allows better transfer and gives

improved results when using long exposure times. 23 The solution consisted of 0.05 g

potassium bromide, 100 g sucrose and 500 ml of distilled water. This solution is then

further diluted 9:1 with distilled water.

In a darkroom using a safelamp with a Kodak Series 1 filter, the stripping film was

floated emulsion side down for three minutes to allow expansion. The sample was dipped

in the solution and raised at a 45 ° angle so that the stripping film smoothly coated the

sample. It was then allowed to dry slowly in a dust-free environment. After drying, it was

stored in a light-tight dry container and exposed for 3-5 days. The sample container was

immersed in liquid nitrogen to minimize the background created by tritium diffusion. The

dcveloping instructions furnished with the AR-10 stripping film were followed:

• 5 minutes in full strength Dektol developer.

° Rinse in slow running water for 30 seconds.

• Fix for 8 minutes in Kodak fixer.

• Rinse in slow running water for 2-3 minutes.

• Dry slowly in a dust-free environment.

238
Observation of the developed sample was done on an optical microscope and SEM.

The underlying microstructure can be seen easier on the optical microscope, however the

distribution of the silver grains is better observed using the energy dispersive x-ray unit

(EDX) on the SEM.

E.3 Results and Discussion

Tritium was successfully charged into many samples in sufficient concentrations to

ensure that some tritium remained after polishing. After many trial and error attempts at

photographic developing and subsequent observation, all experimental obstacles had been

overcome. It was found that placement of the stripping film on the sample and exposure

time were two of the more critical steps. Stretch marks on the emulsion layer during drying

led to many artifacts, as did any dust that may have settled on the sample.

The underlying microstructure was seen quite easily on the optical microscope, and

to observe silver grains, the microscope was focused on the gelatin layer above the sample.

The silver grains were black, random in shape and a few micrometers in diameter. Figure

E.2 shows silver grains uniformly across the sample with no segregation to the eutectic

Y/l('. In order to confirm that these were in fact silver grains, the energy dispersive x-ray

unit (EDX) on the SEM was used. On the SEM the microstructure was almost

indistinguishable through the gelatin layer, and there were serious charging problems.

Despite these problems, it was possible to get a map of the silver distribution using the

EDX unit, as shown in Figure E.3. This confirmed the optical microscopy results that

showed a uniform distribution of silver. It was not possible to correlate the silver

distribution with the underlying microstructure because of the effects of the gelatin layer.

239
ORIGINAL PAGE

BI_CK AND _H!TI_ PHO.r_OGRAP_H

lOgm

(a) (b)

Figure E.2

lOlam

(c) (d)

(a & c) Optical micrograph of eutectic _,/'/observed through the gelatin

layer and (b & d) corresponding Ag layer.

lOlam

240
ORIGINAL PAGE

BLACK AND WHITE PHO[OGRAP.H

Figure E.3

25 I.tm 25 l.tm

(a) (b)

(a) SEM micrograph of microstructure viewed through the gelatin layer.

(b) Corresponding uniform distribution of Ag grains.

The preliminary results differ from the results of Roux, et al., 13 that showed a

strong segregation of the tritium to the eutectic _,/y. The study of Roux, et al., was

performed by the group at the Universit6 Paris-Sud, which has had extensive experience in

tritium autoradiography. It remains unexplained why this study resulted in an apparent

uniform distribution of tritium. A more thorough study needs to be done varying the

charging conditions, outgassing and exposure times. It also would be worthwhile to use

one of the liquid emulsions to facilitate observation on the SEM.

241
References

.

,

.

4.

.

.

.

.

.

12.

13.

14.

15.

16.

17.

G.R. Caskey, Adv. Tech. for Characterizing Hydrogen in Metals, eds. N.F. Fiore

and B.J. Berkowitz, Warrendale, Pa, 1981, 61-75.

N.A. Tiner, T.L. Mackay, S.K. Asunmaa and R.G. Ingersoll, Trans. ASM, 61,

(1968), 195-202.

M. Aucouturier, G. Lapasset and T. Asaoka, Metallography, 11, (1978), 5-21.

J. Chdne and M. Aucouturier, Hydrogen in Metals, eds. I.M. Bernstein and A.W.

Thompson, TMS-AIME, 1980, 583-595.

J.P. Laurent, G. Lapasset, M. Aucouturier and P. Lacombe, Hydrogen in Metals,

eds. I.M. Bernstein and A.W. Thompson, ASM, 1974, 559-573.

M. Aucouturier, Current Solutions to Hydrogen Problems in Steels, eds. C.G.

Interrante and G.M. Pressouyre, ASM, 1982, 407-412.

G.P. Marsh, United Kingdom Atomic Energy Authority AERE-R8560, "An

Assessment of Some Techniques Available for the Local Detection of Hydrogen in

Metals," 1976.

T.D. Le and B.E. Wilde, Current Solutions to Hydrogen Problems in Steels, eds.

C.G. Interrante and G.M. Pressouyre, ASM, 1982, 413-422.

G.L. Downs, Environmental Degradation of Engineering Materials, eds. M.R.

Louthan, Jr., R.P. McNitt and R.D. Sisson, Jr., Virginia Tech. Publ., Blacksburg,

Va., 1981, 425-435.

G.R. Fritsche and B.M. Strauss, Metallography, 4, (1971), 223-230.

H. Ohma, G.P. Tiwari, Y. Iijima and K. Hi.rano, Proc. JIMIS-2, Hydrogen in

Metals, Japan Inst. of Metals, Minakami, Japan, 1980, 229-232.

A.M. Brass, A. Chanfreau and J. Chdne, Hydrogen Effects on Materials Behavior,

eds. N.R. Moody and A.W. Thompson, TMS-AIME, 1989, to be publ.

R. Roux, A.M. Brass and J. Chene, Hydrogen Effects on Material Behavior, eds.

N.R. Moody and A.W. Thompson, TMS-AIME, 1989, to be published.

C.L. Baker, J. Chene, I.M. Bernstein and J.C. Williams, Met. Trans., 19A, (1988),

73-82.

S.tI. Goods, Scripta Met., 20, (1986), 565-569.

J.D. Braun, G.L. Downs and G.W. Powell, Metallography, 4, (1971), 231-242.

G.W. Powell, J.D. Braun, K.F. Chaney and G.L. Downs, Corrosion, 26, (1970),

223-228.

242
18.

19.

22.

P. Lacombe, M. Aucouturier and J. Ch6ne, Hydrogen Embrittlement and Stress

Corrosion Cracking, eds. R. Gibala and R.F. Hehemann, ASM, 1980, 79-102.

J. M.R. Louthan, J. D.E. Rawl and R.T. Huntoon, Corrosion, 28, (1972), 172-

176.

J. Ch6ne, Universite Paris Sud, unpublished research.

Kodak Tech Bits, Spring 1988, "Autoradiography at the Light Microscope Level,"

Eastman Kodak Co., Rochester, NY, 1988.

Kodak Publ. P-64, "Kodak Products for Light Microscope Autoradiography,"

Eastman Kodak Co., Rochester, NY, 1987.

23. Brit. J. of Radiology, 42, (1969), 862-870.

243
APPENDIX F

QUANTITATIVE RESULTS OF REPPICH CRSS MODEL

This appendix will quantitatively examine the applicability of the Reppich model to

PWA 1480. Reppich and co-workers 1-3 have analyzed the increase in the flow stress due

to the presence of coherent, ordered particles as a function of their size and volume

fraction. To apply the model, it must be determined whether the dislocations will be

weakly coupled or strongly coupled, and the critical particle size for this transition, dws, can

be determined using: 2

(5.6)

where w is a term accounting for the elastic repulsion between paired dislocations outside

the particles, T is line tension and y is the APB energy. For similar alloys w = 2.8, 4 and it

has been reported that 7 = 118 rnJ/m 2 in PWA 1480. 5

The constant line tension approximation is inaccurate for these anisotropic alloys,

so a more exact form will be used. Ham 6 has shown that in these types of alloys, edge

dislocations are more resistant to motion than screw dislocations, and Reppich gives the

line tension as: 4

T = K(0) + dO2 .)k4r_J In _.

(5.7)

where L is the distance between precipitates in the [ 111 } plane and can be calculated from

I. = (rt/6f)U2d, where f is the volume fraction of precipitates and d is the precipitate size.

The quantity (K(0) + d2K/d02) is equal to G/2 for edge dislocations, where G is the shear

244
moduhJs. Using f = 0.7, d = 0.45 p.m and b = 2.5 x 10 .8 cm, the line tension for edge

dislocations reduces to:

Gb 2

T = _ (5.8)

3.4

This is slightly different than the constant line tension approximation of Gb2/2 and is close

to other estimates. 4,7 Using this line tension value and G = 54.5 MPa, equation 5.6 gives

a value of 0.18 I.tm for dws. This means that in PWA 1480 with d = 0.45 I.tm, d > dws, and

strong pair coupling occurs, as expected.

According to the Huther-Reppich model, 2 the increase in the CRSS due to overaged, spherical particles that are cut by strongly coupled dislocations is: 2

Twfl/2: d"t _ 1"_1/2

Ax o = 0.86_ 1.28_-'_ )

(5.9)

where b is the Burgers vector and f is the volume fraction of precipitates. The stress to

cause the fh-st dislocation to penetrate the precipitate is: 3

"t 0.72Gb

_A'tp= (5.10) b d

Since Orowan looping does not occur at room temperature because of the high volume

fraction of precipitates, one of the two processes described by (5.9) and (5.10) must be

controlling the CRSS. The cutting stress (5.9) was 103 MPa, and the penetration stress

(5.10) was 449 MPa. The highest of these two will be the controlling strengthening

mechanism, and therefore the CRSS is modeled using the penetration stress.

The Schwarz-Labusch correction 8 must be used because the Huther-Reppich model

is based on a low volume fraction of precipitates. The "reduced particle depth", rlo, is used

245
andforlargeparticlescanbewrittenas:s

(5.11)

A valueof 0.47wasobtainedfor PWA 1480clearly not fulfilling therio<<1 criterion to

uselow volumefraction,Fleischer-Friedeltypemodels.9'1°The interpolationformulato

convertto highvolumefractionmodelsis:8

A'_o = 0.95(1 + Crlo)AXFF (5.12)

where A'_N: is the increase in CRSS due to the presence of a low volume fraction of

particles, as calculated above, and C is a constant close to three for high volume fractions. 4

The increase in the CRSS due to the presence of a high _'olume fraction ofT' precipitates

has been calculated to be 1028 MPa. To obtain the total CRSS for the alloy, the solid

solution strengthening contribution of the matrix, XM, must be incorporated. A variety of

addition rules have been used to calculate the total CRSS in this type of alloy. 11 Linear

addition has met with some success 7'12 and will be used in this study. Using an alloy of

similar composition to the T matrix phase of PWA 1480, A'_ M was determined to be about

115 MPa at room temperature. 13 The total XCRSS then becomes 1143 MPa, which is

considerably higher than the 450 MPa obtained for PWA 1480.

While the Huther-Reppich model does not quantitatively predict the CRSS well, it

does show that initial penetration of the T' precipitate is controlling the CRSS of PWA

1480. There are two reasons why the model did not predict the observed CRSS more

closely. Perhaps the largest inaccuracy is that the Huther-Reppich model only considers

the leading dislocation and does not account for the interaction term between the pair of

dislocations. This explains why the leading term in equation 5.3 is 7/b and not 7_b, as in

other order strengthening theories. 9"11,13,14 This would have reduced the CRSS by almost

h,df, resulting in a reasonable value, i:urther, Ardel111 points out that there is a large

246
spread in the values obtained by different investigators for the constants C and

W 4,7.12,15,16

References

1. W. Huther and B. Reppich, Mat. Sci. and Eng., 39, (1979), 247-259.

2. W. Huther and B. Reppich, Z. Metallk., 69, (1978), 628-634.

3. B. Reppich, Acta Met., 30, (1982), 87-104.

4. B. Reppich, P. Schepp and G. Wehner, Acta Met., 30, (1982), 95-104.

5. M. Dollar and I.M. Bernstein, Superalloys 1988, eds. D.N. Duhl, et al., TMSAIME, 1988, 275-284.

6. R.K. Ham, Trans. Japan Inst. Metals, Vol. 9 supplement, (1968), 52.

7. A.W. Thompson and J.A. Brooks, Acta Met., 30, (1982), 2197-2203.

8. R.B. Schwarz and R. Labusch, J. Appl. Phys., 49, (1978), 5174-5187.

9. L.M. Brown and R.K. Ham, Strengthening Methods in Crystals, eds. A. Kelly and

R.B. Nicholson, Wiley, New York, 1971, 96-112.

10. R.L. Fleischer, Acta Met., 11, (1963), 203-209.

11. A.J. Ardell, Met. Trans., 16A, (1985), 2131-2165.

12. P. Haasen and R. Labusch, 5th Int'l Conf. on Strength of Metals and Alloys, eds. P.

Haasen, V. Gerold and G. Kostorz, Pergamon Press, Oxford, 1979, 639-643.

13. S.M. Copley and B.H. Kear, Trans. TMS-AIME, 239, (1967), 984-992.

14. A. Melander and P.A. Pearson, Metal Sci., 12, (1978), 391-398.

15. E. Nembach, Z. Metallkd., 72, (1981), 401-405.

16. M. Grohlich, P. Haasen and G. Frommeyer, Scr. Met., 16, (1982), 367-370.

247
Nat+or,_l Ae+onatJb(:s _tT_d

Spa(e A(£T}_r/istr AtlO+_

1. Report No 2. Government Accession No.

NASA CR-]85219

4. Title and Subtitle

The Effect of Hydrogen and Microstructure on the Deformation and Fracture i

Behavior of a Single Crystal Nickel-Base Superalloy

7. Author(s)

William S. Walston

9. Performing Organization Name and Address

Carnegie-Mellon University

Pittsburgh, Pennsylvania 15213

t2. Sponsoring Agency Name and Address

National Aeronautics and Space Administration

Lewis Research Center

Cleveland, Ohio r14135-319J

3. Recipient's Catalog No.

5. Report Date

April i 990

6. Performing Organization Code

8. Performing Organization Report No,

None

Report Documentation Page

10. Work Unit No.

553-13-00

11. Contract or Grant No.

NAG3-463

13 Type of Report and Period Covered

Contractor Report

Final

14• Sponsoring Agency Code

15. Supplementary Notes

Project Manager, Robert L. Dreshfield, Materials Division, NASA Lewis Research Center. This report was a

thesis submitted to Carnegie-Mellon University in January 1990 in partial fulfillment of the requirements for the

degree of Doctor of Philosophy in Metallurgical Engineering and Materials Science.

_6 -Kbst;aciA study has been conducted on the efli:cts of internal hydrogen and microstructure on the deformation and fracture of a

single crystal nickel-base superalloy. In particular, room temperature plane strain fracture toughness and tensile tests were

pcrlbrmed on hydrogen-free and hydrogen charged samples of PWA 1480. The role of microstructure was incorporated by

varying the levels of porosity and eutectic y/',(' through hot isostatic pressing and heat treatment. The room temperature beha_,ior of PWA 1480 was unusual because precipitate shearing was not the primary deformation mechanism at all strains. At

strains over 1%, dislocations were trapped in the "f matrix, and an attempt was made to relate this behavior to

conapositional differences between PWA 1480 and other superalloys. Another unique feature of the tensile behavior was

cleavage of the cutcctic y/y', which is believed to initiate the failure process. Fracture occurred on 111 planes and is

likely a result of shear localization along these planes. Elimination of the eutectic y/y' greatly improved the tensile

ductility, but porosity had no effect on tensile properties. Large quantities of hydrogen (1.74 at. %) were gas-phase charged

int,, the material, but surprisingly this was not a function of the amount of porosity or eutectic y/y' present. Desorption

cxperm_ents suggest that the vast majority of hydrogen is at reversible lattice trapping sites. This large, uniform

concentration of hydrogen dranmtically reduced the tensile strain to failure, but only slightly affected the reduction in area.

Further evidence of strain localization due to hydrogen was observed in the deformation structure. Available hydrogen

embrittlemcnt models were examined in light of these results and it was found that the hydrogen enhanced localized

plasticity model can explain much of the tensile behavior. KIc fracture toughness tests were conducted, but it was

necessary to also perform Jic tests to provide valid data. Unexpectedly, the values from these tests did not agree;

further, the crack growth behavior and fractography differed greatly. The eutectic y/y' behaved differently in these

two types.._ of tests, and the possible effect of the unloading 9resent in Jtc tests on this constituent was examined•

-iT_--Key Words (_-ugge_ted g Author(_ 18. Distribution Statement

Supcralloys

Single crystals

Hydrogen effects

Nickel-base alloy

19 Security C assif (of this report)

Unclassi fled

Unclassified - Unlimited

Subject Category 26

20. Security Classif (of this page)

Unclassified

21 No of pages

255

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